Effect of ausforming on the bainitic transformation in ...

Universidad Politécnica de MadridEscuela Técnica Superior de

Ingenieros de Caminos, Canales y Puertos

Effect of ausforming on the bainitictransformation in medium carbon steels

Proyecto de fin de Máster

Adriana Eres-CastellanosMaster en Ingeniería de Materiales

Directores del proyectoCarlos Garcia-Mateo

Francisca García CaballeroFrancisco Gálvez Diaz-Rubio

2018

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AcknowledgementsThis thesis would not have been possible without the inspiration and support ofsome people — my thanks and appreciation to all of them.

I owe my gratitude to my supervisors Dr. Carlos Garcia-Mateo, who has knownhow to answer all my questions without going crazy, and Dra. Francisca G. Ca-ballero. I am forever thankful to my colleagues at CENIM: to Miguel A., David, Javiand Miguel(we’ll miss seeing you around) and Ros.

Of course, this project would not have been possible without all the laboratoriesthat are part of CENIM. Thank you Javi, all the experiments wouldn’t have takenplace without your help. I would also like to thank Sidenor S.A., for providingwith the material that has been used, and to Thyssenkrupp, for allowing to use itsequipment.

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AbstractThe best contribution to the strength of bainitic microstructures is due to the scaleof the bainitic ferrite plates. In order to improve this strength, it is important tounderstand and control the parameters that define the final bainitic microstructure.Among them, the strength of the austenite from where the bainite grows and thedriving force of the transformation are the most influential parameters. Historically,controlling the transformation temperature and the chemical composition has beenthe common way of increasing both, the strength of the austenite and the drivingforce of the transformation.Recently, the viability of increasing both magnitudes has been tested by a processcalled ausforming, that consists on applying plastic deformation prior to the trans-formation, obtaining in this way not only thinner bainitic ferrite plates, but also anacceleration of the bainitic transformation.In this project the viability of this thermomechanical treatment and the influence ofthe parameters that define it (deformation, deformation temperature and transfor-mation temperature) on the bainitic transformation will be tested.

La mayor aportación a la resistencia de las microestructuras bainíticas viene deltamaño de las placas de ferrita. Cualquier mejora de la misma pasa por tanto porel conocimiento y el control de los parámetros que definen la escala final de la mi-croestructura bainítica, entre ellos la resistencia de la austenita a través de la cualcrece la ferrita bainítica y la fuerza motriz de la transformación, directamente rela-cionados con la temperatura de transformación y la composición química del mate-rial.Recientemente, se ha comprobado que la aplicación de deformación plástica previaa la transformación, ausforming, permite conseguir de esta manera no solo placasde ferrita bainítica más delgadas, sino también una aceleración de la transformaciónbainítica.En este trabajo se comprobará la viabilidad de este tratamiento termomecánico, y lainfluencia que los parámetros que lo definen (deformación, temperatura de defor-mación y de transformación), tienen en la transformación bainítica.

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Contents

Acknowledgements iii

Abstract v

1 Introduction and objectives 1

2 State of the art 32.1 Bainitic transformation . . . . . . . . . . . . . . . . . . . . . . . . . . . 32.2 Nanostructured bainitic steels . . . . . . . . . . . . . . . . . . . . . . . 62.3 Ausforming treatments . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

2.3.1 Sorts of ausforming . . . . . . . . . . . . . . . . . . . . . . . . . 11

3 Materials and methods 133.1 Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 133.2 Experimental techniques . . . . . . . . . . . . . . . . . . . . . . . . . . 13

3.2.1 Sample preparation . . . . . . . . . . . . . . . . . . . . . . . . . 13Metallographic abrasive cutting . . . . . . . . . . . . . . . . . 14Mounting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 14Coarse grinding . . . . . . . . . . . . . . . . . . . . . . . . . . . 15Polishing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15Etching . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

3.2.2 Scanning electron microscopy . . . . . . . . . . . . . . . . . . . 15Plate thickness measurements . . . . . . . . . . . . . . . . . . . 16

3.2.3 Optical microscopy . . . . . . . . . . . . . . . . . . . . . . . . . 183.2.4 X-ray diffraction analysis . . . . . . . . . . . . . . . . . . . . . 183.2.5 Hardness tests . . . . . . . . . . . . . . . . . . . . . . . . . . . . 193.2.6 Dilatometer tests . . . . . . . . . . . . . . . . . . . . . . . . . . 20

End of bainitic transformation . . . . . . . . . . . . . . . . . . 22Influence of dilatometer module . . . . . . . . . . . . . . . . . 24

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Dilatometer signal correction because of sample length changesprior to isothermal step . . . . . . . . . . . . . . . . . 25

3.3 Thermodynamic calculation software . . . . . . . . . . . . . . . . . . 293.3.1 Thermo-Calc . . . . . . . . . . . . . . . . . . . . . . . . . . . . 293.3.2 MUCG83 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 30

3.4 Finite elements software . . . . . . . . . . . . . . . . . . . . . . . . . . 30

4 Thermal and thermomechanical treatments 334.1 Selection of non-deformation related variables . . . . . . . . . . . . . 33

4.1.1 Critical temperatures calculation and selection of austenitiza-tion temperature and heating and cooling rates . . . . . . . . 33Theoretical critical temperatures calculation . . . . . . . . . . 33Experimental critical temperatures calculation and selection of

parameters . . . . . . . . . . . . . . . . . . . . . . . . 354.1.2 Selection of isothermal treatments temperatures . . . . . . . . 37

4.2 Selection of deformation-related variables . . . . . . . . . . . . . . . . 384.3 Designed thermomechanical tests . . . . . . . . . . . . . . . . . . . . . 39

5 Results and Discussion 435.1 Austempering treatments . . . . . . . . . . . . . . . . . . . . . . . . . 43

5.1.1 Dilatometer tests . . . . . . . . . . . . . . . . . . . . . . . . . . 435.1.2 Microstructural characterization . . . . . . . . . . . . . . . . . 455.1.3 Plate thickness . . . . . . . . . . . . . . . . . . . . . . . . . . . 455.1.4 X-Ray Diffraction measurements . . . . . . . . . . . . . . . . . 465.1.5 Hardness tests . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47

5.2 Ausforming treatments . . . . . . . . . . . . . . . . . . . . . . . . . . . 485.2.1 Dilatometer tests . . . . . . . . . . . . . . . . . . . . . . . . . . 48

Compression step . . . . . . . . . . . . . . . . . . . . . . . . . . 48Isothermal step . . . . . . . . . . . . . . . . . . . . . . . . . . . 48

5.2.2 Microstructural characterization . . . . . . . . . . . . . . . . . 54Microstructural characterization over the sample . . . . . . . 55Effect of ausforming parameters on final microstructure . . . 60

5.2.3 Plate thickness . . . . . . . . . . . . . . . . . . . . . . . . . . . 625.2.4 X-Ray Diffraction measurements . . . . . . . . . . . . . . . . . 655.2.5 Hardness tests . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67

5.3 Comparison . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 68

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5.3.1 Microstructural characterization . . . . . . . . . . . . . . . . . 695.3.2 Plate thickness . . . . . . . . . . . . . . . . . . . . . . . . . . . 715.3.3 X-Ray Diffraction measurements . . . . . . . . . . . . . . . . . 725.3.4 Hardness tests . . . . . . . . . . . . . . . . . . . . . . . . . . . . 76

6 Conclusions 77

References 79

xi

List of Figures

1.1 Tensile strength and total elongation data for conventional steels andSHSS, adapted from[2] . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

2.1 Summary of the mechanism and microstructure of bainite in steels[19] 42.2 T0 curve[1] . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52.3 NANOBAIN TEM picture, where α stands for bainitic ferrite and γ

represents retained austenite . . . . . . . . . . . . . . . . . . . . . . . 62.4 Heat treatment cycles and resultant microstructures during a multi-

step low-temperature super-bainite transformation. Extracted fromreference[29] . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8

2.5 Sketchs of differerent ausforming routes . . . . . . . . . . . . . . . . . 12

3.1 NEURTEK Brillant 220 Precision Abrasive Cut off Saw . . . . . . . . 143.2 Scanning electron microscope JEOL 6500F . . . . . . . . . . . . . . . . 163.3 Lineal intercept LT measurement . . . . . . . . . . . . . . . . . . . . . 173.4 Distribution fitting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 183.5 Inverted optical microscope Olympus CK10M and Nikon COOLPIX

5700 digital camera . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 193.6 BAHR DIL805 high-resolution dilatometer . . . . . . . . . . . . . . . 213.7 Definition of end of bainitic transformation . . . . . . . . . . . . . . . 233.8 Austempering treatments in order to test the reliability of the com-

pression module . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 243.9 Cooling steps in austempering treatments with different modules . . 253.10 Microstructures obtained by same heat treatment and different modules 263.11 Isothermal steps in austempering treatments with different modules 273.12 Isothermal steps in austempering treatments with and without cor-

rection . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 283.13 Isothermal steps in ausforming treatments before and after plastic

strain correction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29

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4.1 Austempering treatments . . . . . . . . . . . . . . . . . . . . . . . . . 344.2 Ausforming treatments sketch . . . . . . . . . . . . . . . . . . . . . . . 344.3 Thermo-Calc output: Ac1=780 °C ; Ac3=1037°C . . . . . . . . . . . . 364.4 Sketch austenization and quenching treatment . . . . . . . . . . . . . 374.5 Heating step up to 1150 °C . . . . . . . . . . . . . . . . . . . . . . . . . 384.6 Cooling steps at a) 25 °C/s ; b) 15 °C/s . . . . . . . . . . . . . . . . . . 394.7 Evaluation of austenite stress-strain curves sketch . . . . . . . . . . . 404.8 Obtained stress-strain curves for different deformation temperatures 40

5.1 Isothermal steps in austempering treatments with and without cor-rection . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44

5.2 Microstructure obtained by austempering after isothermal held at 300°Cfor 1 hour . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 46

5.3 Isothermal steps in austempering and ausforming treatments . . . . 505.4 Isothermal steps in repeated ausforming treatments . . . . . . . . . . 525.5 Calculation of end of bainitic transformation: Tdef = 520°C; high de-

formation; Tiso=300 °C . . . . . . . . . . . . . . . . . . . . . . . . . . . 535.6 Microstructures obtained by ausforming treatments: Tdef = 520°C;

high deformation; Tiso=300 . . . . . . . . . . . . . . . . . . . . . . . . . 535.7 Strain results obtained from Abaqus simulation, where different col-

ors represent different levels of strain in the sample . . . . . . . . . . 565.8 Microstructures throughout the transverse section Tdef = 400°C; low

deformation; Tiso = 300°C . . . . . . . . . . . . . . . . . . . . . . . . . 585.9 Microstructures throughout the longitudinal section Tdef = 400°C;

low deformation; Tiso = 300°C . . . . . . . . . . . . . . . . . . . . . . . 595.10 Sketch showing the microhardness profile configuration . . . . . . . 605.11 Microhardness measurements at different distances from the centre

of the sample, in the longitudinal section Tdef = 400°C; low deforma-tion; Tiso = 300°C . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61

5.12 Microstructures throughout the longitudinal section depending onthe level of deformation and the Tdef Deformation was applied in thevertical direction Tiso = 300°C . . . . . . . . . . . . . . . . . . . . . . . 63

5.13 Microstructures throughout the transverse section depending on thelevel of deformation and the Tdef Tiso = 300°C . . . . . . . . . . . . . . 64

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5.14 Microstructure obtained by ausforming treatment: Tdef = 400°C; lowdeformation; Tiso=300 . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66

5.15 Microstructures in low-strain regions, throughout the longitudinalsection, classified depending on the level of deformation and the Tdefand compared to austempering microstructure Deformation was ap-plied in the vertical direction Tiso = 300°C . . . . . . . . . . . . . . . . 70

5.16 XRD results of samples deformed at 520°C compared to the austem-pering ones, where the marker shape depends on the deformationtemperature and the marker colour depends on the isothermal tem-perature. Vi and Ci stand for volume fraction and carbon content ofthe phase i, whereasDα and ηα are the ferrite crystallite strain and themicrostrain, respectively . . . . . . . . . . . . . . . . . . . . . . . . . . 73



5.19 Hardness as a function of Vα/t in samples subjected to austemperingtreatments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 76

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List of Tables

3.1 Chemical composition /wt% of the steel used in the present work . . 133.2 XRD measurements and hardness measurements of samples subjected

to austempering treatments with different modules, where D. and C.stand for dilatometer module and compression module, respectively.Vi and Ci stand for volume fraction and carbon content of the phasei, whereas ai and ci are the lattice parameters . . . . . . . . . . . . . . 25

3.3 Maximum and minimum limits imposed on each constituent by MUCG 31

4.1 Results of critical temperatures calculations . . . . . . . . . . . . . . . 334.2 Yield strength of austenite at different temperatures . . . . . . . . . . 394.3 Strain, plastic strain and stress that austenite will be subjected to . . . 414.4 Chosen variables for austempering treatments . . . . . . . . . . . . . 414.5 Chosen variables for ausforming treatments . . . . . . . . . . . . . . . 42

5.1 End of bainitic transformation times for austempering treatments . . 445.2 Plate thickness for austempering treatments . . . . . . . . . . . . . . . 465.3 XRD measurements of samples subjected to austempering treatments 475.4 Hardness measurements of samples subjected to austempering treat-

ments. A load of 1 kg was used. . . . . . . . . . . . . . . . . . . . . . . 475.5 Obtained plastic strains in ausforming treatments . . . . . . . . . . . 495.6 Yield strength of austenite at different temperatures . . . . . . . . . . 495.7 Variable parameters for repeated ausforming treatments . . . . . . . 515.8 End of bainitic transformation times for ausforming treatments . . . 545.9 XRD measurements of samples subjected to ausforming treatments

with 1 and 2 hours isothermal holding. Vi and Ci stand for volumefraction and carbon content of the phase i, whereasDi and ηi are crys-tallite strain and the microstrain, respectively . . . . . . . . . . . . . . 55

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5.10 Hardness measurements of the samples obtained by ausforming treat-ments: Tdef = 520°C; high deformation; Tiso=300. A load of 1 kg wasused. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55

5.11 Plate thickness of ausformed samples . . . . . . . . . . . . . . . . . . 655.12 Ferrite results obtained in XRD measurements in ausforming sam-

ples. Vα andCα stand for volume fraction and carbon content, whereasDα and ηα are crystallite strain and the microstrain, respectively . . . 67

5.13 Austenite results obtained in XRD measurements in ausforming sam-ples. Vγ and Cγ stand for volume fraction and carbon content . . . . 68

5.14 Ausforming samples hardness . . . . . . . . . . . . . . . . . . . . . . 695.15 Plate thickness of all samples . . . . . . . . . . . . . . . . . . . . . . . 71

1

Chapter 1

Introduction and objectives

The quest for new super high strength steels (SHSS) stems from the need ofsimultaneously obtaining the increased performance offered by weight reduction,optimal safety, reduced manufacturing costs and reduced environmental impact.There have been different steel design concepts and innovative processes involvedin the design of the new generation of SHSS with improved strength and ductilityproperties, see Figure 1.1. Notable among them is the design and development ofhigh performance bainitic steels (>1200 MPa UTS with high toughness, tribologi-cal properties, fatigue resistance, etc.) for specific applications, such as carbide-freebainitic rails [1].

FIGURE 1.1: Tensile strength and total elongation data for conventionalsteels and SHSS, adapted from[2]

A recent metallurgical development known as nanostructured or low temper-ature bainite has been found to provide an unrivalled combination of propertieswith yield strength (YS) of 1.4 GPa, ultimate tensile strength (UTS) of 2.2 GPa andabout 30% of ductility [3–9]. One of the difficulties that the achievement of such

2 Chapter 1. Introduction and objectives

microstructures is associated with is the need of high carbon contents, from 0.6 to 1wt.%, to ensure both, low transformation temperatures as well as strong austeniteto give nanoscale plates of bainitic ferrite in residual austenite. The industrializa-tion of such microstructures, nowadays, is limited to niche applications[3, 7, 10]because such high carbon contents do not permit to have a good weldability be-cause of the formation of untempered, brittle martensite in the coarse grained heataffected zones of the joint[11]. In China, there has already been some research di-rected towards achieving nanobainite with lower carbon contents [12–14].

Therefore, the novelty of this microstructure must be transferred to much lowerC contents, 0.4-0.5 wt.%. Several approaches have been tested and they will be re-viewed in Section 2.3. One of the ways to do so is by transformation of deformed(ausformed) austenite, into bainite. This gives two benefits: 1) austenite is strength-ened and 2) the transformation temperatures for the start of bainite and martensite(Bs and Ms) are lowered. Both these factors are beneficial with regard to the refine-ment of the final bainitic ferrite-retained austenite microstructure.

However, not all the parameters that are involved in this thermomechanicaltreatment are fully understood. The objective of this project is to compare the mi-crostructures obtained by the common austempering with the ones obtained by aus-forming, to try to better comprehend how parameters such as deformation temper-ature, deformation or isothermal temperature influence on the bainitic transforma-tion.

3

Chapter 2

State of the art

2.1 Bainitic transformation

The term bainite refers to the product of decomposition of austenite (γ) at tem-peratures between those at which pearlite and martensite form[15]. The transfor-mation occurs via a paraequilibrium nucleation on grain boundaries, meaning thatonly C diffuses, and by a displacive difussionless growth, there is no change in thechemical composition between the parent and product phase. Thermodynamicallythese conditions are expressed by

∆Gm < GN

∆Gγ−→α < −GSB

(2.1)

where ∆Gm stands for the maximum change in free energy accompanying thenucleation under paraequilibrium conditions, ∆GN is the universal nucleation func-tion based on a dislocation mechanism of the kind associated with martensite[16–18] which is defined by the minimum change in free energy that is needed in anysteel in order to nucleate bainite and which is usually expressed as

∆GN = 3.5463 · T (K)− 3499.4Jmol−1 (2.2)

GSB stands for the stored energy of bainite due to its displacive nature (' 400J/mol)and ∆Gγ−→α is the free energy change of the transformation of γ without any changein chemical composition.

The temperature at which both conditions, nucleation and growth, are met isknown as the bainite start temperature BS . As already mentioned, parequilibriumconditions means that only C diffuses, and the concentration ratios of all elements

4 Chapter 2. State of the art

but C should be equal in the bulk material as in the retained austenite (γ) and bainiticferrite (α)[15]:

(χFeχi

)bulk = (χFeχi

)γ = (χFeχi

)α (2.3)

where χFe and χi represent the concentration of iron and other elements but car-bon, respectively[18].

Right after the diffusionless growth of a bainitic ferrite subunit is over, the ex-cess of C is partitioned. Depending on whether the transformation takes placesat a temperature close to BS or at a lower temperature, carbon either diffuses tothe surrounding austenite (Figure 2.1) or diffuses to the surrounding austenite andprecipitates in ferrite. The former process gives rise to the so-called upper bainite,whereas the latter one’s product is known as lower bainite.

FIGURE 2.1: Summary of the mechanism and microstructure of bainitein steels[19]

It has to be noted that nucleation of bainite starts at austenite grain boundaries,but further events can take place via the autocatalytic nucleation near the bainiticplates that have just been formed. It is assumed that the opposite site of a bainiticplate that has just grown provides new potential nucleation sites for autocatalyticnucleation. The dislocations in the austenite near a bainitic plate are thought tostimulate autocatalysis because carbon is assumed to pile up at these regions and

2.1. Bainitic transformation 5

FIGURE 2.2: T0 curve[1]

thereby other austenite–bainite interface regions may contain less carbon which pro-motes nucleation and growth of successive sub-units[20].

The process continues by successive nucleation of subunits until the C concen-tration of the residual austenite reaches the value at which the free energy of bainitebecomes less than that of austenite of the same composition, the transformationthen stops because it is thermodynamically impossible for austenite of such chemi-cal composition to transform to bainite[1]. This is known as the incomplete reactionphenomenon and is perfectly represented by the thermodynamic limit (Figure 2.2),which is the locus of all the points where austenite and ferrite have the same chem-ical composition and the same free energy. By this T0 curve and for a given temper-ature, the carbon content of the austenite for which the transformation will stop canbe calculated.


FIGURE 2.3: NANOBAIN TEM picture, where α stands for bainitic fer-rite and γ represents retained austenite

2.2 Nanostructured bainitic steels

By means of the described bainitic transformation theory it is possible to designsteels that transform at homologous temperatures as low as T/Tm ≈ 0.25, where Tmis the absolute melting temperature[4, 10, 21]. High carbon (0.6-1 wt.%) high silicon(1.5-3 wt.%) steels were designed to produce incredibly fine plates of bainitic ferrite,20-40 nm thick, separated by a percolating network of retained austenite after trans-formation at 200-350°C[7, 10]. No carbides have been ever found because of Siliconadditions that suppress their formation[22]. An example of this microstructure, alsoknown as NANOBAIN, is shown in Figure 2.3.

The extraordinary combination of mechanical properties, with yield strength(YS) of 1.4 GPa, ultimate tensile strength (UTS) of 2.2 GPa, 30% ductility and re-spectable levels of fracture toughness (≈ 51 MPa

√m), had never been achieved be-

fore in bainitic microstructures[5, 23]. Such high strength and hardness are due tothe nano-scale plates of bainitic ferrite, which are so supersaturated in C that thelattice is BCT instead of BCC[24]. This microstructures were demonstrated to be re-liable for wear applications in large components, where a uniform microstructure

2.2. Nanostructured bainitic steels 7

free from residual stresses is required[25].

It has been reported that bainitic ferrite plate thickness depends primarily onthree parameters[26, 27]:

• the strength of the austenite at the transformation temperature

• the dislocation density in the austenite

• the chemical free energy change accompanying transformation

All three factors increase as transformation temperature decreases, so lower bai-nite transformation temperatures, which can be more easily achieved in high carbonsteels because Ms is lower, leads to reductions in bainitic ferrite plates thickness andthus to an improvement of the mechanical properties.

However, the lower the isothermal temperature, the longer the heat treatment(transformation may take up to 20 hours to be done[7, 10] for low temperatures)and, in addition, high carbon contents lead to weldability problems owing to un-tempered and thus brittle martensite in the coarse grained heat-affected zones of thejoints that can easily create cracks. Therefore, taking the concept of NANOBAIN tolower carbon content steels, so that both previously explained problems are workedout, is an important objective.

In order to obtain nanostructured bainite in medium carbon content steels, allor some of those previously named factors which influence plate thickness must beincreased by other means.

Some approaches that have been tried out to transfer the nanoscale bainite con-cept to lower C contents so far are:

• Chemical composition modification, for instance, by introducing significantquantities of Mn and Ni, (2.3% Mn & 5 Ni% wt.%), Bs was lowered in low car-bon content steels (0.1-0.2 wt.%)[28], although the difference between Bs andMs decreased drastically.


However, taking into account that low transformation temperatures will onlylead to a nanostructured bainite if austenite from where bainite grows is suffi-ciently strong, alloying might not always be the solution to obtain nanostruc-tured bainite.

• Heat treatments variations, like the multistep heat treatment performed byWang et al.[29] (Figure 2.4), in which austenite carbon content can be contin-uously increased due to the partial bainitic transformation after each step sothat in the final steps, transformation temperatures can be lower then the Msof the bulk alloy.

FIGURE 2.4: Heat treatment cycles and resultant microstructures dur-ing a multi-step low-temperature super-bainite transformation. Ex-

tracted from reference[29]

• Controlling and decreasing the Prior Austenite Grain Size (PAGS), by decreas-ing the austenitization temperature and time, since the smaller PAGS, the lowerMs[30–34]. The reduction of the PAGS is also associated with the increase inthe bainite transformation kinetics, due to the increase in the number densityof austenite grain surface nucleation sites. Consequently, there is a refinementof the size of the bainitic ferrite plates[21, 35].

2.3. Ausforming treatments 9

• A prior deformation to bainitic transformation, in the so-called ausformingprocess, which will be explained in detail in the following section.

In addition and from the point of view of mechanical properties, it is also de-sirable to reduce the amount of retained austenite with block-morphology becauseit can be less easily carbon enriched and, since Ms (martensite start temperature)decreases with carbon content, it could transform to martensite during cooling toambient temperature. If not, this kind of retained austenite could transform, byTRIP-effect, to martensite during subsequent processing or deformation.

Contrary to what happens with blocky austenite, austenite thin-films that areplaced between subunits of bainitic ferrite are very thermally and mechanically sta-ble, because they are always very supersaturated in carbon and its Ms is below theambient temperature. In addition, as its thickness is about few nanometers, no TRIPeffect takes place[9, 36].

2.3 Ausforming treatments

Ausforming is the treatment in which a prior deformation to bainitic transforma-tion is applied. This deformation can increase both austenite strength and disloca-tion density, by work hardening and plastic deformation, respectively. As both areparameters that play an important role in plate thickness, if optimum parameterswere found, auforming could be the solution to obtain nanobainite in medium andlow carbon steels.

In addition, austenite deformation also increases the amount of nucleation sitesfrom which bainite platelets start to grow. For that reason, if deformation is nottoo high, transformation tends to be accelerated[37], since austenite is not so strongto stop plates growth from happening and the amount of nucleation sites has in-creased substantially. However, for larger deformation, the transformation can beretarded[38, 39] or even stopped due to austenite mechanical stabilization[15].

This owes to the fact that deformation adds a new term in the energy changesequation that describes the chemical energy which is needed to allow a bainitic plate


to grow. When there is also an increase of dislocation density, an additional force tomove the austenite/bainite interface through the forest of dislocations is needed.

∆Gγ−→α < −GSB −∆Gγ−→αS (2.4)

where ∆Gγ−→αS is the force due to the additional dislocation density, which is de-

fined as

∆Gγ−→αS =

µbρ1/2

8π(1− ν)[40] (2.5)

where µ stands for the shear modulus, b is the magnitude of a Burgers vectorand ν is the Poisson’s ratio.

Therefore, although for low strains the amount of nucleation sites increases andthus the transformation is accelerated, for larger strains it starts to be retarded, al-though bainite plates grow thinner, up to a point where there is no enough chemicalenergy, this is, austenite has been stabilized and no bainite will grow.

As the mentioned term increases as a function of dislocation density, it is impor-tant to take into account the relationship between dislocation density and strain ε.For low temperatures dislocation density can be calculated by solving the followingequations

εp = ρbL [40] (2.6)

L =δD

δ +Dε[41] (2.7)

where L is the average distance moved by the dislocations, D is the original grainsize of austenite prior to straining and δ is a coefficient abour equal to 1µm.

However, this expression does not take into account the temperature at whichthe strain is applied, as it is assumed that the temperature is low.

At higher temperatures, a dislocation annihilation phenomenon occurs and thusdislocation density is lower than expected. This phenomenon and the evolution of


dislocation density with strain has been described, for instance, by the Kocks andMecking Model[42], which is based on the following equation

dρ

dε=k1

b

√ρ− fρ (2.8)

where b is the magnitude of the Burgers vector, k1 is the storage coefficient and f isthe dynamic recovery term, which increases with temperature, decreasing disloca-tion density rate thus.

For that reason, effects on either microstructure or mechanical properties at highdeformation temperatures could not be as noticeable as they are at low ones.

2.3.1 Sorts of ausforming

There are three ausforming routes of interest:

• high temperature ausforming (Figure 2.5a), in which the austenite is deformedat high temperatures, although always lower than the non-recrystallizationtemperature (TNR)

• medium temperature ausforming (Figure 2.5b), in which the austenite is de-formed at temperatures that are between the ferritic-pearlitic region and thebainitic region

• low temperature ausforming (Figure 2.5c), in which the austenite is deformedat temperatures that are close to the isothermal temperature, although this isthe most difficult process to perform, due to the high austenite strength atthose temperatures


(A) High-temperature ausforming process

(B) Medium-temperature ausforming process

(C) Low-temperature ausforming process

FIGURE 2.5: Sketchs of differerent ausforming routes

13

Chapter 3

Materials and methods

3.1 Materials

The chemical composition of the steel which has been used in this project andthat will be called 04C3Si in this document is included in Table 3.1. Although isis not added in the corresponding table, there are also trazes of elements such asVanadium, Surfer or Molybdenum, among others. It has been provided by Sidenor(Spain) and it is a medium carbon high silicon steel. While Silicon makes it per-fect to be subjected to heat treatments in order to obtain free carbide bainite[43], allManganese, Chromium and Molybdenum confer hardenability to the steel[44].

3.2 Experimental techniques

3.2.1 Sample preparation

In order to characterize the obtained microstructures, samples must be preparedin a specific way that will be explained subsequently.

TABLE 3.1: Chemical composition /wt% of the steel used in the presentwork

Name C Si Mn Cr

04C3Si 0.4-0.5 3.0-3.5 0.5-1.0 0.5-1.0

14 Chapter 3. Materials and methods

Metallographic abrasive cutting

Specimens were cut by the Automatic Precision Abrasive Cut off Saw Brillant220 which is sold by Neurtek and can be seen in Figure 3.1.

FIGURE 3.1: NEURTEK Brillant 220 Precision Abrasive Cut off Saw

Mounting

Mounting operations are required to improve handling since samples size issmall and sample preparation can be harsh if no mounting is used and to betteridentify the different specimens. Most of characterization equipments have samplecarriers that allow for the same size mountings. Because of that, mounting dimen-sions are set as follows: diameter of 30 mm and thickness lying between 10 and 14mm. Two different kinds of mountings have been used in this project:

• Compression mounting is widely used for metals and it consists of embed-ding either one or multiple samples in bakelite, which is kept for about twentyminutes at specific pressure and temperature. Bakelite is electrically conduc-tive and thus this is the kind of mounting that is used in Scanning ElectronMicroscope.

• Specimens can also be embedded in castables mounting resins, which are twocomponent systems (epoxi and hardener) and are specially recommended forX-Ray Diffraction since they are transparent to the incident beam, they do not

3.2. Experimental techniques 15

block any diffraction from the crystal and they do not produce extra diffrac-tion.

Coarse grinding

Coarse grinding is required to planarize the specimen and to reduce the damagecreated by sectioning. The coarse grinding step is accomplished by decreasing theabrasive grit/particle size sequentially to obtain surface finishes that are ready forpolishing. In this project, all specimens were coarse ground with 600, 1200, 2000 and4000 grit SiC abrasive paper, respectively.

Polishing

The purpose of the polishing step is to remove the damage produced duringcutting and planar grinding. Proper polishing will maintain specimen flatness andretain all inclusions or secondary phases. Polishing was accomplished primarilywith diamond abrasives ranging from 3 micron down to 0.25-micron diamond.

Etching

The purpose of etching is to optically enhance microstructural features such asgrain size and phase features and thus must be only done in case the microstructurehas to be revealed. Nital Etch is the most common metallographic etchant to etchsteels. It can also etch cast irons and some tool steels. The 2% Nital is the most pop-ular percentage of Nital.

Due to the special sensitivity of the obtained microstructure, in order to improvethe looking of the microstructure, specimens were etched with a 2% nital solutionand polished with 1µm diamond paste, etched again and polished with 0.25µm di-amond paste. Finally, the specimen was etched using the same 2% nital solution.This proved to make bainite plates easier to be distinguished.

3.2.2 Scanning electron microscopy

In order to look at the microstructure at high magnification and then seek for dif-ferences among different treatments, the Scanning Electron Microscope (SEM) witha tungsten electron source Schottky JEOL JSM 6500F (Figure 3.2) has been used. It


operates at 1 to 30 kV with a magnification range that lies between 10x and 400,000x,it adopts several detectors that allow to use secondary and backscattered electronimages modes and it allows for X-Ray mapping, electron backscatter diffraction andcathodoluminescende.

FIGURE 3.2: Scanning electron microscope JEOL 6500F

Plate thickness measurements

As plate thickness is directly related to mechanical properties, as explained inSection 2.3, it must be measured in this project. In order to do so, the followingprocedure[45, 46] has been followed:

• The lineal intercept LT in a direction normal to the plate length must be mea-sured in the way shown in Figure 3.3. Lineal intercept must be measured asthe distance between two austenite-ferrite interfaces, always in a perpendicu-lar direction to both.


FIGURE 3.3: Lineal intercept LT measurement

• There must be as many measurements as necessary to fit a lognormal distri-bution, in other words, the logarithm of them must fit a normal distribution(Figure 3.4).

• A correction must be applied in order to obtain the plate thickness t, takinginto account that its relationship with the mean linear intercept (L̄T ) is:

L̄T =πt

2(3.1)

The standard error of t can be obtained by the following expression:

E = ± 2 · σαLπ ·√N

(3.2)

Where σαL stands for the standard deviation of the linear intercept and N refersto the number of measurements.

It’s important to understand that the applied correction is based on the statementthat[47–49] plates are located randomly so it only applies for isotropic microstruc-tures. In case of deformation, anisotropy could nor be reliable.


1.4 1.6 1.8 2 2.2 2.4 2.60

0.1

0.2

0.3

0.4

0.5

log(LT )/nm

Percentfrequency/%

FIGURE 3.4: Distribution fitting

3.2.3 Optical microscopy

Optical microscope allows to look at samples which have been subjected to ther-mal treatments. Although no magnifications higher than 1000x can be reached andthus it has to be combined with higher resolution microscopes, it can be useful tostudy the cleanliness of the material, this refers to the amount of inclusions that canbe found, or to analyze the microstructure for the first time. In this project, an Olym-pus CK10M inverted optical microscope and a Nikon COOLPIX 5700 digital cameraconnected to the microscope have been used. Both are shown in Figure 3.5.

3.2.4 X-ray diffraction analysis

X-Ray Diffraction (XRD) measurements are useful to determine the volume frac-tion of austenite and bainite, as well as the lattice parameters of both phases, takinginto account that bainitic ferrite is considered as tetragonal[24]. From those lat-tice parameters, carbon content of austenite and bainitic ferrite can be obtainedby Dyson & Holmes’[50] and Cohen´s[51] equations, respectively, with an errorof 0.05%. Crystallite size and ferrite microstrain are also given, where the formerone is a measure of the smallest coherent diffraction region (defect free) in a crys-tal (2nm error) and the latter one refers to the no-uniform deformation within thephase(0.2·10−3 error).


FIGURE 3.5: Inverted optical microscope Olympus CK10M and NikonCOOLPIX 5700 digital camera

The XRD measurements were made by a Bruker AXS D8 X-ray diffractometerwith a rotating Co anode X-ray tube as a radiation source, Goebel mirror optics anda LynxEye Linear Position Sensitive Detector for ultra-fast XRD measurements . Inthis procedure, the specimen is simultaneously rotated 360°around its normal axiswhile scanning process. Scanning was carried out in the range from 35°to 135°andthe step between two consecutive measurements was 0.01°. Voltage was 40kV andcurrent was 30mA.

Volume fractions have been calculated according to ASTM E975-13 and their er-ror is 3%. When studying polycrystalline metals, Rietveld,’s is one of the most suit-able methods for lattice parameters calculations from XRD spectra. The methodrequires, for each phase, the prior knowledge of the structural model, which is usu-ally kept fixed during refinement. The mixture pattern is calculated by taking intoaccount the contributing pattern of each phase[52] and the final obtained error is0.001Å. In order to eliminate the effect of texture and collect the correct intensitiesof the phases, harmonic functions were used.

3.2.5 Hardness tests

In order to measure the hardness of different specimens, Vickers hardness testshave been performed. The surface of the material was subjected to a load of 1 kg


applied by a square pyramid with an angle of 136 degrees between opposite faces.The load was always kept for 15 seconds and then both two diagonals of the slop-ing surface of the indentation were measured. Vickers hardness is defined by thefollowing formula:

HV = 1.845 · Fd2

(3.3)

where F is the applied load in kgf and d stands for the arithmetic mean of the twodiagonal values in mm.

3.2.6 Dilatometer tests

Dilatometry is an important technique that can be useful when studying solid-state phase transformations in steels since it allows to know when those transforma-tions take place in real time. It’s based on the fact that a material changes its crys-talline structure when it goes through a phase transformation, varying its specificvolume and thus the sample’s length. Therefore, the temperatures at which a phasetransformation occurs during cooling, heating or isothermal treatment, among oth-ers, can be obtained by dilatometer tests. The study of these transformations can bealso done.

A BAHR DIL805 high-resolution dilatometer (Figure 3.6) was used in this projectto track the isothermal decomposition of austenite into bainitic ferrite and to makesure there is no martensite formation. Sample dimensions were 5 mm in diameter an10 mm length and they were subjected to subsequent thermomechanical treatmentsthat are programmed by inputting such parameters as deformation, temperatureand time, among others.

Temperature is controlled by Type K thermocouples that are welded to the cen-tre of the sample. Heatings and isothermal treatments are carried out in a vacuumchamber below 10−10 Torr in order to prevent the specimen from decarburization oroxidation whereas coolings are performed by controlled Helium flows by a solenoidvalve. Punchers used with the compression module are made out of silicon nitridewhereas the push rods used with the dilatometer module are made out of siliconcarbide.


FIGURE 3.6: BAHR DIL805 high-resolution dilatometer

The dilatometer can provide with different kinds of data, such as the relativechange in length that the specimen undergoes during the treatments, the differentialrelative change in length, time, temperature or stress and strain during compressionsteps, among others.

Since both silicon carbide and silicon nitride have low expansion coefficientscompared to steel’s, if isotropy is assumed, the dilatometric relative change in lengthsignal can be only attributed to the microstructural changes that occur due to thethermal treatments. When no phase change occurs, sample’s length changes by anamount proportional to the original length and the change in temperature:

∆L

L0

= α ·∆T (3.4)

where α (°C-1) is the linear expansion coefficient. However, when a phase changetakes place, the lattice parameters change and this can imply either an expansionor a contraction. The relative change in length vs. time curve obtained during theisothermal step in which bainite is obtained is also known as Avrami’s curve andit is characterized by achieving a plateau after a certain amount of time. Up to thatpoint, both nucleation and growth occur.


It is important to realize that deformation implies anisotropy and thus, dilatome-ter signal might not be as reliable as it is in isotropic cases. In order to better un-derstand how transformation occurs in anisotropic samples, another BAHR DIL805high-resolution dilatometer placed in Germany and own by Thyssenkrupp AG hasbeen used. Besides measuring longitudinal changes, this dilatometer has an addi-tional option to measure radial changes by a laser.

End of bainitic transformation

Since the objective of this project is to compare the effect of different levels ofdeformation at different temperatures, kinetics of bainitic transformation must bewell-known. Therefore, the calculation of the end of the transformation time ob-tained by dilatometer data can be useful to compare the kinetics of the transforma-tion during each of the treatments.

The end of bainitic transformation is defined as the time from which no bainiteis formed anymore. Although this point theoretically corresponds to the time fromwhich the Relative change in length vs. Time curve gets stabilized, it is really hard tofind no variations at all in the dilatometric signal due to noise, so this end time hasbeen redefined as the time at which the differential function takes either 4, 8 or 15percent its maximum value. All the points that correspond to values obtained beforethat maximum differential value is reached must be neglected since it’s known thatthe end of the transformation cannot occur in that period, which is associated withincubation times, nucleation and growth.

All three times obtained by 4, 8 and 15% must be plot and it must be checkedthat there is a negative trend graph, taking then the value of 4% criterion as end oftransformation time. If there is no such a trend, values will be neglected, startingfrom 4%’s value, until graph has the explained trend. In this way we can avoid er-rors provoked by noise or effects of previous thermal steps during the dilatometertest. Figure 3.7 shows this process graphically.

In this project, Avrami’s curves have presented different trends that they usuallyhave, due to anisotropic issues that will be explained in detail in Chapter 5. In those


End of Bain.transf. time

Time/s

Relativechangeinlength/%

max

.diff

.

x% max. diff.

Time

Dif f

eren

tial

/%/s

0 5 10 15

Chosen Endof Bain.transf.

Percentage of maximum differential value /%

End

ofBa

in.t

rans

f.ti

me

/s

FIGURE 3.7: Definition of end of bainitic transformation


cases, the criteria has been applied only taking into account the differential functionin the last part of the experiment that is similar to the typical signal.

Influence of dilatometer module

Two different modules have been used in this project: the dilatometer module,which is more precise, and the compression module, which load can be appliedwith, although there can be some problems regarding accuracy. In order to testhow reliable the dilatometer signal that is given when the compression module isinstalled is, two identical pure isothermal treatments were performed in the puredilatometery module and also in the compression one. The sketch of the treatmentis shown in Figure 3.8.

Tγtγ

5 °C/s

15 °C/s

300 °C/s1 h

25 °C/s

Time /s

Tem

pera

ture

/°C

FIGURE 3.8: Austempering treatments in order to test the reliability ofthe compression module

Although no intermediate phase transformations have been detected in any ofthe cases through dilatometry(Figure 3.9) and all X-Ray Diffraction, hardness tests(Table 3.2) and optical micrographs (Figure 3.10) proved that the microstructure wasthe same, the dilatometric signal during the isothermal step looks rather differentbetween both cases (Figure 3.11). While the signal seems to stabilize from a certain


400 600 800 1,000−0.5

0

0.5

1

1.5

Temperature /°C


Cooling from 990°C to 300°C

0 50 100 150 200 250 300

Temperature /°C

Cooling from 300°C to Troom

— Dilatometer module —Compression module

FIGURE 3.9: Cooling steps in austempering treatments with differentmodules

amount of time when the dilatometric module is used, it keeps increasing for theother one. Therefore, the compression module can be identified as a progressingsystematic error[53].

TABLE 3.2: XRD measurements and hardness measurements of sam-ples subjected to austempering treatments with different modules,where D. and C. stand for dilatometer module and compression mod-ule, respectively. Vi and Ci stand for volume fraction and carbon con-

tent of the phase i, whereas ai and ci are the lattice parameters

Mod. Vγ aγ Cγ Vα aα cα Cα HV1±3/% ±0.001/Å ±0.05/% ±3/% ±0.001 /Å ±0.001 /Å ±0.05/%

D. 16 3.612 0.95 85 2.852 2.876 0.19 578±6C. 19 3.609 0.89 82 2.853 2.875 0.17 584±5

Dilatometer signal correction because of sample length changes prior to isother-mal step

The dilatometer calculates relative change in length as the percentage expansionof the sample compared to its initial length. However, during heating, austeniza-tion and first cooling (prior to isothermal treatment), length undergoes variations


(A) Sample which dilatometer module wasused in

(B) Sample which compression module wasused in

FIGURE 3.10: Microstructures obtained by same heat treatment anddifferent modules

and thus, there can be some error in the Avrami curve associated to the isothermalstep. For that reason, studies about the influence of this change on the signal hasbeen carried out in both treatments with and without prior deformation.

Taking into account that the initial length (l0) that has been input was unknown,it was calculated from the average of the initial lengths that can be obtained fromthe relative value at every point (l0i).

l0 =

∑ni=1 l0in

=

∑ni=1

∆li10·relin

=

∑ni=1

li−l010·relin

(3.5)

where li is the change in length that the sample is subjected to at the momenti, reli is the relative change in length at the moment i and n is the total number ofpoints that the dilatometer has input throughout the test. Then, the corrected changein length after a pure isothermal treatment (relnewi) can be calculated as follows:

relnewi =∆li −∆lbeforeiso

lbeforeiso=

∆li −∆lbeforeisol0 + ∆lbeforeiso

=li − l0∑n

i=1li−l010·relin

+ li − l0(3.6)

where li stands for the length at the point i and lbeforeiso and ∆lbeforeiso are thesample length and the change in length at the first point of the isothermal treat-ment, respectively.


0 500 1,000 1,500 2,000 2,500 3,000 3,5000

0.1

0.2

0.3

0.4

0.5

0.6

0.7

Time/s


Dilatometric moduleCompression module

FIGURE 3.11: Isothermal steps in austempering treatments with differ-ent modules

To measure the accuracy of this correction, the relative error was obtained asaverage of relative errors of every correction:

εT =

∑ni=1 εin

(3.7)

where the relative error is obtained according to the Principle of propagation ofuncertainty:

εi = 2 · ε(li) + ε(l1) + ε(l0) +n∑i=1

(ε(reli) + ε(li) + ε(l0)) (3.8)

In order to test whether the effect of correcting l0 is significant or not, austem-pering treatments that have been performed for this project at 300, 325 and 350 °Chave been corrected. Figure 3.12 shows the signal during isothermal steps with andwithout correction. It is obvious that the dilatometer signal barely changes after be-ing corrected. The error has been calculated in all three cases and it is always about


0.2%. Therefore, the effect of correcting the relative change in length is neglictiblewhen no plastic deformation is applied and it is not necessary to correct it.

0 500 1,000 1,500 2,000 2,500 3,000 3,5000

0.1

0.2

0.3

0.4

0.5

0.6

Time/s


300 °C without correction300 °C with correction



FIGURE 3.12: Isothermal steps in austempering treatments with andwithout correction

However, when applying deformation prior to the isothermal step, the relativeerror gets significant. A calculation to obtain the correction factor by which theAvrami curve should be multiplied in order to be able to compare curves wheredifferent deformations were applied is written below.

relnewi =∆li −∆lbeforeiso

lbeforeiso=

∆li −∆lbeforeisol0 + ∆lbeforedef − εplastic(l0 + ∆lbeforedef )

=

=∆li −∆lbeforeiso

(1− εplastic)(l0 + ∆lbeforedef )≈ ∆li −∆lbeforeiso

l0 · (1− εplastic)=

reli(1− εplastic)

(3.9)

where εplastic stands for the plastic deformation that has been applied and can becalculated from the stress-strain curves that the dilatometer inputs. Therefore, inorder to correct the curves it is enough to divide the relative change in length givenby the dilatometer by (1 − εplastic). Figure 3.13 shows the same dilatometer signal

3.3. Thermodynamic calculation software 29

obtained during isothermal step (after ≈ 16% plastic deformation) before and afterthe above-mentioned correction.

0 500 1,000 1,500 2,000 2,500 3,000 3,500 4,000 4,500 5,0000

5 · 10−2

0.1

0.15

0.2

0.25

0.3

0.35

0.4

0.45

Time/s


Before correctionAfter correction

FIGURE 3.13: Isothermal steps in ausforming treatments before andafter plastic strain correction

3.3 Thermodynamic calculation software

3.3.1 Thermo-Calc

Thermo-Calc is thermodynamic calculation software for tacking mineral equilib-ria problems. It is formed by the application itself and a internally-consistent ther-modynamic dataset. In this project, Thermo-Calc based on the CALPHAD (Com-puter Coupling of Phase Diagrams and Thermo-chemistry) and using the databaseknown as TCFE8 has allowed to calculate the Ac1 and Ac3 temperatures to estimatethe austenization temperature.


3.3.2 MUCG83

MAP_STEEL_MUCG83 is a free powerful program for the pseudo-empirical mod-elling of transformations in steels that has been developed by the Phase Transforma-tions and Complex Properties Group (University of Cambridge). Steel’s chemicalcomposition (in wt%) has to be input, taking into account the eleven most com-mon constituents. There are different upper and botton limits for each of them’scomposition, as shown in Table 3.3. The program supplies with data that allows torepresent a Temperature-Time-Tranformation (TTT) diagram, as well as many otherparameters regarding solid state phase transformations from austenite to marten-site, bainite, Widmanstätten ferrite, allotriomorphic ferrite etc.

The main parameters that are useful for this project and that are given by thispiece of software are the bainite start temperature (Bs) and the martensite start tem-perature (Ms), these are, the higher temperatures at which bainite and martensitestart to form, respectively. Those two values can be used as an estimation of thetemperatures between which the isothermal treatments must be performed in orderto obtain bainite. However, since 04C3Si’s silicon content is slightly higher than theupper limit allowed by the program, it is important to consider that results mightnot be completely accurate and thus those temperatures must be also calculated byempirical equations so that they can be compared before making a decision.

3.4 Finite elements software

Barrelling effect is unavoidable in compression tests and thus, stress and strainare not equally distributed throughout samples. As the effect of both parametershave been studied in the project, it is important to better understand how stress andstrain are distributed along samples which are subjected to barrelling effect. In or-der to do so, Abaqus FEA has been used.

Abaqus FEA is a piece of software suite for finite element analysis and computer-aided engineering which allows to obtain the final state of a part after going throughdeformation or temperature changes, among others. To do so, it is necessary to

3.4. Finite elements software 31

TABLE 3.3: Maximum and minimum limits imposed on each con-stituent by MUCG

Input Order Element Max /wt% Min /wt%

1 Carbon 2.0 0.0012 Silicon 2.5 0.03 Manganese 3.5 0.04 Nickel 3.5 0.05 Molybdenum 1.5 0.06 Chromium 3.5 0.07 Vanadium 1.5 0.08 Cobalt 4.0 0.09 Copper 4.0 0.010 Aluminium 2.0 0.011 Tungsten 4.0 0.0

model the part and input some parameters, such as the material constants or theboundary conditions, both in the initial and the final state.

33

Chapter 4

Thermal and thermomechanicaltreatments

The most simple treatment that bainite can be obtained by is Austempering,whose sketch is shown in Figure 4.1. However, by applying a prior deformationin the so-called Ausforming, whose sketch can be seen in Figure 4.2, the transforma-tion can be either accelerated[37] or retarded[38, 39]. Microstructural characteriza-tion after both treatments will be carried out and compared in this project.

4.1 Selection of non-deformation related variables

4.1.1 Critical temperatures calculation and selection of austeniti-

zation temperature and heating and cooling rates

Theoretical critical temperatures calculation

Critical temperatures were calculated by both specific software and empiricalequations and they are all shown in Table 4.1.

TABLE 4.1: Results of critical temperatures calculations

Empirical formulas Thermo-Calc MUCG_86

Ac1(°C) Ac3(°C) Ms(°C) Bs(°C) Ac1(°C) Ac3(°C) Ms(°C) Bs(°C)

820 888 291 481 780 1037 314 488

Calculations of Martensite Start Temperature (°F), Bainite Start Temperature (°C),Ac1 (°C) and Ac3 (°C) have been respectively made by the different equations that

34 Chapter 4. Thermal and thermomechanical treatments

Tγtγ

HR

CR

Tisotiso

CR-2

Time /s

Tem

pera

ture

/°C

FIGURE 4.1: Austempering treatments

Tγtγ

HR

CR1

Tisotisot r1

Tdef ; ε; ε̇

CR2t r2

CR3

Time /s

Tem

pera

ture

/°C

FIGURE 4.2: Ausforming treatments sketch

4.1. Selection of non-deformation related variables 35

are listed subsequently, where all the chemical compositions are given in wt%. Equa-tions have been chosen so that they are suitable according to steel composition, al-though the calculations have to be contrasted to experimental methods since highSilicon content is sometimes out of bounds. If required, units have been changed sothat all the results are given in Celsius degrees.

Nehrenberg equation[54]:Ms( °F ) = 930− 540 · C − 60 ·Mn− 40 · Cr − 30 ·Ni− 20 · Si− 20 ·Mo

Trzaska equation[55]:Bs( °C) = 771− 231.5 · C − 69 ·Mn− 58.5 · Cr − 31 ·Ni− 23 · Si− 55 ·Mo− 41 · VAndrews equation[56]:Ac1( °C) = 723− 16.9 ·Ni+ 29.1 · Si+ 6.38 ·W − 10.7 ·Mn+ 16.9 · Cr + 290 · AsAndrews equation[56]:Ac3 °C) = 910− 203

√C + 44.7 · Si− 15.2 ·Ni+ 31.5 ·Mo+ 104 · V + 13.1 ·W − 30 ·

Mn+ 11 · Cr + 20 · Cu− 700 · P − 400 · Al − 120 · As− 400 · Ti

Thermo-Calc outputs the molar fraction of each of the phases that is expectedas a function of temperature for a certain composition (Figure 4.3). Ac1 and Ac3temperatures can be taken as those ones at which FCC (austenite) molar fractionstarts to increase and reaches its maximum value, respectively.

Experimental critical temperatures calculation and selection of parameters

Since the theoretical calculations cannot be blindly trusted because Silicon con-tents are sometimes too high for the applied formula, a first dilatometer test wasperformed in order to measure Ac1 and Ac3. Samples were heated at 5°C/s up toTγ=1150°C, held at that temperature for four minutes and then quenched at 180°C/s (Figure 4.4).

Figure 4.5 shows the relative change in length that was observed as a functionof temperature during the heating step. Based on this, we can take Ac1 and Ac3 asthe lower and upper ends of the range where it is observed that the curve does notbehave linearly. The error that previous theoretical calculations have presented canbe attributed to the fact that they assume equilibrium whereas the tests might have


0 200 400 600 800 1,000 1,200

0

0.2

0.4

0.6

0.8

1

Temperature /°C

Mol

arfr

acti

on/

Ac1 and Ac3 temperatures calculation by Thermo-Calc

BCCFCC

FIGURE 4.3: Thermo-Calc output: Ac1=780 °C ; Ac3=1037°C

been affected by kinetic effects.

Therefore, based on these results, the austenization temperature has been de-cided to be Tγ=990°C along the whole project.

It is known that Prior Austenite Grain Size (PAGS) increases with austenizationtemperature and time[30–33] and since PAGS influences Ms[57], it must be mea-sured in samples that were austenized at the same Tγ that has been chosen for thecurrent work.

Critical cooling rate (CCR) is that cooling rate that is low enough so that no otherphases but martensite form during cooling to room temperature. Two different cool-ing rates were tested this time: 25 and 15 °C/s. Because martensite temperature alsodepends on the cooling rate that the sample is quenched at[58], measurements havebeen made by dilatometer tests where cooling rates took both values. Similarly towhat happened when measuring Ac1 and Ac3 temperatures, experimental Ms can

4.1. Selection of non-deformation related variables 37

Tγ=1150°C, tγ=4 min

5°C/s180°C/s

Time /s

Tem

pera

ture

/°C

FIGURE 4.4: Sketch austenization and quenching treatment

be obtained as the temperature at which the relative change in length vs. Tempera-ture curve is not linear anymore. Both cooling steps and their corresponding Ms areplotted in Figure 4.6.

In both cases, no ferrite formation was found during cooling so it can be ensuredthat the CCR is lower than 15 °C/s. In this project, a cooling rate of 15 °C/s will beused in every cooling step before bainitic transformation takes place, since in thisway no other phases such as pearlite or ferrite will form.

4.1.2 Selection of isothermal treatments temperatures

Since the isothermal treatment temperatures that allow to obtain bainite can bechosen among all the temperatures that lie between Ms and Bs and the experimentalvalue of Bs has not been obtained, isothermal treatments have been decided to beperformed at 300 °C, 325 °C and 350 °C,values close to Ms.


0 200 400 600 800 1,000 1,200

0

0.2

0.4

0.6

0.8

1

1.2

1.4

Ac1 = 842°C

Ac3 = 961°C

Temperature/C


FIGURE 4.5: Heating step up to 1150 °C

4.2 Selection of deformation-related variables

Deformation temperatures (Tdef ) were chosen to be 520, 400 and 300°C and dilatome-ter tests to obtain the stress-strain curves at those three temperatures were con-ducted. Samples were heated up at 5°C/s to Tγ=990°C, held at that temperaturefor four minutes and cooled at 15°C/s down to Tdef . After holding at those tem-peratures for 15 s in order to allow the whole sample to stabilize, deformations of50% were applied. The deformation rate was always 0.04 s−1. Figure 4.7 shows thesketch of the previously explained tests.

The obtained stress-strain curves are plot in Figure 4.8. In some cases the loadlimit of the cell was reached and for that reason, data has only been obtained up tostrains smaller than 50%.

From those stress-strain curves, the yield strength of the austenite at differenttemperatures was calculated by the 0.2 offset criterion, which consists of setting YSas the value of stress for which the plastic strain is 0.2%. Results are presented in

4.3. Designed thermomechanical tests 39

0 200 400 600 800 1,000

0

0.5

1

1.5

Ms = 280°C

a)

Temperature/C


200 400 600 800 1,000

0

0.5

1

1.5

Ms = 287°C

b)

Temperature/C


FIGURE 4.6: Cooling steps at a) 25 °C/s ; b) 15 °C/s

TABLE 4.2: Yield strength of austenite at different temperatures

Tdef Yield strength /MPa

520 191.1400 213.0300 223.1

Table 4.2. It can be seen that, as expected, yield strength decreases as Tdef increases.

Taking into account the limit of the dilatometer load cell and the fact that strainmust be high enough so that there is plastic deformation, different levels of defor-mation were decided to be applied on the subsequent ausforming tests dependingon the Tdef . These values, besides their corresponding plastic strains and the stressesthat are required for each of the cases are included in Table 4.3.

4.3 Designed thermomechanical tests

Since one of the objectives of this project is to compare the final microstructureobtained by ausforming with the initial microstructure that can be obtained by sim-ple isothermal holding, firstly some austempering treatments must be conducted,as was explained in Chapter 3. Their sketch can be seen in Figure 4.1 and the cho-sen variables are included in Table 4.4. Although the isothermal time is said to be


Tγ=990°C, tγ=4 min

5°C/s 15°C/s

15s

Tdef ; ε = 0.5ε̇ = 0.04s−1

10s

25°C/s

Time /s

Tem

pera

ture

/°C

FIGURE 4.7: Evaluation of austenite stress-strain curves sketch

0 5 10 15 20 25 30 35 400

200

400

600

800

1,000

1,200

Strain /%

Stre

ss/M

Pa

Tdef=520°CTdef=400°CTdef=300°C

FIGURE 4.8: Obtained stress-strain curves for different deformationtemperatures

4.3. Designed thermomechanical tests 41

TABLE 4.3: Strain, plastic strain and stress that austenite will be sub-jected to

Tdef ε /% εp /% σ /MPa

52010 8.8 415.720 18.8 641.830 28.8 864.4

400 10 9.2 491.220 19.2 820.3

300 10 9.0 628.2

1 hour, some austempering treatments have been repeated for two times the initialisothermal time so that all the tests meet the 4% criteria, assuming that in that waythe transformation will be in the same state in all cases.

TABLE 4.4: Chosen variables for austempering treatments

HR/°C/s

Tγ /°C tγ /min CR/°C/s

Tiso /°C tiso /h CR-2/°C/s

5 990 3 15300

1 25325350

The second cooling rate has been chosen to be 25 °C/s, value for which nomartensite was formed at all when cooling down from Tγ. Taking into accountthat austenite is more enriched after the isothermal step and that Ms temperaturegoes down with carbon content, it can be ensured that no martensite will be foundin any of the treatments.

Regarding ausforming treatments, Figure 4.2 shows the sketch of the ausformingtreatments, whereas the parameters that either will keep fixed or will be changed inorder to study their influence on the microstructure are included in Table 4.5.


TABLE 4.5: Chosen variables for ausforming treatments

(A) Chosen fixed variables for ausforming treatments

HR

/°C

/s

Tγ

/°C

tγ/m

in

CR

/°C

/s

t r1/s

ε̇/s

-1

t r1/s

CR

-2/°

C/s

t iso

/h

CR

-3/°

C/s

5 990 3 15 15 0.04 10 15 1 25

(B) Variable parameters for ausforming treatments

Tdef /°C ε /% Tiso /°C

520

10300325350

20300325350

30300325350

400

10300325350

20300325350

300 10 300

43

Chapter 5

Results and Discussion

5.1 Austempering treatments


The isothermal treatments that are represented in Figure 4.1 were carried out.None of the coolings have shown there was martensite formation during the treat-ment (coolings looked the same as the ones in Figure 3.9, where no expansion takesplace), so the whole microstructure will be formed by bainite and retained austenite.

Figure 5.1 shows the relative change in length vs. time plots during all theisothermal steps where the bainitic transformation takes place. As this transfor-mation is isotropic, it can be assumed that higher dilatometric values correspond tohigher expansions and thus higher bainite volume fractions, although this must bemade sure by XRD measurements.

Under this consideration, it is clear that higher isothermal temperatures leadto lower bainite volume fractions, although transformation occurs faster (becauseit takes shorter times to reach stable values, which are associated with the end oftransformation). Nevertheless, the end of bainitic transformation times have beencalculated based on these curves and they will be discussed below.

Calculation of end of bainitic transformation times

The end of transformation times were calculated as explain in Chapter 3 and theobtained results are included in Table 5.1. They indeed increase for lower isothermaltemperatures, as TTT diagrams predict. It is important to point out that the longest

44 Chapter 5. Results and Discussion

0 500 1,000 1,500 2,000 2,500 3,000 3,5000

0.1

0.2

0.3

0.4

0.5

0.6

Time/s


Tiso=300 °CTiso=325 °CTiso=350 °C

FIGURE 5.1: Isothermal steps in austempering treatments with andwithout correction

time is close to 13 minutes, so isothermal times do not need to be as long as the onesthat were used in this project and they could be shortened if this treatments wereimplemented at an industrial scale.

TABLE 5.1: End of bainitic transformation times for austempering treat-ments

Tiso °C End of B. /s

300 727325 475350 382

5.1. Austempering treatments 45

5.1.2 Microstructural characterization

Figure 5.2 shows one of the obtained microstructures after austempering treat-ments. Along with the rest of the microstructures, different points were concluded:

• Although it was expected that the microstructures of non-ausformed sampleswere purely bainitic, this is, they were composed of ferrite plates and retainedaustenite, in some cases, small amounts of allotromorphic ferrite were found.This is due to a non-rapid enough intermediate cooling which could not beforeseen based on dilatometry results. However, since its amount is so low, incan be ignored when analyzing the characteristics of bainite.

• Carbides and martensite weren’t formed because of silicon addition and austen-ite carbon enrichment, respectively.

• Both Optical microscope and SEM showed that ferrite plates are located ran-domly throughout the sample as PAG are equiaxial. Therefore, all indicationswere that there is no variant selection.

• It was detected blocky austenite that, as known, could be detrimental, since itcould turn into martensite under deformation by TRIP effect. An example ofit is marked in Figure 5.2.

5.1.3 Plate thickness

Plate thickness was measured according to Section 3.2.2 and results are includedin Table 5.2. Lower isothermal temperatures lead to thinner plates, since austenitestrength is higher and this hinders the plate from growing quick, so the microstruc-ture tends to be refined.

Additionally, plates are thicker than the ones obtained in NANOBAIN (Fig-ure 2.3) because carbon content is not as high as the one used for those steels and,since plate thickness is directly related to mechanical properties, the current steel isexpected not to have such a good behaviour as the one found in nanobainitic steels.This matter will be more deeply analyzed in Section 5.3.4.


Blocky austenite

FIGURE 5.2: Microstructure obtained by austempering after isothermalheld at 300°C for 1 hour

TABLE 5.2: Plate thickness for austempering treatments

Tiso °C t/nm

300 47.9±1.6325 62.3±2.8350 92.1±3.2

5.1.4 X-Ray Diffraction measurements

Table 5.3 shows the results obtained from XRD as a function of the isothermaltreatment, where Dxα and ηα stand for ferrite crystallite size and microstrain, re-spectively. According to the T0 curve, volume fraction of bainitic ferrite increaseswith decreasing temperature[59] and austenite must be more carbon enriched too.

However, at those low temperatures it is more difficult for carbon to partitionand austenite thus does not receive as much carbon as it should. Most likely, if weprolonged isothermal times, austenite would be as enriched as T0 curve indicates,this is, sample treated at 300°C would be the most enriched one. However, as this

5.1. Austempering treatments 47

amount is negligible, industrially is not worth to prolong times so much as mechan-ical properties would not change significantly.

Regarding ferrite crystallite size and microstrain, the former one increases andthe latter one decreases as isothermal temperature does. This behaviour is consis-tent with the bainitic transformation itself, the microstructure becoming thinner andwith a higher dislocation density when transformation temperature is lowered[18,27].

TABLE 5.3: XRD measurements of samples subjected to austemperingtreatments

Tiso Vγ Cγ Vα Cα Dxα ηα · 103

°C ±3/% ±0.05/% ±3/% ±0.05/% ±2/nm ±0.2

300 16 0.95 85 0.19 34 2.2325 16 1.03 84 0.17 38 2.0350 20 1.07 80 0.16 48 1.9


Hardness measurements are presented in Table 5.4. As isothermal temperature islower, samples become stronger, as expected since ferrite volume fraction increasesand plate thickness decreases.

TABLE 5.4: Hardness measurements of samples subjected to austem-pering treatments. A load of 1 kg was used.

Tiso °C HV1

300 578±6325 533±4350 502±0


5.2 Ausforming treatments


Ausforming treatments that are represented in Figure 4.2 were carried out. Ashappened in the previous treatments, none of the coolings showed there was marten-site formation during the treatment, so the whole microstructure will be also formedby bainite and retained austenite.

Since, as explained in Section 3.2.6, dilatometer relative change in length mustbe corrected when deformation has been applied by multiplying it by 1

1−εp , first ofall the applied plastic deformation must be obtained. For that reason, compressionstep will be firstly analyzed, followed by the isothermal step dilatometer signal.

Compression step

To obtain the plastic deformation that has been applied, compression step stress-strain curves must be analyzed, taking plastic strain as the subtraction of the strainassociated with the elastic region from the maximum strain value. Plastic strainsobtained for each of the tests are included in Table 5.5. As can be seen, really appliedplastic deformations differ from the originally selected values, although since for thesame conditions of deformation, they are not really scattered and then can be worthto compare those conditions. In order to easily name the tests which have similardeformation, those with strains that lie between 8.9% and 9.2% will be referred toas low deformation ones, whereas medium and high deformation ones will strainsthat lie between 16% and 16.6% and between 22.5% and 23.6€%, respectively.

Yield strengths at each temperature have also been recalculated taking into ac-count the previous tests and the currently explained ones. Their values are pre-sented in Table 5.6, which evinces that, yield strength increases as temperature de-creases.

Isothermal step

Figure 5.3 shows the relative change in length vs. time plots during all theisothermal steps where the bainitic transformation takes place. Although Avrami


TABLE 5.5: Obtained plastic strains in ausforming treatments

Tdef /°C Tiso /°C ε /% εp /%

520

300 9.7 8.9325 9.7 9.1350 9.7 9.0

300 17.0 16.4325 16.9 16.1350 16.5 16.0

300 23.1 22.5325 24.2 23.5350 24.5 23.6

400

300 9.9 9.3325 9.9 9.0350 9.9 9.2

300 17.6 16.6325 16.9 16.2350 16.9 16.1

300 300 9.4 8.8

curves are tipically always positive when samples are isotropic, in this case the de-formation has resulted to make samples anisotropic and thus, transformation doesnot really have to happen in all directions in the same way.

In isotropic cases, the shear components of the individual shape deformations ofthe large number of variants which form along any dimension tend to cancel out ona macroscopic scale, whereas the dilatational ones average, provoking an isotropicexpansion. If there is a texture in the sample, as the one that comes from austenite

TABLE 5.6: Yield strength of austenite at different temperatures

Tdef Yield strength /MPa

520 192.6 ± 8.2400 213.6 ± 10.3300 215.6 ± 10.6


deformation in this case, differences between strains throughout different axis canbecome apparent.[60]

low deformation medium deformation high deformation

Tiso =300°C Tiso =325°C · · · · · Tiso =350°C

0 1,000 2,000 3,000−1

−0.8

−0.6

−0.4

−0.2

0

0.2

0.4

0.6

Tdef =520°C

Time/s


0 1,000 2,000 3,000

Tdef =400°C

Time/s

0 1,000 2,000 3,000

Tdef =300°C

Time/s

FIGURE 5.3: Isothermal steps in austempering and ausforming treat-ments

This effect has been mostly reported when applying stress during the isothermalheld[61–64] and it is explained to have to do with a crystallographic variants favour-ing. Variant selection has been also proved to be present in the transformation fromtextured austenite to bainite[37, 65] and, therefore, this could lead to anisotropictransformation along transverse and longitudinal axis.


Two experiments were repeated in a different dilatometer that allows to mea-sure the radial expansions during the test by a laser. In this way, the explained effectcould be proved and it could be made sure that the curves weren’t due to machineerrors. The variables of the repeated experiments, besides their obtained final plas-tic strains, are included in Table 5.7, whereas the relative change in length in bothtransverse and longitudinal direction is shown in Figure 5.4.

It is clear that indeed there is an expansion in the radial direction while the longi-tudinal one gets contracted due to the previous texture. The kink in the longitudinalsignal (specially noticeable in Subfigure 5.4a) probably means that, at the beginninganisotropy plays an important role in the transformation, so most of the plates growin the radial direction, whereas from a certain point of time, since no so much spaceis left, they start to grow in a more random way and that’s why the longitudinalsignal becomes less negative.

TABLE 5.7: Variable parameters for repeated ausforming treatments

Tdef/°C

ε /% Tiso /°C εp /%

520 30.25 325 29.51400 10.43 350 9.24

Calculation of end of bainitic transformation times

The end of transformation times were calculated as explain in Chapter 3 and theobtained results are included in Table 5.8. However, for dilatometer signals that donot have the common Avrami’s curve trend, like the ones with kinks, times seem tobe excessive and not to correspond to the real time.

For example, after a 23.7% plastic deformation at 520°C and an isothermal stepat 300°C for 1 hour, dilatometry showed that there was no end for the 4% criteria, soisothermal time should be longer. The treatment was repeated for a longer isother-mal time, where the end of bainitic transformation time was calculated to be 5141 s,as can be seen in Figure 5.5.


0 1,000 2,000 3,000−1

−0.8

−0.6

−0.4

−0.2

0

0.2

0.4

0.6

0.8

1

Time/s


RadialLongitudinal

(A) Tdef = 520°C; high deformation;Tiso = 325°C

0 1,000 2,000 3,000

Time/s

(B) Tdef = 400°C; low deformation;Tiso = 350°C

FIGURE 5.4: Isothermal steps in repeated ausforming treatments

Both samples were taken to X-ray diffraction, where volume fractions and lat-tice parameters (and thus carbon content of each of the phases) were obtained andthey are included in Table 5.9. Taking into account the equipment errors associatedto the XRD measurements, it is clear that the transformation was done before onehour, even though the end of transformation time couldn’t be calculated. Althoughhardness measurements seem to indicate that the microstructure is slightly different(Table 5.10), micrographs show that it is the same in both cases (Figure 5.6).

Therefore, although the end of transformation time calculations seem reliable forpure isothermal treatments, they cannot be really trusted in cases where the com-pression module has been used in the dilatometer. This is probably due to the errorassociated to the compression module, as happened when testing its reliability in


0 1,000 2,000 3,000 4,000 5,000−0.4

−0.3

−0.2

−0.1

0

0.1

0.2

EndofB.t=5141 s

Time/s


FIGURE 5.5: Calculation of end of bainitic transformation:Tdef = 520°C; high deformation; Tiso=300 °C

(A) 1 hour isothermal step (B) 2 hour isothermal step

FIGURE 5.6: Microstructures obtained by ausforming treatments:Tdef = 520°C; high deformation; Tiso=300


TABLE 5.8: End of bainitic transformation times for ausforming treat-ments

Tdef °C Deformationlevel

Tiso °C End ofB. /s

520

low300 2782.8325 1631350 2999

medium300 5954.4325 1987350 4118

high300 5140.8325 2797350 234

400

low300 273.6325 162350 108

medium300 190.8325 166350 115

300 low 300 137

Chapter 3.


Microstructures obtained after ausforming treatments turned out to be ratherdifferent that the previous ones. Strain and the temperature at which this strainwas imposed seem to have a big influence over platelets distribution over the sam-ple. The parameter that links both mentioned ones is the dislocation density, whichincreases as deformation temperature decreases and strain increases, respectively.However, it is also important to take into account that if, dislocation density is toohigh, mechanical stabilization of austenite might occur.

Microstructural characterization was performed for all conditions and the mostimportant found out aspects will be discussed subsequently. As strain is applied


TABLE 5.9: XRD measurements of samples subjected to ausformingtreatments with 1 and 2 hours isothermal holding. Vi and Ci stand forvolume fraction and carbon content of the phase i, whereas Di and ηi

are crystallite strain and the microstrain, respectively

tiso Vγ Cγ Vα Cα/h ±3/% ±0.05/% ±3/% ±0.05/%

1 21 0.97 79 0.172 16 1.15 85 0.16

TABLE 5.10: Hardness measurements of the samples obtained by aus-forming treatments: Tdef = 520°C; high deformation; Tiso=300. A load

of 1 kg was used.

tiso /h HV1

1 564±122 591±0

in one specific direction, microstructures have been analysed for both transverseand longitudinal sections in every case. They indeed present differences that willbe explained in more detail below. In addition, as proper lubrication is difficult toobtain during compression tests, microstructural changes over samples will be alsoanalysed, taking into account that strain changes due to barrelling effect and thismight affect how transformation occurs.

Microstructural characterization over the sample

Due to barrelling effect, there is a strain distribution over the austenitic samplebefore bainitic transformation that affects the final microstructure, as mechanicalfree energy associated with strain changes depending on dislocation density. In or-der to better understand how strain, and thus, dislocations, are distributed through-out samples, and then see how mechanical free energy change actually affects thetransformation, FEM simulations were performed by Abaqus FEA. A sample withthe same dimensions that the ones used for the dilatometer tests, this is, a 5-mm-diameter, 10-mm-long cylinder has been modelled. Since only the strain distribu-tion and not its real values are required for this comparison, several aspects havebeen assumed:

• both dilatometer punchers were simplified by being modelled as shells


(A) Transverse section of the sample (B) Longitudinal section of the sample

FIGURE 5.7: Strain results obtained from Abaqus simulation, wheredifferent colors represent different levels of strain in the sample

• it was assumed that the strain distribution would look similar regardless ma-terial properties, so only one simulation where material properties at 400 °C(obtained from dilatometer tests in Section 4.2) were input was done

• the coefficient of friction was assumed, being its only requirement to be highenough to provoke a noticeable barrelling effect

• only one deformation level was simulated, so that it could be seen where thetheoretically applied strain is really reached and which are the different areaswhere this strain is much lower than expected so that it could be comparedwith final microstructure

Therefore, simulation was performed for 10% deformation at 400 °C and the ob-tained strain results can be seen in Figure 5.7. Due to barrelling effect, strain is inho-mogeneously distributed along the sample. While this effect can be barely noticedthroughout the transverse section, where microstructure and hardness are expectedto be the same, there is a stronger effect in the longitudinal section. Strain is almostnegligible in the points which are directly in contact with the punchers, whereas ittakes its maximum value, which is closer to the applied deformation, in the centralpoints.

In order to study whether the effect of strain is significant or not, the differentregions that correspond to different strain levels in both transverse and longitudi-nal sections were put under the microscope. Microhardness profiles were also per-formed.


This study has been performed in several samples with different ausformingconditions, assuming that the strain distribution is the same, despite strain valueschange. The same effect is always found, although its not so noticeable for samplesdeformed at the highest temperatures because of dislocation annihilation. There-fore, while the mechanical energy added by deformation is proportional to strainwhen deformation temperature is low enough, at higher temperatures dislocationdensity can be reduced because dislocations are more prone to rearrange in the crys-tal structure (as explained in Section 2.3).

For that reason, the differences in microstructure and hardness will be only shownfor the sample deformed up to ≈10% at 400°C and then isothermally held at 300°C.

Figures 5.8 and 5.9 show the different obtained microstructures depending onthe area and the section. As the former image proves, same levels of deformationsare directly related to similar microstructures, since the rest of the energies that formpart of the bainitic transformation remain the same in the whole sample and the onlyenergy that varies along the sample is the one that is originated by the deformation.The later figure is an example of how different microstructure can look like depend-ing on the area that is put under the microscope. While microstructure in Subfig-ure 5.9a is totally affected by deformation, presenting rather oriented bainite platesand a lack of blocky austenite, Subfigure 5.9c looks quite similar to microstructuresobtained by pure isothermal held (Figure 5.2), in which growth takes place in noparticular order.

A microhardness profile was also perfomed in the studied sample, microhard-ness was measured from regions near the no-strain edge to the centre, as Figure 5.10.A load of 100g was applied for 15 seconds and seven measurements were realizedin each point. Figure 5.11 shows the obtained results with their corresponding stan-dard error as a function of the distance from the centre to the point where the mea-surement was made. Although, due to the low load that had to be chosen to be ableto make so many measurements, results are rather scattered, they are enough to ob-serve that hardness increases as points are closer to the centre. This might owe tothe fact that those regions are more deformed, so austenite became stronger prior to


transformation and it could thus have to do with a plate thickness decrease. In ad-dition, this could also be indicative of a dependency between more ordered bainiticmicrostructures, in which not so much blocky austenite seems to be present, and agreater hardness.

Therefore, from now on, all the microstructural characterization have been car-ried out in the central parts of both transverse and longitudinal sections so that theycan all be compared among them. Moreover, hardness measurements have been

(A) Microstructure in the centre of the sample, transverse section

(B) Microstructure close to the edge of the sample, transverse section

FIGURE 5.8: Microstructures throughout the transverse sectionTdef = 400°C; low deformation; Tiso = 300°C


(A) Microstructure in the centre of the sample, highest strain longitudinal section

(B) Microstructure close to the lateral edge of the sample, medium strain, longitudinal sec-tion

(C) Microstructure close to the top edge of the sample, lowest strain, longitudinal section

FIGURE 5.9: Microstructures throughout the longitudinal sectionTdef = 400°C; low deformation; Tiso = 300°C


FIGURE 5.10: Sketch showing the microhardness profile configuration

always made in the transverse section, where hardness is assumed to take its maxi-mum values and not to be too scattered because the strain level does not change.

Effect of ausforming parameters on final microstructure

Microstructures in both tranverse and longitudinal sections have been lookedunder the microscope for all ausforming samples.To simplify, and because microstruc-tures found at different isothermal temperatures are rather similar among themif the rest of the parameters remain constant, only micrographs of samples whichwere isothermally held at 300 °C will be shown, although it must be clear that allwhich will be commented also applies to Tiso of 325 and 350 °C. The fact that mi-crostructures are slightly more refined, a slight additional amount of ferrite formsand austenite is slightly more enriched when samples have been held at lower tem-peratures will be then disregarded in this subsection, although these aspects will bediscussed in more detail in Subsections 5.2.3 and 5.2.4.

Longitudinal section

Figure 5.12 shows the microstructure that was found in the central part of thelongitudinal sections of the samples (compression had been applied vertically), de-pending on the deformation and the temperature which that deformation was ap-plied at. First of all, microstructures deformed at the highest temperature are muchmore disordered. This can have happened because dislocation density increased by


500 520 540 560 580 600 6200

0.5

1

1.5

2

2.5

3

3.5

HV 0.1

Vert

ical

dist

ance

toth

ece

ntre

/mm

FIGURE 5.11: Microhardness measurements at different distances fromthe centre of the sample, in the longitudinal section

Tdef = 400°C; low deformation; Tiso = 300°C

deformation is lower at higher temperatures, at which dislocations are more able tomove and reorganize.

Second of all, at lower deformation temperatures, plates seem to have formedalong specific directions, which are always at about ±45°. Plates thus cross one an-other at a 90°angle (Subfigure 5.9). This behaviour has already been published byGong W. et at[66], who concluded that a strong variant selection had taken placeand that this variant selection was related to slip systems operated in austenite atthe deformation. In the current case, further study would be needed.

Moreover, at the lowest deformation temperature, this effect, although still ap-preciable, is less noticeable than at the medium temperature and plates seem to growin only one direction. It can have to do with mechanical stabilization: austenite wastoo deformed that bainitic ferrite could not grow as freely as it did at samples de-formed at 400°C.


Finally, because of bainite growth preference, blocky austenite seems to be presentmuch less frequently at lower deformation temperatures than in the other left case,which can make mechanical properties improve for the those cases as TRIP effectcannot take place in retained austenite.

Transverse section

Figure 5.13 also shows the microstructures depending on the deformation andthe temperature which that deformation was applied at, although this time the sec-tion that was put under the microscope was the transverse one. The same trend thatwas discussed in the previous subsection is found, although there are some differ-ences.

First of all, compression direction does not promote any special growth direc-tion in the transverse section, so plates are not placed in any preferred direction.Secondly, specially at lower deformation temperatures, a higher amount of platesgrow in the same direction within every grain, so sheaves are bigger up to the pointthat they sometimes span whole grains. As explained before, this leads to loweramounts of blocky austenite.

In addition, owing to deformation, there are some pancake grains (one of themcan be clearly seen on the top left part of Subfigure 5.13f). Although they can bemore likely to appear in the micrographs that regard to the longitudinal section,because of the deformation direction, since a grain changes its shape depending onthe stress that is transferred from the adjacent grains, they can also be found in thesection perpendicular to the applied load.


Table 5.11 shows plate thickness that were measured in all the ausforming spec-imens in the transversal section. It is clear that plates are thinner as isothermaltemperature decreases since the yield strength of austenite is higher at lower tem-peratures[67].


(A) Tdef = 520°Chigh deformation

(B) Tdef = 520°Cmedium deformation

(C) Tdef = 520°Clow deformation

(D) Tdef = 400°Cmedium deformation

(E) Tdef = 400°Clow deformation

(F) Tdef = 300°Clow deformation

FIGURE 5.12: Microstructures throughout the longitudinal section de-pending on the level of deformation and the Tdef

Deformation was applied in the vertical direction Tiso = 300°C







(F) Tdef = 300°Clow deformation

FIGURE 5.13: Microstructures throughout the transverse section de-pending on the level of deformation and the Tdef

Tiso = 300°C


TABLE 5.11: Plate thickness of ausformed samples


Tiso °C t /nm

520

low300 46.9± 2.2325 50.4±2.2350 69.4±4.6

medium300 61.4±2.1325 64.0±2.7350 53.4± 1.4

high300 54.7±3.3325 59.8±3.8350 64.4± 1.7

400

low300 56.7±3.0325 63.7±2.6350 67.1±3.2

medium300 64.5±3.4325 84.7±3.9350 86.7±4.5

300 low 300 41.4±1.9

Since it was surprising that there seems not to be any trend between dislocationdensity ρ(Tdef , ε) and plate thickness, microstructures in both sections in the samesamples were compared (Figure 5.14). At a glance, it is obvious that the lineal inter-cepts LT from which plate thickness is obtained (Section 3.2.2) do not measure thesame. This can be explained by the fact that plates are located in specific directionsand the calculated plate thickness based on the stereographic corrections is not asreliable as it is when plates have grown randomly.


Tables 5.12 and 5.13 include XRD measurements made in ausforming samples,in both ferrite and austenite, respectively. Regarding volume fractions, it is clearthat larger amounts of austenite are found on samples that were deformed at higherlevels, problably because of its mechanical austenitization, which prevents bainitic


(A) Transverse section (B) Longitudinal section

FIGURE 5.14: Microstructure obtained by ausforming treatment:Tdef = 400°C; low deformation; Tiso=300

ferrite from growing. However, this effect is more pronounced in samples whichhad been subjected to isothermal holdings at higher temperatures (Tiso=350°C), atwhich the chemical driving force is less representative and then the added mechan-ical driving force is more significant in the transformation.

In addition, as happened in the austempering samples, higher isothermal tem-peratures lead to smaller ferrite volume fractions because chemical driving force forbainitic transformation is lower.

Austenite carbon content was expected to decrease with plastic strain since a de-crease in the amount of ferrite that is formed means that there is less carbon availableto partition towards austenite and because carbon tends to go towards the defects.Regarding crystallite size, it tends to decrease as deformation is larger and isother-mal temperature is lower, since dislocation density is higher and they can movemuch more difficult so it’s not possible for them to redistribute properly. Takinginto account the error, microstrain barely changes with deformation, although it de-creases as isothermal temperature is higher. The trend should be the opposite as theone observed for crystallite size, this is, higher plastic strains constrain the latticemore so microstrain should increase with deformation. The fact that austenite car-bon contents and microstrains do not follow the expected trend could be related tothe difficulty to measure because of applied deformation. The lattice distortion pro-voked by deformation could makee latice parameters (and thus carbon contents),


TABLE 5.12: Ferrite results obtained in XRD measurements in ausform-ing samples. Vα and Cα stand for volume fraction and carbon content,whereas Dα and ηα are crystallite strain and the microstrain, respec-

tively

Tdef Deformation Tiso Vα Cα Dxα ηα · 103

/°C level /°C ±3/% ±0.05/% ±2/nm ±0.2

520

low300 79 0.17 36 2.1325 77 0.16 43 1.9350 74 0.15 46 1.7

medium300 82 0.18 29 2.0325 81 0.17 26 1.7350 73 0.15 71 1.6

high300 79 0.17 26 2.0325 73 0.16 28 1.9350 67 0.14 33 1.7

400

low300 76 0.17 36 2.2325 79 0.17 42 2.0350 71 0.16 50 1.7

medium300 74 0.17 30 1.8325 74 0.17 37 1.8350 69 0.16 37 1.5

300 low 300 74 0.18 43 2.2

crystallite size and microstrain not to be reliable.


Table 5.14 shows ausforming samples hardness, classified according to the threevariables that had been modified in the tests. As happened in austempering treat-ments, samples get harder as isothermal temperature increases. Concerning de-formation, it seems that, for same isothermal temperature, hardness slightly de-creases with deformation. This could happen because ferrite volume fraction hasa more significant influence over hardness than plastic deformation. Therefore, forhighly deformed specimens, although austenite has hardened because of deforma-tion, since not such high amount of ferrite is formed as in less deformed samples,


TABLE 5.13: Austenite results obtained in XRD measurements in aus-forming samples. Vγ and Cγ stand for volume fraction and carbon con-

tent

Tdef Deformation Tiso Vγ Cγ°C level °C ±3/% ±0.05/%

520

low300 21 0.91325 23 1.07350 26 1.04

medium300 18 1.04325 19 1.08350 24 1.05

high300 21 0.97325 27 1.03350 34 1.02

400

low300 24 1.20325 21 1.20350 29 1.24

medium300 26 1.28325 26 1.28350 31 1.26

300 low 300 26 1.17

hardness decreases.

5.3 Comparison

In the current chapter, both austempering and ausforming samples will be com-pared according to different factors. Dilatometer results cannot be compared since,as proved, anisotropy is present in the latter case and then the found expansionscannot be assumed to be proportional to ferrite volume fractions, as is done whenno deformation is applied prior to transformation.

5.3. Comparison 69

TABLE 5.14: Ausforming samples hardness


Tiso °C HV1

520

low300 588 ±10325 561 ±0350 513 ±18

medium300 588±5325 574±12350 509±8

high300 564 ±12325 527 ±17350 493 ±4

400

low300 551 ±8325 521 ±4350 504 ±8

medium300 521 ±8325 481 ±20350 496 ±8

300 low 300 606 ±5


As dislocation density increases (this is, deformation temperature is lower ordeformation is higher), higher anisotropy is found in samples and then microstruc-tures seem to be ordered in a specific way. For that reason, two effects are found:

• Microstructures found in samples where deformation temperature is higherare more similar to the ones where no deformation was applied, since disloca-tion density is lower, as can be seen in Figures 5.12 and 5.13.

• All microstructures found in low-strain areas (this is, those areas which arecloser to the punchers), look similar than the ones found in no-deformationsamples, although this similarity is more pronounced in samples deformed at520°C. Figure 5.15 shows the comparison among microstructure at those low-strain regions and the austempering microstructure.







(F) no deformation (G) Tdef = 300°Clow deformation

FIGURE 5.15: Microstructures in low-strain regions, throughout thelongitudinal section, classified depending on the level of deformation

and the Tdef and compared to austempering microstructureDeformation was applied in the vertical direction

Tiso = 300°C

5.3. Comparison 71


Table 5.15 includes all plate thickness values for austempering and ausform-ing treatments. It looks like deformation makes plate thickness be less sensitive toisothermal temperature since there are more pronounced differences in plate thick-ness among austempering samples than among ausforming ones.

TABLE 5.15: Plate thickness of all samples


Tiso °C t /nm

zero300 47.9±1.6325 62.3±2.8350 92.1±3.2

520

low300 46.9± 2.2325 50.4±2.2350 69.4±4.6

medium300 61.4±2.1325 64.0±2.7350 53.4± 1.4

high300 54.7±3.3325 59.8±3.8350 64.4± 1.7

400

low300 56.7±3.0325 63.7±2.6350 67.1±3.2

medium300 64.5±3.4325 84.7±3.9350 86.7±4.5

300 low 300 41.4±1.9

Although, contrary to what was expected, deformation does not seem to directlyimply plate thickness decrease and in some cases plates even broaden when de-formation has been applied, this might be due to the bad use of the stereographiccorrection factor, which only applies for isotropic cases, as explained in Section 3.2.2.


However, it is important to point out that, even though the real plate thicknessvalues are still not known, they are always in the same order as the ones obtainedby isothermal treatments; they all have nanometric scale and this microstructure canbe thus still considered nanostructured. By better understanding how the bainitictransformation from deformed austenite works, plate thickness could be decreased.


Figures 5.16, 5.17 and 5.18 show how different parameters obtained by XRDchange as deformation increases. Figures show the results obtained in tests wheredeformation had been applied at 520, 400 and 300°C, respectively, all contrastedwith the non-deformed status.

Results have been divided depending on the Tdef because the behavior com-pletely changes since, at higher temperatures, it’s harder for dislocation density toincrease and, once dislocations have been created, they are freer to redistribute alongthe lattice. Each of the parameters will be discussed below.

First of all, regarding austenite and ferrite volume fraction, two effects are mostlyfound: austenite volume content increases as isothermal temperature and defor-mation do. The former effect can be explained by the fact that chemical drivingforce increases as isothermal temperature decreases and then ferrite is more proneto grow. The latter one is explained by mechanical austenite stabilization. The moredeformed austenite is prior to bainitic transformation, the more stabilized it is andthen it is harder for bainite platelets to grow.

Second of all, carbon contents must be also discussed. In all cases, taking intoaccount the error, it can be assumed that ferrite carbon contents do not change de-pending on how previously deformed austenite was. Concerning austenite carboncontents, it seems that two different behaviors are present depending on Tdef . Asexplained in Section 5.2.4, lattice distortion caused by deformation makes austenitecarbon content values, crystallite sizes and microstrains not to be reliable.

5.3. Comparison 73

Tdef = 520°C no deformation

Tiso = 300°C Tiso = 325°C Tiso = 350°C

0 5 10 15 20 25

20

30

εp/%

Vγ/%

0 5 10 15 20 25

0.9

1

1.1

εp/%

Cγ/%

0 5 10 15 20 25

70

80

90

εp/%

Vα/%

0 5 10 15 20 25

0.1

0.15

0.2

0.25

εp/%Cα/%

0 5 10 15 20 25

30

40

50

εp/%

Dxα/nm

0 5 10 15 20 25

1.5

2

2.5

εp/%

η α·1

0−3

FIGURE 5.16: XRD results of samples deformed at 520°C compared tothe austempering ones, where the marker shape depends on the defor-mation temperature and the marker colour depends on the isothermaltemperature. Vi and Ci stand for volume fraction and carbon contentof the phase i, whereas Dα and ηα are the ferrite crystallite strain and

the microstrain, respectively


4 Tdef = 400°C no deformation


0 5 10 15

20

30

εp/%

Vγ/%

0 5 10 15

1

1.2

εp/%

Cγ/%

0 5 10 15

70

80

90

εp/%

Vα/%

0 5 10 150.1

0.15

0.2

0.25

εp/%

Cα/%

0 5 10 15

30

40

50

εp/%

Dxα/nm

0 5 10 15

1.5

2

2.5

εp/%

η α·1

0−3



5.3. Comparison 75

� Tdef = 300°C no deformation


0 2 4 6 8

15

20

25

30

εp/%

Vγ/%

0 2 4 6 8

0.9

1

1.1

1.2

εp/%

Cγ/%

0 2 4 6 870

75

80

85

εp/%

Vα/%

0 2 4 6 8

0.15

0.2

0.25

εp/%Cα/%

0 2 4 6 8

35

40

45

εp/%

Dxα/nm

0 2 4 6 8

2

2.2

2.4

εp/%

η α·1

0−3





Figure 5.19 shows hardness as a function of Vα/tα for both austempering and aus-forming treatments. As can be seen, there is a trend between both parameters, prov-ing that ferrite volume fraction increases hardness, whereas it decreases as platesare wider. There does not seem to be any relation between the level of deformationor the temperature at which this deformation occurs and all three previous values.

0.7 0.8 0.9 1 1.1 1.2 1.3 1.4 1.5 1.6 1.7 1.8

460

480

500

520

540

560

580

600

620

Vα/t /%/nm

HV

1

AustemperingAusforming

FIGURE 5.19: Hardness as a function of Vα/t in samples subjected toaustempering treatments

77

Chapter 6

Conclusions

• Contractions were detected during the isothermal step in the longitudinal di-rection in some experiments where austenite had been previously deformed.Expansions detected in the radial direction proved that deformation provokessuch an anisotropy that bainitic transformation occurs in certain directions andfor that reason even contractions can be detected

• In austempering treatments, higher isothermal temperatures lead to faster bainitictransformations, although also to thicker bainitic plates and lower ferrite vol-ume fractions, which is directly related to lower hardness

• As strain is not homogeneously distributed along the sample after being de-formed, microstructure change. In order to compare the resulted microstruc-tures obtained after different ausforming treatments, it must be looked in thecentral part of both transveral and longitudinal sections, since the central sam-ple area is the one where strain takes its maximum value throughout the sam-ple

• After ausforming it seems that, microstructures along the longitudinal sectionseem to have some similarities, although they are more noticeable at mediumdeformation temperatures. Plates seem to be oriented in specific directions,crossing one another at 90°angle. This trend seems to be directly related todeformation and for that reason the effect is not so apparent for the highestdeformation temperature. It also seems that, at low deformation temperatures,plates grow in only one direction instead of two

• In ausforming treatments, when both deformation temperature and appliedplastic strain keep constant, the trend among the isothermal temperature and

78 Chapter 6. Conclusions

all plate thickness, ferrite volume fraction and hardness is the same than wasfound in isothermal treatments

• All isothermal treatments and ausforming ones presented a nanostructuredmicrostructure

• Ausformed samples have smaller amounts of ferrite as deformation increases,due to mechanical stabilization

• Hardness is directly proportional to the rate between ferrite volume fractionand ferrite plate thickness

79

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