Microstructure and Transformation
Kinetics in Bainitic Steels
Ph.D. thesis
Kinetics in Bainitic Steels
Proefschrift
ter verkrijging van de graad van doctor aan de Technische Universiteit Delft,
op gezag van de Rector Magnificus prof. dr.ir. J.T. Fokkema, voorzitter van het College voor Promoties,
in het openbaar te verdedigen op dinsdag 2 december 2008 om 10.00 uur
door
Natalia Vadimovna LUZGINOVA
Master of Science (Physics) Tomsk State University, Russia
Copromotor: dr.ir. J. Sietsma Samenstelling promotiecommissie: Rector Magnificus, voorzitter
Prof.dr.ir. L.A.I. Kestens, Technische Universiteit Delft, promotor Dr.ir. J. Sietsma, Technische Universiteit Delft, copromotor
Prof.dr.-ing. W. Bleck, RWTH Aachen, Aachen, Deutschland
Prof.dr. P.J. Jacques, Université catholique de Louvain, Louvain-la-Neuve, Belgique Prof.dr. R. Boom, Technische Universiteit Delft
Prof.dr. I.M. Richardson, Technische Universiteit Delft
Dr. L. Zhao, Materials innovation institute, Delft, the Netherlands
Dr. L. Zhao heeft als begeleider in belangrijke mate aan de totstandkoming van het proefschrift bijgedragen.
Keywords: hyper-eutectoid steels, multiphase steels, phase transformations, thermodynamics, microstructure evolution
ISBN: 978-90-77172-414
Copyright © 2008 by N.V. Luzginova
All right reserved. No part of the material protected by this copy right notice may be reproduced or utilized in any form or by any means, electronical or mechanical, including photocopying, recording or by any information storage and retrieval system, without written permission from the author.
Contents
1
General Introduction ... 1
1.1 Hyper-eutectoid steel... 2
1.2 Thermal treatment of hyper-eutectoid steel ... 3
1.3 Outline of this thesis... 6
1.4 References ... 9
2
Experimental... 11
2.1 Materials... 12
2.2 Dilatometry and Heat Treatment... 16
2.3 Optical Metallography and Microhardness Measurements ... 18
2.4 Electron Microscopy... 20
2.5 X-Ray Diffraction ... 21
2.6 Vibrating Sample Magnetometry... 23
2.7 DICTRA simulations... 24
2.8 References ... 26
3
Experimental Characterization of Fe-C-Cr Steel ... 27
3.1 Introduction... 28
3.2 Transformation kinetics and morphology of lower bainite... 29
3.3 Retained austenite ... 35
3.4 Thermal stability of retained austenite... 37
3.5 Conclusions ... 44
3.6 References ... 45
4
Modeling of Lower Bainite Formation in Fe–C–Cr Steel... 47
4.1 Introduction... 48
4.2 Reconstructive model... 50
4.3.1 Nucleation rate... 60
4.3.2 Overall kinetics for bainite formation ... 61
4.4 Discussion ... 63
4.4.1 K
1, K
2, and t
0parameters in Quidort and Brechet’s
reconstructive model... 64
4.4.2 λ and κ parameters in Van Bohemen and Sietsma’s
displacive model... 67
4.5 Conclusions ... 74
4.6 References ... 76
5
Effect of Alloying Elements on the Spheroidization Process in
Hyper-eutectoid Steels... 79
5.1 Introduction... 80
5.2 Design of the spheroidization heat treatment... 82
5.3 Microstructural observations and hardness measurement... 83
5.4 Quantitative analysis of spheroidized microstructures... 89
5.5 Discussion ... 91
5.6 Conclusions ... 95
5.7 References ... 98
6
Effect of Alloying Elements on Cementite Dissolution in
Hyper-eutectoid Steels... 99
6.1 Introduction... 100
6.2 DICTRA simulations... 101
6.3 Experimental observations... 111
6.4 Conclusions ... 119
6.5 References ... 120
7
Effect of Alloying Elements on Lower Bainite Formation in
Hyper-eutectoid Steels ... 121
7.1 Introduction... 122
7.2 Experimental Results ... 123
7.3 Modelling Results... 132
7.4 Conclusions ... 141
7.5 References ... 142
Summary... 143
Samenvatting... 149
Publications ... 155
Acknowledgements... 157
Curriculum Vitae... 159
1 General Introduction
In this chapter, the description of hyper-eutectoid steel in general and SAE 52100 steel in particular is given in Section 1.1. In Section 1.2 the thermal treatments are discussed in order to obtain the desired microstructure and properties of hyper-eutectoid steels. The scope of the thesis is presented in Section 1.3.
1.1 Hyper-eutectoid steel
Hyper-eutectoid steel is a steel with a carbon concentration greater than the eutectoid content (Figure 1.1), which will strongly depend on the concentrations of other alloying elements. It should be noted that Figure 1.1 shows a quasi-binary Fe–C phase diagram calculated for steel with 1.5 wt.% Cr, the A1–temperature line splits
up, depending on the other alloying elements. Upon cooling of hyper-eutectoid steel from the fully austenitic region into the austenite and cementite (γ+Fe3C) region, first
cementite will start to nucleate and grow along the austenite grain boundaries. This cementite is called pro-eutectoid cementite, as it forms before the eutectoid reaction takes place. Upon further cooling more cementite will be formed and the composition of the austenite will reach the eutectoid composition, and as the temperature is lowered through the eutectoid temperature, all remaining austenite of eutectoid composition will be transformed into pearlite. Pearlite with pro-eutectoid cementite along the prior austenite grain boundaries is usually the initial microstructure in hyper-eutectoid steels after casting and forming.
800 900 1000 1100 1200 1300 0.0 0.5 1.0 1.5 2.0 Carbon content, wt.% Temper at ure, K γ γ γ γ γ γ + α γ γ γ γ Fe3C Pearlite α + Fe3C Proeutectoid Fe 3C Eutectoid Fe3C Eutectoid composition A1
Figure 1.1. Schematic representation of the microstructures for a quasi-binary iron–carbon alloy of hyper-eutectoid composition (Fe – 1.0 wt.% C – 1.5 wt.% Cr), as it is cooled from the fully austenitic region to below the eutectoid temperature [1].
The popular bearing steel SAE 52100 (1.01 wt.% C - 1.36 wt.% Cr - 0.32 wt.% Mn - 0.25 wt.% Si) is one example of hyper-eutectoid steel. SAE 52100 steel possesses many characteristics favorable to the production of tools, dies, precision components, through-hardened bearings [2, 3], and it is relatively low in cost. For these applications, SAE 52100 steel has excellent fatigue properties [4-6], high compressive/tensile strength, high hardenability and hardness as well as a low level of both solid and gas inclusions. Such excellent properties cannot be achieved with the pearlitic microstructure obtained directly after casting and forming. In order to develop the desired properties in this steel a special heat treatment is required. In the next section the description of the entire heat treatment process for hyper-eutectoid steel is presented.
1.2 Thermal treatment of hyper-eutectoid steel
As discussed in Section 1.1 the initial microstructure of hyper-eutectoid steel consists of pearlite and pro-eutectoid cementite along the prior austenite grain boundary (Figure 1.1). Such a pearlitic microstructure has a poor machinability, which is considered to be a disadvantage for industrial applications. In order to overcome this problem and to reduce the hardness of the material before machining and further hardening treatment, a spheroidization treatment of cementite particles,
i.e. soft annealing should be performed (Figure 1.2, I). Two types of spheroidization
treatment are often used:
(i) Subcritical spheroidization below the A1–temperature (Figure 1.1), which is
mainly applied for hypo-eutectoid steels. During subcritical annealing of steels with an initial pearlite structure, the cementite lamellae in pearlite break up into spheroids driven by the reduction in surface energy [7, 8].
(ii) Intercritical spheroidization above the A1–temperature (Figure 1.1), which is
mainly applied for hyper-eutectoid steels in order to spheroidize and to partially dissolve the grain boundary cementite [9-12]. During intercritical spheroidization an
incomplete dissolution of cementite occurs at the austenitisation annealing temperature and upon slow cooling austenite with fine cementite particles transforms into a mixture of ferrite and spheroidized cementite by the Divorced Eutectoid Transformation (DET) reaction.
In the present work the main focus is on the intercritical spheroidization treatment. The spheroidization annealing of hyper-eutectoid steels is of significant interest not only for industrial application but also for the new insight that can be gained on the spheroidization mechanism. Although various studies address the principles of the intercritical spheroidization and successful empirical recipes have been developed for certain alloys, many aspects regarding the mechanism of intercritical spheroidization, the controlling parameters and the effect of alloying elements remain uncertain.
Figure 1.2. Schematic representation of the heat treatment processes for hyper-eutectoid steels.
The final desired properties of hyper-eutectoid steels are obtained after intercritical austenitisation (Figure 1.2, II), when not all of the spheroidized cementite
is dissolved, followed either by an isothermal formation of lower bainite (Figure 1.2, III) or by quenching to room temperature to form martensite. The initial microstructure before austenitisation consists of ferrite and spheroidized cementite obtained by the intercritical spheroidization process discussed earlier. The presence of incompletely dissolved cementite after austenitisation has a beneficial effect on the rolling contact fatigue life of bearing steels [13]. By properly controlling the amount of dissolved cementite the required composition of austenite can be obtained in order to achieve a high hardness and yield strength of the product. Furthermore, the dissolution process can significantly influence the subsequent bainite hardening [14]. For instance, the presence of a cementite volume fraction of 0.03–0.05 can prevent an excessive austenite grain growth during austenitisation. It should be noted that cementite dissolution in austenite has been extensively studied in the literature, especially for the bearing SAE 52100 steel with 1.5 wt.% Cr [13-20]. However, the influence of different alloying elements like cobalt and aluminum, as well as the effect of different chromium contents on the austenitisation parameters, has not been given much attention.
Many applications of SAE 52100 bearing steel require that the steel is heat-treated to obtain a lower-bainitic microstructure as the final product (Figure 1.2, III). By creating the lower bainite microstructure in this steel the advantageous mechanical properties, such as excellent fatigue life, high strength and hardness, as well as greater toughness than in fully martensitic steels can be achieved. It should be noted that in hyper-eutectoid steels the lower-bainitic microstructure can only be produced isothermally, since during continuous cooling of hyper-eutectoid steels other transformation products can be formed like upper bainite and pearlite, the presence of which will be a disadvantage for the mechanical properties of the hardened components of bearings. Thus, the main disadvantage associated with the production of steels with a lower bainite microstructure is that it is a time-consuming process. In order to reduce the production time without loss of the desired
mechanical properties, a better understanding of bainite formation and a suitable model of its kinetics are required.
The additions of alloying elements might significantly affect the bainite formation kinetics in steels. For instance, Cr, the main substitutional alloying element in SAE 52100 steel is a strong carbide-forming element, which can therefore be expected to have a strong influence on the lower bainite formation. It is shown in the literature [21] that even a small amount of chromium retards the reaction of austenite decomposition into bainite compared with chromium-free steels. Beside this negative retardation effect Cr has many positive effects in terms of hardenability, spheroidization [9], and the resistance to decarburization [2]. The additions of other alloying elements might be beneficial in order to accelerate the bainite formation in Cr-containing steels by the increase of the free energy change accompanying the austenite to ferrite transformation [22, 23].
1.3 Outline of this thesis
In this work the attention has been focused on the microstructure evolution and the phase transformation kinetics in hyper-eutectoid steels, in a commercial SAE 52100 bearing steel and 7 model alloys with different concentrations of chromium, cobalt and aluminum, but with the carbon content of model alloys being the same as of a commercial SAE 52100 steel (1 wt.% C).
Chapter 2 describes the experimental equipment and the simulation software extensively used throughout this thesis to study microstructure and transformation kinetics in hyper-eutectoid steels. An overview of the material compositions studied in this thesis is presented and a brief introduction of the experimental techniques is given. In the present work dilatometry, optical metallography and X-ray diffraction analysis were used to follow the phase transformation during different heat treatments. A thermo-magnetic technique was used to study the evolution and the thermal stability of retained austenite, and electron microscopy to reveal the details
of the microstructural morphologies. The application of the DICTRA software [24] to simulate the cementite dissolution kinetics during the austenitisation and spheroidization process is discussed.
Chapter 3 focuses on the experimental characterization of the lower bainitic microstructure of hyper-eutectoid SAE 52100 steel (1.01 wt.% C - 1.36 wt.% Cr - 0.32 wt.% Mn - 0.25 wt.% Si). The microstructure and the kinetics of isothermal formation of lower bainite, and the evolution and thermal stability of retained austenite in SAE 52100 steel, is investigated using dilatometry, optical microscopy, electron microscopy, X-ray diffraction and a thermo-magnetic technique.
Chapter 4 presents two different physical approaches to model the formation of lower bainite in high carbon and chromium SAE 52100 steel. In the first model, a reconstructive approach is used. Nucleation of bainitic laths is considered in the general framework of the classical nucleation theory and a diffusion-controlled growth model is used. In the second model, displacive growth of bainitic ferrite is assumed, where the change in the bainite volume fraction is governed by the nucleation rate. Model calculations are compared to the experimentally obtained lower bainite fractions for SAE 52100 steel. The advantages and disadvantages of the proposed models are discussed, and the appropriate model is chosen for the description of the overall isothermal lower bainite formation in high-carbon steels.
In Chapter 5 the effect of alloying elements on the cementite spheroidization process in hyper-eutectoid steels is investigated experimentally and theoretically. A spheroidized structure in high carbon steel is obtained using an intercritical spheroidization process, after which during the slow cooling of austenite with fine cementite particles a divorced eutectoid transformation (DET) reaction occurs. A criterion for the occurrence of the DET reaction, as opposed to pearlite formation, is defined for the Cr-containing steels, and a reasonable agreement is found between the criterion and the experimental results. This DET criterion is further extended for steels with other alloying elements, like Co, Al and Mn.
Chapter 6 shows the effect of Cr, Co and Al as alloying elements on the cementite dissolution during austenitisation in hyper-eutectoid steels with 1 wt.% C. The dissolution of cementite is investigated with dilatometry, optical microscopy and scanning electron microscopy. The austenitisation process parameters are chosen from the results of DICTRA simulations, where the experimentally observed initial size of cementite particles is taken into account. A comparison between results calculated with DICTRA and experimental results for the kinetics of cementite dissolution in hyper-eutectoid steels is discussed.
Finally, in Chapter 7 an investigation of the effect of alloying elements on lower bainite formation in hyper-eutectoid steels is performed using dilatometry, scanning electron microscopy and X-ray diffraction measurements. The displacive model is successfully applied to describe the kinetics of lower bainite formation, including the effects of both isothermal transformation temperatures and alloying elements.
1.4 References
1. W.D. Callister: Materials Science and Engineering: An Introduction, 7th edition,
Wiley, New York, 2006, p. 298.
2. J.M. Beswick: Met. Trans. A, 1987, vol. 18A, pp. 1897–1906. 3. Y.B. Gou, C.R. Liu: J. Manuf. Sci. Eng., 2002, vol. 124, pp. 1–9.
4. G.E. Hollox, R.A. Hobbs, J.M. Hampshire: Wear, 1981, vol. 68, pp. 229–240. 5. F.C. Akbasoglu, D.V. Edmonds, Met. Trans. A, 1990, vol. 21A, pp. 889–893.
6. J.M. Hampshire, J.V. Nash, G.E. Hollox, in: J.J.C. Hoo (Ed.), Rolling Contact Fatigue
Testing of Bearing Steels, ASTM (American Society for Testing Materials),
Philadelphia, 1982, pp. 47–66.
7. S. Chattopadhyay, C.M. Sellars: Metallography, 1977, vol. 10, pp. 89–105.
8. D. Hernandez–Silva, R.D. Morales, J.G. Cabanas–Moreno: ISIJ Int., 1992, vol. 32, pp. 1297–1305.
9. J.D. Verhoeven: Met. Mater. Trans. A, 2000, vol. 31A, pp. 2431–2438.
10. G.M. Michal, M.D. Novak: Austenite Formation and Decomposition, eds. E.B. Damm, M.J. Merwin, Minerals, Metals and Materials Society, Warrendale, PA, 2003, pp. 397–413.
11. W. Hewitt: Heat Treatment of Metals, 1982, vol. 3, pp. 56–62.
12. T. Oyama, O.D. Sherby, J. Wadsworth, B. Walser: Scripta Met., 1984, vol. 18, pp. 799–804.
13. C.A. Stickels: Met. Trans. A, 1974, vol. 5, pp. 865–874.
14. L. Zhao, F.J. Vermolen, A. Wauthier, J. Sietsma: Met. Mat. Trans. A, 2006, vol. 37, pp. 1841–1850.
15. J.M. Beswick: Met.Trans. A, 1978, vol. 18A, pp. 1897–1901. 16. J.M. Beswick: Met.Trans. A, 1984, vol. 15A, pp. 299–306. 17. K. Nilsson: Trans. ISIJ, 1971, vol. 11, pp. 149–156.
18. J. Epp, H. Surm, O. Kessler, T. Hirsch: Acta Mater., 2007, vol. 55, pp. 5959–5967. 19. E.L. Brown, G. Krauss: Met. Trans. A, 1986, vol. 17A, pp. 31–36.
20. C.A. Stickels, A.M. Janotik: Met. Trans. A, 1980, vol. 11A, pp. 467–473.
21. E.S. Davenport, E.S. Bain, N.J. Kearny: Trans. Met. Soc. AIME, 1930, vol. 90, pp. 117–154.
22. C. Garcia-Mateo, F.G. Caballero, H.K.D.H. Bhadeshia: ISIJ Int., 2002, vol. 43, pp. 1821–1825.
23. M. De Meyer, D. Vanderschueren, B.C. De Cooman: ISIJ Int., 1999, vol. 39, pp. 813–822.
2 Experimental
Chapter 2 gives a description of the experimental equipment and the simulation software extensively used throughout this thesis to study microstructure and transformation kinetics in hypereutectoid steels. An overview of the materials studied in this thesis is presented in Section 2.1. In Sections 2.2 through 2.6 a brief introduction is given to the experimental techniques. Dilatometry (Section 2.2), Optical Metallography (Section 2.3) and X-ray Diffraction analysis (Section 2.5) were used to follow the phase transformation progress during different heat treatments. Electron Microscopy (Section 2.4) was used to investigate the microstructure morphologies, and Vibrating Sample Magnetometry (Section 2.6) to study the evolution and the thermo-stability of retained austenite. The background on the DICTRA calculations used to simulate cementite dissolution during the austenitisation process is discussed in section 2.7.
2.1 Materials
Materials studied in the present work were a commercial SAE 52100 steel (as a base material) and 7 model high-carbon alloys. The composition of SAE (AISI) 52100 steel, in the literature also known as 100Cr6 (Germany), GCr15 (China), SUJ-2 (Japan), EN-31 (UK), ШХ15 (Russia), is listed in Table 2.1. The as-received SAE 52100 steel had a microstructure consisting of a ferrite volume fraction of 0.85 and a spheroidized-cementite volume fraction of 0.15. The above mentioned spheroidized microstructure was obtained (after casting and cold forming) by a soft-annealing treatment, which was austenitisation at 1093 K for one hour, slow cooling to 963 K at a rate of 10 K/hour, and air cooling to room temperature (Figure 2.1 (a)).
Table 2.1. Alloy composition of SAE 52100 steel in wt.%.
Fe C Si Mn Cr Ni Cu Mo Al S P
bal. 1.01 0.25 0.32 1.36 0.16 0.12 0.04 0.03 <0.02 <0.01 Besides SAE 52100 steel 7 model high-carbon alloys were studied in the present thesis. The model steel compositions are listed in Table 2.2, where 1.5Cr steel has a similar composition as the commercial SAE 52100 steel (Table 2.1). All alloys were manufactured at the Corrosion and Metals Research Institute, Sweden by chill casting under inert conditions using high–purity alloying metals, resulting in ingot dimensions of 40×40×160 mm3. After casting, a chemical analysis of each ingot was
made and the ingots were further treated as shown in Figure 2.1 (b). Hot isostatic pressing (HIP) was performed at a temperature of 1420 K and under a hydrostatic pressure of 100 MPa for 4 hours, followed by furnace cooling to 1070 K at an average rate of 12 K/min and cooling to room temperature at an average rate of 35 K/min, to obtain a pore-free and homogenized structure. The microstructure of all model alloys after HIPing consisted of pearlite and pro-eutectoid cementite at the prior austenite grain boundaries.
300 600 900 1200 0 2 4 6 8 10 12 14 Time, hours Tempera tur e, K 10 K/hour Austenitization for 1 hour (a) 300 600 900 1200 1500 0 1 2 3 4 Time, hours Tempera tur e, K 5 12 K/min Austenitization and HIP
for 4 hours
35 K/min 1420 K
1070 K
(b)
Figure 2.1. (a) – a soft-annealing treatment (cementite spheroidization) of SAE 52100 steel. (b) – the HIP treatment of the model high-carbon alloys after casting.
All thermodynamic calculations for the investigated alloys were performed using the ThermoCalc software (TCCR version, TCFE2 database) [1]. Figure 2.2 presents the quasi-binary Fe-C phase diagrams for steels with different Cr (Figure 2.2 (a)), Co (Figure 2.2 (b)) and Al (Figure 2.2 (c)) contents. Three-phase regions (ferrite, austenite and cementite), which split A1 into two lines, are observed in the phase
diagrams for all steels. In this work lower and upper A1–temperatures are presented
950 1000 1050 1100 1150 0 0.2 0.4 0.6 0.8 1 1.2 Carbon content, wt.% Te m p era tur e, K A1 A '1 0.5Cr 3.5Cr 2.5Cr 1.5Cr α+θ γ+θ γ+α γ (a) 950 1000 1050 1100 1150 0 0.2 0.4 0.6 0.8 1 1.2 Carbon content, wt.% Te m p era tur e, K A1 A '1 1Co - 1.5Cr 2Co - 1.5Cr α+θ γ+θ γ+α γ (b) 950 1000 1050 1100 1150 1200 1250 0 0.5 1 1.5 2 2.5 3 Carbon content, wt.% Temper at ur e, K A1 A '1 1Al - 1Co - 1.5Cr γ γ+θ γ+α+θ α+θ γ+α (c)
Figure 2.2. Quasi-binary Fe-C phase diagrams for the investigated alloys (α - ferrite, γ – austenite, θ – cementite). (a) – 0.5Cr, 1.5Cr, 2.5Cr, 3.5Cr steels, (b) – 1Co-1.5Cr and 2Co-1.5Cr steels, (c) – 1Al-1Co-2Co-1.5Cr steel.
Characteristic temperatures and eutectoid compositions from the phase diagram are listed in Table 2.3 and it can be seen that the addition of alloying elements changes the phase equilibria for these steels significantly. The effect of alloying elements on the phase equilibria will be discussed in detail in the Chapters 5 and 6.
Table 2.2. Alloy composition of model high carbon steels in wt.%.
Steel name Fe C Si Mn Cr Co Al 0.5Cr bal. 1.04 0.25 0.30 0.53 -- -- 1.5Cr bal. 1.05 0.25 0.34 1.44 -- -- 2.5Cr bal. 1.04 0.27 0.31 2.39 -- -- 3.5Cr bal. 1.02 0.27 0.30 3.38 -- -- 1Co-1.5Cr bal. 1.05 0.26 0.32 1.36 1.02 -- 2Co-1.5Cr bal. 1.04 0.25 0.31 1.36 2.05 -- 1Al-1Co-1.5Cr bal. 1.06 0.25 0.31 1.38 0.98 1.04
Table 2.3. Characteristic temperatures and eutectoid compositions.
Steel name A1, K A’1, K Acm, K weutc , wt.%
0.5Cr 1005 1010 1135 0.68 1.5Cr 1010 1015 1170 0.57 2.5Cr 1015 1025 1200 0.46 3.5Cr 1020 1030 1220 0.37 1Co-1.5Cr 1015 1025 1160 0.58 2Co-1.5Cr 1020 1030 1170 0.59 1Al-1Co-1.5Cr 1055 1160 1280 0.68
2.2 Dilatometry and Heat Treatment
A Bähr 805 dilatometer was used to study the length change (dilatation) of the specimen during a heat treatment. The monitoring of the dilatation is a commonly used method to study phase transformations in steels. The cylindrical massive specimens for dilatometry experiments were prepared with a size of 10 mm in length and 5 mm in diameter. A specimen is placed in the dilatometer between two quartz rods with a thermocouple spot-welded in the middle of specimen in order to control the temperature. In the dilatometer the specimen is heated by induction. Sufficiently high cooling rates (up to 85 Ks-1) are obtained by helium gas quenching. A
description of the phase transformations investigated in the present work is presented in Chapter 1.
In the experiments to study the spheroidization process specimens were heated to the austenitisation temperature Taus = 1040–1110 K at a rate of 120 K/min (Figure
2.3 (a)). After holding for 2 hours at the austenitisation temperature the specimens were cooled at a cooling rate of 15 K/hour to 10 K above the A′1 temperature,
followed by cooling at 5 K/hour to 955 K, and further air cooling to room temperature. The use of the dilatometer enables the recording of the change in length and thus the progress of the phase transformations during heating, cooling and isothermal holding can be followed (Figure 2.3 (b)). In order to perform the spheroidization heat treatment of a large number of specimens a box furnace was used. To prevent oxidation and decarburization of the material the specimens were placed in quartz tubes filled with helium and sealed.
In the experiments to study the kinetics of cementite dissolution during austenitisation, specimens were heated to the higher austenitisation temperature
Taus = 1090–1170 K at a rate of 120 K/min and austenitized for different times
950 1000 1050 0 5 10 15 Time, hours Tempera tur e, K A'1 Taus=1040 - 1110 K 15 K/hour 5 K/hour 2 K/s (a) 60 70 80 90 100 110 120 0 5 10 Time, hours Cha n ge in Le n gth, μm 15 80 90 100 110 120 0.05 0.10 0.15 Ferrite transformation Austenitization (b)
Figure 2.3. (a) – an example of a spheroidization heat treatment and (b) – the corresponding change in length.
In the experiments to study the microstructural evolution and the kinetics of lower bainite formation (Figure 2.4, II) specimens after austenitisation at Taus = 1090–
1170 K for 30 min were quenched to the bainite holding temperatures TLB = 480–
570 K, and annealed for different times (0–120 min), followed by quenching to room temperature. The dilatation was recorded both during austenitisation and during the lower bainite formation heat treatments (Figure 2.4 (b)).
All specimens for the further analysis described in Sections 2.3–2.6 were cut from the dilatometry specimens.
300 500 700 900 1100 1300 0 20 40 60 80 100 120 140 160 Time, min Temper at ure, K TLB= 483 - 573 K Taus= 1093 - 1173 K (I) (II) (a) -20 20 60 100 140 0 20 40 60 80 100 120 140 160 Time, min Chang e in Leng th, μm (I) (II) (b)
Figure 2.4. (a) – an example of (I) the austenitisation heat treatment followed by (II) the formation of lower bainite. Dashed lines show the interrupt quenching after partial transformation. (b) – the corresponding change in length.
2.3 Optical Metallography and Microhardness Measurements
A metallographic examination of each specimen was made with optical microscopy. The microstructures were quantitatively analyzed using AnalySIS Image Processing Software. To obtain a good contrast for the optical analysis of the spheroidized microstructure and the microstructure after partial cementite dissolution during the austenitisation process, the specimens were pre-etched in 5%Nital followed by Klemm’s tint etching (50 ml of saturated aqueous sodium thiosulfate solution and 1 g of sodium disulfide). After etching cementite appears in white, and ferrite and martensite appear in black or dark brown, enabling a reliable setting for the threshold value for further quantitative analysis of the size and the volume fraction of cementite (Figure 2.5).
(a) 0.E+00 5.E+04 1.E+05 2.E+05 0 30 60 90 120 150 180 210 240 Grey value Num b er of pi xel s Black White Cementite Ferrite Threshold (b)
Figure 2.5. (a) – an example of a spheroidized microstructure after Klemm’s etching, (b) – grey value distribution of the microphotograph of a spheroidized microstructure.
The prior austenite grain size was determined after etching in a saturated picric acid solution with additions of HCl and the “Teepol” wetting agent [2]. The analysis of the austenite grain size was based on the concept of the equivalent diameter, where the equivalent diameter of an austenite grain is the diameter of the circle that contains the same area as the austenite grain.
To observe the lower bainite microstructures, specimens were etched with 2% Nital for 10 s; after the etching retained austenite and martensite were unaffected (light), whereas lower bainite appeared black.
The Vickers hardness of all specimens was measured using the Buehler automatic microhardness testing system OmniMent MHT 7.0 Rev.1 with a load of
1 kg. Every hardness value presented in this work is the average of at least five measurements.
2.4 Electron Microscopy
Scanning Electron Microscopy (SEM) measurements of SAE 52100 steel were performed at TU Delft with a JSM-6500F Field Emission Scanning Electron Microscope to characterize the overall morphology. Energy Dispersive X-ray Spectroscopy (EDS) was used for chemical analysis of cementite particles. All micrographs were obtained using a beam of 15 keV electrons. The microstructures were examined after etching for 10 s in 2% Nital solution.
Scanning Electron Microscopy (SEM) measurements of the 7 model hypereutectoid steels (Table 2.2) were performed at Corus with a Zeiss Ultra 55 Field Emission Gun Scanning Electron Microscope to characterize the overall morphology. The microscope was equipped with an in-lens electron optic system, which allows an optimal recovery of secondary electrons and results in enhanced resolution. Specimens were hot mounted in Polyfast resin, which is electrically conductive with low emission in the vacuum chamber during examination. All micrographs were obtained using a beam of 15 keV electrons. The microstructure details were examined after etching for 5 s in 1% Nital solution.
Transmission Electron Microscopy (TEM) measurements were performed at TU Delft using Philips CM30T microscope operating at 300KV. The bright field (BF) and the selected area electron diffraction (SAED) techniques were used in order to characterize the microstructure of lower bainite in SAE 52100 steel. Thin foils were prepared for TEM study as follows: (1) specimen was cut by a diamante saw in slices of approximately 1 mm thickness, (2) cut foils were ground and polished to 0.1 mm thickness, (3) circular pieces of the foils were cut and placed on the TEM copper ring and the final thinning was performed by conventional Ion Milling technique using Gatan 691 Precision Ion Polishing System (PIPS).
2.5 X-Ray Diffraction
X-ray diffraction measurements were carried out at room temperature on a Bruker D8-Advance diffractometer equipped with a Vantec Position Sensitive Detector (PSD). CoKα radiation was used and 2θ scans were performed with step
time of 0.6 s and step size of 0.025º. 2θ values were ranged from 40° to 130°, containing four ferrite, four austenite and a set of cementite peaks. Typical XRD scans are shown in Figure 2.6.
The EVA software suite (DIFFRACplus Evaluation Package, version 2.2) was
used to analyze the diffraction peaks. The volume fraction of retained austenite was determined from the integrated intensities of austenite and ferrite peaks using the method described in [3] with:
(
)
(
)
∑
∑
∑
+ + = α γ γ θ γ γ γ α α α γ γ γ γ n n hkl hkl hkl hkl n hkl hkl f R I n R I n R I n f R 1 1 1 1 1 1 (2.1)where fθ is the volume fraction of all carbides in the material, I and γhkl are the
integrated intensities of austenite and ferrite peaks, respectively; nγ and nα are the numbers of {hkl} lines for which the integrated intensities have been measured;
and are theoretical intensities [
αhkl I γhkl R αhkl R 3] presented in Table 2.4.
Table 2.4. Theoretical line intensities (R-values) for the ferrite and austenite phases in
steel for Co radiation (λCo=1.79021Å) [3].
{hkl}phase {111}γ {200}γ {220}γ {311}γ {110}α {200}α {211}α {220}α
0 50 100 40 60 80 100 120 2θ (degrees) Intensity (CPS) {110}α {200}α {211}α {220}α θ θ (a) 0 50 100 40 60 80 100 120 2θ (degrees) Intensity (CPS) {1 11 }γ {110}α {20 0}γ {2 00 }α {31 1}γ {22 0}γ {21 1}α {22 0}α θ θ (b)
Figure 2.6. Typical diffraction spectra of SAE 52100 steel. (a) – a soft-annealed specimen (only diffraction peaks of ferrite (α) and cementite (θ) are observed), (b) – a specimen annealed for 45 minutes at 503 K and quenched to room temperature (diffraction peaks of ferrite (α), cementite (θ) and austenite (γ) are present). The intensity is shown in counts per second (CPS).
The austenite lattice parameters (aγ) were calculated from the {311}γ austenite
diffraction peak [4]. The carbon content of austenite (w ) in wt.% was calculated γc
γ γ 3.555 0.44 c
a = + w (Å), (2.2)
which is considered to most reliably describe the variation of the retained austenite parameter with carbon content [5]. The room temperature lattice parameter of 3.555 Å is given for pure Fe (austenite).
2.6 Vibrating Sample Magnetometry
Cylindrical specimens for the magnetization measurements with a size of 2 mm in length and 2 mm in diameter were machined from dilatometry specimens using an electro-discharging machine (EDM). All magnetic measurements were performed in a LakeShore 7307 Vibrating Sample Magnetometer (VSM). Before experiments the VSM was calibrated with a standard NIST nickel specimen. A typical magnetization curve at room temperature for as-received SAE 52100 steel is presented in Figure 2.7.
To obtain the saturation magnetization values, Ms, the high-field part of the
magnetization curve was fitted to the equation described in [6]:
2 1 s a b M M H H ⎛ ⎞ = ⎜ − − ⎝ ⎠⎟ , (2.3)
where M is the magnetization at the applied magnetic field H, Ms the saturation
magnetization, a and b the fitting parameters.
To study the thermal stability of retained austenite the thermal cycles from 300 K to 1173 K (high temperature magnetic experiment) and from 300 K to 10 K (low temperature magnetic experiment) under a constant magnetic field of 0.79×106 A/m
(1 Tesla) were performed. Stepwise heating and cooling were carried out and the magnetization was measured one minute after the set temperature was reached. The high temperature magnetic experiments were performed in a High Temperature
Oven (Model 73034) at a heating rate of 5 K/min and a cooling rate of 10 K/min. The low temperature magnetic experiments were performed in a Closed Cycle Refrigerator (Model 73018) at cooling and heating rates of 10 K/min.
-1.5 -0.5 0.5 1.5 -1.5 -0.5 0.5 1.5 Field, A/m x 10-6 Ma gnetiza tion, A/m x 10 -6
Figure 2.7. Magnetization curve for as-received SAE 52100 steel.
2.7 DICTRA simulations
During the austenitisation processes described in Section 2.2 of this chapter cementite dissolution in austenite occurs. In order to simulate the cementite dissolution the DICTRA [1, 7] software package was used. In the DICTRA simulations the cementite particle was spherical and the initial compositions of cementite and austenite were inherited from the spheroidized cementite and ferrite. It was assumed that during heating ferrite rapidly transforms into austenite without any cementite being dissolved. This assumption is approximative, but it has been applied before in the literature [8] and is considered to be realistic in case of high carbon steels and sufficiently high austenitisation temperatures. It can be envisaged that the initial stage of the dissolution process does actually partly occur during the heating step. In DICTRA local equilibrium is assumed at the moving
austenite/cementite interface. It should be noted that in DICTRA the dissolution of only one spherical particle in a spherical volume element (austenite) is actually simulated (Figure 2.8), which has been shown to be a good assumption [9] to describe the change in volume fraction of cementite during the austenitisation process.
Figure 2.8. Schematical representation of the system used in the DICTRA simulations. θ is
cementite, γ is austenite, Rtotal is a total radius of the system, Rθ is the cementite particle
radius, Rγ is the austenite radius.
The initial particle size was set to the average particle size observed experimentally after the spheroidization heat treatment. The volume fractions of the phases were determined by assuming that only the substitutional elements contributed to the system volume and the initial state (the austenite size and the initial compositions of phases) was obtained from ThermoCalc at a temperature of
A1 – 10 K. The effect of the surface tension was neglected, because the particle
2.8 References
1. ThermoCalc & DICTRA software: http://www.thermocalc.com. 2. J. van der Sanden: SKF Report, 1978, № NL77M524.
3. C.F. Jatczak, J.A. Larson, S.W. Shin: Retained Austenite and Its Measurements by
X-Ray Diffraction, SAE Inc., Warrendale, 1980, p. 12.
4. B.D. Cullity: Elements of X-Ray Diffraction, Addison–Wesley Inc., Reading, 1978, p. 359.
5. N. Ridley, H. Stuart, L. Zwell: Trans. Met. Soc. AIME, 1969, vol. 245, pp. 1834– 1836.
6. J.W. Cahn, P. Haasen: Physical Metallurgy, Elsevier, Amsterdam, 1983, p. 2558. 7. A. Borgenstam, A. Engstöm, L. Höglund, J. Ägren: J. Phase Equilibria, 2000, vol. 21,
pp. 269–280.
8. M. Hillert, K. Nilsson, L-E. Törndahl: J. Iron and Steel Inst., 1971, vol. 209, pp. 49– 66.
9. Z-K. Liu, L. Höglund, B. Jönsson, J. Ägren: Met. Trans. A, 1991, vol. 22A, pp. 1745– 1752.
3 Experimental Characterization of
Fe-C-Cr Steel
In this chapter, the kinetics of isothermal formation of lower bainite and the evolution and thermal stability of retained austenite in SAE 52100 steel, 1.01C-1.36Cr-0.32Mn-0.25Si (wt.%), is investigated with dilatometry, optical microscopy, electron microscopy, X-ray diffraction and thermo-magnetic measurements. It is demonstrated that an increase in carbon content of austenite with bainitic holding time occurs, as a result of which the retention of a significant amount of austenite at room temperature takes place in SAE 52100 steel. The thermal stability of retained austenite is investigated. The temperature at which retained austenite starts to decompose to ferrite and carbides upon heating varies with bainitic holding time. The transformation of austenite to martensite during cooling to 10 K is found to be not complete, and a large amount of austenite remains untransformed.
3.1 Introduction
High-carbon and chromium steels, for instance SAE 52100 steel, are widely used in the bearing industry due to a combination of excellent fatigue life and high strength. Most of the applications of SAE 52100 bearing steel require that the steel is heat-treated to obtain either a martensitic or a lower-bainitic microstructure as the final product. The important advantage of the production of lower bainite over martensite is to gain greater toughness at the same hardness level. A hard martensitic or lower bainitic microstructure leads to a reasonable bearings lifetime due to high strength and high resistance to fatigue [1, 2]. When bearings are used in water-containing environments, however, the martensitic microstructure is sensitive to hydrogen-induced cracking, since twin boundaries in high carbon martensite are susceptible to hydrogen adsorption and crack nucleation. As an alternative, lower-bainite microstructures are produced, which give a similar fatigue life as a martensitic one under good lubricant conditions, whereas in water-containing environments fully lower-bainitic bearings show an increased fatigue life [2, 3]. When the service conditions of bearings are in the presence of water the lower bainitic microstructure is therefore often desired despite the fact that the manufacture of the steel with lower-bainitic microstructure is more expensive due to an extra time-consuming isothermal holding.
Bainitic or martensitic microstructures are often obtained together with retained austenite (γR). The amount and the morphology of γR in bearings is an important
issue. Austenite can often be retained in two forms: blocky and film types [4, 5]. The former is relatively unstable, which significantly influences the dimensional stability of the material [6]. The film austenite, located between bainitic ferrite plates, is very fine and stable due to the small dimensions and the carbon enrichment. Such a fine microstructure consisting of bainitic ferrite and film austenite gives an excellent combination of strength and toughness in high Si steels [5] and, moreover, the retained austenite islands can act as additional obstacles for crack propagation [7].
The austenite decomposition into lower bainite results in the formation of a microstructure consisting of ferritic plates and carbides within ferrite. To create a lower-bainitic microstructure the heat treatment of SAE 52100 steel consists of a partial austenitisation at temperatures of 1123–1143 K followed by the isothermal holding at lower-bainitic temperatures (483–573 K). It should be noted that in this steel the lower-bainitic microstructure cannot be produced by continuous cooling without the formation of other transformation products like upper bainite and pearlite, the presence of which will be a disadvantage for the mechanical properties of the hardened components of bearings. The mechanical properties of SAE 52100 steel after lower-bainitic treatment, including fatigue life, are well studied in the literature [1-3]. However, the details of the kinetics of lower bainite formation in hyper-eutectoid steel and the role of retained austenite, which is of essential importance for the use and application of the material, have not been reported in much detail. To obtain a better understanding of the lower bainite formation a dilatometry study together with metallographic observations, electron microscopy and X-ray diffraction analysis has been performed (Section 3.2). The evolution of retained austenite (Section 3.3) with bainitic holding time and its thermal stability (Section 3.4) upon heating and cooling has been studied using X-ray diffraction and thermo-magnetic measurements.
3.2 Transformation kinetics and morphology of lower bainite
In order to study the morphology and the evolution of the microstructure in SAE 52100 steel during lower bainite hardening, partial transformation was performed in the temperature range 480–570 K for different times, followed by quenching to room temperature. The experimental transformation–temperature–time diagram for the lower bainite formation in SAE 52100 steel is shown in Figure 3.1.450 500 550 600 0.1 1 10 100 1000 Time, min Temperature, K Mstart=470 K austenite + 5% cementite lower bainite + 5% cementite 2% 50% 90%
Figure 3.1. TTT diagram for lower bainite formation in SAE 52100 steel after austenitisation at 1133 K for 30 min.
It should be noted that the initial microstructure of SAE 52100 steel after austenitisation and before bainitic hardening consists of austenite and 5% spheroidized cementite [8]. The bainite fractions are calculated from dilatometry results using the lever rule and confirmed by optical microscopy. The martensite starting temperature is obtained from dilatometry results and equals 470 K, which is very close to the estimated value Mstart = 460 K using Andrew’s empirical equation
[9]. The reaction rate of lower bainite formation is strongly dependent on the isothermal holding temperature, at higher transformation temperatures the formation of lower bainite is faster than at lower temperatures.
A general view of a lower bainitic microstructure in hypereutectoid steel after partial transformation for 10 and 20 min at 533 K is shown in Figure 3.2, where black needles are lower bainite (LB), white particles are spheroidized cementite (θ) and the rest is a mixture of martensite (α’) and retained austenite (γR). Although the optical
microscopy can be successfully used to estimate the volume fraction of lower bainite, it is not able to reveal the individual plates of lower bainite, for which the
high-resolution scanning and transmission electron microscopy have been used in the present study.
(a) (b)
Figure 3.2. Optical metallography of the lower bainitic microstructure. (a) – 533 K for 10 min (total fraction of lower bainite is 0.15), (b) – 533 K for 20 min (total fraction of lower bainite is 0.75). Black needles are lower bainite (LB), small white particles correspond to the primary spheroidized cementite (θ), the grey matrix is formed by martensite (α’) and
retained austenite (γR).
Lower bainite formation is a decomposition of austenite into non-lamellar aggregates of ferrite and cementite. Due to the low transformation temperatures at which lower bainite is formed carbon cannot easily diffuse away into the austenite from the firstly formed supersaturated ferritic plate. Therefore, the only way to reduce the carbon content of bainitic ferrite is carbide precipitation. Usually the precipitation of either cementite (θ) or ε– and η–carbides within ferrite is expected in lower bainitic microstructures [10], which will depend on the composition of the steel and the transformation time and temperature. Figure 3.3 shows a typical complex microstructure of lower bainite in hardened SAE 52100 steel after partial transformation at different temperatures. Spherical particles in all micrographs are
spheroidized cementite particles, which are undissolved during intercritical austenitisation.
(a) (b)
(c) (b)
Figure 3.3. Lower bainitic microstructure revealed by scanning electron microscopy. (a) – 533 K for 5 min (lower bainite fraction is 0.02), (b) – 553 K for 10 min (lower bainite fraction is 0.50), (c) – 573 K for 10 min (lower bainite fraction is 0.80), (b) – 553 K for 20 min
(lower bainite fraction is 0.90). θ corresponds to the primary spheroidized cementite, θLB to
carbide in lower bainite, α’ to martensite and γR to retained austenite. Black arrows show
It can be seen that lower bainite first nucleates at the prior austenite grain boundary (Figure 3.3, black arrows) as a thin ferritic plate followed by the precipitation of carbides within ferrite. As transformation progresses (Figure 3.3 (b–d)) additional plates nucleate both at the sides of the original plates and intragranularly. It is shown that the isothermal transformation temperature in the investigated temperature range has no effect on the microstructural morphology. At all temperatures the formation of lower bainite is observed.
During quenching to room temperature after partial transformation to lower bainite not all of the residual austenite is transformed to martensite. The retained austenite is very difficult to be distinguished from martensite by means of SEM, but it can be seen that retained austenite films tend to be trapped between lower bainitic plates (Figure 3.3 (d)) at high transformed fractions, when no martensite is formed upon quenching. In the next sections the detailed analysis of the evolution and thermal stability of retained austenite in SAE 52100 steel will be presented in more detail.
In Figure 3.4 bright-field transmission electron microscopy shows the morphology of lower bainite at high magnification. It can be seen that ferritic plates with the thickness of 0.2–0.3 μm contain elongated rod–like carbide particles. In order to identify the structure of the carbides observed here, selected area diffraction was used and the measured spacings dhkl are compared with the theoretical values. The
analysis reveals that carbides in the ferritic plates have the orthorhombic structure of cementite, and precipitate in a single variant within a given ferritic plate. Cementite particles are inclined at about 60o to the longest side of the ferritic plate. The
precipitation of neither ε– nor η– carbides has been detected in SAE 52100 steel in the present study.
(b)
(a)
(c)
Fe3C
Figure 3.4. (a) – Bright field transmission electron microscopy, specimen transformed at 503 K for 45 min (lower bainite fraction is 0.70). (b) and (c) – Diffraction pattern taken
3.3 Retained austenite
The volume fraction of retained austenite in SAE 52100 steel after different bainite holding times/temperatures and quenching to room temperature is determined by XRD, and shown in Figure 3.5. It can be seen that the fraction of austenite retained at room temperature in SAE 52100 steel increases with the partial decomposition of residual austenite into lower bainite up to a lower bainite fraction of about 0.75. During the lower bainitic heat treatment the fraction of austenite gradually decreases, but it becomes more stable with respect to martensite formation during quenching to room temperature, which explains the maximum of retained austenite observed at a high volume fraction of lower bainite (Figure 3.5).
0.00 0.03 0.06 0.09 0.12 0.15 0.18 0.0 0.2 0.4 0.6 0.8 1.0
Volume fraction of lower bainite
Volume fraction of reta ine d au st eni te 483 K 503 K 533 K 553 K 573 K
Figure 3.5. Retained austenite volume fraction vs. volume fraction of lower bainite. The calculated error is due to counting statistics.
Usually the amount of retained austenite in the material will strongly depend on the reaction temperature and alloy composition. For instance, in high silicon steels (or steels with high concentrations of other alloying elements that retard the precipitation of cementite from austenite and ferrite) after upper bainite formation significantly more austenite is retained at ambient temperature than in low Si steels.
The retention of the significant amount of austenite at room temperature in SAE 52100 steel is associated with the effect of alloying elements, like chromium and carbon. It has been shown in [8] that the chromium distribution over the austenite grain in SAE 52100 steel is not homogeneous after cementite dissolution at 1133 K, which affects the austenite stability with respect to bainite formation.
The measurements of the lattice parameter of the retained austenite in SAE 52100 steel after partial formation of lower bainite at different temperatures indicate that there is a carbon enrichment of retained austenite (Figure 3.6).
0.7 0.8 0.9 1.0 1.1 1.2 0.0 0.2 0.4 0.6 0.8 1.0
Lower bainite volume fraction
Ca rbon co nte n t, wt.% 483 K 503 K 533 K 553 K 573 K
Figure 3.6. Carbon content of the retained austenite as a function of lower bainite volume fraction. The solid square represents the carbon content of austenite calculated from the mass balance, when the volume fraction of lower bainite is zero. Dashed line shows the expected carbon enrichment of austenite when no martensite forms on quenching.
It means that not all carbon precipitates in cementite within the ferritic plates, and carbon diffuses into the austenite without precipitation of carbides either. It should be noted that overestimation of the carbon content is likely to occur at low fractions of bainite due to the formation of martensite upon quenching. The magnitude of this effect can be estimated for a lower bainite fraction of zero. The austenite carbon content calculated from the mass balance between ferrite and austenite (Figure 3.6,
square) is lower than one obtained from the XRD analysis by more than 0.2 wt.%. The values for austenite lattice parameter at lower bainite formation temperatures is lower than the lattice parameter of retained austenite measured at room temperature, which is an indication that actually more significant carbon enrichment of austenite occurs during lower bainite formation than observed from the XRD analysis (Figure 3.6, dashed line).
3.4 Thermal stability of retained austenite
In the previous section the evolution of the volume fraction of retained austenite in SAE 52100 steel is investigated. A thermal stability of the retained austenite is also a very important issue for the dimensional stability of the bearings [6]. In this section the stability range of retained austenite upon cooling and heating is investigated using thermo-magnetic experiments. The thermal cycle from room temperature to 1173 K and back is performed for specimens with lower bainitic microstructure obtained after different holding times at 503 K. Figure 3.7 shows the temperature dependence of the magnetization at a constant magnetic field of 0.79×106 A/m for two specimens: (a) soft-annealed and (b) annealed for 45 minutes at
503 K. The microstructure of the soft-annealed specimen consists of ferrite and a fraction of 0.05 of spheroidized cementite, whereas the specimen annealed for 45 min at 503 K is partially transformed to lower bainite and contains a significant fraction of retained austenite. It is clear that the magnetization of the soft-annealed material (Figure 3.7 (a)) decreases with increasing temperature; this is due to the increase in thermal precession at the atomic level. The magnetization approaches zero at about 1050 K, above this temperature the ordered magnetic structure disappears, and the material becomes paramagnetic [11]. The observed hysteresis of cooling and heating curves occurs due to the phase transformation. As an example of typical M(T) behavior the temperature dependence of the magnetization for a material containing retained austenite, the specimen annealed for 45 min at 503 K is shown in Figure 3.7
(b) and (c). An increase of magnetization is observed during heating in the range 485 K–550 K, which is mainly due to the increase of the volume fraction of ferro-magnetic phases, i.e. decomposition of retained austenite. At temperatures above 480 K the cementite becomes paramagnetic, and the difference in magnetization during heating and cooling can be directly related to the volume fraction of retained austenite, although the possible formation of the intermediate carbides can affect the magnetization, as will be discussed below.
It should be noted that all specimens showing this increase in magnetization present an interesting phenomenon: an “overshoot” of the heating curve with respect to the cooling curve is observed at about 575 K (Figure 3.7 (b) and (c)), which is known to be the Curie temperature of ε-carbides [12]. The observed overshoot can likely be attributed to the formation of ε-carbides during the decomposition of retained austenite, as reported in [13]. The intermediate ε-carbide could be further transformed to ε’- (also known as η-) carbides and to cementite [13].
0.00 0.40 0.80 1.20 1.60 250 450 650 850 1050 Temperature, K Ma gnetiza tio n, A/m x 10 6 heating cooling (a)
0.00 0.40 0.80 1.20 1.60 250 450 650 850 1050 Temperature, K Mag n etization, A/m x 10 6 heating cooling (b) 1.10 1.20 1.30 1.40 475 575 675 775 Temperature, K Ma gnetiz ation, A /m x 1 0 6 Δ MC MH (c)
Figure 3.7. The temperature dependence (high temperature cycle (300 K → 1133 K → 300
K)) of magnetization under constant field 0.79 x 106 A/m for: (a) soft-annealed specimen
and (b) specimen transformed for 45 min at 503 K. (c) a close-up of (b) in order to illustrate the calculation of thermal stability of retained austenite from the high temperature magnetic measurements.
The decomposition process of γR is estimated using the following proposed
( ) 1 ( ) Δ R H γ C M T f M T = − + , (3.1)
for T < T
( )
Δ . M and H MC are magnetization values obtained at the sametemperature from the heating and cooling curves, respectively. Δ is the maximum overshoot between heating and cooling curves, for more detail see Figure 3.7 (c). It is assumed that the decomposition of austenite starts with the formation of ε-carbides,
and the change in magnetization value during heating is a result of the increase in volume fraction of ferro-magnetic phases, on the one hand, and the decrease in the saturation magnetization of ε-carbides, on the other hand. In addition to the volume fraction of γR at room temperature, these observations give important information
about the thermal stability of austenite. The temperature dependence of the volume fraction of retained austenite calculated using Equation (3.1) is shown in Figure 3.8 and summarized in Figure 3.9.
0.00 0.05 0.10 0.15 0.20 380 430 480 530 580 Temperature, K V olume fra ction of austenite 0 min 20 min 45 min 60 min 90 min 120 min b
Figure 3.8. Temperature dependence of retained austenite fraction obtained from the high temperature magnetic measurements.
One can see that the starting temperature of austenite decomposition varies with bainitic holding time, and is in the range from 480 K (45 min at 503 K) to 540 K (120 min at 503 K). The end temperature of austenite decomposition is 580 K for all holding times (Figure 3.9).
450 500 550 600 0 20 40 60 80 100 1 Time, min Temper at ur e, K 20 end start
Figure 3.9. Start (triangles) and end (circles) temperatures of retained austenite decomposition upon heating vs. bainite holding time.
Figure 3.10 shows the temperature dependence of the magnetization at a constant magnetic field of 0.79×106 A/m for a thermal cycle from room temperature
to 10 K and back for the soft-annealed material and the specimen annealed for 45 min at 503 K. It is shown that the magnetization of the soft-annealed material increases with decreasing temperature, which is due to the increase in saturation of the ferromagnetic phases. Cooling and heating curves are the same for the soft-annealed material because it contains only ferrite and cementite. For the heat-treated specimens an additional increase in magnetization is observed during cooling due to the increase in the volume fraction of ferromagnetic phases during cooling, i.e. retained austenite transforms to martensite. The maximum difference between cooling and heating curves is observed for the specimen that has been annealed for
45 min at 503 K. A significant amount of austenite (0.09) is found to be stable down to 10 K. 1.00 1.10 1.20 1.30 1.40 1.50 1.60 0 50 100 150 200 250 300 Temperature, K Ma gne tiz ation, A/m x 10 6 cooling heating soft-annealed 45 min at 503 K
Figure 3.10. The temperature dependence (low temperature cycle (300 K → 10 K → 300 K)
of magnetization under constant field of 0.79 x 106 A/m for soft-annealed material and
specimen annealed for 45 min at 503 K.
The essential information that can be obtained from the experiments described here is the thermal stability of austenite upon cooling (γ → α') and upon heating (γ → α + carbides). The austenite fraction as a function of temperature is calculated upon heating (Equation 3.1) and cooling [14] for the specimen with the highest amount of retained austenite at room temperature, and is shown in Figure 3.11. It is observed that all retained austenite is decomposed into ferrite and cementite (possibly with the formation of intermediate carbides) during heating, whereas during cooling to 10 K a large amount of austenite remains untransformed.
One can see from Figure 3.10 and Figure 3.11 that the transformation of retained austenite to martensite stops at around 110 K. This is likely an indication of the presence of two types of retained austenite with different thermal stability against martensite transformation. That is, stable film-type austenite is retained between
bainitic plates, while relatively less stable blocky austenite is retained between sheaves of bainite [4, 5]. The thermal stability of retained austenite is understood in terms of the carbon content [9] and the size of austenite grains [15, 16].
0.00 0.04 0.08 0.12 0.16 0.20 0 100 200 300 400 500 600 Temperature, K Vo lum e f ractio n of au st enit e γ --> α' γ --> α + carbides
Figure 3.11. Temperature dependence of volume fraction of retained austenite in specimen held for 45 min at 503 K.
Both types of retained austenite are enriched with carbon to some extent during the bainite formation at 503 K (Figure 3.6), since austenite contains about 0.25 wt.% of silicon acting as an inhibitor for carbide precipitation. As the blocky type of retained austenite has usually larger grain size than the film-type austenite [15], the carbon enrichment in film-type austenite is expected to be more than in the blocky type [4, 5]. Furthermore, the small grain size of film-type retained austenite leads to insufficient nucleation sites for martensite transformation and it thus also increases the stability of retained austenite significantly. The film-type retained austenite remains untransformed during cooling to 10 K. The temperature range in which the austenite is stable, i.e. neither transforms to ferrite and carbides during heating nor to
martensite during cooling, is determined to be 220 K, from 230 K to 450 K after a lower bainitic heat treatment at 503 K for 45 min.
3.5 Conclusions
In this chapter, the experimental characterization of the kinetics of lower bainite formation and its morphology in SAE 52100 bearing steel, as well as the evolution and thermal stability of retained austenite is investigated with optical and electron microscopy, X-ray diffraction and thermo-magnetic measurements. A significant fraction of austenite is retained in the material. It has been demonstrated that the maximum of retained austenite volume fraction occurs as a combination of the increasing carbon concentration in the austenite and of the decreasing volume fraction of the residual austenite at bainite formation temperatures.
R fγ
The thermal stability of austenite upon cooling and heating is investigated. The temperature range in which the austenite is stable, i.e. neither transforms to ferrite and carbides during heating nor to martensite during cooling, is from 230 K to 450 K for the specimens held for 45 min at 503 K. The temperature at which retained austenite starts to decompose to ferrite and carbides upon heating varies with the bainitic holding time and is in the range from 480 K (45 min at 503 K) to 540 K (120 min at 503 K). The end temperature of austenite decomposition is 580 K for all holding times. The transformation of austenite to martensite during cooling till 10 K is not complete, the transformation stops at 110 K, which is an indication of the presence of the very stable film-type austenite.
