Low Temperature Austenite Decomposition in Carbon Steels

Albin Stormvinter Doctoral Thesis

KTH Royal Institute of Technology School of Industrial Engineering and Management Department of Materials Science and Engineering Division of Physical Metallurgy SE-10044 Stockholm, Sweden

Stockholm 2012

Albin Stormvinter Low Temperature Austenite Decomposition in Carbon Steels

KTH Royal Institute of Technology, School of Industrial Engineering and Management, Department of Materials Science and Engineering, Division of Physical Metallurgy, SE-10044 Stockholm, Sweden

ISRN KTH/MSE--12/19--SE+METO/AVH ISBN 978-91-7501-449-4

Akademisk avhandling som med tillstånd av Kungliga Tekniska högskolan i Stockholm framlägges till offentlig granskning för avläggande av teknologie doktorsexamen 27 september 2012, kl. 10.00 i sal F3, Lindstedsvägen 26, Kungliga Tekniska högskolan, Stockholm.

Printed by Universitetsservice US-AB, Stockholm, Sweden

To Emma

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Abstract

Martensitic steels have become very important engineering materials in modern society. Crucial parts of everyday products are made of martensitic steels, from surgical needles and razor blades to car components and large-scale excavators. Martensite, which results from a rapid diffusionless phase transformation, has a complex nature that is challenging to characterize and to classify. Moreover the possibilities for modeling of this phase transformation have been limited, since its thermodynamics and kinetics are only reasonably well understood. However, the recent development of characterization capabilities and computational techniques, such as CALPHAD, and its applicability to ferrous martensite has not been fully explored yet.

In the present work, a thermodynamic method for predicting the martensite start temperature (Ms) of commercial steels is developed. It is based mainly on information on Ms from binary Fe-X systems obtained from experiments using very rapid cooling, and Ms values for lath and plate martensite are treated separately. Comparison with the experimental Ms of several sets of commercial steels indicates that the predictive ability is comparable to models based on experimental information of Ms from commercial steels.

A major part of the present work is dedicated to the effect of carbon content on the morphological transition from lath- to plate martensite in steels. A range of metallographic techniques were employed: (1) Optical microscopy to study the apparent morphology; (2) Transmission electron microscopy to study high-carbon plate martensite; (3) Electron backscattered diffraction to study the variant pairing tendency of martensite. The results indicate that a good understanding of the martensitic microstructure can be achieved by combining qualitative metallography with quantitative analysis, such as variant pairing analysis. This type of characterization methodology could easily be extended to any alloying system and may thus facilitate martensite characterization in general.

Finally, a minor part addresses inverse bainite, which may form in high-carbon alloys. Its coupling to regular bainite is discussed on the basis of symmetry in the Fe-C phase diagram.

Keywords: Carbon steels, Electron backscattered diffraction, Martensite, Microscopy, Microstructure, Thermodynamic modeling.

Preface

In the present doctoral thesis I would like to share my knowledge gained during the past four years, as a PhD student in physical metallurgy at the Department of Materials Science and Engineering at KTH Royal Institute of Technology, Sweden. The work has mainly been carried out within the martensite formation project (Supplements 1-5) which runs within the Hero-m research center, although Supplement 6 addresses inverse bainite and hence belongs to the bainite formation project within Hero-m. A large part of the characterization work on martensite is the result of two research stints conducted by the present author: at INP-Grenoble (France) in the spring of 2010 and at Tohoku University (Sendai, Japan) in the summer of 2011. In particular this doctoral thesis will discuss the following challenges, which are related to low temperature austenite decomposition in carbon steels:

 Thermodynamically based prediction of the martensite start temperature (Ms)  Quantitative microstructure characterization: The effect of carbon content on the martensite morphology with emphasis on the transition from lath to plate martensite.  Inverse bainite in hypereutectoid Fe-C alloys.

The outline of this thesis is as follows; In the first chapter, introduction, the topic low temperature austenite decomposition and the aims and purposes related to the three challenges listed above are introduced. In Chapters 2-4 my intention has been to condense selected works on three important subjects related to this work: Phase transformations, Characterization techniques, and Thermodynamics. I make no claim that these chapters would advance the understanding of low temperature austenite decomposition further; instead they should be viewed as either an introduction to the different topics for the novice or as an interpretation of the dense literature on the subjects done by the present author. The readers are referred to the cited works for more details. The expert reader may in fact skip Chapters 2-4 and move on to Chapters 5-7 where I summarize the results from the supplements and discuss them in relation to existing literature.

Albin Stormvinter Stockholm, 2012-08-01

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Appended papers

I. Investigation of Lath and Plate Martensite in a Carbon Steel A. Stormvinter, P. Hedström and A. Borgenstam Solid State Phenomena, 172-174, (2011), p. 61-66.

II. Thermodynamically Based Prediction of the Martensite Start Temperature for Commercial Steels A. Stormvinter, A. Borgenstam and J. Ågren Metallurgical and Materials Transactions A, Online first, 2012, DOI: 10.1007/s11661-012-1171-z

III. A transmission electron microscopy study of plate martensite formation in high-carbon low alloy steels A. Stormvinter, P. Hedström and A. Borgenstam Submitted manuscript

IV. Effect of Carbon Content on the Orientation Relationship between Austenite and bct- Martensite in Fe-C Alloys resolved by Electron Backscattered Diffraction A. Stormvinter, G. Miyamoto, T. Furuhara, and A. Borgenstam Submitted manuscript

V. Effect of Carbon Content on the Variant Pairing of Martensite in Fe-C alloys A. Stormvinter, G. Miyamoto, T. Furuhara, P. Hedström and A. Borgenstam. Submitted manuscript

VI. On the Symmetry Among the Diffusional Transformation Products of Austenite A. Borgenstam, P. Hedström, M. Hillert, P. Kolmskog, A. Stormvinter and J. Ågren Metallurgical and Materials Transactions A, 42(6), (2010), p. 1558-1574.

* The appended papers will be referred to as [1–6] in the present thesis

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My contribution to the appended papers I. Major part of: literature survey, experimental work, and writing. II. Major part of: literature survey, modeling work, and writing. III. Major part of: literature survey, experimental work, analysis, and writing. IV. Major part of: literature survey, experimental work, analysis, and writing. V. Major part of: literature survey, experimental work, analysis, and writing. VI. Part of experimental work, contributed to discussions and took part in the writing.

Other related reports not included in the thesis A. On the Three-Dimensional Microstructure of Martensite in Carbon Steels Peter Hedström, Albin Stormvinter, Annika Borgenstam, Ali Gholinia, Bartlomiej Winiarski, Philip J. Withers, Oskar Karlsson and Joacim Hagström Proceedings of the International Conference on 3D Materials Science 2012.

Parts of this work have been presented at the following international conferences 1. “Investigation of the Transition from Lath to Plate Martensite in Fe-C”, Albin Stormvinter, Peter Hedström and Annika Borgenstam, Solid-Solid Phase Transformations in Inorganic Materials (PTM2010), 6 - 11 June 2010, Avignon, France. 2. “A Phenomenological Model for the Prediction of Martensite Start Temperature in Steels”, Albin Stormvinter and Annika Borgenstam, The TMS Annual Meeting & Exhibition 2011 (TMS2011), 27 February - 3 March 2011, San Diego, USA. 3. “Thermodynamically Based Prediction of the Ms Temperature for Commercial Steels”, Albin Stormvinter and Annika Borgenstam, International Conference on Martensitic Transformations 2011 (ICOMAT2011), 4 - 9 September 2011, Osaka, Japan. 4. “Martensitic Meso- and Nanostructures in High-Carbon Low-Alloyed Steels”, Albin Stormvinter, Peter Hedström, and Annika Borgenstam, The TMS Annual Meeting & Exhibition 2012 (TMS2012), 11 - 15 March 2012, Orlando, USA.

Table of Contents

Abstract ...... v Preface ...... vii 1 Introduction ...... 1 1.1 Aim of the present work ...... 2 2 Low temperature austenite decomposition ...... 5 2.1 Martensite ...... 6 2.1.1 Lattice orientation relationship ...... 7 2.1.2 The habit planes ...... 8 2.1.3 Martensite – A hierarchic structure ...... 8 2.2 Bainite and inverse bainite ...... 10 3 Characterization techniques ...... 13 3.1 Optical microscopy ...... 13 3.2 Electron microscopy ...... 15 3.2.1 Scanning electron microscopy ...... 15 3.2.2 Transmission electron microscopy ...... 18 4 Thermodynamics ...... 21 4.1 Measuring the martensite start temperature ...... 23 4.2 Predicting the martensite start temperature...... 24 5 Martensite characterization in carbon steels ...... 29 5.1 High-carbon martensite investigated by TEM and ASTAR ...... 35 5.2 EBSD – Variant pairing tendency of martensite in Fe-C alloys ...... 41 5.3 Transition between lath and plate martensite...... 51 6 Discussion on inverse bainite ...... 53 7 Thermodynamically based prediction of the martensite start temperature ...... 57 8 Concluding remarks and future prospects ...... 61 9 Acknowledgments ...... 63 10 Bibliography ...... 65

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1 Introduction

Steels are by far the most widely-used metallic material and worldwide the crude steel production reached 1,527 megatons (Mt) for the year of 2011 [7]. The world production has almost doubled in ten years (851 Mt crude steel in 2001) and it is the demand from fast developing countries, in particular China, that drives this expansion [7]. This ever-growing demand for steel is certainly due to its attractive combination of high strength and high toughness with a low cost and easy recycling. These material properties make steel ideal for most engineering applications, although unlocking these inherent properties may not be trivial. It is the basic insight in physical metallurgy that allows for tailoring of steel properties, such as strength and toughness, by tuning alloying content and heat treatment processes [8].

The vast majority of engineering steels belong to the category carbon steels, these are steels that are often loosely classified according to carbon content. The manufacturing of carbon steels relies on controlling the decomposition of austenite, the high-temperature allotrope of iron. Austenite is not a stable phase at ambient temperature for carbon steels so it will decompose into ferrite and cementite on cooling. Figure 1.1 displays the well-known Fe-C phase diagram, which may be used to predict this decomposition. The phase diagram concept is best applied for phase transformations with fairly rapid kinetics, i.e. at relatively high temperatures, since diffusion is needed to reach equilibrium. However, the present work is dedicated to low temperature austenite decomposition. This might occur by some diffusional process as in the case of bainite formation. On the other hand, it might also be a diffusionless transformation as in the case of martensite formation. It all depends on the alloying content and the processing. Regardless whether the austenite decomposition results in bainite or martensite, phase transformations at low temperatures offer many challenges from a scientific point of view. These challenges include, but are not limited to: (1) Assessing the thermodynamics of important alloying systems, since there are very few experimental data available at such low temperatures. (2) Controlling and analyzing the kinetics of the phase transformations. (3) Characterizing and analyzing the complex microstructures which the phase transformations give rise to.

Figure 1.1: The Fe-C phase diagram calculated by Thermo-Calc [9], [10]. Solid lines represent the metastable phase diagram with cementite. Dotted lines represent the stable phase diagram with graphite.

1.1 Aim of the present work

The major part of this work has been dedicated to the effect of carbon on the martensitic microstructure in carbon steels. This is usually referred to as “well-known”, even though there is no standardized quantitative method to objectively evaluate this microstructure. Recent advances in electron microscopy have enabled quantitative characterization based on crystallographic information. Such techniques are electron backscatter diffraction (EBSD) [11] in the scanning electron microscope (SEM) and automatic phase and orientation mapping (ASTAR) [12], [13] in the transmission electron microscope (TEM). Here, these techniques have been applied to martensite formed in Fe-C alloys with the purpose to develop an objective method to characterize the change in the microstructure with the carbon content.

The second challenge addressed in this work is the prediction of the martensite start temperature

(Ms) as a function of alloying content, which has gained a lot of attention over the past century. It is quite understandable since it is a key measure for alloy design. Often a pure empirical approach has been used to model the alloy effect on Ms, although other approaches have been attempted as well such as neural network models. The problem with both these approaches is that the predictability may not be sufficient when the full composition range is considered. Therefore several researchers have attempted a more fundamental approach, based on thermodynamics, for predicting Ms. With a thermodynamic approach the idea is to model the transformation barrier and how it depends on alloying content and other parameters, such as prior austenite grain size (PAGS) or prior deformation. Such transformation barrier model would then be coupled to a reliable thermodynamic database to calculate the Gibbs free energy of austenite and martensite as a function of alloying content and temperature. It is this approach that has been attempted in the present work.

Finally, a minor part of this work is devoted to the formation of inverse bainite in high-carbon alloys. This is a subject that has not received so much attention as it might have deserved in the literature. Here the decomposition of high-carbon austenite into inverse bainite is characterized and classified by optical microscopy (OM) and by SEM. The results from the experimental study are discussed in relation to formation of ordinary bainite, with the main purpose to strengthen the evidence for bainite being a diffusion controlled transformation.

2 Low temperature austenite decomposition

Low temperature austenite decomposition in carbon steels may be better described by the word commonly used for the process, which also explains its purpose: hardening or quench hardening. The sole purpose of this process is to produce a microstructure that provides strength and hardness to the final product. In principle the processing may seem trivial: rapid quenching from the austenitic regime to complete transformation into martensite or bainite. Nevertheless, when studied in more detail several technological and scientific challenges arise, of which many can be attributed to the fine microstructure. However, before discussing these challenges in more detail a few words on the subject hardening should be addressed.

Iron began to replace bronze as the most common metallic material for tools and weapons some 3000 years ago [7]. To craft iron into a hatchet or a sword that could hold an edge and withstand long wear was not common knowledge, instead it was an art of craftsmanship. The ability to hold an edge is of course a direct consequence of the quench hardening process, which has to be carried out with accuracy. At this time there was for obvious reasons no method to study the process in more detail. So even though hardening had been applied for a few millennia it was not until the introduction of modern steelmaking and later the birth of metallography that the hardening process could be studied more in-depth. The hardened microstructure of steels was first described by Osmond and he named it after another pioneer metallographer: Adolf Martens [8]. Thus, the hard phase of steels was named martensite.

Ferrous martensite may be produced from austenite by providing sufficient cooling so all diffusional transformations are omitted, such as formation of ferrite or perlite. It may be convenient to introduce the concept of hardenability which is defined as the “susceptibility to hardening by rapid cooling” or as “the property, in ferrous alloys, that determines the depth and distribution of hardness produced by quenching” [9]. In general, an increase in alloying content permits a decrease in quenching rate, i.e. increases the hardenability.

2.1 Martensite

The martensitic microstructure has a high level of complexity and to resolve its true nature is a challenge. It is also known that the microstructure varies with composition and especially the carbon content has a strong influence. A common practice is to distinguish between two major types of ferrous martensite, namely, lath- and plate martensite [14], [15]. This is a general classification and several researchers recognize morphologies such as surface martensite [16], butterfly martensite [17], [18], thin-plate martensite [19], [20], and lenticular martensite [21], [22].

A martensitic transformation is defined by the following characteristics: (a) displacive; (b) diffusionless; and (c) kinetics and morphology are dominated by the strain energy arising from the shear-like displacements [23], [24]. Martensitic transformations can occur in many types of metallic and nonmetallic crystals, minerals, and compounds [25]. A displacive phase transformation is a structural change in the solid state which occurs by coordinated shifts of atoms [23]. This coordinated movement of atoms is sometimes referred to as a military phase transformations, which is in contrast to civilian diffusion-based phase changes that occur by uncorrelated jumps of individual atoms across an essentially incoherent interface [25]. A characteristic feature of the martensitic transformation is the shape change that may produce a surface relief if the transformation occurs near a free surface. In 1924, Bain [26] discussed the nature of the martensitic transformation in steels. He recognized that a contraction of about 20 percent along one of the austenite <001> axes and an expansion of about 12 percent along two of the <011> axes perpendicular to the chosen <001> axis transformed the austenite lattice into a ferrite bcc lattice. Hence, the fcc lattice of austenite could as well be viewed as a bct lattice with the axial ratio 1.414. This lattice correspondence has been named as the Bain correspondence and may be illustrated by Figure 2.1, which is reproduced from [27]. The Bain correspondence demonstrates that the bcc lattice could be generated from the fcc lattice through a homogeneous distortion. Nevertheless, it lacks some of the fundamental features of a martensitic transformation, for instance it does not produce an invariant plane. This plane is the interface between the martensite and the parent phase and should be undistorted and unrotated, and hence becomes the martensite habit plane. The Bain strain alone does not fulfill the criteria to produce such a plane, and therefore is not an invariant plane strain [28].

Figure 2.1: Bain correspondence in the transformation of (a) austenite into (b) martensite. Annotation: ○, Fe atom; x, positions available for carbon. Reproduced from Guo et al. [29].

Some 30 years after Bain proposed his theory, works by Wechsler et al. [30] and by Bowles and Mackenzie [31], [32] established a more solid understanding of martensite formation. These works applied the principles of matrix algebra to martensitic transformations and established what is now known as the phenomenological theory of martensite crystallography (PTMC) [33]. The PTMC was the first martensite theory that could predict experimentally observed habit planes, orientation relationships, and macroscopic distortions from knowledge only of the crystal structures of the parent and product phases. The PTMC thus introduced a mathematical and logical route between experimental observation and theoretical considerations [27].

2.1.1 Lattice orientation relationship A characteristic feature of the martensitic transformation is the lattice orientation relationship (OR) between the parent and product phases [27]. Pioneers Kurdjumov and Sachs [34] studied a carbon steel with 1.4C* by X-ray pole figure analysis and suggested the well-known K-S OR:

(111)γ//(011)α and [-101]γ//[-1-11]α.

The K-S OR states that the closed packed planes (CPP) along with the closed packed directions (CPD) of martensite (product phase) and austenite (parent phase) are parallel. Other commonly quoted OR are the Nishiyama-Wasserman (N-W OR) [35], [36] and the Greninger-Troiano (G-T OR) [37]. These were both determined from studies on high-nickel ferrous alloys containing large amount of retained austenite using X-ray diffraction techniques. More recently, Kelly et al. [38] and

* Compositions are given in mass% if not stated otherwise

Zhang and Kelly [39] have studied martensite formed in low carbon steels by TEM and measured orientation of thin layers of austenite retained between laths using convergent beam Kikuchi line diffraction patterns (CBKLDP) to determine the OR, they report:

(111)γ // (101)α and [1-10]γ 2.5°±2° from [1-1-1]α.

Zhang and Kelly [39] pointed out that CBKLDP has an accuracy of ±0.5° and can thereby very accurately determine the local OR between parent and product phase, given that sufficient amounts of retained austenite are present. Therefore this method cannot be applied to low-carbon martensite containing low amounts of retained austenite.

2.1.2 The habit planes In the early 1940s Greninger and Troiano [40] studied the orientation habit of martensite formed in plain carbon and nickel steels. They concluded that lenticular martensite did not form along any low- indices plane of austenite. Instead they reported that the orientation habit of martensite abruptly changed from a {4 4 10}γ to a {4 10 18}γ habit at roughly 1.4C. Later Greninger and Troiano [37], [41] extended their analysis of the orientation relationships and were the first to suggest a double shear mechanism, that later developed into the PTMC [30–32]. In carbon steels the martensite morphology changes with carbon content and so does the observed habit plane. For lath martensite it is close to {111}γ or {557}γ, but changes to {225}γ or {259}γ for plate martensite [27].

2.1.3 Martensite – A hierarchic structure The K-S OR predicts that 24 unique crystallographic martensite variants may develop from a single parent austenite grain, when annealing twins are disregarded. For lath martensite, these 24 variants may in turn be divided into four groups, each consisting of six variants with a common habit plane. These four groups have been named as “packets” [42–44]. By adopting the nomenclature used by Morito et al. [42] the four packets are made up by the following variant groups: V1-V6, V7-V12, V13-V18, and V19-V24. The packet may be further subdivided into three “blocks”, which each consists of two martensite variants with a low degree of misorientation. For instance, the three blocks of packet V1-V6 are: V1/V4, V2/V5, and V3/V6 [42]. The unique variants, for example V1, have been named as sub-blocks and are in fact the martensite unit that contains the individual laths. The lath has been described as a needle-like martensite unit, a few microns long and less than a

8 micron in diameter [45]. The substructure of the laths consists of dislocations arranged in a typical cellstructure, which are a result of accommodation deformation due to the large shape strain [46]. Adjacent laths may have a misorientation of a few degrees, but are in principal of the same crystallographic martensite variant. With the lath as the smallest building block, lath martensite may be regarded as hierarchic structure, which becomes: parent grain – packet – block – sub-block – individual lath.

For plate martensite, which is found in high-carbon and high-alloyed grades with low martensite start temperature, it is not at all evident if it exists a hierarchic structure comparable to what has been shown for lath martensite. When studied by optical microscopy (OM) plate martensite appears as individual units of various sizes. The unit size is limited by the parent grain size and by preceding martensite plates, which inhibits growth of a fresh unit. Some characteristic features of plate martensite are: (1) the midrib, (2) transformation twins, (3) deformation twins, (4) dislocation arrays, and (5) transverse micro-cracks [47], [48]. Characterization work done on plate martensite have commonly addressed the irrational habit planes [40] and remarkably few studies have been made on variant selection. One of the few studies were made by Okamoto et al. [19] who characterized martensite formed in a Fe-30.70Ni-0.28C alloy. By following Okamoto et al. [19] it may be concluded that a nucleation event in ferrous alloys with a low Ms temperature will result in the growth of a “plate group”, which consists of four unique crystallographic martensite variants. The principle plate group, which is referred to variant V1, becomes: V1, V16, V17, and V6 (again by adopting the nomenclature used in [42]). In total six unique plate groups may form from a single parent austenite grain.

The regime between these two extreme morphologies of martensite, i.e. lath- and plate martensite, is often referred to as a mixed regime with a gradual transition from the former to the later. This morphological transition has been schematically represented in Figure 2.2. However, this is more or less a crude simplification since it has been shown that the morphological transition is gradual over the entire composition range [49], [50]. In a recent review on the martensite morphology by Maki [51] he pointed out that the factors that determine the morphology and substructure of martensite remain poorly defined.

Figure 2.2: Schematic representation of the transition from lath to plate martensite with increasing carbon content.

As the carbon content is increased the Ms decreases and the martensite morphology changes from lath to plate martensite. In between these two extremes it is said to be a “mixed” regime.

2.2 Bainite and inverse bainite

The decomposition of austenite into ferrite and cementite in the temperature range below that for pearlite results in a microstructure that has been named as bainite [52]. Bainite, like pearlite, is a eutectoid transformation product and may be further subdivided into upper- and lower bainite. Both these transformation structures are acicular and ferrite is the leading phase in the main growth direction during the transformation. In bainite, the ferrite has a tendency to form as plate-like units and grow with an orientation relationship with the austenite [25]. In the case of upper bainite, carbon will diffuse away from the growing ferrite into the austenite and eventually the carbon content reaches such a level that cementite nucleates and grows. The shape of the cementite is usually as a broken lamella in between the ferrite units. At temperatures where lower bainite forms the diffusion of carbon is slow, which makes the carbides to precipitate in more close relationship with the ferrite. The resulting carbide dispersed ferrite becomes very fine and may in fact be hard to tell apart from tempered martensite. Another feature that bainite share with martensite is that it produces a surface relief effect [53]. This is probably one of the reasons for the uncertainty regarding the mechanism by which bainite grows: If it forms by a rapid, diffusionless shear mechanism or if it is controlled by

10 carbon diffusion. However, the surface relief should not be taken as a proof of a diffusionless nature, since it has been shown that Widmanstätten ferrite [54] as well as Widmanstätten cementite [55] gives a surface relief. Both these transformations are without a doubt diffusion controlled to their nature. The diffusionless hypothesis was initially proposed by Zener [56] and suggests that a bainite unit or a sheaf grows by short consecutive steps [57]. Each of them so rapid that there is not enough time for carbon to diffuse. The carbides are then formed in a subsequent step from the supersaturated bainitic ferrite, which involves carbon diffusion. The final product, here lower bainite, may be seen from Figure 2.3. Macroscopically the growth rate of bainite is rather slow and could as well be controlled by carbon diffusion. This hypothesis, that the nature of bainite growth is diffusion controlled, has long been supported by Hillert, who in 1957 [58] proposed that there should be a symmetry among the eutectoid transformation products of austenite in the Fe-C system. He suggested that pearlite would form if ferrite and cementite behaves as equal partners as the eutectoid transformation progresses. Bainite forms when ferrite is the leading phase, and the ferrite constituent is identical to Widmanstätten ferrite. Hillert also introduced, for symmetry reasons, a third transformation product that would have cementite as the leading phase and he named it as inverse bainite. Also for inverse bainite the cementite, as the leading phase, should be identical to Widmanstätten cementite.

Figure 2.3: Lower bainite formed at 618 K (345 °C) in alloyed steel with 0.69C. 10,000 times magnification. Reproduced from Oblak and Hehemann [57].

3 Characterization techniques

The science discipline of examining and determining the constitution and structure of metals has been named metallography. The father of metallography is Henry Clifton Sorby who was the first person to examine correctly polished and chemically etched metal samples under the microscope in the 1860s. Today the discipline covers the examination of metals over a wide range of length scales, from millimeter scale down to the atomic scale [52].

3.1 Optical microscopy

The classical metallographic tool is the optical microscope (OM), which allows for accurate metallographic examinations of metallic materials of various shapes and sizes. Before the OM investigation, specimens are usually mechanical polished and chemical etched [52]. This will produce a proper surface and thus allow for examination of microstructural features which may arise from solidification or solid state transformations. Features that might be studied are for instance:

 Grains (size and shape) and grain boundaries  Phases and phase interfaces  Annealing twins  Composition gradients that affect the microstructure  Inclusions (distribution, size and shape)  Precipitates (distribution, size and shape)  Imperfections (cracks and notches)  Surface structures (Corrosion attacks, fracture surfaces, and cracks)

Figure 3.1 displays an OM micrograph on transformation structures in a high-carbon alloy. This figure well illustrates the capabilities of accurate metallography using OM. From the figure three different features can quite easily be distinguished, those are: spiky nodules, plate-like units formed in zig-zag patterns, and a matrix phase. Since the heat treatment process and alloy composition is known in this case, these three features may be recognized as inverse bainite, plate martensite, and retained austenite, respectively. However, there are limitations of the OM technique; those are first and foremost the limited depth of field and the spatial resolution, which is of the same order of magnitude as the wavelength of light, i.e. features smaller than a few tens of a micron cannot be observed. Moreover, the limited depth of field at high magnification makes it difficult to picture features which are non-planar [52]. The limited depth of field might not be critical for studying

13 martensite and inverse bainite on a polished surface, as seen from Figure 3.1, but to resolve the fine microstructural features in more detail is beyond the resolution of OM.

Figure 3.1: Microstructure of a 1.67C carbon steel isothermally heat treated at 400 °C for 1 minute after a 10 minutes austenitization at 1100 °C. The dark etching spiky nodules are inverse bainite formed during the isothermal hold. Apart from inverse bainite the figure shows plate martensite (dark gray) formed during the final quenching to room temperature. The matrix is retained austenite. The specimen was prepared by mechanical polishing followed by etching in Nital.

3.2 Electron microscopy

Even though the OM is a versatile characterization tool, its limited spatial resolution makes it insufficient to study the nanostructure of materials. Therefore, it becomes necessary to apply electron microscopy techniques to resolve such details of the materials structure. By using modern SEM and TEM it is possible to achieve a resolution better than 1 nm and thus resolve the very fine structural features. In addition, it is possible to determine the local crystallography as well as the local chemical composition using electron microscopes, which will add substantial information to the characterization work.

3.2.1 Scanning electron microscopy The SEM is primarily used to study the surface, or near surface, structure of bulk specimens. Electrons are generated from a source and then focused on to the surface by condenser lenses. The electron source has usually been of the tungsten filament thermionic emission type, although field emission gun (FEG) sources are growing in popularity since they may achieve higher resolutions and higher probe currents [59]. The electrons generated from the source will generate various secondary emissions after interaction with the specimen surface. These emissions and corresponding contrast mechanisms have been listed in Table 3.1:

Table 3.1: Overview of contrast mechanisms, detectors, and typical lateral and depth resolution [52]. Lateral Depth of Detected signal Type of detector Information resolution information Surface topography, Secondary electrons Scintillator/photomultiplier 5–100 nm 5–50 nm compositional contrast Compositional contrast, Solid-state detector or Backscattered electrons surface topography, crystal 50–1000 nm 30–1000 nm scintillator/photomultiplier orientation, magnetic domains Complementary contrast to Same as Same as No external detector Specimen current backscattered plus secondary backscattered backscattered necessary electron signal electrons electrons Semiconductor detector Characteristic X-rays (energy-dispersive) or Element composition, element 500–2000 nm 100–1000 nm (primary fluorescence) crystal/proportional counter distribution (wavelength-dispersive) Detection of nonmetallic and Cathodoluminescence Photomultiplier ...... semiconductive phases

All the detectable signals listed in Table 3.1 may be used for materials characterization. For instance, secondary electrons excel in detecting the surface topography and are more or less insensitive to atomic number. Instead the contrast depends strongly on the angle between the incident beam and the specimen surface. By applying a suitable etchant, chemical or electrolytic, interesting features may be revealed and could then easily be viewed using the SEM in secondary electron mode. Figure 3.2 gives an example on such image, here carbide precipitations as a result of martensite tempering are revealed.

Backscattered electrons (BSE) are electrons that escape the specimen due to elastic scattering. These electrons are more energetic than the secondary electrons and their yield have a strong correlation to the average atomic number. Hence, the backscattered electrons carry information on the local composition and their signal may be used to produce an image with compositional contrast. In addition, backscattered electrons carry crystallographic information since the yield also is dependent on the orientation of a crystal with respect to the incident beam [59]. This effect, known as electron channeling, is usually much weaker when compared to the atomic number contrast. Figure 3.3 displays a martensitic steel where the microstructure has been revealed by channeling contrast in the SEM.

Figure 3.2: Plate martensite formed in a high-carbon steel quenched to 200 °C after austenitization at 1000 °C for 10 minutes. Carbide precipitation is clearly visible inside the large martensite unit. These carbides precipitated during a 1 minute tempering at the quenching temperature. The figure is a SE image using an in-lens detector and acquired using the following setup: 5.0 kV, WD 5 mm, x10,000. The specimen was prepared by mechanical polishing followed by electrolytic etching.

Figure 3.3: Displays plate martensite formed in a high-carbon steel. The figure is a BSE image acquired using the following setup: 10.0 kV, WD 2.7 mm, x1000. The specimen was prepared by mechanical polishing.

A SEM technique that is rapidly developing is the electron backscattered diffraction technique (EBSD). It is a fully automated quantitative material characterization technique based on mapping the crystallography of the investigated specimen. EBSD relies on the detection of Kikuchi patterns, which arise from subsequent elastic scattering of inelastically scattered electrons [59]. To obtain an increased amount of electrons, adding to the Kikuchi pattern, the incident beam should be tilted about 70° relative to the analyzed surface. To be successful the surface should have a mirror finish and be essentially free from deformation, for instance by using a vibration-assisted final polishing. The electron beam is scanned across the specimen while the position (x- and y-coordinates) and Kikuchi patterns are recorded simultaneously. The next step in the EBSD analysis is to index the collected Kikuchi patterns; this is commonly done using a Hough transform [11]. For every mapped- point the best possible solution, i.e. phase and crystallographic orientation, is calculated from the recorded Kikuchi data. Recent EBSD systems may even select the appropriate phases automatically, given that a reliable database including the current phases is available. The resulting dataset will contain crystallographic data for every successfully analyzed point within the mapped area of the analyzed 2-D cross-section. Commercial EBSD software includes numerous post-processing possibilities, for example the collected dataset can be viewed as a phase colored map or an inverse pole figure (IPF) colored map as seen in Figures 3.4a and 3.4b respectively. The reader who wants further insight in the SEM and EBSD techniques are recommended to the excellent monographs by Goldstein et al. [60] and by Schwartz et al. [11], respectively.

Figure 3.4: (a) Phase colored EBSD map of a high-carbon steel where martensite is colored in red and retained austenite is colored in blue. (b) An inverse pole figure colored EBSD map of martensite that corresponds to (a), the retained austenite is here colored in gray. The EBSD dataset is 45 by 35 microns and acquired using an OXFORD system and a 50nm step size.

Besides generating characteristic electrons the beam specimen interaction also gives rise to X-rays. Among these there are the characteristic X-rays, which are a result of ionization of the atoms in the material. An electron of the primary beam may knock an electron out from an inner shell and thus leave an electron vacancy. This vacancy will immediately be replaced by an electron from an outer shell. As this other shell electron transfers to the new lower energy level it releases some excess energy. This energy can be lost as an X-ray photon which in turn is characteristic for every element. Therefore, these X-rays may be used to analyze the local or average chemical composition. However, there are some overlaps in the X-ray energy spectra which may be difficult to overcome.

3.2.2 Transmission electron microscopy The transmission electron microscope builds on the same principle as the scanning electron microscope; that is interaction between the generated electron beam and the specimen. The central part of the TEM is the vertical vacuum chamber. At its top the electron gun is located, which accelerates electrons to a voltage of about 40-400 kV. Below the electron gun are two or more condenser lenses with an aperture present in between. The lenses together with the aperture are used to change the character of the incident beam, which may be defocused to a parallel beam or focused to a convergent beam. It all depends on the desired analyzing conditions. Below the lenses lies the specimen chamber. As the TEM relies on interaction of transmitted electrons the specimen chamber and holder are a crucial part of the equipment. The holder should allow the operator to move the

18 specimen as well as to tilt it to rather large angles. The electrons that have interacted with the specimen should also be able to leave the specimen chamber. In addition, X-rays that may have been generated should be permitted to leave as well. After the specimen the objective lens and the projector lenses are located. These lenses are used to produce the image. There are several types of images that may be produced in the TEM, these include: bright field (BF) images, dark field (DF) images, high angle annular dark field (HAADF) images, selected area electron diffraction (SAED) images, and convergent beam electron diffraction (CBED) images [59]. It is the objective aperture that is used to select either the undeflected (bright field) or the diffracted (dark field) beam for the imaging. The magnified image is often recorded using a charge-coupled device (CCD) array camera. An example of a BF image is given in Figure 3.5, which displays plate martensite in a high-carbon steel.

A technique that has many similarities with SEM-EBSD is the automatic phase and orientation mapping for TEM (ASTAR) [12], [13]. Instead of determining the phase and crystal orientation from collected Kikuchi patterns, ASTAR is based on template matching of electron diffraction spot patterns acquired in the TEM. The templates, i.e. theoretical diffraction patterns, are calculated with the crystal structure as input parameter. The commercial ASTAR technique also contains quality indicators such as “reliability”, which is comparable to confidence index found in the TSL-EBSD system. ASTAR is explained in more detail in Rauch and Dupuy [12] and Rauch et al. [13].

Figure 3.5: Displays plate martensite formed in a high-carbon steel. The figure is a BF image acquired using an acceleration voltage of 300 kV and x6000 magnification. The specimen was prepared by electrolytic polishing.

4 Thermodynamics

Today computational thermodynamics are a fundamental engineering tool for materials design. Coordinated international research using CALPHAD (CALculation of PHAse Diagrams) techniques has provided the engineering society with reliable databases of important alloying systems, such as carbon steels. Computational thermodynamics rely on accurate descriptions of the Gibbs energies for the pure elements and relevant compounds of the material of interest. From these modeled Gibbs energies, the equilibrium state of the system, i.e. steel grade, may be calculated for a certain alloy content, temperature and pressure. Moreover, it is possible to study the relative difference in Gibbs energy of stable and metastable phases as a function of temperature. Regarding martensite formation in carbon steels, it is of interest to know the T0 temperature [56], the temperature where the Gibbs energies of austenite and ferrite phases with the same composition are equal, i.e.

. At temperatures above T0, austenite is more stable than ferrite, with the same composition, and below T0 the vice versa is true. Evidently, the total energy of the system would decrease if austenite would transform into ferrite of the same composition below T0. Hence, there is a positive driving force for ferrite formation, which may be defined as:

[Eq. 4.1]

In practice, when using computational thermodynamics, martensite may be treated as supersaturated ferrite and Equation 4.1 thus becomes relevant for martensite formation as well. Ordinary ferrite forms in a diffusional manner with a different composition than the parent austenite even above T0 but still inside the austenite-ferrite two-phase field of the phase diagram. The compositions of austenite and ferrite, locally at the migrating phase-interface, are close to those of the equilibrium tie line in the phase diagram, i.e. not much driving force is needed for the interface to migrate. Most of the driving force is used to change the composition by diffusion. However, when regarding the experimental Ms temperature for pure iron, reported as 818 K (545 °C) [61], the calculated available driving force is considerable. Figure 4.1 displays how the available driving force increases with decreasing temperature and reaches 1250 J/mol at the experimental Ms. It should be remembered though that only the chemical part of the change in Gibbs energy is considered in this calculation. In general, the Gibbs energy of a martensitic phase is higher than that of ferrite, since coherency strain energy, surface energy, and defect energies needs to be accounted for [62], [63]. These effects may be 21 difficult to quantify and therefore it may be more useful to define the requirement of chemical driving force to start the martensitic transformation. This may be thought of as a transformation barrier, which will hereafter be denoted . The barrier is essentially different from , which is the available chemical driving force at a given temperature. The available driving force is a thermodynamic property, defined in Equation 4.1, which depends on the temperature and alloy composition. The barrier, on the other hand, depends on the details of the transformation mechanisms, and will thus account for the elastic and plastic energies stored in the material after the martensitic transformation. For each steel, the available driving force varies with temperature, and MS is reached when the available driving force is equal to .

Commercial multicomponent steel databases, such as TCFE6 [64], do not contain a metastable martensitic phase. Instead, the user may treat martensite as supersaturated ferrite when extracting thermodynamic data from the database. This may lead to some concerns, for instance how to treat the ordering of carbon atoms [56]. The ordering occurs as a consequence of the very rapid martensitic reaction giving no time for carbon atoms to perform diffusive jumps. In the ferrite structure there are three times more interstitial sites than in austenite but only one-third of them corresponds to the interstitial sites where carbon may be located in the austenite. Thus the carbon atoms are not randomly distributed in the fresh martensite, since the positions are inherited from the austenite. The ordering of carbon atoms was first discussed by Zener [56] and is therefore sometimes referred to as Zener ordering. Ordering of carbon atoms has a very small effect on the thermodynamic properties of the martensite at low carbon contents. On the other hand, it will have a dramatic impact at high carbon contents. Therefore, this effect should be taken into account when evaluating the driving force for martensitic transformations in high-carbon alloys. For example, this has been done in an evaluation by Fisher [65].

Figure 4.1: Displays the driving force for martensite formation as a function of temperature for pure iron. The experimental Ms temperature is taken from Morozov et al. [61] and the Gibbs energies for martensite (ferrite) and austenite was calculated in the Thermo-Calc system [9], [64].

4.1 Measuring the martensite start temperature

Several methods have been proposed to predict the martensite start temperature for commercial steels. These include empirical-, neural network-, and thermodynamic methods. They have in common that they rely on experimental Ms temperatures in some sense. It is interesting to notice that there is no standard method to evaluate the experimental Ms temperature. Instead several methods have been used, for example metallography [66–68], dilatometry [69–72], magnetometer [72–74], and thermal arrest measurements [61], [75–78]. Often the experimental uncertainty of the method used for determining Ms has not been evaluated. In addition, there are uncertainties related to the effect of the PAGS, which seems to affect Ms significantly [79–81]. Therefore it is reasonable to believe that the uncertainty for experimental Ms data is rather large. In a paper by Ghosh and Olson [82] they suggested it to be ±40 K. On the other hand, Mirzayev et al. [83] claim that the accuracy of their measurements is not worse than ±15 K.

4.2 Predicting the martensite start temperature

Many modern engineering steels rely on the hard phase of martensite, for example: transformation induced plasticity (TRIP)-assisted steels [84], quenched and partitioned steels [85], and maraging steels [86]. Therefore, it is no surprise that the prediction of Ms and especially its dependence of alloying content has gained a lot of attention over the years. An early attempt to rationalize information on Ms in commercial steels was made by Payson and Savage in 1944 [87] who proposed an empirical expression that could be used for predicting the Ms temperature of new steels:

[K], [Eq. 4.2]

with alloying content given in mass percent [mass%]. The original equation gave Ms in Farenheit but has here been recast to give Ms in K. A large number of similar empirical expressions have later been proposed [67], [88–90] and they have been reviewed a number of times, for instance by Wang et al. [91], by Skrotzki and Hornbogen [92], and by Sourmail and Garcia-Mateo [93]. These empirical expressions may have a very practical usage, although they are limited to a certain composition range.

Figure 4.2: Variation of driving force with Ms temperature in binary iron alloys. Reproduced from Raghavan and Antia [94].

A more fundamental approach is to apply a thermodynamic method to model Ms. This can be done by evaluating the barrier for the diffusionless austenite to martensite transformation at the measured

Ms temperatures and study how it varies with composition. Raghavan and Antia [94] attempted such analysis. However, they did not use experimentally determined Ms values in their analysis, instead they accepted the improved Andrews’ [90] equation evaluated by Kung and Rayment [95]. Raghavan and Antia [94] calculated the driving force at Ms for some binary and ternary systems, within the composition range covered by the equation, as a function of composition. Their results are presented in Figure 4.2 and illustrate how they represented the variation of driving force at Ms in binary alloys with straight lines. In their analysis they concluded that for multicomponent alloys it did not appear feasible to develop a simple equation for the required driving force at Ms as a function of composition. However, they suggested that a more useful result would be to derive a correlation between the required driving force and Ms by statistical fitting. By calculating the driving force at Ms for a total of 1152 arbitrarily selected low alloy steels within a given composition range they obtained a dataset for their statistical analysis. A least-squares fit of all 1152 compositions yielded the following linear equation:

( ) [J/mol], [Eq. 4.3] Another example of a thermodynamic approach is the expression proposed by Hsu and Chang [96], obtained by studying information from Fe-C alloys:

( ( )) [J/mol], [Eq. 4.4]

where xc denotes mole fraction of carbon. This expression was intended for use only above xc = 0.02. The last term was supposed to give the temperature dependence of the effect of shear and the second term was supposed to account for the solution hardening of carbon. The expression could just as well be written as:

( ) [J/mol], [Eq. 4.5]

Wang et al. [97] have also derived an expression for the required driving force as function of temperature. They selected 104 engineering steel compositions and calculated the driving force thermodynamically at estimated Ms temperatures, which were obtained from a trained neural network [98]. They fitted a straight line to the calculated driving forces and obtained:

( ) [J/mol], [Eq. 4.6]

When taking the carbon ordering energy into account they obtained a slightly different expression:

( ) [J/mol], [Eq. 4.7]

It should be noted that Wang et al. [97] did not find it necessary to include the effect of alloying elements explicitly in their expressions.

A very extensive study was made by Ghosh and Olson in 1994 [99], [100]. They plotted the driving force, evaluated at the experimental Ms for various alloying contents, xi, as function of that content. Assuming that the effect of an alloying element on the required driving force should be similar to the effect on solution hardening they fitted the data with a parabolic curve,

[J/mol], [Eq. 4.8]

where K1, which should hold for pure Fe, was treated as independent of temperature. The second parameter, , was evaluated for a large number of alloying elements and assuming that the effects were additive they could then predict the Ms values for a large number of steels and obtained reasonable agreement with experimental values.

Ghosh and Olson [82], [101] later refined their method by introducing a new multicomponent database, specially evaluated for martensitic transformations, together with a complex method for representing the critical driving force for starting the martensitic transformation. In their paper they concluded that the accuracy of Ms predictions, in ternary and multicomponent systems, was satisfactory if an experimental uncertainty of ±40 K was accepted.

So, even though a lot of work has been invested in predicting Ms there is still no general method readily available, that one can apply easily. However, it seems like a semi-physical approach that combines state-of-the-art thermodynamics with key experiments might be the best option to establish a more general method of predicting Ms.

5 Martensite characterization in carbon steels

It is well-known that the microstructure of ferrous martensite is affected by the alloying content [15], [27], [102]. Even though the morphological change may seem obvious, it is still very challenging to make a qualitative as well as quantitative classification of the observed microstructure. Often microstructure characterization relies on the experience of the operator, i.e. the metallographer. It may seem unfair to thus suggest that microstructure characterization is to some extent subjective, since there are quantitative indicators as well. Those are for instance hardness, strength, and other measurable material properties, although they all would be indirect measurements.

As discussed in Chapter 2 the martensite morphology changes from lath-like to plate-like as the alloying content is increased. This progressive change in martensite morphology can be represented by the Figures 5.1a-5.1e, which display specimens with an increasing carbon content. The figures also illustrate the complexity of the martensitic microstructure, which consists of individual martensite units forming together in hierarchic structures. Many of the microstructural features are beyond the magnification range of the OM and therefore only a qualitative understanding of the microstructural transition is realistic from this series of OM micrographs.

Figure 5.1: Martensite formation in steels with increasing carbon content, all specimens were mechanical polished and etched with 3% Nital. (a) Interstitial free steel (Fe-1.48Mn-0.0024C), (b) Fe-0.35C, (c) Fe-0.75C, (d) Fe-1.05C, and (e) Fe-1.80C. The figure is reproduced from [5].

A very useful metallographic technique, which has not been used so extensively in the past decade, is to study the microstructural change in a specimen with a composition gradient. For instance, Borgenstam et al. [103] successfully applied the gradient technique to study the critical growth temperature for martensite and how it was related to Ms. More recently Stormvinter et al. [1] applied this technique to a carbon steel to qualitatively study the microstructural transition from lath to plate martensite with increasing carbon content. The steel composition is listed in Table 5.1. After decarburization using hydrogen, a carbon gradient was obtained close to the surface, as seen from the wavelength dispersive X-ray spectroscopy (WDS) measurement displayed in Figure 5.2a. They also measured the hardness versus carbon content and concluded that it agreed well with previously reported hardness numbers [45], [74], [104], [105], as seen from Figure 5.2b.

Table 5.1: Composition of the carbon steel studied in [1], alloying content is given in mass%. Carbon Chromium Silicon Manganese 0.20 – 1.67 0.47 0.27 0.23

Figure 5.2: (a) Carbon profile close to the surface, ~0.01-5 mm, after decarburization measured with WDS and simulated with DICTRA using the actual conditions. (b) Hardness vs. carbon content [1], [45], [74], [104], [105]. The figure is reproduced from [1].

Stormvinter et al. [1] used a quenching and tempering procedure when characterizing the martensitic microstructure in the carbon gradient specimens. This procedure made it possible to separate the first martensite units that formed, close to the Ms temperature, from units formed during the final quenching to room temperature (RT). Moreover, any position in a heat-treated specimen could be related to the local carbon content by using the information in Figure 5.2a and hence it was possible to see how the martensite formation close to Ms was affected by the carbon content.

Figures 5.3a-5.3c are from carbon gradient experiments [1] and show typical transformation structures found in carbon steels: lath martensite (Figure 5.3a), plate martensite (Figure 5.3c), and martensite formation in the transition regime between these two extremes (Figure 5.3b). The tempered martensite appears dark while the martensite formed during quenching to RT appears bright. Close to the eutectoid composition the martensite microstructure displays characteristic features for both lath and plate martensite. For instance, in Figure 5.3b a large martensite plate is visible to the right that has grown to a considerable size. Moreover, it can be seen how new martensite units have been autocatalyzed on this primary unit. The autocatalysis of new martensite units on a primary plate is better seen in Figure 5.4, taken from a Fe-1.80C binary alloy (the same alloy as seen in Figure 5.1e). Apart from these plate-like martensite units, Figure 5.3b shows small martensite units without typical features of plate martensite and could as well be lath martensite.

The gradual morphologic transition of martensite in the composition range 0.6 to 1.1C can be seen in Figures 5.5a-5.5f. It is realized that the individual martensite units become more defined and midribs, characteristic for plate martensite, are visible in Figures 5.5e and 5.5f. It should further be noted that the change in hardness also appears to be continuous over this composition range, as seen from Figure 5.2b.

Figure 5.3: OM micrographs of the two main types of martensite and the transition region in between; (a) Quenched and tempered at 623 K, 100-300 µm, ~0.32C. (b) Quenched and tempered at 473K, 1000-1100 µm ~0.81C. (c) Quenched to 298 K, >5000 µm, ~1.67C. The figure is reproduced from [1] and the micrometer range refers to the distance from the surface.

Figure 5.4: Tendency for autocatalysis of new martensite units on a primary martensite unit in a Fe-1.80C binary alloy.

Stormvinter et al. [1] performed SEM-BSE imaging of the carbon gradient specimens, which are shown in Figures 5.6a and 5.6b. In these BSE images the martensite units become visible through crystallographic contrast, and hence a change in lattice orientation will be reflected in the images. Figure 5.6b is centered on a martensite unit, which is also indicated with a white arrow in Figure 5.6a, and it is clearly seen that this unit contains two lattice orientations. This is most probably due to the internal transformation twinning that occurs when plate martensite grows from the austenite, although the twin-relation cannot be confirmed in this imaging mode.

The morphological transition observed in specimens with a carbon gradient by Stormvinter et al. [1] agrees well with the commonly accepted view of martensite, i.e. lath martensite forms in the composition range 0 to 0.6C and above 1.0C acicular plate martensite is the dominant morphology [47]. The transition between these two morphologies is usually observed within the composition range 0.6 to 1.0C [14], [47], [106]. Maki [51] has emphasized that the substructure and crystallography of martensite units, which forms in this composition range, is fairly complicated. Moreover, Maki et al. [107] have shown that the morphology of lath martensite changes with the carbon content and Greninger and Troiano [40] have observed a change in the morphological appearance of plate martensite in Fe-C alloys. For 1.5 to 1.8C, they found plate-type martensite with clearly visible midribs, which form in zig-zag patterns. Below this composition range they reported a sudden change in morphological appearance of the plate martensite, since two plate-like units seemed to be emanating from a point and grow with an obtuse angle with much less prominent midribs. This latter morphology has later been described as butterfly martensite [17], [49] and the former as lenticular martensite [49].

Figure 5.5: Martensite morphology of a sample cooled to RT without tempering, observed in OM: (A) 900-1000 µm, ~0.76C; (B) 1100-1200 µm, ~0.83C; (C) 1300-1400 µm, ~0.92C; (D) 1500-1600 µm, ~0.97C; (E) 1700-1800 µm, ~1.09C; (F) 1900-2000 µm, ~1.11C. The figure is reproduced from [1] and the micrometer range refers to the distance from the surface.

Figure 5.6: (a) Martensite morphology observed in SEM-BSE, 800-1000 µm ~0.75C; (b) High magnification of the area marked by the arrow in (a). The figure is reproduced from [1].

5.1 High-carbon martensite investigated by TEM and ASTAR

Classic TEM has been used to study martensite since the early 1960s, and it was clear already from the pioneering works on medium- and high-carbon ferrous martensite that the microstructure and substructure of plate martensite is complex [45], [108]. Characteristic features that has been reported are for instance: (a) twins, (b) planar defects, (c) dislocation arrays, (d) midrib region, and (e) micro- cracks [47], [48]. Pioneers Kelly and Nutting [45] were the first to make examinations of ferrous martensite using thin-foils in the TEM when they studied martensite in two high-carbon low alloy steels with 1.0 and 1.4C respectively. They reported that both steels displayed plate-like martensite units, a few microns long and less than a micron thick. Within these units they observed nano-sized twinning, which had a 15 to 100 Å interspacing, and reported the twin-plane to be {112}α.

Furthermore, they found some retained austenite in the 1.4C steel and concluded that the {112}α twins were produced from {110}γ planes. Later, Oka and Wayman [48] studied two high-carbon low alloy steels with 1.28 and 1.82C. The martensite units found in both these steels also contained

{112}α transformation twins, but in the 1.82C steel {101}α planar defects were observed. They suggested that the {101}α planar defects were accommodation distortion imposed by the growing martensite.

There have in fact been surprisingly few studies following the works by Kelly and Nutting [45] and by Oka and Wayman [48] on plate martensite in high-carbon low alloy steels. On the contrary, when regarding martensite formation in general a vast number of TEM studies have been reported. In recent years there has been a significant development of characterization capability, for instance automated crystallographic techniques as TEM-ASTAR. Therefore, Stormvinter et al. [3] selected two high-carbon low alloyed steels with the ambition to combine ASTAR with classical TEM to provide a comprehensive basis for discussing the development of plate martensite in these alloys.

The alloys chosen by Stormvinter et al. [3] for TEM observations contained 1.20 and 1.67C respectively. Figure 5.7a is from the 1.67C alloy and shows a large plate-like martensite unit on the right-hand side with a distinct midrib in the vertical direction. On the left-hand side of this large martensite unit, it can be seen that smaller martensite units have developed in characteristic zig-zag patterns, which are surrounded by heavily dislocated retained austenite. Stormvinter et al. [3] confirmed by SAED (Figure 5.7c) that the large martensite unit to the right was internally twinned.

The twins were of {112}α-type and can be seen from Figures 5.7b and 5.7d, which are DF 35 micrographs. The extension of these twins appears limited by a discrete crystallographic plane, as can be seen from the figures. This limiting plane was determined as {101}α and may in fact be accommodation distortion of either fault- or twin-type.

Figure 5.7: TEM micrographs of the 1.67C steel. (a) On the right hand side in the BF micrograph a large plate martensite unit is seen with a midrib in the vertical direction. On the left hand side, areas of retained austenite are present, which seem to contain a high amount of dislocations. (b) DF micrograph centered on the midrib of the large plate martensite unit on the right-hand side in Figure 5.7a. There is strong diffraction close to the midrib. (c) SAED from the midrib area with the 200α-twin spot used for DF indicated. (d) High magnification DF micrograph corresponding to the central part of Figure 5.7b. It is seen that the strong diffraction is due to nano-sized transformation twinning along the midrib. The figure is reproduced from [3].

Another example of plate martensite in the 1.67C alloy, with transformation twinning on the midrib, is shown in Figures 5.8a and 5.8b. In addition to classic TEM, Stormvinter et al. [3] attempted ASTAR on these martensite units. The result of the ASTAR is displayed in Figure 5.9 as an IPF colored map. When comparing the ASTAR results to the TEM imaging it is realized that the nano- size transformation twinning was not resolved by ASTAR, since only a small area (Indicated by (1)) shows a crystal orientation that deviates from the primary orientation of unit A. This can easily be understood from the fact that the size of the twins was in the order of the probe-size (20nm). However, by decreasing the probe-size to 10nm and analyze only a small fraction, close to the midrib, of martensite unit A the transformation twinning could be better reproduced by ASTAR. This result is seen from the insert in Figure 5.9. Stormvinter et al. [3] used a conventional electron source when conducting the ASTAR; however, by using a FEG source the probe-size could be further decreased and it is thus likely that the spatial resolution would be good enough to accurately resolve the transformation twinning. Nevertheless, it will still be very challenging to simultaneously resolve transformation twinning in multiple martensite units, since only twins aligned parallel to the electron beam may be accurately analyzed using ASTAR. The reason is simply that for nonparallel twins the incident electron beam will pass through multiple twins and thus give rise to double diffraction. This will make the recognition of the diffraction pattern impossible since it does not correspond to an actual crystal orientation. Still ASTAR can be a very useful technique for studying martensite formation, since it allows for orientation and phase mapping in materials with large deformations: i.e. martensitic steels. For example, the martensite-martensite orientation relationships may be calculated from the ASTAR data. In the study by Stormvinter et al. [3] this was done for the martensite units annotated A, B and C in Figure 5.9.

Figure 5.8: TEM micrographs of the 1.67C steel. (a) BF micrograph of martensite units surrounded by retained austenite. The midrib of the central unit has been indicated by an arrow. In addition, two kink couplings with adjacent units have been indicated by arrows. (b) High-magnification DF micrograph on the midrib of the martensite unit indicated by the arrow in a. The figure is reproduced from [3].

Figure 5.9: IPF colored map of martensite obtained by ASTAR. The area corresponds to the Figure 5.8a where it is shown in BF. The insert displays an IPF colored map on the midrib of the martensite unit indicated by the arrow in Figure 5.9a. The figure is reproduced from [3].

Based on the microstructural observations in 1.20 and 1.67C steels Stormvinter et al. [3] presented a schematic on the development of martensite in high-carbon low alloy steels, as seen from Figure 5.10. They concluded that development of plate martensite seems to occur as follows: (A) In the first stage of the transformation, the growth occurs by formation of {112}α twins. (B) As the martensite unit continues to grow there is a build-up of elastic stresses, which are represented by the dark-blue area in the surrounding matrix. These elastic stresses may be partly relieved and become more uniform by various accommodation components. The accommodation may occur by dislocation motion, in the martensite unit or in the surrounding austenite. The area of the martensite units, which are mainly affected by these accommodations, is colored light-green in Figure 5.10. If the transformation temperature is low, and as found in the 1.67C steel, accommodation may occur through {101}α twinning/faults, which have been illustrated with black lines. (C) The martensite transformation progresses by autocatalytic nucleation of new martensite units, which may have a different shape strain and orientation relationship to the parent phase. Commonly these new units belong to the same plate group and make a well-defined coupling to the preceding unit, which may be of kink- or wedge-type. The martensite units from the same plate group in Figure 5.10 “(C)” may be viewed as a “cluster”. Hence, it may be the result of a single nucleation event with successive autocatalysis of new martensite units.

It is realized that the formation temperature and alloying content is important to consider when discussing the development of plate martensite. This may be illustrated by the variety of plate martensitic microstructures found in the literature, i.e.: butterfly- [17], [18], lenticular- [21], [22], and thin-plate martensite [19], [109]. On the other hand, when the development of plate martensite is seen from a more general perspective some common principles may be distinguished: (1) Initial growth of a midrib, which has a completely twinned substructure. (2) Transverse thickening of the midrib, given that plastic accommodation can occur by either deformation twinning or dislocation formation. (3) Subsequent autocatalytic nucleation of other martensite units from the same plate group. The variety of observed microstructures may be explained by the conditions for these steps to occur, which most likely are governed by temperature and alloying content.

Figure 5.10: Schematic on the development of martensite in high-carbon low alloy steels. (A) Represents the first stage of transformation, the growth occurs by formation of {112}α twins. (B) As the martensite unit continues to grow there is a build-up in the elastic stresses, which are represented by the dark-blue area in the surrounding matrix. The light-green area belongs to the martensite crystal and has been affected by accommodations. However, if the transformation temperature is low, accommodation may occur through {101}α twinning/faults, which have been illustrated with black lines. (C) The martensite transformation progresses by nucleation of new martensite units, which have a different orientation relationship to the parent phase. These new units may form a kink coupling or a zig-zag pattern with the preceding unit. The figure is reproduced from [3].

5.2 EBSD – Variant pairing tendency of martensite in Fe-C alloys

In the previous sections the emphasis has been on qualitative martensite characterization of the transition from lath to plate martensite [1] and the development of plate martensite in high-carbon steels [3]. In this final section on martensite characterization the emphasis is on quantitative martensite characterization by EBSD. The discussion is based on two of the appended supplements; of which one introduces the methodology and the feasibility for studying tetragonal bct-martensite [4] and the second presents the effect of carbon content on the variant pairing tendency of martensite in binary Fe-C alloys [5].

As discussed in Chapter 3 (Section 3.2.1) EBSD is an analysis technique that relies on detection of the local crystallography of a specimen. The collected data can be used to determine the phase as well as the crystal orientation. Recently, Miyamoto et al. [110] developed a novel method to accurately determine the OR between austenite and martensite that is based only on the martensite crystal orientations acquired by EBSD [110]. Stormvinter et al. [4] extended this methodology to study martensite formation with large tetragonality formed in Fe-C alloys. For this purpose they selected high-purity binary alloys having the nominal compositions of 0.35, 0.75, 1.05, 1.80C and an interstitial free (IF) steel. The chemical compositions of the alloys together with the Ms temperature, calculated according to Stormvinter et al. [2], are listed in Table 5.2.

Figure 5.11a displays IPF colored martensite units in the 1.80C alloy, which all belong to a single parent austenite grain with no annealing twins. To interpret the martensite microstructure solely based on IPF colored maps is rather difficult. However, these martensite units display a very characteristic pole figure (PF) pattern as seen from Figure 5.11b. The figure displays a standard

<001>α-BCC stereographic projection of experimental data points, which have been superimposed by the calculated nodes obtained by numerical fitting [110]. This pattern can be compared to the well- known K-S OR [34]. In agreement with the K-S OR, Figure 5.11b displays 24 unique martensite orientations, which may be seen from the clustering of 8 martensite variants around every <001>γ.

This circle pattern around every <001>γ is called a Bain circle [111].

Moreover, Stormvinter et al. [4] discussed how the tetragonality of high-carbon martensite would influence the EBSD data analysis, since in general martensite is treated as bcc rather than bct when conducting EBSD analysis. They first attempted indexing of martensite as tetragonal bct on the

1.80C alloy, since it has the highest c/a ratio. Figures 5.11c and 5.11d from the 1.80C alloy display

[001]α-BCT and [010]α-BCT/[100]α-BCT pole figures of the martensite units displayed in Figure 5.11a when indexed as tetragonal bct instead of cubic bcc. As can be seen from Figure 5.11c, the [001]α-BCT crystal direction could be separated from the [010]α-BCT and [100]α-BCT crystal directions. Figure 5.11d on the other hand contains data belonging to both [100]α-BCT and [010]α-BCT since they are indistinguishable.

Taking Bain correspondence ([1-10]γ // [100]α-BCT, [110]γ // [010]α-BCT, [001]γ // [001]α-BCT) into account, it can be understood that only [001]α-BCT should be contained in the Bain circle. Stormvinter et al. [4] attempted bct indexing on the 0.35, 0.75 and 1.05C alloy as well, and the success rate of bct indexing, can be seen in Figure 5.11e. The success rate is defined as a ratio between the number of data points, whose [001]α-BCT is contained in the Bain circle, and the total number of data points. It is clearly seen that the success rate decreases with decreasing tetragonality and reaches 1/3 at 0.35C, which would be expected from mathematical statistics for the random case. Thus, it could be concluded that accurate indexing of martensite as tetragonal bct was only feasible for the 1.80C alloy of the five alloys studied. In addition, Stormvinter et al. [4] showed that bct indexing of the EBSD data has a minor effect on the OR analysis and thus the simplified bcc indexing is feasible for all carbon contents.

Table 5.2: Chemical composition, calculated Ms and PAGS for the studied alloys. C Mn B Fe Calc. Ms PAGS 0.0023 1.48 0.0027 Bal. 726K 320μm 0.345 0.02 - Bal. 672K 200μm 0.74 <0.003 - Bal. 531K 370μm 1.06 <0.003 - Bal. 472K 300μm 1.80 0.004 - Bal. 363K 250μm

Figure 5.11: (a) Displays an IPF colored map of martensite units from one single parent austenite grain, in the

1.80C alloy. Only data points with a confidence index higher than 0.2 are displayed. (b) <001>α-BCC PF of martensite units visible in (a). Experimental data-points have been superimposed by calculated nodes for <001>α-BCC and <001>γ, which were obtained by fitting to the former. Hence, white circles and black crosses are for martensite and austenite, respectively. (c) PF [001]α-BCT and (d) PF [010]α-BCT/[100]α-BCT of martensite units visible in (a) when indexed as tetragonal bct. (e) Variation of success rate of bct indexing as a function of carbon content and tetragonality. The figure is reproduced from [4].

Figure 5.12: The observed change in orientation relationship between parent and product phase with increasing carbon content expressed as deviation angle between austenite and martensite CPP and CPD. The figure is reproduced from [4].

However, when calculating the deviation angle between CPP and CPD of martensite and austenite they found it necessary to consider the martensite tetragonality. This may be understood from the fact that the CPP and CPD will rotate when going from a cubic to a tetragonal lattice, unlike the principal cube axes that will stay parallel. Comparing the deviation angle between CPP and CPD of the parent and product phase is a common way to express the OR. In Figure 5.12 the effect of carbon content on the OR, as reported by Stormvinter et al. [4], is displayed and it is seen how the OR shift from close to G-T OR towards K-S OR as the carbon content increases. Moreover, these results were found consistent with previously reported OR on low-carbon [110] and high-carbon steels [34].

There may be several factors affecting the OR, such as lattice parameters of martensite and austenite, austenite strength for accommodation of shape strain and substructure in martensite arising from lattice invariant shear (LIS) mode. Oka and Okamoto [112] made theoretical predictions of the OR for an 1.8C alloy by using phenomenological theory of martensitic crystallography (PTMC) with the assumption of two different LIS systems (112)[-1-11]α and (101)[10-1]α, which correspond to the

ORs shown as A and B in Figure 5.12, respectively. (112)[-1-11]α is the ordinary LIS for (112)α twin and it appears in most of plate martensite in ferrous alloys while (101)[10-1]α corresponds to the

(101)α twin reported only in high carbon martensite [13]. The prediction by Oka and Okamoto [112], A and B in Figure 5.12, illustrate the significant effect of LIS on OR. It is interesting that the measured ORs for high carbon alloys agree with the OR B whereas ORs for lower carbon martensite are rather close to OR A. Stormvinter et al. [4] did however not investigate the mechanism for the observed change in OR in more detail.

Instead, Stormvinter et al. [5] used the experimentally determined OR to investigate the effect of carbon content on the variant pairing tendency of martensite formed in these five alloys. Variant pairing tendency implies quantification of the nearest neighboring of martensite units. In most cases the number of crystallographically unique martensite-martensite neighbors is limited to 24. The misorientation, and thus boundary character, of such neighbors may be calculated from the experimentally determined OR. An EBSD dataset, for example as seen from Figure 5.13, may contain several thousands of martensite-martensite boundaries and is therefore ideal for variant pairing determination. Stormvinter et al. [5] showed how the quantified variant pairing can be interpreted and transposed into knowledge of the martensite microstructure.

Figure 5.13: IPF colored EBSD map taken from the IF steel. The figure is reproduced from [5].

Figure 5.14: Variant pairing frequency diagrams which display the length fractions of inter-variant boundaries for (a) IF steel, (b) 0.35C alloy, (c) 0.75C alloy, (d) 1.05C alloy, and (e) 1.80C alloy. The error bars represent the standard deviation of three measurements, each acquired from an area of 800x265µm. Reproduced from [5].

The results from the analysis [5] of martensite-martensite boundaries are presented in Figures 5.14a- 5.14e, and show the length fraction of variant boundaries in relation to variant 1. As seen from the figures there is a clear dominance of intra close-packed plane group (CP group) variant pairing at lower carbon contents. Moreover, there is a shift from V1/V4 dominance towards V1/V2 dominance as the carbon content is increased. However, the variant pairing tendency of martensite formed in the 1.80C alloy is in strong contrast, as seen from Figure 5.14e, to that of the other alloys. For this high-carbon alloy, pairing of V1 with variants V2, V6, V16 and V17 seems to be more frequent, with a dominance of V1/V16 pairing. This tendency for V1/V16 pairing is also observed in the 0.75 and 1.05C alloys, although not as pronounced as in the 1.80C alloy.

Figure 5.15: IPF colored EBSD map taken from the IF steel. The map shows an area with a dominant martensite packet and blocks and sub-blocks have been labeled with its corresponding martensite variants. EBSD data from the region enclosed by the yellow box was used for the analysis by standard stereographic projection. Misorientation boundaries are colored as follows: (White) 4°-13° and (Black) > 19°. The figure is reproduced from [5].

As seen from Figure 5.14a, martensite formed in the IF steel is dominated by V1/V4 pairing. This particular variant pair, V1/V4, has previously been associated with sub-block boundaries of low misorientation [42], [113], [114]. The packet and block martensite microstructure is rather obvious in the IF steel, as seen from Figure 5.1a. Another noticeable feature of the IF steel is seen from Figure 5.15; the block boundaries are few when compared to the sub-block boundaries. In fact, the large block-size of the IF steel may be seen from Figures 5.1a and 5.13 as well. Hence, as the sub-block boundaries, as seen from Figure 5.15, dominates the microstructure it is realized that the V1/V4 boundary also is the most frequently observed, as seen from Figure 5.14a. The smaller block- size/sub-block-size ratio for the 0.35C explains why the V1/V4 variant pairing is not found to be as dominant as for the IF steel, although it displays martensite formed in packets and blocks as seen from Figure 5.1b. In the 0.35C alloy all martensite units that belong to the CP group shows up in the variant pair frequency analysis. Nevertheless, the sub-block boundaries, V1/V4, are the most frequent.

The 0.75C alloy, with the close to eutectoid composition, does not display a typical lath martensitic microstructure, since packets and blocks cannot be easily identified as seen from Figure 5.1c. However, when the EBSD data is used to color martensite units according to their CP group association the packets become more evident, as seen from Figure 5.16a. Moreover, the martensite units, a few microns in size, seem to be separated primarily by high-angle boundaries in this alloy. Figure 5.16b displays an IPF colored map of a limited number of such martensite units, taken from a

47 single CP group (CP1 in Figure 5.16a). These martensite units have been labeled with their corresponding martensite variant, V1 through V6, and boundaries are colored as follows: Yellow – Twin (V1/V2; V3/V4; V5/V6), White – Low angle (V1/V4; V2/V5; V3/V6), and Black – High angle (All other combinations). High angle boundaries, preferably twin boundaries (yellow), seem to dominate in this alloy, although there are a considerable fraction of low angle (white) boundaries as well. This is also seen from the variant pairing analysis, displayed in Figure 5.14c. The results for the 0.75C alloy may be compared with the results for a 0.61C alloy by Morito et al. [42]. Both results display variant pairing of martensite units that belong to the same CP group. However, no tendency for block or sub-block formation is observed in either of the alloys. These similarities indicate that an increase in carbon content from 0.61 to 0.75C does not affect the structure significantly. The observed strong peak of the V1/V2 variant pairing in Figure 5.14c is an important observation. If the martensite in the 0.75C alloy may be classified as lath martensite, it indicates that a lower formation temperature would promote V1/V2 variant pairing. This temperature dependence of variant pairing has recently been observed by Takayama et al. [114] in low-carbon bainite. They found that a lowered formation temperature promotes V1/V2 pairing. This variant pairing is unique since it is the only pairing that is twin-related. In addition, a recent EBSD work [115] by Naraghi et al. on athermal martensite formed in stainless steels also supports this observation of twin-related variant pairing at low transformation temperatures.

Figure 5.16: Displays EBSD data from the 0.75C (a) Martensite units from a single prior austenite grain colored according to CP group belonging. (b) IPF coloring of martensite units from a single CP group. The martensite units have been labeled with their corresponding crystallographic variant. The boundaries are colored as follows: Yellow – Twin (V1/V2; V3/V4; V5/V6), White – Low angle (V1/V4; V2/V5; V3/V6), and Black – High angle (All other combinations). The figure is reproduced from [5].

In the 1.80C alloy the variants V1, V6, V16 and V17, which all belong to the same plate group, tend to be paired more frequently, as seen from Figure 5.14e. The dominance of plate group variant pairing is supported by the previous studies on variant selection of plate martensite [19], [116], [117]. The strong coupling between the plate group martensite variants may also be seen from Figure 5.17. Here the EBSD-data from the 1.80C alloy is displayed as an IPF colored map, where the martensite variants are colored as follows: V1 – Orange, V6 – Dark blue, V16 – Pink, and V17 – Light blue. According to Okamoto et al. [19], variant pairs V1/V6, V1/V16 and V1/V17 correspond to wedge-, kink-, spear-type respectively. They reported that wedge and spear variant pairs are favored due to self-accommodation, while kink type is not. Therefore, Stormvinter et al. [5] supposed that the observed dominance of V1/V16 pairing (kink type) could be due to the autocatalytic plate formation.

The results from the variant pairing analysis [5] of 1.05C displayed in Figure 5.14d may be interpreted as a transition between CP group pairing and plate group pairing. This conclusion agrees well with the observed microstructure obtained using OM, as seen from Figure 5.1d. In the figure it can be seen that large plate-like martensite units, with transverse micro-cracks, are embedded in a matrix of smaller martensite units.

Figure 5.17: IPF colored EBSD map taken from the 1.80C alloy. EBSD data from the region enclosed by the yellow box was used for the analysis by standard stereographic projection. Martensite variants are IPF colored as follows: V1 – Orange, V6 – Dark blue, V16 – Pink, and V17 – Light blue and the white regions correspond to retained austenite. It should be noted how the transverse thickening of the primary plates seems to occur by autocatalytic nucleation of new martensite units on the former. The figure is reproduced from [5].

Based on their results, Stormvinter et al. [5] presented three schematics to illustrate the observed change of martensite morphology with increasing carbon content. These schematics are presented in Figures 5.18a-5.18c. Martensite formation in the IF steel is represented by Figure 5.18a, where a prior austenite grain has transformed into a martensitic packet and block structure. In this figure one CP group has been highlighted and it contains three blocks, which are colored according to their corresponding Bain group association. The blocks contain sub-blocks, which are separated by boundaries colored in light gray. Figure 5.18b displays formation of CP groups in the Fe-0.75C alloy. The microstructure within every CP group would be very similar to Figure 5.16b, i.e. preferably V1/V2 twin-type boundaries between the individual martensite variants. The absence of blocks and sub-blocks seems to be the characteristic feature of this type of medium carbon lath martensitic microstructure. The well-known characteristic hierarchy of lath martensite is in strong contrast to the plate martensite formed in the Fe-1.80C alloy displayed in Figure 5.18c. Here the development of martensite into plate groups is illustrated. This occurs by subsequent autocatalysis of variants which belong to the same plate group. Every plate group contains four martensite variants and in total there may be six plate groups present in one prior austenite grain (when disregarding annealing twins). However, to make the microstructural interpretation easier, only one out of the six plate groups are displayed here. This also seems reasonable since it is not clear from the present analysis how the individual plate groups will be distributed within the prior austenite grain.

Figure 5.18: Schematics of martensite formation in (a) IF steel, (b) Fe-0.75C alloy and (c) Fe-1.80C alloy. The region composed by variants belonging to the same Bain group is represented by the same color. The figure is reproduced from [5].

5.3 Transition between lath and plate martensite

In the present thesis it has been demonstrated how different analyzing tools may provide information to get a more fundamental understanding of the martensite microstructure. This includes: (1) OM and classic SEM to analyze apparent morphology of the martensite units. (2) Classic TEM to get detailed information on the substructure as well as the possibility to do more detailed crystallographic analysis. (3) Automated crystallographic techniques, such as EBSD and ASTAR, to analyze orientation relationships as well as martensite-martensite boundary characteristics, i.e. variant pairing tendency. When interpreting the combined knowledge gathered [1], [3–5] it seems clear that the complexity of the martensite morphology has to be resolved in multi- length scale approach, where all these techniques are combined. Moreover it also seems rather clear that the transition from lath martensite towards plate martensite is gradual and that the two morphologies makes good reference points for interpreting a material in the mixed regime. In addition, the recently developed variant pairing analysis may prove a good compliment to traditional metallography, since it allows for quantification of boundary character, and thus martensite character. This is well illustrated in Figure 5.19 that shows how the variant pairing tendency changes with the carbon content. This EBSD technique has a clear advantage over ASTAR, since it can be applied to an area large enough to be representative for the whole specimen. On the other hand, ASTAR has its advantages since it can very well reproduce the local microstructure and allows for on-line interchanging to classic TEM. In conclusion, by combining the strengths from various characterization techniques it is possible to get a good understanding of the morphological transition in ferrous martensite with increasing carbon content, although it may still prove very challenging and time-consuming to perform the characterization.

Figure 5.19: Variant pairing tendency given as length fraction of variant boundary versus carbon content and based on the experimental results from the present work. The figure is reproduced from [5].

6 Discussion on inverse bainite

As mentioned in Chapter 2 (Section 2.2), there has been a long debate concerning the nature of the bainite formation and whether it can be regarded as diffusionless or diffusion controlled. The conflict has been on-going since the 1940s when Zener [56] presented a comprehensive review of the mechanisms involved in the austenite decomposition. The experimental results by Ko and Cottrell [53], that bainite grows with a surface relief, were taken as a strong indication of a diffusionless nature. However, as pointed out by Hillert [118] the surface relief is an indication of the displace nature of the transformation, i.e. the fcc lattice is transformed into the bcc lattice in a military fashion. It does not imply that such transformation omits diffusion of carbon, especially for a transformation with slow growth rate as for instance the bainite transformation.

Recently, Borgenstam et al. [6] studied the decomposition of high-carbon austenite with the purpose to test the symmetry among the diffusional transformation products in the Fe-C alloys. The term symmetry was here used to define the relation between the eutectoid decomposition products as follows: “Pearlite is the result of cooperative growth of ferrite and cementite where the two phases behave as equal partners. It is favored by a close-to-eutectoid composition. Bainite forms with ferrite as the leading phase in the main growth direction, which is identical to the growth direction of Widmanstätten ferrite. Bainite is favored by a low carbon content. Inverse bainite forms with cementite as the leading phase in the main growth direction, which is identical to the growth direction of Widmanstätten cementite. Inverse bainite is favored by a high carbon content.”. The concept of inverse bainite was first introduced by Hillert in 1957 [58] but has been discussed only a few times during the elapsed 50 years, mainly by Aaronson and co-workers [119–121]. Evidently, the diffusionless mechanism cannot be applied to inverse bainite, since it is not possible to imagine the growth of a low-carbon cementite plate that later enriches in carbon by the onset of a diffusional process.

Borgenstam et al. [6] based their study on a large number of micrographs from isothermal treatments of a high-carbon steel. Below the A1 temperature they observed austenite decomposition into pearlite in a Fe-1.67C-0.22Mn-0.27Si-0.47Cr. However, below 773 K (500 °C) pearlite was replaced by an acicular eutectoid with cementite as the leading phase. An example of these “spiky nodules” can be seen from the micrograph by Modin and Modin [122] presented in Figure 6.1.

Figure 6.1: Spiky nodules of inverse bainite formed after 60 s at 723 K (450 °C) in a steel with 1.65C. OM micrograph at 600 times magnification. Reproduced from Modin and Modin [122].

Figure 6.2: Acicular eutectoid (inverse bainite) between plates of cementite, formed after 1 s at 823 K (550 °C) in steel with 1.34C. The arrows mark inverse bainite ahead of the cementite plate. TEM micrograph on plastic replica at 12,000 times magnification. Reproduced from Kinsman and Aaronson [119].

As the transformation temperature was lowered further the “spikes” of the nodules became less pronounced and the eutectoid structure appeared more “fan-like” or columnar. At even lower transformation temperatures the eutectoid structure again appeared spiky with protruding tips. However, Borgenstam et al. [6] could show that this was in fact a gradual shift from cementite being

54 the leading phase of the hypereutectoid structure towards ferrite as the leading phase. This may be illustrated by Figure 6.2 that shows how cementite is dominating the growth of inverse bainite in the upper temperature range. In contrast, Figure 6.3 displays a hypereutectoid structure, transformed at much lower temperature, where the growth has been dominated by the ferrite. From the figure it can be seen how the shorter carbide units have been nucleated and grown from a primary ferrite plate. Hence, the carbides are formed subsequently. This implies that as the isothermal transformation temperature is lowered the growth of the hypereutectoid transformation structure shifts from a cementite dominance (Figure 6.2) towards a ferrite dominance (Figure 6.3), i.e. there is a transition from growth of inverse bainite towards ordinary bainite as the transformation temperature is lowered. This is not so unlikely when taking the asymmetry of the Fe-C phase diagram into account.

As can be seen from Figure 6.4, the extrapolated A3 and Acm show that the supersaturation with respect to ferrite increases more rapidly than to the supersaturation with respect to cementite. The supersaturation may be interpreted as a driving force and hence the nucleation and growth of ferrite becomes more likely as the temperature decreases.

Figure 6.3: Two thin spikes with a central plate of ferrite (white) and shorter plates of cementite (dark). Formed after 1 h at 573 K (300 °C) in steel with 1.67C. Etched for carbides. Additional etching in nital was applied in the insert. Two arrows mark the same platelet of cementite. SEM at 12,000 times magnification. Reproduced from Borgenstam et al. [6].

Figure 6.4: Extrapolated solubility lines for ferrite and cementite in Fe-C austenitee [9], [64]. The figure is reproduced from [6].

Most of the changes of microstructure of the eutectoid products in high-carbon steels with temperature are gradual and may be explained by the asymmetry of supersaturation with respect to ferrite and cementite at decreasing temperature. Moreover, Borgenstam et al. [6] showed that inverse bainite, which forms by a diffusion controlled process, has many similarities to ordinary bainite. For instance, they both form with a shape change that gives rise to a surface relief. However, the bainite controversy will probably survive until a key experiment accepted by the proponents of both sides can resolve the true nature of the transformation mechanism.

7 Thermodynamically based prediction of the martensite start temperature

Some years ago, Borgenstam and Hillert [123] reviewed a comprehensive amount of experimental Ms temperatures of binary and ternary ferrous alloys and made some noticeable distinctions in their final assessment: (a) They used experimental Ms values directly and took them mainly from studies using very rapid cooling; (b) They distinguished between data for two kinds of martensite yielding separate plateaux in the cooling curves, where plateau III was supposed to belong to lath martensite and plateau IV to plate martensite; (c) Results for Fe-Co were included which was an important addition because it expanded the temperature range to above the value for pure Fe. They also proposed the following expressions graphically for the requirement of driving force to start a martensitic transformation into either lath or plate martensite:

( ) [J/mol], [Eq. 7.1]

( ) [J/mol], [Eq. 7.2]

The concept of two separate Ms temperatures, one for lath- and one for plate martensite, is originating from the works by Morozov et al. [61], [75] who studied how the decomposition of austenite in pure Fe and in Fe with 0.01C was affected by the cooling rate, from moderate to extremely rapid cooling. They observed four transformation plateaux, as seen in Figure 7.1, where plateaux III and IV were identified as caused by lath and plate martensite, respectively. According to that information one should distinguish between the Ms temperatures for lath and plate martensite, supposedly 818 K (545 °C) [61] and 693 K (420 °C) [75] for pure iron.

Figure 7.1: Beginning of the γ → α transformation in iron as a function of cooling rate, redrawn from Morozov et al. [75]

In an attempt to construct a sound thermodynamic method to predict Ms of commercial steels, Stormvinter et al. [2] accepted that the two expressions, Equations 7.1 and 7.2, for the necessary driving force to form lath and plate martensite by Borgenstam and Hillert [123] holds for pure Fe.

They then referred experimental data on Ms for binary alloys, Fe-X(= C, Cr, Mn, Ni), from rapid cooling experiments [70], [74], [76], [78], [124–127] to either plateau III or IV. From this information they calculated the driving force at Ms and the difference from Equations 7.1 and 7.2 was accepted as an effect of the alloying element and was evaluated as proportional to the content of each alloying element. A linear superposition law was assumed for the combined effect of alloying elements on the driving force, yielding the following expressions to be used for predicting the Ms temperature of commercial steels (compositions given in mole fraction):

( ) ( )

, [J/mol] [Eq. 7.3] ( )

( ) ( )

, [J/mol] [Eq. 7.4] ( )

In the analysis by Stormvinter et al. [2] Zener ordering of carbon atoms was taken into account, since it will have a dramatic impact on the Gibbs energy of the fresh martensite at high carbon contents.

In order to use Equations 7.3 and 7.4 to predict Ms of a commercial steel, it is necessary to find the temperature where the available driving force for the alloy composition equals the required driving force, i.e. the transformation barrier. To validate their proposed thermodynamic method for prediction of Ms, they applied it to a number of datasets available in the rich martensite literature

[67], [68], [128], [129]. Both Equations 7.3 and 7.4 were used and the highest predicted Ms temperature was accepted for each steel. For steels where Equation 7.3 (lath) gave the highest predicted Ms, open symbols were used in the figures, while filled symbols were used when Equation 7.4 (plate) gave the highest predicted Ms. For example, Figure 7.2 displays a comparison of the predicted Ms by the model and the experimentally determined Ms by Grange and Stewart [67] and by Steven and Haynes [68].

Figure 7.2: Comparison between experimental Ms by Grange and Stewart [67], and by Steven and Haynes [68] and predicted Ms temperatures. The figure is reproduced from [2]. Open symbols: lath martensite and filled symbols: plate martensite.

The thermodynamic model proposed by Stormvinter et al. [2] has been used to develop a simple and user-friendly software for calculation of Ms. This software is for the moment exclusive for the Hero- m research center partners and may therefore not be distributed outside Hero-m yet. A screenshot of the user-interface can be seen in Figure 7.4. It is the belief of the present author that this software will be developed further, for instance to include the effect by PAGS on Ms.

Figure 7.3: User-interface of the computer program for calculating Ms developed within the Hero-m center.

8 Concluding remarks and future prospects

The martensite transformation in steels has a substantial amount of complexity. In addition, it is also one of the most important phase transformations to account for in process development and materials design. Some important aspects of the martensitic transformation have been thoroughly discussed in the present thesis. Those include: (1) The quantification of the martensite microstructure and how that may be used to classify martensitic steels. (2) The feasibility of developing a predictive model for Ms that is based on the current state-of-the-art thermodynamics. Both these subjects may have a true industrial significance, since it may aid engineers in alloy design. However, there are still challenges left before the full potential of both these methods can be fully realized. A number of these challenges will be listed here:

 The limited amount of experimental Ms data from binary and ternary alloys that is relevant for steels. It is evident that additional experimental data would make further improvements of the

thermodynamic model as developed here [2], since it in principle relies on experimental Ms data from binary and ternary alloys. However, it is important to keep in mind that such data should contain information of experimental conditions, for instance cooling rate and PAGS. Very often this kind of additional data is left out when the final experimental data is presented.

 Improved thermodynamic databases, which are optimized for low temperatures. It is likely that the never-ending increase in computational power will further fuel the capabilities of doing computer based experiments. Ab initio and related techniques have already proved that they can provide reliable data that may be extremely difficult to obtain through an ordinary experiment. This new information may be used together with experimental data to optimize thermodynamic databases to handle low temperature phase transformations, as the martensitic for instance.

 Improvement of the commercially available EBSD software. For instance, to include easy-to-use post-processing software to analyze orientation relationships and variant pairing. It is most probable that EBSD analysis will become more sophisticated since this is a recent developed area. Especially if one considers the growing popularity of FEG-SEMs equipped with EBSD detectors.

 There has been a recent interest in three-dimensional (3D) characterization of materials, for instance TMS (The Minerals, Metals and Materials Society) hosted an international conference (3DMS in Seven Springs, Pennsylvania) dedicated to the topic in the past summer (2012). It is likely that 3D characterization will add substantial information to the structure-property relationships of materials. 3D characterization includes, among other techniques, 3D EBSD that allows for detailed analysis of the crystallography of a finite volume. 3D EBSD is already commercially available and may be applied to study morphology, boundary character, and variant grouping of martensitic steels.

 The growth kinetics of martensite may not yet be categorized as “well-known”. Obviously the rapid transformation rate has made it difficult to study how the transformation progresses. On the other hand, synchrotron sources and detector capabilities have improved significantly since most of the experiments were conducted. It is therefore likely that it would be feasible to study the growth kinetics using a synchrotron source to gain more insight to this phase transformation. Steels forming lath martensite would probably be the best candidate to choose, since experiments indicate that the growth rate can be several orders of magnitude lower when compared to the growth rate of plate martensite.

9 Acknowledgments

First of all I would like to express my sincere gratitude to my principal supervisor Assoc. Prof. Annika Borgenstam and co-supervisor Prof. John Ågren for giving me this opportunity to do research on some very interesting aspects of physical metallurgy. I do much appreciate how you encourage me to find and develop the researcher within me, although always prepared with guidance if needed. I would also like to express my deepest gratitude to my second co-supervisor, Dr. Peter Hedström, for non-depleting enthusiasm as well as professional assessment of my work.

I would like to acknowledge Prof. Yves Brechét and his co-workers at INP-Grenoble for a superb spring in 2010. Also Prof. Tadashi Furuhara and his co-workers should be sincerely acknowledged for accepting me into their lab in the summer of 2011. A special thanks to Dr. Goro Miyamoto for his never-ending patience in my struggles to learn crystallography of martensite.

I would also like to acknowledge my colleagues at KTH for providing a welcoming, professional and inspiring environment for research. A special thanks to Dr. Lars Höglund for being very helpful with computer related challenges.

I would also like to thank the VINN Excellence Center Hero-m, VINNOVA, Swedish Steel Industry, and KTH Royal Institute of Technology for providing finance for performing this work. As well as Stiftelsen Axel Hultgren fond, Ångpanneföreningens Forskningsstiftelse, Stiftelsen Bergshögskolans jubileumsfond, Knut och Alice Wallenbergs stiftelse, and Stiftelsen De Geerska fonden for providing travel grants.

Finally, I would like to express my deepest gratitude to my family for their support and encouragement. I am also in debt to all my friends for giving me such a wonderful time; you know who you are (Ingen nämnd – Ingen glömd).

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