The Influence of Austenite Grain Size on Hot Ductility of Steel

99
NOTE This online version of the thesis may have different page formatting and pagination from the paper copy held in the University of Wollongong Library. UNIVERSITY OF WOLLONGONG COPYRIGHT WARNING You may print or download ONE copy of this document for the purpose of your own research or study. The University does not authorise you to copy, communicate or otherwise make available electronically to any other person any copyright material contained on this site. You are reminded of the following: Copyright owners are entitled to take legal action against persons who infringe their copyright. A reproduction of material that is protected by copyright may be a copyright infringement. A court may impose penalties and award damages in relation to offences and infringements relating to copyright material. Higher penalties may apply, and higher damages may be awarded, for offences and infringements involving the conversion of material into digital or electronic form.

Transcript of The Influence of Austenite Grain Size on Hot Ductility of Steel

Page 1: The Influence of Austenite Grain Size on Hot Ductility of Steel

NOTE

This online version of the thesis may have different page formatting and pagination from the paper copy held in the University of Wollongong Library.

UNIVERSITY OF WOLLONGONG

COPYRIGHT WARNING

You may print or download ONE copy of this document for the purpose of your own research or study. The University does not authorise you to copy, communicate or otherwise make available electronically to any other person any copyright material contained on this site. You are reminded of the following: Copyright owners are entitled to take legal action against persons who infringe their copyright. A reproduction of material that is protected by copyright may be a copyright infringement. A court may impose penalties and award damages in relation to offences and infringements relating to copyright material. Higher penalties may apply, and higher damages may be awarded, for offences and infringements involving the conversion of material into digital or electronic form.

Page 2: The Influence of Austenite Grain Size on Hot Ductility of Steel

The Influence Of Austenite Grain Size

On Hot Ductility Of Steels

A thesis submitted in fulfilment of the requirements

for the award of the degree

MASTER OF ENGINEERING BY RESEARCH

From

UNIVERSITY OF WOLLONGONG

By

Suk-Chun Moon, B.Eng (Met.)

DEPARTMENT OF MATERIALS ENGINEERING

2003

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1

CERTIFICATION

I, Suk-Chun Moon, declare that this thesis, submitted in fulfilment of the requirements for the

award of Master of Engineering by Research, in the Department of Materials Engineering,

University of Wollongong, is wholly my own work unless otherwise referenced or acknowledged.

The document has not been submitted for qualifications at any other academic institution.

Suk-Chun Moon

December 2003

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ACKNOWLEDGEMENTS

I wish to express my gratitude to Prof. Rian Dippenaar who has supervised my work. I deeply

appreciate his prudent regard on all things. He always encouraged me saying “Excellent”. His

words always made me confident.

I express my gratitude to POSCO and its staff for giving me this opportunity to study abroad

and supporting unsparingly. I also acknowledge Mr. Shin-Eon Kang (POSCO Technical Research

Laboratories) for providing the experimental materials.

I am deeply indebted to Mr. Bob DeJong for his support with the hot tensile tests on the

GLEEBLE 3500. I also thank Mr. Ron Marshall for preparing my specimens and Mr. Greg

Tillman for assistance in metallographic preparation. Many thanks to Dr. Dominic Phelan and Mr.

Mark Reid for their assistance with my experiments.

I thank Rev. Tae-Joo Lee who helped my family to adapt easily when we arrived in Australia in

cold windy winter.

I am most grateful to my parents for their deep concern for me and my family and for their

support over such a long distance.

Finally, my most severe thanks to my wife, Kwang-Young and my son, Joon-Young, for their

patience, loving support and encouragement to finish my task and giving me good heart and love.

Also, thanks to my unborn baby who will bring ‘big surprise’ sooner or later.

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ABSTRACT

The high temperature ductility of steel increases as the grain size is decreased. However, in

microalloyed steels this interdependence of grain size and hot-ductility is generally less

pronounced because of the overriding effect of precipitation of carbo-nitrides of alloying elements

at grain boundaries. Nevertheless, many types of crack have been shown to be associated with

coarse austenite grains, and the tendency to crack is reduced when the formation of coarse

austenite grains is prevented by the use of suitable secondary cooling. To our knowledge, a

systematic study, to isolate the contribution of grain size on hot-ductility has not been reported and

hence, under these circumstances, it is important to define clearly the contributing role of grains

size to hot-ductility and therefore it is necessary to separate the effect on hot ductility of

microalloying precipitates from that of grain size.

The materials used for study were Fe-0.05%C, Fe-0.18%C and Fe-0.45%C alloys, prepared in

an experimental facility. Carbon was the only alloying element deliberately added to the melt and

tramp elements and impurities were kept to as low a value as experimentally possible. A

GLEEBLE 3500 thermomechanical simulator was used to conduct hot-tensile tests and from these

results the relationship between austenite grain size and hot-ductility could be determined. The

hot-ductility tensile tests in this study were performed in the temperature range 1100 to 700�.

Hot-tensile test specimens were either solution treated or melted in-situ (direct cast) and cooled

to the test temperature in order to simulate the microstructural characteristics of commercially as-

cast slab which contains coarse austenite grains at the base of oscillation marks. In order to obtain

different grain sizes, the specimens were solution treated at temperatures between 1100� and

1350� for 10minutes. The specimens were then rapidly cooled to the test temperature, in the

range 1100� to 700� at a rate of 200�min-1 and then pulled to fracture at a low strain rate of

7.5×10-4s-1. Specimens that were cast in-situ (referred to as the ‘direct casting condition’) were

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cooled at rates of 100�min-1 and 200�min-1 respectively in order to study the effect of cooling

rate on hot ductility.

It was found that grain size increased almost linearly with increasing solution treatment

temperature in all the Fe-C alloys and in case of specimens solution treated at 1350� an average

austenite grain size of ~4mm in diameter was obtained. Increasing the grain size resulted in

ductility loss under all testing conditions. The existence of a ductility trough between Ar3 and Ae3

temperature was considered to be due to the formation of deformation induced ferrite as evidenced

by a constant peak stress region between these two temperatures. However, convincing

experimental evidence of the ductility trough extending beyond the Ae3 temperature well into the

austenite was found. The mechanism of this interesting, and important observation for low carbon

alloy below 0.3%C has not been explored as yet and at present it has to be assumed that it is

related to the occurrence of grain boundary sliding because in the high austenite temperature

region the grain boundary sliding is favored in a coarse grained structure.

For specimens cast in-situ, the largest grains were found in the Fe-0.18%C alloy for both

cooling rates. This important observation is attributed to the higher austenitizing temperature of a

Fe-0.18%C alloy compared to that of the other Fe-C alloys studied. Moreover, the formation of

columnar austenite grains were observed in this alloy on cooling, whereas equi-axed grains were

formed in the other two Fe-C alloys. By such columnarization, the surface cracking susceptibility

of the peritectic grade steel will be accelerated. The hot-ductility of the Fe-C alloy of near-

peritectic composition was the lowest of the alloys studied and the inferior ductility in this alloy is

attributed to the coarse grain size and columnar shape of the grains. The Fe-0.45%C alloy had the

smallest grain size at any specimens cast in-situ but the hot-ductility was much lower than would

have been expected in an alloy of such small grain size. This much reduced ductility may be due

to increased grain boundary sliding.

At a higher cooling rate of the in-situ melted specimens, smaller grains were produced in all the

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Fe-C alloys resulting in ductility improvement. This grain refinement obtained at higher cooling

rates have important implications for near-net shape casting operations such as thin-slab casting or

strip-casting where much higher cooling rates are realized than in the conventional casting process,

if the factors related with precipitation are removed.

An important insight derived at through this study was that enlarged grain contributed more to

reduced ductility at high temperature than did cast structure, at least under the pertaining

experimental conditions. This observation has important practical implications because it means

that efforts in industry could be concentrated on reducing the chances of forming large austenite

grains, such as at the roots of oscillation marks due to a decreased cooling rate, without undue

regard to the effect of these measures on cast structure.

This study has provided convincing new experimental evidence of the extremely detrimental

effect of large austenite grains on hot-ductility in plain carbon steels. The very large columnar

shaped grains that can form in alloys of near-peritectic composition is particularly disconcerting.

However, the influence of AlN precipitation on austenite grain boundaries on hot-ductility was not

studied and it is recommended that this important topic should be included in subsequent

investigation. The experimental data on these Fe-C alloys, provided in this study, may now be

used to benchmark and further analyse the much larger body of information on low-alloyed steels

available in the literature.

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TABLE OF CONTENTS CHAPTER 1 - INTRODUCTION 1

CHAPTER 2 - LITERATURE REVIEW 4 2.1 Continuous casting process 4

2.1.1 Outline of continuous caster 4

2.1.2 Near net shape casting technology 5

2.2 Slab cracking 9

2.2.1 Features of Cracking 9

2.2.2 Coarse prior austenite grains 11

2.3 Hot ductility test 14

2.3.1 Simulation of straightening operation during continuous

casting 14

2.3.2 Suitability of the hot tensile test to the problem of

transverse cracking 15

2.4 Fracture mechanisms 16

2.4.1 Region of embrittlement 17

2.4.1.1 Embrittlement by strain concentration and

microvoid coalescence at grain boundaries 17

2.4.1.2 Grain boundary sliding 21

2.4.2 High ductility, low temperature region 23

2.4.3 High ductility, high temperature region 23

2.4.4 Hot ductility behaviour of steels 24

2.4.4.1 Plain C-Mn and C-Mn-Al steels 24

2.4.4.2 C-Mn-Al steels with high Al and N levels 25

2.4.4.3. Microalloyed steels 26

2.5 Factors influencing hot ductility 27

2.5.1 Grain Size 27

2.5.2 Precipitation 29

2.5.3 Composition 32

2.5.3.1 Sulphur 32

2.5.3.2 Carbon 35

2.5.4 Cooling rate 38

2.5.5 Strain rate 38

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2.5.6 Thermal history 39

CHAPTER 3 - EXPERIMENTAL 41 3.1 Preparation of specimens 41

3.2 Hot ductility test 42

3.2.1 GLEEBLE3500 42

3.2.2 Thermomechanical cycles 43

3.2.3 Measurement of hot ductility 45

3.2.4 Correction of stress-strain curve 45

3.2.5 Determination of GLEEBLE setting temperature for melting 47

3.3 Metallography 49

CHAPTER 4 - RESULTS 50 4.1 Hot ductility curves 50

4.1.1 Fe-0.05%C alloys 50

4.1.2 Fe-0.18%C alloys 52

4.1.3 Fe-0.45%C alloys 53

4.2 Stress strain curves 55

4.2.1 Fe-0.05%C alloys 55

4.2.2 Fe-0.18%C alloys 57

4.2.3 Fe-0.45%C alloys 58

4.3 Austenite grain size 60

CHAPTER 5 - DISCUSSION 67 5.1 Grain growth 67

5.2 Ductility troughs 69

5.3 Influence of grain size on hot ductility 74

5.3.1 Reduction in area 74

5.3.2 Position and width of ductility trough 76

5.4 Influence of carbon content on hot ductility 77

5.5 Influence of cooling rate following direct-casting on hot ductility 81

5.6 Comparison between as-cast condition and solution treatment

condition 82

5.7 Practical implications of the experimental findings 83

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CHAPTER 6 - CONCLUSIONS 85

BIBLIOGRAPHY 88

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LIST OF FIGURES

Fig.2.1 Schematic diagram of typical continuous casting machine [1] 4

Fig.2.2 Schematic diagram of (a) conventional continuous casting and hot-rolling (b) Thin slab-hot

rolling (TSHR) (c) Strip-casting 6

Fig.2.3 Practical examples of thin-slab casting techniques 9

Fig.2.4 Widespread crazing and fine transverse cracks at oscillation marks on the as-cast surface of

a line pipe steel slab (top side). Etched in hot HCL. Magnification not specified. CD :

Casting Direction [4] 10

Fig.2.5 Crazing around a transverse crack at the base of an oscillation mark on the as-cast top

surface of a 0.20%C steel slab. Etched in hot HCL [4]. 10

Fig.2.6 Section through the surface across a system of transverse cracks through the coarse-grained

zone. Prior austenite grain boundaries white (ferrite-decorated). Etched in nital.

Magnification not specified [6]. 11

Fig.2.7 Grain growth of austenite during continuous cooling. The specimens were remelted at

1580�, cooled to a given temperature at a rate of 0.28�s-1, and then quenched into water

[7]. 12

Fig.2.8 Formation of surface cracks due to blown grain during casting [4] 14

Fig.2.9 Schematic diagram of a ductility curve defining the three characteristic regions of hot-

ductility [1] 17

Fig.2.10 Schematic diagram showing mechanism for transformation induced intergranular failure

[1] 18

Fig.2.11 Microstructure at room temperature of steel tested at 800�, showing intergranular failure

associated with a thin layer of grain boundary ferrite [17] 19

Fig.2.13 Schematic illustration of intergranular microvoid coalescence of Nb-bearing steels. a-c

depict deformation in austenite above the Ar3 temperature, d-f depict deformation in the

(austenite + ferrite) region 21

Fig.2.14 Schematic models showing the formation of wedge cracks by grain boundary sliding: The

arrows indicate the sliding boundaries and the sense of translation [23] 22

Fig.2.15 Intergranular microvoid coalescence type fracture in C-Mn-Al steel, solution treated at

1350�, tested at 850� at strain rate of 10-3s-1 [16] 24

Fig.2.16 (a) C-Mn-Al steel showing flat facets on fracture surface, (b) enlarged view of a showing

lack of voiding around MnS inclusions [16] 25

Fig.2.17 Hot ductility curves for C-Mn steels at various grain sizes (a) 0.19wt%C steel (b)

0.65wt%C steel [32] 28

Fig.2.18 Influence of (a) particle size and (b) interparticle spacing on hot ductility of Nb-containing

steels, solution treated at 1330�, cooled to a test temperature of 850�, and fractured at a

strain rate of 3×10-3s-1 [1] 30

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Fig.2.19 Effect of sulphur content on the minimum reduction of area for two cooling rates of 1°Cs-1

and 4°Cs-1 [51] 33

Fig.2.20 Schematic illustration showing the effect of C on surface cracking of continuous cast slabs

[7] 35

Fig.2.21 Effect of C content on austenite grain size and calculated value of RA (a) Relationship

between austenite grain size and C content. Specimens (0.35%Si-1.5%Mn-0.05%Nb)

were remelted at 1580�, cooled to 900� at a rate of 5�s-1 and then quenched. (b)

Calculated values RA which were deduced from the relationship between grain size and RA

assuming that the specimens were deformed at 800� at a strain rate of 0.83×10-3s-1 [7] 36

Fig.2.22 Influence of C and Mn on the width of the ductility trough [39] 37

Fig.3.1 Geometry of GLEEBLE specimen 41

Fig.3.2 Schematic diagram of the GLEEBLE testing arrangement 43

Fig.3.3 Schematic diagram of thermomechanical cycles for hot ductility tests under the conditions

of (a) Solution treatment and (b) Direct casting 44

Fig.3.4 Geometry of sample after fracture 45

Fig.3.5 Measured tensile force during tensile deformation 46

Fig.3.6 Stress-strain curve. (a) curves obtained from GLEEBLE (b) modified curves 46

Fig.3.7 The results of preliminary experiments for the determination of the apparent melting point,

(a) Fe-0.05%C alloy (b) Fe-0.18%C alloy (c) Fe-0.45%C alloy 48

Fig.4.1 Hot ductility curves for Fe-0.05%C alloys under (a) solution treatment condition (b) direct

casting condition 51

Fig.4.2 Sample geometry after fracture tested at 850� under solution treatment condition 52

Fig.4.3 Hot ductility curves for Fe-0.18%C alloys under (a) solution treatment condition (b) direct

casting condition 53

Fig.4.4 Hot ductility curves for Fe-0.45%C alloys under (a) solution treatment condition (b) direct

casting condition 54

Fig.4.5 Stress-strain curves at different test temperatures for Fe-0.05%C alloys solution treated at

(a) 1100� (b) 1200� (c) 1350� 55

Fig.4.6 Stress-strain curves at different test temperatures for Fe-0.05%C Alloys under direct casting

condition at cooling rate (a) 100�min-1 (b) 200�min-1 56

Fig.4.7. Peak stress as a function of test temperature for Fe-0.05%C alloys under (a) solution

treatment condition (b) direct casting condition 56

Fig.4.8 Stress-strain curves at different test temperatures for Fe-0.18%C alloys solution treated at

temperature (a) 1100� (b) 1200� (c) 1350� 57

Fig.4.9 Stress-strain curves at different test temperatures for Fe-0.18%C alloys under direct casting

condition at cooling rate (a) 100�min-1 (b) 200�min-1 57

Fig.4.10. Peak stress as a function of test temperature for Fe-0.18%C alloys under (a) solution

treatment condition (b) direct casting condition 58

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Fig.4.11 Stress-strain curves at different test temperatures for Fe-0.45%C alloys solution treated at

temperature (a) 1100� (b) 1200� (c) 1350� 58

Fig.4.12 Stress-strain curves at different test temperatures for Fe-0.45%C alloys under direct

casting condition at cooling rate (a) 100�min-1 (b) 200�min-1 59

Fig.4.13 Peak stress as a function of test temperature for Fe-0.45%C alloys under (a) solution

treatment condition (b) direct casting condition 59

Fig.4.14 Cross-section of a GLEEBLE specimen of a Fe-0.05%C alloy. Etched in nital (diameter

of specimen = 10mm) 61

Fig.4.15 Cross-section of a GLEEBLE specimen of a Fe-0.18%C alloy. Etched in saturated picric

acid based etchant (diameter of specimen = 10mm) 62

Fig.4.16 Cross-section of a GLEEBLE specimen of a Fe-0.45%C alloy. Etched in saturated picric

acid based etchant (diameter of specimen = 10mm) 63

Fig.4.17 The distribution of austenite grain size of a Fe-0.05%C alloy under (a) solution treatment

condition (b) direct casting condition 64

Fig.4.18 The distribution of austenite grain size of a Fe-0.18%C alloy under (a) solution treatment

condition (b) direct casting condition 65

Fig.4.19 The distribution of austenite grain size of a Fe-0.45%C alloy under (a) solution treatment

condition (b) direct casting condition 66

Fig.5.1 Grain size as a function of solution treatment temperature 68

Fig.5.2 Grain size as a function of cooling rate under direct casting condition 69

Fig.5.3 Hot-ductility curves for Fe-0.05%C alloys for solution treated specimens having various

grain sizes 70

Fig.5.4 Hot-ductility curves for Fe-0.18%C alloys for solution treated specimens having various

grain sizes 70

Fig.5.5 Hot-ductility curves for Fe-0.45%C alloys for solution treated specimens having various

grain sizes 73

Fig.5.6 Relationship between tensile properties and reciprocal of austenite grain size (D) at

different test temperatures in solution treated specimen (a) Fe-0.05%C (b) Fe-0.18%C (c)

Fe-0.45%c alloy 75

Fig.5.7 Relationship between tensile properties and reciprocal of austenite grain size (D) at

different test temperatures under direct casting condition for (a) Fe-0.05%C (b) Fe-0.18%C

(c) Fe-0.45%c alloy 75

Fig.5.8 Relationship between minimum RA value and reciprocal of austenite grain size (D) for

different Fe-C alloys in the solution treatment condition 76

Fig.5.9 (a) Position of ductility trough, (b) width of trough as a function of grain size under solution

treatment condition 77

Fig.5.10 Hot ductility curves from solution treatments at (a) 1100� (b) 1200� (c) 1350� 78

Fig.5.11 Effect of C content on austenite grain size and RA value under direct casting condition for

cooling rate (a) 100�min-1 (b) 200�min-1 79

Fig.5.12 Hot ductility curves for Fe-0.18%C alloy from two different thermal conditions 82

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CHAPTER 1 - INTRODUCTION

The continuous casting of steel was introduced commercially around 1960. Through numerous

efforts to solve problems related to solidification phenomena, many developments and

improvement in technology were made resulting in large-scale production and high yield ratios.

However, one of the problems still pertaining to this process has been the occurrence of cracks on

the slab surface such as transverse cracking, crazing, corner cracking and star cracking. In many

cases these cracks remain even after rolling with high reduction in strip thickness. The

introduction of thin-slab casting and hot-direct-rolling (HDR), where surface inspection prior to

rolling is not possible, elevated the need to ensure the elimination of these cracks and the

improvement of surface quality if defect-free rolling is to be achieved. Maintaining a defect free

slab surface has now become a prime requirement for the economic production of HDR of steel

sheet.

The straightening operation of the continuously cast strand is carried out when the slab surface

temperature is in the range 1000� to 700�. These surface temperatures unfortunately coincide

with the temperature range in which steel exhibits a ductility minimum as measured in laboratory

hot-tensile tests [1]. Although numerous studies have been devoted to establishing a relationship

between transverse cracking and the hot-ductility trough as well as the factors that effect this

relationship, only a few researchers [7, 32, 39] have reported the hot-ductility behavior of the steel

in relation to the characteristics of austenite grain growth.

Generally, the high temperature ductility of steel increases as the grain size is decreased.

Refining the grain size leads to a reduction in both the depth and width of the ductility trough [32].

On the other hand, microalloyed steel generally shows little indication that grain size has a

significant influence on hot ductility. This observation is mainly due to the overriding effect of the

precipitation of AlN or Nb(C,N) at grain boundaries. Nevertheless, in the straightening operation

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of continuous casting strands, many types of cracks are well associated with coarse austenite

grains, and convincing evidence has been provided that the tendency to crack is reduced when the

formation of coarse austenite grains is prevented by the use of suitable secondary cooling [6].

Under these circumstances, it is important to define clearly the role of grains size in controlling

hot ductility. In order to do this, it is necessary to separate the effect on hot ductility of

microalloying precipitates from that of grain size.

It is generally conceded that the early stages of solidification and subsequent high temperature

phase transformations profoundly effect cast structure, and therefore also hot-ductility during the

straightening operation, but conclusions have mostly been drawn from indirect evidence because

of the difficulty of conducting reliable experiments at these high temperatures.

In the present study, a GLEEBLE3500 thermo-mechanical simulator has been used to evaluate

the ductility of three different Fe-C alloys. The relationship between austenite grain size and hot-

ductility of specimens in the temperature range 1100� to 700�, was obtained through hot tensile

tests under conditions similar to that of as-cast slab which contain coarse austenite grains at the

base of oscillation marks.

For the purposes of this investigation, plain carbon steels have been prepared in a laboratory

furnace. Carbon was the only alloying element deliberately added to the melt and tramp elements

and impurities were kept to as low a value as was experimentally possible. The rationale behind

this approach is to attempt to isolate the effect of austenite grain size on hot-ductility in the

absence of complicating factors such as the precipitation of alloy carbides and the presence of low

melting point impurities such as iron sulphides.

An attempt was made to prepare pure Fe-C alloys but aluminum had to be added as deoxidizer

and some ingress of nitrogen could not be avoided. Hence, as will be shown in Table 3.1, there is a

distinct possibility that AlN can form in the alloys prepared for this investigation. The influence of

AlN grain boundary precipitates on hot-ductility is well documented and no attempt was made in

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this study to isolate the contribution of AlN grain boundary precipitation to hot-ductility.

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CHAPTER 2 - LITERATURE REVIEW

2.1 Continuous casting process

2.1.1 Outline of continuous caster

A typical continuous caster is shown schematically in Fig.2.1 [1]. A tundish which is supplied

with molten steel from a ladle feeds it into the mold, which is oscillating and water-cooled. Mold

oscillation is carried by a hydraulically driven mold oscillation system to ensure surface quality in

the cast slab and to prevent sticking of the solidifying strand to the mould. This oscillation is

responsible for the formation of oscillation marks on the surface of the strand. From mold to

secondary cooling zone, molten steel is solidified into slab and cooled.

There are many stresses involved in this process ; Friction between solidified shell and mold

wall, ferrostatic pressure inside the shell, thermal stresses on the strand surface and bending

stresses at the straightening point. Whenever a tensile stress is present when the steel shell is in a

region of low ductility, there is the likelihood that crack can form.

Fig.2.1 Schematic diagram of typical continuous casting machine [1]

2.1.2 Near net shape casting technology

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Near net shape casting is advanced process technology that has been a major focus of recent

research and development efforts. The technology includes the methods by which liquid steel is

continuously cast into shapes more nearly approximating the final dimensions to be achieved in

the hot finishing stage of production. By doing so, the construction cost of plant and the

production cost of product can be reduced significantly [2]. These features are illustrated in

Fig.2.2. In this figure, the process routes of conventional slab continuous casting, thin-slab casting

and continuous strip-casting are compared.

Direct linkage or merger of two or more consecutive process steps is a powerful means by

which production cost in the steel manufacturing process can be reduced. Thin-slab hot rolling

(TSHR) and strip-casting are the latest examples of process developments in which these

principles are applied.

In the TSHR process, continuous casting and hot-strip rolling are integrated so that the

roughing mill is eliminated, whereas in the strip-casting process, the hot-strip rolling process is

eliminated completely. The thickness of thin slabs ranges between 50 and 150mm, whereas the

thickness of strips produced by strip-casting is about 2-6mm. Efforts are still being made to

develop a technology known as DSC for casting of slabs of plain carbon steels with thickness

between thin-slab casting and strip-casting by utilizing the horizontal twin belt concept.

6 0 0 m

S la b(2 0 0 ~ 2 5 0

m m )

R e h e a tin gF u rn a c e

R o u g h in gM ill

F in is h in gM ill C o ilin g

6 0 0 m

S la b(2 0 0 ~ 2 5 0

m m )

R e h e a tin gF u rn a c e

R o u g h in gM ill

F in is h in gM ill C o ilin g

(a) Conventional slab continuous casting and hot-rolling

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T h in S lab(50 ~8 5m m )

C o il B oxC o iling

200~300m

T hin S lab(50 ~8 5m m )

C o il B oxC o iling

T h in S lab(50 ~8 5m m )

C o il B oxC o iling

200~300m

(b) Thin-slab hot rolling (TSHR)

1 0 0m

C o ilin gF in a l S trip(0 .7 ~ 3 .0 m m )

C a s t S trip(1 .4 ~ 6 m m )

1 0 0m1 0 0m

C o ilin gF in a l S trip(0 .7 ~ 3 .0 m m )

C a s t S trip(1 .4 ~ 6 m m )

(c) Strip-casting

Fig.2.2 Schematic diagram of (a) conventional continuous casting and hot-rolling (b) Thin slab-hot rolling

(TSHR) (c) Strip-casting

Recently many steel making industrials have invested in commercial mini mills that employ

thin-slab casting technology, thereby deriving technical and economic benefit from this lower

cost, streamlined alternative for producing many types of hot-rolled sheets. Also, this process can

be rendered environmentally friendly by selecting the electric-arc furnace route of steelmaking

which uses recycled scrap and directly reduced iron as a raw material.

There are a variety of thin slab-casting techniques currently being used in the world. Fig.2.3

shows several examples of this technology.

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(a) Thyssen Krupp Stahl AG, Germany (CSP : Compact Strip Production) [3]

(b) Aceria Compacta, Spain (CSP) [3]

(c) AST-Terni, Italy (CSP) [3]

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(d) Cremona, Italy (ISP) [3]

(e) POSCO #1, South Korea (ISP : In-line Strip Production) [47]

(f) Algoma, Canada (FTSC : Flexible Thin Slab Casting)

(g) CORUS Ijmuiden, Netherlands (DSP : Direct Sheet Plant) [3]

Fig.2.3 Practical examples of thin-slab casting techniques

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2.2 Slab cracking

2.2.1 Features of Cracking

Transverse cracking, crazing and star cracking are linked by an important common feature.

Szekeres [4] suggested that in all these cases the cracks follow the boundaries of exceptionally

large prior-austenite grains. On the surface of the as-cast product, the diameters of these grains

may vary between 1 and 4mm, and in some instances they are even larger, An example is shown

in Fig.2.4 [4]. At some locations the large grains may extend to a depth beyond 6mm. Turkdogan

[5] referred to these abnormally large grains as “blown grains”. The diameter of blown grains

tends to be greater at the base of oscillation marks as shown in Fig.2.5 [4]. Schmidt and Josefsson

[6] also provided evidence of blown grains and found that transverse cracks occurred only in areas

where blown grains are present, Fig.2.6 [6]. They suggested that the large grains may be caused by

secondary recrystallization. Similarly, Maehara et al. [7] maintained that control of the austenite

grain size should be the first priority in preventing surface cracking.

Fig.2.4 Widespread crazing and fine transverse cracks at oscillation marks on the as-cast surface of a line

pipe steel slab (top side). Etched in hot HCL. Magnification not specified. CD : Casting Direction [4]

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Fig.2.5 Crazing around a transverse crack at the base of an oscillation mark on the as-cast top surface of a

0.20% C steel slab. Etched in hot HCL [4].

Fig.2.6 Section through the surface across a system of transverse cracks through the coarse-grained zone.

Prior austenite grain boundaries white (ferrite-decorated). Etched in nital. Magnification not specified [6].

2.2.2 Coarse prior austenite grains

Maehara et al. [7] provided convincing evidence that austenite grain size increases very rapidly

in the temperature range of 1450 to 1350� when specimens are re-melted in-situ and

continuously cooled, Fig.2.7 [7]. Moreover, steel containing 0.16% C had significantly larger

grains than other plain carbon steels under similar heat treatments. Because the surface

temperature of most strands are below 1300� at the point of mold exit, blown grains must

develop while the surface region is still within the mold.

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Molten steel level fluctuations in the mold give rise to deep transverse depressions and

oscillation marks. Wolf [8] suggested that large prior austenite grains form in such depressions

because of locally reduced cooling rates as a result of lack of contact between the solidified shell

and mold wall.

Fig.2.7 Grain growth of austenite during continuous cooling. The specimens were remelted at 1580�,

cooled to a given temperature at a rate of 0.28�s-1, and then quenched into water [7].

Surface cracking mechanism based on the existence of blown grains during casting is

schematically illustrated in Fig.2.8 [4]. Stage� represents newly solidified grains on the mold

wall with very small grain size. The grain diameter at the surface is 500 um or less.

Stage� represents blown austenite grains that have grown to a size several times larger than

those shown in stage�. In this case, the surface temperature is probably above 1350�.

In stage�, liquid copper (or a Cu alloy), if present at the scale-matrix interface, penetrates the

blown grain boundaries and allows microcracks to initiate. However, if liquid copper is not

present, solid state sulfides may precipitate on the blown grain boundaries and be the weakening

agent responsible for microcrack formation.

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In stage�, the temperature is low enough to allow precipitation of nitrides, e.g., AlN, Nb(C,N),

or V(C,N). Subsequently, or simultaneously, the nucleation and growth of proeutectoid ferrite can

weaken the grain boundaries and cause a significant loss of ductility. The size of the virgin

austenite grains has a profound effect on the nature of proeutectoid ferrite that precipitates and

grows along the grain boundaries. With blown grains, the proeutectiod ferrite forms in a

continuous, film-like fashion. When the austenite grain size is small, the ferrite grains are more

equi-axed and discontinuous. It is far easier for a crack to propagate along a continuous ferrite film

as opposed to a discontinuous one. Because AlN precipitates far more rapidly in ferrite than in

austenite, the film can be weakened further by AlN precipitation.

In stage�, the top side of the strand experiences tension during the straightening operation.

Thus, any microcracks that are aligned with decorated blown grain boundaries are easily extended,

or new ones develop. If the strand surface temperature at the straightener is above the temperature

at which proeutectoid ferrite forms, this stage may actually precede stage�.

Fig.2.8 Formation of surface cracks due to blown grain during casting [4]

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2.3 Hot ductility test

2.3.1 Simulation of straightening operation during continuous casting

There are several test methods by which the unbending operation during continuous casting

may be simulated. These techniques include: hot-bend testing, hot-compression testing, torsion

testing and hot-tensile testing. The hot-bend test most closely simulates the unbending operation

but it is difficult to quantify the severity of surface cracks that form during such a test [9,10]. The

hot-compression test is carried out on a flanged sample where the hoop strain corresponding to the

first appearance of a crack is taken as a measure of hot-ductility [11]. Torsion testing produces

large strains and the difficulties in interpreting the fracture appearance after failure makes this test

unsuitable [12]. The most popular test for studying the problems of transverse cracking is the

simple hot-ductility tensile test.

2.3.2 Suitability of the hot tensile test to the problem of transverse cracking

It is clear that laboratory hot-ductility tensile tests do not precisely simulate the straightening

operation in continuous casting, A major disparity is the degree of straining involved. In the

straightening operation, it is at most only 1~2% [9], whereas, in a hot-ductility test, the fracture

strains are in the range 5~100%. Thus, the mechanisms which pertain to tensile tests do not

necessarily coincide with those associated with transverse cracking. Notwithstanding these

objections, the simple hot tensile test has proven to be the most popular test for the study of

transverse cracking on laboratory scale. Generally, tests are carried out in a protective atmosphere

using a servo-hydraulic load frame equipped with either a furnace, or an induction heater. The

GLEEBLE machine has also proven to be popular for the investigation of hot-ductility because of

its ability to melt samples and its versatility in simulating thermal cycles. In this instance, heating

is effected by electrical resistance, which has the advantage that there is no practical limit to the

rate of heating, and temperature gradients can be kept to a minimum [1].

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Samples are usually heated to a temperature above the solution temperature of the microalloy

precipitates, both to dissolve all these particles and to produce a coarse grain size reminiscent of

the continuously cast microstructure before the unbending operation. The sample is cooled at the

rate experienced by the surface of the strand during the continuous casting operation and is

strained at rates between 10-3 and 10-4s-1 [9,13]. These strain rates are specifically chosen to

simulate the strain rates associated with straightening operation. More sophisticated simulations of

the continuous casting operation involve actual melting the samples, either by induction or

electrical resistance, in a quartz tube placed over the mid span region to retain the liquid. In-situ

melting can be combined with the complex cooling patterns that are experienced in the secondary

cooling zone before straightening, although unfortunately, this technique has rarely been used.

Some investigators [14,15] have used total elongation to fracture as a measure of the ductility,

and it can provide useful information regarding the role played by dynamic recrystallisation in

influencing the ductility. The majority of researchers have used reduction in area (RA) at fracture

to provide quantitative information on the fracture strain. Although the RA at the onset of fracture

is probably a better measure of ductility in relation to transverse cracking, the total RA has been

generally used because it is easier to measure. This measurement has the advantage that it is

independent of the fracture geometry of the sample.

2.4 Fracture mechanisms

The general characteristics of ductility as a function of temperature are shown in Fig. 2.9 [1].

Three distinct regions may be

(i) a high ductility, low temperature (HDL) region

(ii) a deep ductility trough, indicating a region of embrittlement

(iii) a high ductility, high temperature (HDH) region.

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It is instructive to briefly discuss the deformation and fracture characteristics of material

subjected to tensile stress in each of these regions.

Fig.2.9 Schematic diagram of a ductility curve defining the three characteristic regions of hot-ductility [1]

2.4.1 Region of embrittlement

The ductility region is invariably associated with intergranular fracture, the fracture facets on

microscale, being either covered with fine dimples or microvoids, or they are smooth. Two distinct

fracture mechanisms may be deduced from this difference in appearance of the fracture surface.

Preferential deformation in regions close to grain boundaries initiates voids at grain boundary

inclusions or precipitates, which leads to intergranular failure via microvoid coalescence and in

this instance fine dimples and microvoids are expected to be seen on a fracture surface. However,

grain boundary sliding in the single phase austenite region followed by wedge cracking would

result in a smooth fracture surface [1].

2.4.1.1 Embrittlement by strain concentration and microvoid coalescence at grain

boundaries

Two distinctly different microstructural features may lead to strain concentrations at austenite

grain boundaries.

▪ thin ferrite films

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▪ precipitate free zones

Plain carbon as well as low-alloyed steels have been shown to be susceptible to intergranular

fracture in the temperature region in which unbending of the continuously cast strand is performed

[1]. For this reason, it is pertinent to further discuss the likely microstructural features that may

lead to a mechanism of intergranular fracture.

Ferrite films Intergranular failure may occur when the austenite to ferrite transformation has

partially occurred and a thin film of ferrite (~5-20� thick) has formed around austenite grains.

Such a situation is depicted in Fig.2.10 [1]. The comparative ease of dynamic recovery in ferrite

translates into a low flow stress compared to austenite, and therefore to strain concentration in the

ferrite film. This strain concentration leads to ductile voiding, generally at MnS inclusions located

at austenite grain boundaries.

Fig.2.10 Schematic diagram showing mechanism for transformation induced intergranular failure [1]

Strain induced ferrite can be formed at temperatures above the Ar3 temperature (the

austenite/ferrite transformation start temperature at a constant cooling rate), and often as high as

the Ae3 temperature (the austenite/ferrite transformation start temperature under equilibrium

condition) when the tensile test is conducted at these temperatures [16,17], Fig.2.11 [17]. Between

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the Ar3 and Ae3 temperatures, the thickness of the ferrite film forming around austenite grains does

not change significantly with temperature. Hence, also in this case a thin ferrite layer will form

around austenite grains. At test temperatures below Ar3 the ferrite film thickens rapidly and the

ductility is fully recovered when approximately 50% ferrite is present before the tensile test is

conducted.

Fig.2.11 Microstructure at room temperature of steel tested at 800�, showing intergranular failure associated

with a thin layer of grain boundary ferrite [17]

Although deformation-induced ferrite can be formed quite readily during straining in the course

of hot tensile testing, there is as yet, no convincing evidence of the presence of strain-induced

ferrite in coarse grained steels at the low strains (~2%) applied during straightening. Essadiqi and

Jonas [18] provided limited evidence that strain induced ferrite can be produced in a fine grained

(~25�) low C, Mo containing steel, deformed to a true strain of 0.016 at a low strain rate of

7.4×10-4s-1. This strain and strain rate are similar to those undergone during the straightening

operation, but no such evidence has been provided for coarse grained steels.

Precipitate free zones on grain boundaries In Nb-containing steels that have been solution

treated prior to cooling to the test temperature, precipitation of Nb(C,N) occur in austenite during

deformation. These carbon-nitrides usually precipitate on austenite grain boundaries and such

precipitation is frequently accompanied by the formation of relatively weak precipitate free (and

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carbon depleted) zones (PFZs) on both sides of the austenite grain boundaries (500nm wide) [19].

Fine matrix precipitation can also take place, leading to significant matrix strengthening. The

microstructural situation is then similar to the presence of soft films of deformation-induced ferrite

on austenite grain boundaries, and microvoid coalescence fractures are frequently observed. In this

case however, void formation takes place at the microalloy precipitates (Nb(C,N) and AlN when

Nb and Al are both present). This fracture process is schematically shown in Fig.2.13 [20]. When

deformation is induced in the single phase austenite region, grain boundary sliding is also likely to

be instrumental in the embrittling process, most probably contributing to crack propagation.

Fig.2.13 Schematic illustration of intergranular microvoid coalescence of Nb-bearing steels. a-c depict

deformation in austenite above the Ar3 temperature, d-f depict deformation in the (austenite + ferrite) region

2.4.1.2 Grain boundary sliding

Grain boundary sliding followed by crack propagation occurs in austenite because it displays

limited dynamic recovery [1] giving rise to high flow stresses and work hardening rates. These

high stresses, in turn, prevent the accommodation, by lattice deformation, of the stresses built up at

triple points or grain boundary particles, and this series of events may lead to intergranular failure

by the nucleation of grain boundary cracks. This rupture mechanism is usually associated with

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creep, the latter occurring at strain rates typically below 10-4s-1. However, fractures characteristic

of failure initiated by grain boundary sliding are frequently found at the strain rate generally used

in hot tensile testing (10-3s-1). Furthermore, Ouchi and Matsumoto [21] have observed grain

boundary sliding at strain rates as high as 10-1s-1 in a 0.054%Nb steel strained in tension at 900�.

Therefore much insight may be gained from a study of the creep literature about the factors that

influence intergranular crack nucleation and growth at high temperatures [22].

Traditionally, intergranular creep defects have been classified as either ‘grain edge’ (r-type, r for

rounded), or ‘grain corner’ (w-type, w for wedge) cavities. Both types of cracks are observed in

samples tensile tested at strain rates in the range 10-3s-1 to 10-4s-1, and both require grain boundary

sliding to nucleate cracks. The models proposed for the formation of w-type cracks are illustrated

in Fig.2.14 [23]. These models are important at high temperatures, where other embrittling

mechanisms, associated with precipitates and the thin ferritic films, are not viable.

Fig.2.14 Schematic models showing the formation of wedge cracks by grain boundary sliding: The arrows

indicate the sliding boundaries and the sense of translation [23]

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2.4.2 High ductility, low temperature region

In the high ductility, low temperature (HDL) region, which coincides with a relatively high

volume fraction of ferrite, the strain is no longer concentrated in a thin ferrite film at austenite

grain boundaries. Furthermore, the difference in strength between austenite and ferrite decreases

with decreasing temperature, thus increasing plastic strain in the austenite and, more importantly,

decreasing the strain in the ferrite [24]. The concentration of strain at the grain boundaries is thus

minimized, and high ductilities are observed. In Ferrite, dynamic recovery, which is a softening

process that operates at all strains, readily takes place [25].

Generally, the ductility is very good when high percentages of ferrite are present in the

microstructure, in the vicinity of 700� [16,17]. At this temperature, recovery in the ferrite takes

place with ease, the subgrain size is large, and the flow stress is low. Thus, ferrite flows readily at

triple points to relieve stress concentrations, therefore discouraging the initiation of cracks.

2.4.3 High ductility, high temperature region

One obvious reason for the improvement in ductility in the high temperature region is the

absence of the thin ferrite films. In this region failure occurs either by grain boundary sliding or

through strain concentration in the PFZ. However, higher temperatures also lead to less

precipitation in the matrix and at the grain boundaries. Finally, increased temperatures lead to

lower flow stresses via increased dynamic recovery so that stress concentrations at the crack

nucleation sites are reduced [1].

At higher temperature such as in the HDH region, grain boundary migration may occur, leading

to increased ductility. Cracks can grow intergranularly during the early stages of deformation, and

become isolated within the grains as a result of grain boundary migration. The original cracks are

then distorted into elongated voids, until final failure occurs by necking between these voids.

One way to achieve a high driving force for grain boundary migration is by dynamic

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recrystallisation. It is therefore not surprising that the HDH region has been observed to coincide

with the onset of dynamic recrystallisation [26,27].

2.4.4 Hot ductility behaviour of steels

2.4.4.1 Plain C-Mn and C-Mn-Al steels

For steel with carbon contents below 0.3% and low in Al and N, fractography of samples which

have failed intergranulary reveals that the coarse grain surfaces (~30� grain diameter) are covered

by cavities [16], Fig.2.15. These cavities are caused by microvoid coalescence within the thin

ferrite films. The increased ductility at lower temperatures in the HDL region therefore

corresponds to an increased volume fraction of ferrite. On the other hand, the increased ductility at

higher temperature in the HDH region is attributed to the absence of ferrite films on austenite

grain boundaries as well as the possibility that grain boundary

Fig.2.15 Intergranular microvoid coalescence type fracture in C-Mn-Al steel, solution treated at 1350�,

tested at 850� at strain rate of 10-3s-1 [16]

migration may isolate small cracks and hence prevent the formation of cracks exceeding the

critical crack length [1].

2.4.4.2 C-Mn-Al steels with high Al and N levels

At the high temperature end of the trough, fractured samples display intergranular failure

(Fig.2.16) with flat facets and no evidence of microvoid coalescence or ferrite formation [16,17].

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Embrittlement is therefore by grain boundary sliding in austenite, accompanied by crack

nucleation at triple points. In case of the high Al-N steels, it is likely that AlN precipitates are

formed on the austenite grain boundaries, pinning the boundaries and allowing the cracks formed

by grain boundary sliding to join up, as well as encouraging void formation [1].

Fig.2.16 (a) C-Mn-Al steel showing flat facets on fracture surface, (b) enlarged view of a showing lack of

voiding around MnS inclusions [16]

Lowering the tensile test temperature to below Ae3 will introduce deformation induced ferrite,

giving rise to embrittlement by microvoid coalescence. Thus, in these steels, the ductility trough is

a result of both grain boundary sliding and microvoid coalescence in the ferrite. A further

reduction in temperature increases the volume fraction of ferrite and the recovery of ductility.

Since the high temperature side of the trough occurs in the single phase austenite region,

ductility recovery in the HDH region is by grain boundary migration, aided by the dissolution of

AlN precipitates.

2.4.4.3. Microalloyed steels

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In microalloyed steels precipitation occurs at austenite grain boundaries, often generating

precipitation free zones and thus leading to embrittlement in the single phase austenite region by a

combination of microvoid coalescence and grain boundary sliding or shear. Fine precipitates may

pin the grain boundaries, allowing the cracks to join up. Thus, intergranular fractures at the high

temperature end of the trough are of mixed character, containing flat facets as well as coalesced

microvoids. Lowering the temperature increases the amount of intergranular microvoid

coalescence, until the fractures are entirely due to microvoid coalescence in deformation-induced

ferrite. The mechanisms of fracture in the HDH and HDL regions are essentially those pertaining

to high Al-N steels, described above [1].

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2.5 Factors influencing hot ductility

2.5.1 Grain Size

The high temperature ductility increases as the grain size is decreased. When the failure is

intergranular, refining the grain size affects crack growth via :

A. The decrease in the crack aspect ratio, which controls stress concentration at the crack

tip [30].

B. The difficulty in propagating smaller cracks formed by sliding through triple points [30].

C. The increase in the specific grain boundary area (for a given volume fraction of

precipitate), which reduces the precipitate density on the grain boundaries [11, 21].

D. The reduction in the critical strain for dynamic recrystallisation by increasing the

number of grain boundary nucleation sites [31].

The influence of grain size on the hot ductility of C-Mn steel having 0.19 and 0.65wt%C is

illustrated in Fig.2.17 [32]. The grain sizes before deformation were varied by heating cylindrical

samples to temperatures of 925-1330� and holding for 15minutes using Instron tensile tester and

GLEEBLE machine. And then the samples were tested to failure at a strain rate of 3×10-3s-1 after

cooled to the test temperature of 550-950�. For the 0.19wt%C steel, in the finer grained steels

with grain size below 180 �, the ductility trough starts close to the Ar3 temperature when films of

ferrite form around the stronger austenite grains. It is difficult to observe deformation induced

ferrite because of the much higher austenite to ferrite transformation rate resulting from the fine

austenite grain size [17]. The narrow trough in a fine grained steel is then a consequence of the

rapid increase in volume fraction of the ferrite which forms when the temperature is lowered to

below the Ae3. In the coarse grained steels, on the other hand, deformation induced ferrite can

have a pronounced influence on hot ductility over a wide range of temperatures leading to a wide

and deep ductility trough. In this case deformation can raise the Ar3 temperature to almost the Ae3

temperature. The volume fraction of deformation induced ferrite is always small in the coarse

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grained steels, often over a wide temperature range from the Ae3 to the Ar3 (undeformed) because

there is insufficient deformation away from the boundary regions to increase the volume fraction

of ferrite significantly.

Fig.2.17 Hot ductility curves for C-Mn steels at various grain sizes [32]

Refining the grain size has an even greater influence on the hot ductility of the 0.65wt%C steel

as shown in Fig.2.17 (b). Intergranular fracture in coarse grained steel occurs by grain boundary

sliding in the austenite resulting in a very wide ductility trough. Raising the carbon level increases

the activation energy for dynamic recrystallization, and hence, more grain boundary sliding than

grain boundary migration is expected to account for deformation. It is probable that intergranular

failure in these steels now only occur below the Ar3 temperature.

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2.5.2 Precipitation

The effect of precipitates on hot ductility depends strongly on their size distribution and

interparticle spacing. These characteristics are controlled by composition and the thermomechanical

history of the cast strand. Fig.2.18 [1] shows the influence of particle size and interparticle spacing of

Nb(C,N) precipitates at the austenite grain boundaries on the hot ductility of C-Mn-Nb-Al steel. The

results of the surface examination of plates commercially produced from continuously cast slabs are

also shown in the figure. Plates that have been rejected because of the presence of surface cracks

contained Nb(C,N) precipitates with mean particle sizes of less than 14nm and interparticle spacings

of less than 60nm. Mintz et al. [13] proposed that the excessive surface cracks observed in slabs that

contained fine precipitates less than 14nm, pins the boundaries and when deformation occurs by

grain boundary sliding in the austenite, cracks are allowed to join up. Microvoid coalescence failures

are also encouraged by an increase in the precipitate or inclusion density at the austenite grain

boundaries, these being preferential sites for void initiation, and hence when the interparticle spacing

is low, typically less than 64nm, excessive cracks can occur.

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Fig.2.18 Influence of (a) particle size and (b) interparticle spacing on hot ductility of Nb-containing steels,

solution treated at 1330�, cooled to a test temperature of 850�, and fractured at a strain rate of 3×10-3s-1 [1]

The precipitation of AlN, Nb(C,N), and V(C,N) in austenite is accelerated by deformation. This

acceleration of precipitation in deformed austenite compared the rate of precipitation in

undeformed austenite at the same temperature results from the favorable nucleation sites, such as

dislocation networks and vacancy clusters, provided by deformation [1].

Both Nb(C,N) and VN can precipitate rapidly during testing at a strain rate <10-1s-1, and hence

the precipitation of these carbides occurring during the test can have a very important influence on

the experimentally observed hot ductility and transverse cracking. For steels which are solution

treated and cooled to the test temperature, Nb extends the ductility trough to higher temperatures

than V because the maximum rate of Nb(C,N) precipitation in undeformed austenite occurs at

950�, whereas it is about 885� for VN. Al steels have even narrower ductility troughs because

the maximum rate of AlN precipitation occurs at 815� [33,34].

Under cooling rates typically experienced in continuous casting, AlN will not precipitate in

austenite either statically or dynamically, because nucleation is inhibited. However AlN can easily

precipitates in ferrite, and thermal cycling through the ferrite/austenite transition encourages such

precipitation. Although the bulk cooling rate of a continuously cast slab is roughly constant,

temperature cycling that occurs at the strand surface can accelerate AlN precipitation. Depending

on the secondary cooling conditions, it is possible for the temperature to fall periodically below

the Ar3, particularly at slab corners. This can be very detrimental to the ductility, because AlN

precipitates mainly at austenite grain boundaries [35].

In steels containing aluminum and vanadium, the precipitation of carbo-nitrides that occur

before deformation (the static precipitate) is more detrimental than those formed during testing,

although they are coarser, because they form only at grain boundaries. By contrast, in Nb-

containing steels, dynamic precipitation occurs more readily and reduces hot ductility significantly

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28

[15,27]. Moreover, if such steels are solution treated before deformation so as to take carbides and

nitrides into solution, very fine Nb(C,N) is formed on the austenite boundaries and within the

matrix. Conversely, if the test is done under direct casting conditions, the amount of Nb available

to form fine deformation-induced Nb(C,N) precipitation will be reduced because coarse niobium

carbo-nitrides precipitate interdendriticly during solidification [36].

Under solution treated testing conditions, Ti additions to steel are most effective in maintaining

hot-ductility and reducing or eliminating the ductility trough because TiN and TiN-rich

precipitates form at high temperatures close to the solidus and tend to be coarse and randomly

distributed having a high enough volume fraction to restrain grain growth at high temperatures

(~1350�) [1]. The benefits of Ti addition therefore result from the refinement of grain size and the

ability to combine preferentially with N, thereby preventing the formation of AlN or Nb(C,N)

precipitates. In industrial continuous casting, however, Ti addition does not significantly influence

the grain size obtained during cooling after solidification [37]. However, the knowledge about the

influence of Ti appears to be incomplete as laboratory investigations do not always agree with

operational experiences.

2.5.3 Composition

2.5.3.1 Sulphur

The effect of sulphur on hot ductility depends largely on the test conditions. For steels solution

treated at 1330�, it is the amount of S which can be redissolved at 1330� and subsequently

reprecipitated as sulphides in a fine form at grain boundaries that is most detrimental to ductility

[1, 38, 39, 40]. Dissolution of sulphides allows S to segregate to grain boundary precipitate/matrix

interfaces as well as austenite grain boundaries. This leads to enhanced nucleation of microvoids

and a loss in ductility. The amount of sulphur that redissolves depends on the manganese content

of the specimen [39]. Using the solubility data of Turkdogan [41], for a steel containing 1.4% Mn,

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the amount of sulphur redissolved at 1330°C is ~0.001%. Accordingly, once the sulphur content

becomes high enough to reach the maximum dissolvable amount at the solution treatment

temperature, the hot ductility behaviour will be independent of the sulphur level.

Under direct casting conditions, it is the total sulphur level that is important for controlling hot

ductility [1, 38, 39]. Fig.2.19 [51] illustrates the effect of sulphur on RA at two cooling rates, 1°Cs-

1 and 4°Cs-1, where the ductility decreased as the sulphur content is increased. Abushosha [42]

found that increasing the sulphur level caused lower ductility for C-Mn-Al steels and C-Mn-Al-Nb

steels under direct casting conditions. Increasing the sulphur level increases the volume fraction of

sulphides formed on solidification, reducing ductility. The reduction in ductility is related to the

increase in the volume fraction of sulphides at the interdendritic boundaries. The interdendritic

boundaries later form the austenite grain boundaries; as a consequence sulphide particles are

located at austenite grain boundaries. These sulphide inclusions enhance intergranular failure in

austenite by preventing either austenite grain boundaries from migrating or encourage voiding

during grain boundary sliding [1, 39]. Sulphide inclusions are also instrumental in accelerating

intergranular failure in thin, deformation induced, ferrite films formed at lower temperatures.

Fig.2.19 Effect of sulphur content on the minimum reduction of area for two cooling rates of 1°Cs-1 and

4°Cs-1 [51]

In direct cast steels that have been reheated, segregation that occured during solidification will

Page 43: The Influence of Austenite Grain Size on Hot Ductility of Steel

30

influence the amount of sulphur that can redissolve at the holding temperature. The regions

surrounding the sulphides are depleted in Mn, so a higher volume fraction of sulphur can be

redissolved on reheating [39]. Hence, if specimens are in the as-cast condition and reheated, they

may have a lower ductility than if the specimen were only reheated to the solution treatment

temperature.

Sulphur can impair hot ductility by forming precipitate free zones (PFZs), allowing strain to

concentrate in the weaker regions (where there are no precipitates) in the vicinity of austenite

boundaries. In specimens exhibiting poor ductility, PFZs have been found responsible for

encouraging intergranular failure. Dense precipitation at the prior austenite grain boundaries and

an extremely fine dispersion of sulphides in the matrix are observed in these steels [1, 38].

It is generally recommended that sulphur levels be kept to a minimum to avoid transverse

cracking. Reducing sulphur levels reduces the volume fraction of sulphides available for

precipitation on solidification. Calcium treatment of liquid steel has been shown to be beneficial

for improving hot ductility by modifying the sulphides. Sulphur is bound in these modified

sulphides and hence, they show little ability to redissolve at 1330°C, thus, reducing the amount of

free sulphur available for precipitation. Calcium additions to liquid steel will also reduce the total

amount of sulphur in the steel by removal of sulphur in the slag (important for direct cast

specimens) and reduce precipitation in this way [1, 42].

Page 44: The Influence of Austenite Grain Size on Hot Ductility of Steel

31

2.5.3.2 Carbon

The carbon content largely determines hot cracking susceptibility of low alloy steels during

continuous casting, as shown in Fig.2.20 [7]. Maximum cracking susceptibility is found in the

peritectic composition region, 0.10 to 0.16%C. Hot-tensile tests on reheated specimens fail to

show this carbon dependence, and the cause of the ductility loss in the medium C steels cannot be

ascribed merely to the uneven solidification on the surface of slabs due to the peritectic reaction.

The carbon content dependence can be ascribed to microstructural changes during solidification. If

intergranular failure in austenite is the dominant failure mode, then hot-ductility of direct-cast

specimens will depend largely on the austenite grain size. This is described in Fig.2.21 [7].

Because austenite forms at high temperatures in the peritectic region, large grains develop due to

rapid grain growth. For this reason the largest austenite grain sizes are generally found in steels

with carbon contents close to the peritectic composition [38].

Fig.2.20 Schematic illustration showing the effect of C on surface cracking of continuous cast slabs [7]

Page 45: The Influence of Austenite Grain Size on Hot Ductility of Steel

32

(a) Relationship between austenite grain size and C content. Specimens (0.35%Si-1.5%Mn-0.05%Nb) were

remelted at 1580�, cooled to 900� at a rate of 5�s-1 and then quenched.

(b) Calculated values RA which were deduced from the relationship between grain size and RA assuming that

the specimens were deformed at 800� at a strain rate of 0.83×10-3s-1

Fig.2.21 Effect of C content on austenite grain size and calculated value of RA [7]

Austenite grain growth is largely retarded by the presence of a small amount of a second phase

and especially by the presence of fine grain boundary precipitates. Thus the grain size will mainly

be determined by the austenitizing temperature, and will be a maximum in steels with the

peritectic composition. In steels of hypo-peritectic composition, the second phase is delta-ferrite.

In steels of hyper-peritectic composition, however, a liquid phase will be present up to much lower

temperatures [7].

Although microalloyed steels exhibit poor ductility in the temperature range 1000-700�, a

ductility trough is present even in plain C-Mn steels. The hot-ductility behaviour observed by

Crowther and Mintz [6] was a strong function of the carbon content. For steels containing more

than 0.3%C, troughs up to ~200� wide were obtained having minimum RA values close to 30%.

Page 46: The Influence of Austenite Grain Size on Hot Ductility of Steel

33

In steels with carbon contents lower than 0.3%, minimum RA values were found but the ductility

troughs were only about 50-100� wide. It has been shown that the width of the trough decreases

with decreasing carbon and manganese contents, Fig.2.22 [39]. The trough in the hot-ductility

curve of low-carbon C-Mn steels is mainly a result of the presence of thin films of deformation-

induced ferrite which form around the austenite grain boundaries. In higher carbon steels, grain

boundary sliding in the austenite can be responsible for the loss in ductility at temperatures above

Ae3.

Fig.2.22 Influence of C and Mn on the width of the ductility trough [39]

Page 47: The Influence of Austenite Grain Size on Hot Ductility of Steel

34

2.5.4 Cooling rate

Generally, increasing the cooling rate in the range 25 to 240�min-1, results in lower ductility

for most types of steel. In most cases, the decrease in ductility with increasing cooling rate is

ascribed to either the formation of finer precipitates or finer inclusions [40].

For C-Mn steels, a finer MnS distribution in the ferrite film surrounding the austenite grains, as

well as a reduction in thickness of the ferrite film, can lead to the deterioration in ductility at

increased cooling rates [40]. In the case of C-Mn-Al steels, the deterioration in ductility is due to

finer AlN precipitation and/or a finer sulphide inclusion distribution [40, 43]. For C-Mn-Al-Nb

steels, an increase in cooling rate can lead to a larger amount of Nb being held in solution,

resulting in an increase in finer, more detrimental, strain-induced Nb(C,N) precipitation [43].

It is pertinent however, to point out that although a reduction of the cooling rate of the strand

may increase the resistance to transverse cracking through coarsening of the precipitates and a

reduction in the thermal stresses in the strand, an increased cooling rate may result in ductility

improvements through grain refinement [1].

2.5.5 Strain rate

With respect to straightening operations where the strain rate varies between ~10-3-10-4s-1, an

increase in the strain rate significantly improves hot ductility. The fracture appearance changes

from intergranular to ductile by an increase in the strain rate [1, 38, 40]. This improvement of hot

ductility with strain rate may be attributed to the following [1, 21, 44]:

. Insufficient time to allow for strain induced precipitation,

. Less grain boundary sliding,

. Insufficient time for the formation of voids near the precipitates or inclusions at grain

boundaries

. Prevention of the formation of deformation-induced ferrite

Page 48: The Influence of Austenite Grain Size on Hot Ductility of Steel

35

In commercial operations changes in casting speed can only increase the strain rate by a factor

of two and hence, it is not possible to change the cracking behavior by a change in casting steed

through its influence on strain rate alone but a change in casting speed can still modify the

temperature distribution and thus have an indirect effect on cracking behavior [1, 45].

2.5.6 Thermal history

The actual thermal cycle at the surface of a continuously cast steel slab is very complex

because of the temperature oscillation introduced by the alternate impingement of water sprays

into the slab and contact with the rolls [35]. Computer simulations [46] have shown that there is

very large temperature drop just below the mold, where the surface temperature reaches 600-700�.

The surface is then reheated by heat transfer from the interior of slab to over 1000�, after which it

is cooled more steadily along the rest of the strand.

This cyclic behavior of the surface temperature has been incorporated into tensile tests to

simulate actual continuous casting. The results showed that, if the temperature drops below the

final test temperature during cycling, increased fine precipitation at the test temperature reduces

ductility. This has been demonstrated for C-Mn-Al and C-Mn-Nb-Al steels [17, 36, 46].

Furthermore, if the temperature falls below the transformation so that ferrite is present,

precipitation is further enhanced due to the lower solubilities of nitrides and carbonitrides in ferrite

than in austenite. This is again expected to be detrimental to hot-ductility.

Page 49: The Influence of Austenite Grain Size on Hot Ductility of Steel

36

CHAPTER 3 - EXPERIMENTAL

3.1 Preparation of specimens

The specimens used in this investigation were prepared in a laboratory scale vacuum induction

melting furnace. Table 3.1 shows the chemical composition of these specimens. In order to

eliminate the overriding effect of precipitation of alloying element at the austenite boundaries

masking any influence of grain size, tramp element and impurities were kept as low a level as

possible. Nevertheless, Al was used as deoxidizer and some ingress of N occurred. Solidified

casting ingots were reheated to 1200� and then hot-rolled into 15mm thick plates using a

laboratory scale rolling mill. Cylindrical GLEEBLE specimens, with 10mm diameter and 115mm

in length were machined from these plates and the details of such specimens are shown in Fig.3.1.

Table 3.1 Chemical composition of specimens

Elements, wt%

Spec

imen

s

C P Mn Si S Ni,Cr,Mo

,Cu,Sn Al Nb Ti, V N

A 0.05 0.002 <0.01 <0.005 0.002 <0.002 0.016 <0.001 <0.003 0.0015

B 0.18 0.002 <0.01 <0.005 0.002 <0.002 0.034 <0.001 <0.003 0.0016

C 0.45 0.002 <0.01 <0.005 0.002 <0.002 0.025 <0.001 <0.003 0.0016

115 mm

10 mm

Thermocouple Type-R(Pt-Pt13%Rh)

Fig.3.1 Geometry of GLEEBLE specimen

3.2 Hot ductility test

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37

3.2.1 GLEEBLE3500

GLEEBLE3500 is a thermomechanical simulator in which the samples are heated by electrical

resistance. Sample cooling can be achieved by several methods such as simple cooling by

anvil/grips or by any combination of air/inert gas/water quenching. Testing can be done in either

ambient or inert atmosphere or alternatively under vacuum. The testing is done fully automatic, or

with a limited manual control. It is possible to deform under uniaxial tensile, uniaxial compressive

or plane strain compressive conditions at the desired temperature. The time – temperature –

deformation program is written on a desk-top computer through the use of table programming.

The table program is converted automatically to an executable script program by the machine

software. Script programs are executed, not table programs. Machine capabilities, as specified by

the manufacturer, are:

. Heating rate up to 3000�s-1.

. Cooling rate up to 600�s-1.

. Ram speed up to 1000mms-1 (which affords a very high strain rate).

. Maximum Load 10 tons.

A schematic diagram of the GLEEBLE testing arrangement is shown in Fig.3.2. In order to

control the temperature of a specimen, a type-R thermocouple (Pt / Pt-13%Rh) is spot-welded

onto the specimen at the middle of the span.

Page 51: The Influence of Austenite Grain Size on Hot Ductility of Steel

38

Fig.3.2 Schematic diagram of the GLEEBLE testing arrangement

3.2.2 Thermomechanical cycles

Fig.3.3 shows the thermomechanical cycles imposed during hot ductility tests. In order to attain

different grain sizes, specimens were solution treated at pre-determined temperatures between

1100� and 1350� for 10minutes in order to allow sufficient grain growth, Fig.3.3 (a). The

specimens were then cooled to the test temperature, in the range 1100� to 700� at a rate of

200�min-1. After holding for 1 minute at the test temperature, the specimens were pulled to

fracture at a low strain rate of 7.5×10-4s-1, simulating the strain rate obtained during straightening

of cast slab. Tensile tests were conducted at temperature intervals of 50� in the temperature range

1100� to 700�. Tests were carried out in a vacuum of approximately 10-3 atm.

In order to simulate direct casting conditions, specimens were melted and then cooled to the

test temperature (Preliminarily experiments conducted to determine the melting point will be

discussed in Section 3.2.5). In order to prevent the specimen from collapsing during melting,

quartz crucibles covering the middle part of GLEEBLE specimen were used to contain the melted

zone of the specimen. Two different cooling rates, 100�min-1 to simulate conventional continuous

casting and 200�min-1 for the simulation of thin-slab continuous casting, were used to study the

effect of cooling rate on hot ductility, Fig.3.3 (b).

Page 52: The Influence of Austenite Grain Size on Hot Ductility of Steel

39

For the purposes of metallographic examination, specifically to determine grain size, samples

were subjected to the same heat treatments and water quenched after cooling to the test

temperature but without applying deformation.

10m in

1m in

7.5¡ ¿10-4s-1

200¡ É/m in

10¡ É/s

Tim e

1100-1350¡ É

700-1050¡ É

(a) Solution treatment

20sec

1m in

7.5¡ ¿10-4s-1

200¡ É/m in

10¡ É/s

M elting

700-1050¡ É

100¡ É/m in

Tim e

(b) Direct casting

Fig.3.3 Schematic diagram of thermomechanical cycles for hot ductility tests under the conditions of (a)

Solution treatment and (b) Direct casting

3.2.3 Measurement of hot ductility

Reduction in area at fracture has been a most popular method for assessing hot-ductility. Fig.3.4

Page 53: The Influence of Austenite Grain Size on Hot Ductility of Steel

40

shows the geometry of sample after fracture. Reduction in area at fracture was calculated as

follows:

RA (%) = (A0�A1) / A0 × 100 (3.1)

Where: A0 = cross-sectional area before test

A1 = cross-sectional area after fracture

A1 was calculated by an average of four measurements.

A1 A0A1 A0

Fig.3.4 Geometry of sample after fracture

3.2.4 Correction of stress-strain curve

The load-curve sensed by the load-cell during a tensile test indicates that there is a residual

force remaining after the specimen has been fractured as shown in Fig.3.5. Both Friction between

the ram and chamber and software resetting errors seemed to cause this effect. It is therefore

necessary to correct the load obtained from GLEEBLE load cell every time. The net tensile force

was obtained by shifting the load curve down (or up) by the amount of the residual load. After

doing this, modified stress-strain curves could be obtained. Fig.3.6 shows an example of this

correction showing that exactly the same result was obtained from two different experiments using

the same steel grade specimens under the same test conditions.

Page 54: The Influence of Austenite Grain Size on Hot Ductility of Steel

41

1000 1200 1400 1600 18000

50

100

150

200

250

300

350

Net tensile force

Mesured tensile forcefrom load cell

Forc

e (k

gf)

Time (sec)

Fig.3.5 Measured tensile force during tensile deformation

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

-10

0

10

20

30

40

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

-10

0

10

20

30

40

2

1

(b)(a)

Strain

2

1

Stre

ss (M

Pa)

Strain

Fig.3.6 Stress-strain curve. (a) curves obtained from GLEEBLE (b) modified curves

Page 55: The Influence of Austenite Grain Size on Hot Ductility of Steel

42

3.2.5 Determination of GLEEBLE setting temperature for melting

There exists a radial thermal gradient in GLEEBLE specimens due to heat loss at the specimen

surface. This heat loss is mainly due to radiative heat loss under vacuum conditions. Heat loss by

convection in vacuum is reported to be about 0.1% of the total heat loss [48]. In the case of

specimen temperatures above 1400�, this temperature difference between the surface and center

of a sample is approximately 100°C [49].

Because only surface temperature of a sample is measured in a GLEEBLE test and this

temperature is used to control the heating current, it was necessary to determine the apparent

melting point of the three specimen types used before a melting experiment could be conducted.

Several preliminary experiments were conducted to determine an accurate estimate of this

apparent melting point. Samples within a quartz crucible in its center span was heated up to

1500� a very slow rate of 0.2�sec-1 under a small compressive force of 8kgf. During heating the

stroke (the relative distance between the two rams) as well as the temperature was recorded. This

experiment was repeated 2-3 times. Fig.3.7 shows the effects measured. Specimen starts to melt

from the center of specimen due to thermal gradient. The stroke increases or may become stable

until some outer part of specimen begin to melt. As melting of outer zone is initiated, the specimen

seems to collapse resulting in a decrease of the stroke (because there is no resisting power against

the applied compressive force). Sometimes the thermocouple detached just before or after the

stroke collapse which resulted in the loss of experimental control. The GLEEBLE setting

temperatures for the melting experiments were set at 1440�, 1440� and 1430� for the Fe-

0.05%C, Fe-0.18%C and Fe-0.45%C alloys respectively.

Page 56: The Influence of Austenite Grain Size on Hot Ductility of Steel

43

200 250 300 350 400 450

0.56

0.58

0.60

0.62

0.64

Time (sec)

Stro

ke (m

m)

1340

1360

1380

1400

1420

1440

1460

1442?

Temperature (?

)

¡ É

¡É

(a) Fe-0.05%C alloy

350 400 450 500 550 600 6500.6

0.7

0.8

0.9

1.0

1.1

Time (sec)

Stro

ke (m

m)

1360

1380

1400

1420

1440

1460

1444?

Temperature (?

)

¡ É

¡É

(b) Fe-0.18%C alloy

350 400 450 500 550 6000.64

0.66

0.68

0.70

0.72

0.74

1435?

Time (sec)

Stro

ke (m

m)

1360

1380

1400

1420

1440

1460

Temperature (?

)

¡ É

¡É

(c) Fe-0.45%C alloy

Fig.3.7 The results of preliminary experiments for the determination of the apparent melting point, (a) Fe-

0.05%C alloy (b) Fe-0.18%C alloy (c) Fe-0.45%C alloy

Page 57: The Influence of Austenite Grain Size on Hot Ductility of Steel

44

3.3 Metallography

In order to establish the relationship between the heat treatment condition and austenite grain

size, specimens were heat-treated within the GLEEBLE. Heat treated specimens were then cut

into appropriate sizes using an Accutom precision cutting machine. Samples were polished up to

1� finishing.

In order to delineate prior austenite grain boundaries, the following etching technique was used.

Ten drops of hydrochloric acid and ten drops of teepol were added to 70ml of saturated picric acid.

The etchant was heated to 67°C and polished samples were immersed for 10-15 minutes in this

solution. In some cases, nital (2.5% nitric acid) was used as an etching agent.

Page 58: The Influence of Austenite Grain Size on Hot Ductility of Steel

45

CHAPTER 4 - RESULTS

4.1 Hot ductility curves

In this chapter, the temperature range in which a ductility trough is found, is defined as the

range in which the RA value is below 40%.

4.1.1 Fe-0.05%C alloys

Fig.4.1 shows the reduction in area (RA) as a function of test temperature for Fe-0.05%C alloys.

For the solution treatment condition, at temperatures below 800�, almost the same RA value is

found. On the other hand, at temperatures above 900� higher solution treatment temperatures

resulted in lower RA values. In the temperature range 830-1020� a ductility trough was found for

specimens solution treated at 1350�. The minimum RA values are 15.7% at 850�, 40.1% at 900�

and 65.6% at 800�, for solution treatment temperatures 1100�, 1200� and 1350� respectively.

For direct cast condition, the RA values at a cooling rate of 200�min-1 were slightly higher than

that for 100�min-1 at all test temperatures.

0%

20%

40%

60%

80%

100%

650 700 750 800 850 900 950 1000 1050 1100

Test temperature (�)

RA

(%)

1100� , solution1200� , solution1350� , solution

(a) solution treatment

Page 59: The Influence of Austenite Grain Size on Hot Ductility of Steel

46

0%

20%

40%

60%

80%

100%

650 700 750 800 850 900 950 1000 1050 1100

Test temperature (�)

RA

(%)

100�/min, melting

200�/min, melting

(b) direct casting condition

Fig.4.1 Hot ductility curves for Fe-0.05%C alloys under (a) solution treatment condition (b) direct casting

condition

When specimens were solution treated at 1100� or 1200� and tested at 850�, it was not

possible to obtain RA values because neck formation did not always coincide with the center of

the specimen in axial direction as shown in Fig.4.2.

Fig.4.2 Sample geometry after fracture tested at 850� under solution treatment condition

4.1.2 Fe-0.18%C alloys

Fig.4.3 shows RA curves as a function of test temperature for Fe-0.18%C. When specimens

were solution treated, the ductility trough was present at 780-860�, 770-880� and 760-960� for

Page 60: The Influence of Austenite Grain Size on Hot Ductility of Steel

47

solution treatment condition temperature 1100�, 1200� and 1350� respectively. At test

temperatures below 800�, the curves show similar RA values for all solution treatment

temperatures. As in the case of Fe-0.05%C alloy, the higher solution treatment temperatures cause

poor ductility when the test temperature is above 850� especially for specimens solution treated

at 1350�. The minimum RA value was 11.0% at 850�, 13.2% at 850� and 32.3% at 800� for

solution treatment temperature 1100�, 1200� and 1350� respectively.

For the direct casting condition, the curves showed the minimum ductility at 800� for both

cooling rates.

0%

20%

40%

60%

80%

100%

650 700 750 800 850 900 950 1000 1050 1100Test temperature (� )

RA

(%)

1100� , solution1200� , solution1350� , solution

(a) solution treatment

Page 61: The Influence of Austenite Grain Size on Hot Ductility of Steel

48

0%

20%

40%

60%

80%

100%

650 700 750 800 850 900 950 1000 1050 1100Test temperature (�)

RA

(%)

100� /min, melting200� /min, melting

(b) direct casting condition

Fig.4.3 Hot ductility curves for Fe-0.18%C alloys under (a) solution treatment condition (b) direct casting

condition

4.1.3 Fe-0.45%C alloys

The measured reduction in area (RA) as a function of temperature for Fe-0.45%C alloys is

shown in Fig.4.4. Unlike the other alloys, the RA value peaked at 700� in both heat treatment

conditions.

In case of solution treatment condition, the ductility trough began at 850�, 890� and 970� for

solution treatment temperature 1100�, 1200� and 1350� respectively. Ductility was not

recovered at temperatures below 650�. At all test temperatures, a higher solution treatment

temperature showed the lower RA value.

For direct cast condition, the RA values at a cooling rate of 200�min-1 were slightly higher than

that for 100�min-1 in all test temperatures.

Page 62: The Influence of Austenite Grain Size on Hot Ductility of Steel

49

0%

20%

40%

60%

80%

100%

600 650 700 750 800 850 900 950 1000 1050

Test temperature (� )

RA

(%)

1100� , solution1200� , solution1350� , solution

(a) solution treatment

0%

20%

40%

60%

80%

100%

600 650 700 750 800 850 900 950 1000 1050

Test temperature (� )

RA

(%)

100� /min, melting200� /min, melting

(b) direct casting condition

Fig.4.4 Hot ductility curves for Fe-0.45%C alloys under (a) solution treatment condition (b) direct casting

condition

Page 63: The Influence of Austenite Grain Size on Hot Ductility of Steel

50

4.2 Stress-strain curves

4.2.1 Fe-0.05%C alloys

The stress-strain tensile test curves for Fe-0.05%C alloys are shown in Fig.4.5 and Fig.4.6

under solution treatment condition and direct casting condition respectively. At lower test

temperatures, once the maximum stress was reached, the stress dropped rapidly to failure. But, at

higher test temperatures, more ductility was retained. Fig.4.7 shows peak stress as a function of

test temperature. By and large, the peak stress decreases with increasing test temperature as

expected, except in the region 800-950� for solution treatment condition and 800-900� for direct

casting condition.

0.0 0.1 0.2 0.3 0.4 0.5 0.60

10

20

30

40

50

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.70

10

20

30

40

50

0.0 0.1 0.2 0.3 0.4 0.5 0.60

10

20

30

40

50

1000¡ É

750¡ É

800¡ É

700¡ É

Stre

ss (M

Pa)

Strain

1050¡ É

1000¡ É

800¡ É

950¡ É

900¡ É750¡ É

700¡ É

Strain

950¡ É

1100¡ É

1050¡ É

1000¡ É

870¡ É

800¡ É

830¡ É

750¡ É

700¡ É

Strain

(a) 1100� (b) 1200� (c) 1350�

Fig.4.5 Stress-strain curves at different test temperatures for Fe-0.05%C alloys solution treated at (a) 1100�

(b) 1200� (c) 1350�

Page 64: The Influence of Austenite Grain Size on Hot Ductility of Steel

51

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.70

10

20

30

40

50

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.70

10

20

30

40

50

1000¡ É

800¡ É

900¡ É

700¡ É

Stre

ss (M

Pa)

Strain

1000¡ É

800¡ É

900¡ É

700¡ É

Strain

(a) 100�min-1 (b) 200�min-1

Fig.4.6 Stress-strain curves at different test temperatures for Fe-0.05%C alloys under direct casting condition

at cooling rate (a) 100�min-1 (b) 200�min-1

0

20

40

60

650 700 750 800 850 900 950 1000 1050 1100 1150

Test temperature(¡ É)

Peak

stre

ss (M

Pa)

1100¡ É, solution1200¡ É, solution1350¡ É, solution

Ae3

0

20

40

60

650 700 750 800 850 900 950 1000 1050 1100 1150

Test temperature(�)

Peak

stre

ss (M

Pa)

100� /min, melting200� /min, melting

(a) solution treatment condition (b) direct casting condition

Fig.4.7. Peak stress as a function of test temperature for Fe-0.05%C alloys under (a) solution treatment

condition (b) direct casting condition

Page 65: The Influence of Austenite Grain Size on Hot Ductility of Steel

52

4.2.2 Fe-0.18%C alloys

The stress-strain curves for tensile tests of Fe-0.18%C alloys are shown in Fig.4.8 and Fig.4.9

under solution treatment conditions and direct casting conditions respectively. Fig.4.10 shows

peak stress as a function of test temperature. All these curves show the same trend as Fe-0.05%C

alloys. The flattened region in peak stress value are at 800-900� for solution treatment condition

and 700-800� for direct casting condition.

0.0 0.1 0.2 0.3 0.4 0.5 0.60

10

20

30

40

50

60

0.0 0.1 0.2 0.3 0.4 0.5 0.60

10

20

30

40

50

60

0.0 0.1 0.2 0.3 0.4 0.5 0.60

10

20

30

40

50

60

1000¡ É

950¡ É

900¡ É

850¡ É800¡ É

750¡ É

700¡ É

Stre

ss (M

Pa)

Strain

1050¡ É

1000¡ É

950¡ É 900¡ É

850¡ É800¡ É

750¡ É

700¡ É

Strain

850¡ É

800¡ É

950¡ É

1000¡ É

900¡ É

750¡ É

700¡ É

Strain (a) 1100� (b) 1200� (c) 1350�

Fig.4.8 Stress-strain curves at different test temperatures for Fe-0.18%C alloys solution treated at

temperature (a) 1100� (b) 1200� (c) 1350�

0.0 0.1 0.2 0.3 0.40

10

20

30

40

50

60

0.0 0.1 0.2 0.3 0.40

10

20

30

40

50

60700¡ É

900¡ É

800¡ É

Stre

ss (M

Pa)

Strain

1000¡ É900¡ É

800¡ É

700¡ É

Strain (a) 100�min-1 (b) 200�min-1

Fig.4.9 Stress-strain curves at different test temperatures for Fe-0.18%C alloys under direct casting condition

at cooling rate (a) 100�min-1 (b) 200�min-1

Page 66: The Influence of Austenite Grain Size on Hot Ductility of Steel

53

0

20

40

60

80

650 700 750 800 850 900 950 1000 1050 1100

Test temperature(¡ É)

Peak

stre

ss (M

Pa)

1100¡ É, solution1200¡ É, solution1350¡ É, solution

Ae3

0

20

40

60

80

650 700 750 800 850 900 950 1000 1050 1100Test temperature(� )

Peak

stre

ss (M

Pa)

100� /min, melting200� /min, melting

(a) solution treatment condition (b) direct casting condition

Fig.4.10. Peak stress as a function of test temperature for Fe-0.18%C alloys under (a) solution treatment

condition (b) direct casting condition

4.2.3 Fe-0.45%C alloys

The stress-strain curves of tensile tests of Fe-0.45%C alloys are shown in Fig.4.11 and Fig.4.12

under solution treatment condition and direct casting condition respectively. Fig.4.13 shows peak

stress as a function of test temperature. As the carbon content increases, peak stress values

increase. No significant flattened regions in peak stress were observed, unlike the other alloys.

0.0 0.1 0.2 0.3 0.4 0.5 0.60

10

20

30

40

50

60

70

80

90

0.0 0.1 0.2 0.3 0.4 0.5 0.60

10

20

30

40

50

60

70

80

90

0.0 0.1 0.2 0.3 0.4 0.5 0.60

10

20

30

40

50

60

70

80

90

Stre

ss (M

Pa)

Strain

900¡ É

800¡ É

700¡ É

900¡ É

850¡ É

800¡ É

750¡ É

700¡ É

Strain

850¡ É

800¡ É

900¡ É

750¡ É

700¡ É

Strain

(a) 1100� (b) 1200� (c) 1350�

Fig.4.11 Stress-strain curves at different test temperatures for Fe-0.45%C alloys solution treated at

temperature (a) 1100� (b) 1200� (c) 1350�

Page 67: The Influence of Austenite Grain Size on Hot Ductility of Steel

54

0.0 0.1 0.2 0.3 0.4 0.50

10

20

30

40

50

60

70

80

90

100

0.0 0.1 0.2 0.3 0.4 0.50

10

20

30

40

50

60

70

80

90

100

1000¡ É

900¡ É

850¡ É

800¡ É

750¡ É

700¡ É

Stre

ss (M

Pa)

Strain

800¡ É

900¡ É

850¡ É

750¡ É

700¡ É

Strain

(a) 100�min-1 (b) 200�min-1

Fig.4.12 Stress-strain curves at different test temperatures for Fe-0.45%C alloys under direct casting

condition at cooling rate (a) 100�min-1 (b) 200�min-1

0

20

40

60

80

100

650 700 750 800 850 900 950 1000 1050

Test temperature(¡ É)

Pea

k st

ress

(MP

a)

1100¡ É, solution1200¡ É, solution1350¡ É, solution

Ae3

0

20

40

60

80

100

120

650 700 750 800 850 900 950 1000 1050

Test temperature(� )

Pea

k st

ress

(MP

a)

100� /min, melting

200� /min, melting

(a) solution treatment condition (b) direct casting condition

Fig.4.13 Peak stress as a function of test temperature for Fe-0.45%C alloys under (a) solution treatment

condition (b) direct casting condition

Page 68: The Influence of Austenite Grain Size on Hot Ductility of Steel

55

4.3 Austenite grain size

In order to determine the grain size, samples were subjected to the same heat treatments

without applying deformation. Specimens were then water quenched following cooling into the

two-phase region so that ferrite could form on the austenite grain boundaries. Specimens were

alternatively etched in a picric acid based solution or nital. Fig.4.14, Fig.4.15 and Fig.4.16 show

the cross-section of GLEEBLE heat-treated specimen for Fe-0.05%C, Fe-0.18%C and Fe-0.45%C

alloys respectively. The diameter of each cross-sectional specimen is 10mm.

In determining the mean grain size, generally known measuring method such as linear intercept

technique could not be used because the grain sizes were too large. Individual grain areas were

measured using software UTHSCS IMAGE TOOL and the grain size was calculated using the

equivalent diameter to account for different morphologies.

The distributions of austenite grain size are shown in Fig.4.17, Fig.4.18 and Fig.4.19. In these

graphs each measured grain sizes is plotted as a function of the distance between the specimen

center and the grain center. For specimens having large grains, in order to get more data two or

three cross-sections of specimen were used in obtaining the graph.

Due to the radial thermal gradient present in GLEEBLE specimens, there is grain size

differences between the center and surface regions as shown in these graphs. This difference

becomes larger as the solution treatment temperature is increased. However, in the case of direct

casting, this difference is smaller than in the case where specimens were solution treated. For the

direct casting condition, when the specimen is melted, the radial thermal gradient seems to

decrease, resulting in a more homogeneous grain size.

Page 69: The Influence of Austenite Grain Size on Hot Ductility of Steel

56

Fig.4.14 Cross-section of a GLEEBLE specimen of a Fe-0.05%C alloy. Etched in nital. For the sake of

clarity the position of grain boundaries were photo-enhanced in the photographs ‘Direct casting’ (diameter

of specimen = 10mm)

1100�

1200� Solution

treatment

1350�

Direct

casting

100�min-1

200�min-1

Page 70: The Influence of Austenite Grain Size on Hot Ductility of Steel

57

Fig.4.15 Cross-section of a GLEEBLE specimen of a Fe-0.18%C alloy. Etched in saturated picric acid based

etchant. For the sake of clarity the position of grain boundaries were photo-enhanced in the photographs

‘200�min-1-Direct casting’ (diameter of specimen = 10mm)

1100�

1200� Solution

treatment

1280�

1350�

Direct

casting

100�min-1

200�min-1

Page 71: The Influence of Austenite Grain Size on Hot Ductility of Steel

58

Fig.4.16 Cross-section of a GLEEBLE specimen of a Fe-0.45%C alloy. Etched in saturated picric acid based

etchant (diameter of specimen = 10mm)

1100�

1200� Solution

treatment

1350�

Direct

casting

100�min-1

200�min-1

Page 72: The Influence of Austenite Grain Size on Hot Ductility of Steel

59

0

2

4

6

0 1 2 3 4 5

Distance from center (mm)

Gra

in s

ize

(mm

)

1100�, solution1200�, solution1350�, solution

(a) solution treatment

0

2

4

6

0 1 2 3 4 5

Distance from center (mm)

Gra

in s

ize

(mm

)

100�/min, melting

200�/min, melting

(b) direct casting

Fig.4.17 The distribution of austenite grain size of a Fe-0.05%C alloy under (a) solution treatment condition

(b) direct casting condition

Page 73: The Influence of Austenite Grain Size on Hot Ductility of Steel

60

0

2

4

6

0 1 2 3 4 5

Distance from center (mm)

Gra

in s

ize

(mm

)

1100�, solution1200�, solution1280�, solution1350�, solution

(a) solution treatment

0

2

4

6

0 1 2 3 4 5

Distance from center (mm)

Gra

in s

ize

(mm

)

100�/min, melting

200�/min, melting

(b) direct casting

Fig.4.18 The distribution of austenite grain size of a Fe-0.18%C alloy under (a) solution treatment condition

(b) direct casting condition

Page 74: The Influence of Austenite Grain Size on Hot Ductility of Steel

61

0

2

4

6

0 1 2 3 4 5

Distance from center (mm)

Gra

in s

ize

(mm

)

1100�, solution1200�, solution1350�, solution

(a) solution treatment

0

2

4

6

0 1 2 3 4 5

Distance from center (mm)

Gra

in s

ize

(mm

)

100�/min, melting

200�/min, melting

(b) direct casting

Fig.4.19 The distribution of austenite grain size of a Fe-0.45%C alloy under (a) solution treatment condition

(b) direct casting condition

Page 75: The Influence of Austenite Grain Size on Hot Ductility of Steel

62

CHAPTER 5 - DISCUSSION

5.1 Grain growth

In order to study the influence of grain size on hot ductility, it is necessary to determine the

representative grain size resulting from the thermal history of each specimens used in this study.

The cross-sections shown in the previous chapter reveal a single plane, possibly masking the

cross-section that actually contains the largest grain. For this reason, and also because of the grain

size difference between center and surface, it would be incorrect to determine the average grain

size by merely averaging the grain sizes in the plane of the polished surface. In an attempt to

overcome this problem, only selected grains which occupied 40% of the cross-sectional area of the

specimen were used in determining the ‘average’ grain size. Although this is an arbitrary choice, it

gives a better measure of the effective or true grain size.

Fig.5.1 shows grain size as a function of solution treatment temperature. In this graph, the

maximum grain sizes as well as the mean size of the selected grains are included. The grain size

increases almost linearly with increasing solution treatment temperature. It is surprising that in the

case of a solution treatment at a temperature of 1350� the average grain size reached 4mm in

diameter in all steel grades. In the Fe-0.05%C alloy, a maximum grain size of almost 6mm was

obtained. Other researches [7, 32] using Nb containing steel and plain C-Mn steel under similar

thermal conditions obtained grain sizes of 0.4-1.5mm in diameter. This observation implies that

alloying elements or alloying compounds pin the grain boundaries and prevent grain boundary

migration at these high solution treatment temperatures.

Carbon content seems to have little effect on grain size at any given solution treatment

temperature. On the other hand, when specimens are melted in-situ, there is a significant influence

of carbon on grain size as shown in Fig.5.2. In the in-situ melted specimens the largest grains were

found in the Fe-0.18%C alloy.

Inspection of the Fe-C phase diagram reveals that in the steel of peritectic composition, Fe-0.18%C,

Page 76: The Influence of Austenite Grain Size on Hot Ductility of Steel

63

austenite forms on cooling at a much higher temperature than either the Fe-0.05%C alloy or the Fe-

0.45%C alloy (almost 100� higher). Since grain growth is a sharp function of temperature, austenite

grains will grow much larger in any given time in a steel of peritectic composition. Therefore the grain

size in the steel of peritectic composition is expected to be largest at any given cooling rate. That this

argument holds, is bourne out by the grain sizes obtained at different cooling rates. At a cooling rate of

100�min-1, the specimen spends more time in the high temperature region than at a cooling rate of

200�min-1 and hence the grain size is larger. The grain size of the Fe-0.05%C alloy is also relatively

large because delta-ferrite grains grow very large in the single delta-ferrite phase on cooling following

solidification. The austenite grain size is determined, in large measure by the delta-ferrite size and

hence, large austenite grains form [50].

0

2

4

6

1050 1100 1150 1200 1250 1300 1350 1400

Solution treatment temperature (¡ É)

Gra

in s

ize

(mm

)

0.05%C 0.05%C0.18%C 0.18%C0.45%C 0.45%C

Average(selected) M axim um

Fig.5.1 Grain size as a function of solution treatment temperature

Page 77: The Influence of Austenite Grain Size on Hot Ductility of Steel

64

0

1

2

3

4

5

100¡ É/min 200¡ É/min

Gra

in s

ize

(mm

)0.05%C 0.05%C0.18%C 0.18%C0.45%C 0.45%C

Average (selected) Maximum

Fig.5.2 Grain size as a function of cooling rate under direct casting condition

5.2 Ductility troughs

Fig.5.3 and Fig.5.4 shows the hot-ductility curves for solution treated specimens having various

grain sizes for Fe-0.05%C and Fe-0.18%C alloys respectively. The Ae1 temperature (the eutectoid

transformation temperature under equilibrium condition) and Ae3 temperature (the austenite/ferrite

transformation start temperature under equilibrium condition) which are calculated using software

MTDATA are also shown on the graph. As shown in these figures, as the grain sizes increase, the

ductility trough extends towards higher temperatures and even beyond the Ae3 temperature. The

reason for the existence of a low ductility region just below the Ae3 temperature may be explained

as follows. According to Crowther [16] and Cardoso [17], deformation induced ferrite can be

formed at temperatures above the Ar3 temperature (the austenite/ferrite transformation start

temperature at a constant cooling rate), and often as high as the Ae3 temperature when the tensile

test is conducted at these temperatures. The comparative ease of dynamic recovery in ferrite

translates into a low flow stress compared to

Page 78: The Influence of Austenite Grain Size on Hot Ductility of Steel

65

0%

20%

40%

60%

80%

100%

650 700 750 800 850 900 950 1000 1050 1100

Test temperature (¡ É)

RA

(%)

0.7mm

1.6mm4.3mm

Ae1

Ae3

Fig.5.3 Hot-ductility curves for Fe-0.05%C alloys for solution treated specimens having various grain sizes

0%

20%

40%

60%

80%

100%

650 700 750 800 850 900 950 1000 1050 1100

Test temperature (¡ É)

RA

(%)

0.4mm1.4mm3.8mm

Ae1

Ae3

Fig.5.4 Hot-ductility curves for Fe-0.18%C alloys for solution treated specimens having various grain sizes

austenite, and therefore to strain concentration in the ferrite film. This strain concentration leads to

ductile voiding, generally at void nucleation sites located at austenite grain boundaries and this

results in a loss of ductility. Coarsening the grain size causes the temperature of the start of the

ductility trough to increase up to the Ae3 temperature. However, there is no evidence in the

literature that deformation induced ferrite can form beyond the Ae3 temperature for the steels

having a carbon content of less than 0.3%. The extension of ductility trough beyond the Ae3

temperature as shown in Fig.5.3 and Fig.5.4, could possibly be ascribed to the occurrence of grain

Page 79: The Influence of Austenite Grain Size on Hot Ductility of Steel

66

boundary sliding, although evidence in the literature suggests this mechanism of failure is only

favored in higher carbon steel containing more than 0.3%C [16]. In coarse grained steels, it is

difficult for dynamic recrystallisation to occur due to the decrease in the number of grain boundary

nucleation sites where dynamic recrystallisation initiates, thereby decreasing the possibility of

ductility improvement via grain boundary migration [1, 45].

On the other hand, AlN precipitates can play a role in widening the ductility trough into the

single phase austenite region by encouraging grain boundary sliding. It is likely that AlN

precipitates are formed on the austenite grain boundaries, pinning the boundaries and allowing the

cracks formed by grain boundary sliding to join up, as well as encouraging void formation. It is

recognized that the precipitation of AlN on grain boundaries can have a significant influence on

hot-ductility. However, in this study, it was not possible to study this aspect in detail. It is therefore

recommended that the effect of AlN precipitation in these alloys be distinguished from the effect

of grain size on hot-ductility by fractography and transmission electron microscopy.

It is interesting to compare the position of the ductility trough relative to temperature to the

temperature dependence of the peak stress obtained in high-temperature tensile tests. The peak

stress vs. temperature curves for the Fe-0.05%C and Fe-0.18%C alloys are shown in Fig.4.7 and

4.10 respectively. On lowering the temperature of the tensile test, the strength of austenite

increases down to a temperature of about 950� in the case of the Fe-0.05%C alloy and 900� in

the case of the Fe-0.18%C alloy. At test temperatures between 950� and 750�, the peak stress is

lower than what it would have been for pure austenite and higher than for pure ferrite, evidently

because the test temperature falls within the two-phase (ferrite + austenite) region and the mixture

of ferrite and austenite is softer than austenite. However, this softening occurs at a temperature at

least 50� above the Ae3 temperature. The formation of ferrite between the Ar3 and Ae3

temperature during deformation is a well known phenomenon and such ferrite is usually referred

to ‘deformation induced ferrite’.

Page 80: The Influence of Austenite Grain Size on Hot Ductility of Steel

67

It is more difficult to explain the softening behavior above the Ae3 temperature. One possibility

is that the temperature, measured in the surface of the tensile specimen, is not fully representative

of the bulk temperature of the specimen in the region where deformation occurs. Hence that the

temperatures boundaring the two-phase region are displaced. There is clearly a radial temperature

distribution in a hot-tensile test specimen in the GLEEBLE arrangement. However, experimental

evidence suggests that the temperature measured on the surface of the specimen in the area of neck

formation is actually lower than the temperature in the center. This temperature difference also

becomes higher at higher test temperature. It therefore seems improbable that the extended

‘softening temperature region’ is merely due to a temperature displacement of the two-phase region.

In the Fe-0.45%C alloy, the two phase region is much narrower and little evidence is found of

the influence of the two-phase region on the peak stress except in the specimen containing very

large grains as shown in Fig.4.13.

Fig.5.3 shows a very strong dependence on grain size of ductility as well as the temperature

range of the ductility trough for the Fe-0.05%C alloy. The ductility trough becomes deeper and

wider with an increase in grain size and the reduction in ductility in the larger-grained specimens

at temperatures above the Ae3 temperature can clearly not be attributed to the formation of ferrite.

In specimens with a grain size of 4.3mm, a significant reduction in ductility occurs on lowering

the test temperature from 1100� to 1000�. For this temperature range the structure is clearly

austenitic and the reduction in ductility must be attributed mainly to a grain size effect. The

dependence of ductility on temperature of the Fe-0.18%C and Fe-0.45%C alloys show similar

behavior but the ductility troughs become wider with an increase in carbon content. These aspects

are analysed more quantitatively below.

Fig.5.5 shows hot ductility curves for the Fe-0.45%C alloy with various grain sizes. The low

ductility above Ae3 and between Ae3 and Ar3 seems to be due to grain boundary sliding and

deformation induced ferrite respectively as in the case of the other alloys. But at 700�, there is a

Page 81: The Influence of Austenite Grain Size on Hot Ductility of Steel

68

slight increase in ductility. The drop in ductility below a temperature of 700� probably arises from

the fact that a relatively large amount of pearlite forms in this alloy compared to the other alloys at

temperatures below Ar1 (the eutectoid transformation temperature at a constant cooling rate). Until

the Ar1 temperature is reached, ductility will increase due to an increase in the volume fraction of

ferrite. Once the Ar1 temperature is reached, pearlite can form and the strength of the matrix is

increased. In addition, the presence of a thin film of ferrite at the grain boundaries, leads to even

more strain concentration at the grain boundaries and the ductility decreases even further [40].

0%

20%

40%

60%

80%

100%

600 650 700 750 800 850 900 950 1000 1050Test temperature (¡ É)

RA

(%)

0.7mm2.6mm4.1mm

Ae3Ae1

Fig.5.5 Hot-ductility curves for Fe-0.45%C alloys for solution treated specimens having various grain sizes

(The Ar1 and Ar3 temperature are not shown in the figure)

5.3 Influence of grain size on hot ductility

5.3.1 Reduction in area

The relationship between tensile properties and the reciprocal of austenite grain size (D) at

several test temperatures is shown in Fig.5.6 and Fig 5.7 for the solution treatment conditions and

direct casting conditions respectively. The ductility generally decreases with increasing grain size

under all testing conditions notwithstanding the fact that the matrix strength was essentially

independent of the grain size at any given test temperature.

These results are in good agreement with previous studies in which Nb-containing steels [7] and

plain C-Mn steels [32] were used. However in these studies, the steels contained relatively small grains

Page 82: The Influence of Austenite Grain Size on Hot Ductility of Steel

69

(smaller than 1.5mm in diameter).

In the case of direct casting, the change in grain size results from the difference of cooling rate.

At the higher cooling rate there is less time available for grain growth in the high austenite

temperature region. Hence, the grain size decreases with an increase in cooling rate and a

concomitant ductility improvement results as shown in Fig.5.7.

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.40

20

40

60

80

1000

10

20

30

40

50

60

RA

(%)

1/D (mm-1)

900¡ É

1000¡ É

900¡ É

1000¡ É

Pea

k S

tress

(MP

a)

0.0 0.5 1.0 1.5 2.0 2.5 3.

1/D (mm-1)

900¡ É

900¡ É

1000¡ É

1000¡ É

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6

800¡ É

800¡ É

900¡ É

900¡ É

1/D (mm-1) (a) Fe-0.05%C (b) Fe-0.18%C (c) Fe-0.45%C

Fig.5.6 Relationship between tensile properties and reciprocal of austenite grain size (D) at different test

temperatures in solution treated specimen (a) Fe-0.05%C (b) Fe-0.18%C (c) Fe-0.45%c alloy

Page 83: The Influence of Austenite Grain Size on Hot Ductility of Steel

70

0.30 0.35 0.40 0.40

20

40

60

80

1000

10

20

30

40

50

60

70

RA

(%)

1/D (mm-1)

1000¡ É

1000¡ É

900¡ É

900¡ ÉPe

ak S

tress

(MPa

)

0.25 0.26 0.27 0.28 0.29 0.300

0

0

0

0

00

0

0

0

0

0

0

0

1/D (mm-1)

900¡ É

900¡ É

1000¡ É

1000¡ É

0.35 0.40 0.45 0.50 0.55

850¡ É

850¡ É

800¡ É

800¡ É

1/D (mm-1) (a) Fe-0.05%C (b) Fe-0.18%C (c) Fe-0.45%C

Fig.5.7 Relationship between tensile properties and reciprocal of austenite grain size (D) at different test

temperatures under direct casting condition for (a) Fe-0.05%C (b) Fe-0.18%C (c) Fe-0.45%C alloy

Fig.5.8 shows the effect of austenite grain size on the minimum RA value for the solution

treatment condition. It is clear that with increasing grain size minimum RA value decrease for all

steel grades. Above all, this is obvious in case of the Fe-0.05%C alloys, showing a much higher

RA value at the same grain size. It is generally known that the lower carbon contents exhibit the

more ductile behavior.

0.0 0.5 1.0 1.5 2.0 2.50

10

20

30

40

50

60

70

80

1/D (mm-1)

0.05%C 0.18%C 0.45%C

Min

imum

RA

(%)

Fig.5.8 Relationship between minimum RA value and reciprocal of austenite grain size (D) for different Fe-

Page 84: The Influence of Austenite Grain Size on Hot Ductility of Steel

71

C alloys in the solution treatment condition

5.3.2 Position and width of ductility trough

As a measure of the position of a ductility trough, the temperature at minimum RA value was

taken except in the case of very wide ductility troughs where the temperature was chosen to be at

the center of the trough. This information is summarized in Fig.5.9 as a function of grain size.

Also Included is the width of trough taken at 40% of RA value. With increasing grain size, the

position of the trough as well as the width of the trough are increased.

0 1 2 3 40

50

100

150

200

250

300

800

850

900

950

0.05%C 0.18%C 0.45%C

Width of trough (¡

É)

Grain size (mm)

0.05%C 0.18%C 0.45%C

Position of ductility trough (¡

É) (a)

(b)

Fig.5.9 (a) Position of ductility trough, (b) width of trough as a function of grain size under solution

treatment condition

Page 85: The Influence of Austenite Grain Size on Hot Ductility of Steel

72

5.4 Influence of carbon content on hot ductility

Fig.5.10 shows hot ductility curves for the alloys of different carbon contents at the same

solution treatment temperature. It is clear that at any given solution treatment temperature the

position and shape of the ductility curve vary with carbon content. As shown in Fig.5.8 and

Fig.5.9 (a), the carbon content affected the position of trough as well as the depth of trough under

solution treatment condition.

0%

20%

40%

60%

80%

100%

600 650 700 750 800 850 900 950 1000 1050 1100

Test temperature (�)

RA

(%)

0.05%C0.18%C0.45%C

`

(a) 1100�

0%

20%

40%

60%

80%

100%

600 650 700 750 800 850 900 950 1000 1050 1100

Test temperature (� )

RA

(%)

0.05%C0.18%C0.45%C

`

(b) 1200�

Page 86: The Influence of Austenite Grain Size on Hot Ductility of Steel

73

0%

20%

40%

60%

80%

100%

600 650 700 750 800 850 900 950 1000 1050 1100Test temperature (� )

RA

(%)

0.05%C0.18%C0.45%C

`

(c) 1350�

Fig.5.10 Hot ductility curves from solution treatments at (a) 1100� (b) 1200� (c) 1350�

In order to investigate the influence of carbon content on hot ductility under direct casting

conditions, the austenite grains size and the variation of the corresponding RA value were plotted

against % C as shown in Fig.5.11. The Fe-0.18%C alloy exhibits the largest grain size at both

cooling rates and the large grains in this Fe-C alloy results from the higher austenitizing

temperature compared to that of other Fe-C alloys as discussed earlier. The loss of ductility seems

to be directly related to an increase in grain size and the RA value for Fe-0.18%C alloy has the

lowest value. Although the Fe-0.45%C alloy had the smallest grain size, it was still very brittle.

This observation may be explained

Page 87: The Influence of Austenite Grain Size on Hot Ductility of Steel

74

0.0 0.1 0.2 0.3 0.4 0.50

20

40

60

80

1000

1

2

3

4

800¡ É 900¡ É 1000¡ É

RA

(%)

C (%)

Gra

in s

ize

(mm

)

0.0 0.1 0.2 0.3 0.4 0.50

20

40

60

80

1000

1

2

3

4

800¡ É 900¡ É 1000¡ É

RA

(%)

C (%)

Gra

in s

ize

(mm

)

(a) 100�min-1 (b) 200�min-1

Fig.5.11 Effect of C content on austenite grain size and RA value under direct casting condition for cooling

rate (a) 100�min-1 (b) 200�min-1

by grain boundary sliding becoming dominant as the carbon content increases. Crowther and

Mintz [16] showed that increasing the carbon content to above 0.3% in a coarse grained steel

(~300�) causes intergranular failure to occur by grain boundary sliding in the austenite, resulting

in a very wide ductility trough. Increasing the carbon content was found to increase the activation

energy for dynamic recrystallization, and hence to encourage more grain boundary sliding and

linkage of cracks.

Inspection of the cross-section of heat treated samples in the direct-casting condition shown in

Fig.4.14, Fig.4.15 and Fig.4.16 reveals interesting information. Columnar austenite grains were

formed in the Fe-0.18%C alloy compared to the equi-axed grains observed in the lower or higher

C alloys. The formation of columnar grains in the Fe-C alloy close to the peritectic composition

may be related to the larger temperature gradient during cooling of specimens of this composition

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75

or to the occurrence of the peritectic reaction followed by the peritectic transformation to austenite

[7]. This observation has great practical significance because the susceptibility of slabs to surface

cracking can be significantly accelerated since the effective grain size affecting intergranular

fracture can be taken as the length of the columnar grains rather than the average grain size [7].

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76

5.5 Influence of cooling rate following direct-casting on hot ductility

As discussed in the literature review an increase in the cooling rate generally results in reduced

ductility in most types of steel. In most cases, the decrease in ductility with increasing cooling rate

is ascribed to either the formation of finer precipitates or finer interdendritic inclusions [40]. The

decrease in ductility is the result of the overriding effect of precipitation of alloy carbo-nitrides on

the austenite grain boundaries, thereby masking any influence of grain size on ductility. Because

plain carbon steels in which tramp elements and impurities were kept very low were used in this

study the effect of alloying element precipitation was minimized. When a specimen is cooled from

the molten state, the lower cooling rate allows time for grains to grow in the high austenite

temperature region. Consequently, an increased cooling rate may result in ductility improvements

through grain refinement. For this reason, in terms of grain size refinement, a higher cooling rate

may be an advantage in near-net shape casting such as thin-slab casting or strip-casting processes.

As shown in Fig.5.2, with an increase in cooling rate from 100�min-1 to 200�min-1, the grain

size decreased and accordingly, ductility is improved in all steel grades and at most test

temperatures as shown in Fig.4.1 (b), Fig.4.3 (b) and Fig.4.4 (b).

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77

5.6 Comparison between as-cast condition and solution treatment condition

It is interesting to note that the ductility of the as cast specimens was sometimes better than that

of solution treated specimens as shown in Fig.5.12. It is well known that an as-cast structure is

inferior to reheated structures with respect to high temperature mechanical properties due to the

differences in microstructure. Generally, the austenite grain structure of reheated slabs has a

smaller grain size than that of the as-cast structure. However, it follows from the data displayed in

Fig.5.12 that the coarse grained structure of the solution treated steel exhibits much poorer hot-

ductility than the as-cast structure in the temperature range 820� to 1000�.

0%

20%

40%

60%

80%

650 700 750 800 850 900 950 1000 1050

Test temperature (� )

RA

(%)

1350� , solution200� /min, melting

Fig.5.12 Hot ductility curves for Fe-0.18%C alloy from two different thermal conditions

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78

5.7 Practical implications of the experimental findings

This study was undertaken to probe the influence of grain size on hot-ductility primarily

because evidence in the literature implicated reduced hot-ductility in continuously cast slabs to the

presence of large grains occurring at the roots of oscillation marks [4, 6, 7]. Convincing

experimental result was found that an increased austenite grain size results in a deepening as well

as a widening of the ductility trough observed in the temperature range 700� to 1100�. The

detrimental effect of grain size was most severe in alloys of near peritectic composition and

moreover, in specimens directly cast in a GLEEBLE arrangement, austenite grains were of

columnar nature compared to equi-axed structures observed in the other alloys. A further finding

of great significance was that the presence of large austenite grains seemed to be more detrimental

to hot-ductility in the austenitic region in the solution treated specimens than an as-cast structure,

at least under the pertaining experimental conditions.

It is pertinent to relate these findings to continuous casting practice. Very large austenite grains

have been observed at the roots of oscillation marks [4] in commercially produced continuously

cast slab, especially in steels of near-peritectic composition. These large austenite grains form

when the peritectic transformation occurs and the thin solidifying shell shrinks from the mould

locally. The concomitant lowering of the heat extraction rate results in this thin part of the shell

being less efficiently cooled and hence, more time is allowed for austenite grain growth to occur.

The observation that columnar grains form in the GLEEBLE samples that have been melted in-

situ, compounds the problem because in the presence of such columnar grains, it is not the average

grain size that will determine the surface crack susceptibility, but the size of the extended

columnar grains.

Although large austenite grains are formed in the mould, the detrimental effect thereof is

mainly realized during unbending where the temperature corresponds to that of the ductility

trough. During unbending the surface of the slab is subjected to tensile stresses and if large grains

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79

exist at the roots of oscillation marks, the surface crack susceptibility is substantially increased.

In order to prevent the formation of large austenite grains, remedial action need to be taken in

the mould. Attempts need to be made to ensure that the volume contraction resulting from the

peritectic transition does not cause large differentials in the cooling rate of the thin solidifying

shell. The judicious selection of mould flux that will equalize, at least in part, the heat extraction

along the surface of the strand is of the essence and some success has already been achieved in

this regard. A better understanding of the mechanism of the peritectic reaction and quantitative

information on the rate of the peritectic transformations are also required if a strategy is to be

designed to prevent the formation of these excessively large, or ‘blown-out’ grains. Fortunately

great strides are currently being made in pursuit of this goal [52].

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CHAPTER 6 - CONCLUSIONS

1. The grain size of the three Fe-C alloys studied, increases almost linearly with increasing

solution treatment temperature. A solution treatment temperature of 1350� rendered an average

austenite grain size of ~4mm in diameter in all three Fe-C alloys. This coarsening of grain size is

attributed to the absence of alloying elements or impurity elements which can restrict grain

growth at this high solution treatment temperature.

2. In the case where the specimens were melted in-situ (direct casting condition), the largest grains

were found in the Fe-0.18%C alloy at both cooling rates used. This observation can be explained

by the higher temperature at which the first austenite forms on cooling in the Fe-0.18%C alloy

than in the others, so that the grains have a better chance to grow at high temperature.

3. Columnar austenite grains were formed in the Fe-0.18%C alloy on cooling from the molten

state, whereas equi-axed grains were formed in the lower or higher C alloys.

4. Increasing grain size resulted in a loss of ductility, widening and deepening the ductility trough

under all test conditions. The existence of a ductility trough after the Ae3 temperature is reached

on cooling, may be due to the formation of deformation induced ferrite. The extension of the

ductility trough to temperatures higher than the Ae3 may possibly be ascribed to the occurrence

of grain boundary sliding. There is a possibility that the precipitation of AlN on austenite grain

boundaries can contribute to the loss of ductility but this aspect was not studied in detail.

5. The peak stress as a function of test temperature in hot tensile tests showed a flattening of the

peak stress in the ferrite-austenite duplex region. In the Fe-0.05%C and Fe-0.18%C alloys, the

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81

Ae3 temperature falls within this flattened region, indicating that deformation induce ferrite can

form at least up to the Ae3 temperature. The extension of this flattened region beyond the Ae3

temperature can not be explained as yet.

6. Because there is an almost linear relationship between peak stress and temperature for the Fe-

0.45%C alloy, it seems that the main embrittling mechanism in this steel is grain boundary

sliding.

7. In the case of the Fe-0.45%C alloy, the loss of ductility below 700� is most probably due to the

formation of a relatively large amount of pearlite at temperatures below Ar1 compared to the

other alloys. Once the Ar1 temperature is reached, pearlite can form and the strength of the

matrix is increased. In addition, the presence of a thin film of ferrite at the grain boundaries,

leads to even more strain concentration at the grain boundaries and the ductility decreases even

further.

8. With increasing grain size the position of the ductility trough as well as the width of trough is

increased.

9. In the case where specimens were cast in-situ (direct casting condition), the higher cooling rate

allows less time for grains to grow in the high austenite temperature region and hence, ductility

improves.

10. For the specimens solution treated to render different grain sizes, the position of the ductility

trough as well as the depth of the trough decreases as the carbon content increases. This

observation correlates with the tendency of decreasing the temperature range of the flattened

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82

region in the peak stress curve as a function of test temperature with increasing carbon content.

This reflects the characteristics of two phase region, i.e., the higher carbon content represent the

lower Ar3 temperature and narrowing two phase region.

11. Under direct casting conditions, the Fe-0.18%C alloy showed the largest grain size at both

cooling rates and consequently the RA value for this alloy represents the lowest value amongst

the three alloys.

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83

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