Physical Metallurgy of Thermomechanical Treatment of Structural Steels

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Physical Metallurgy of Thermomechanical Treatment of Structural Steels Mazanec, Karel.; Mazancova, E. Cambridge International Science Publishing 1898326436 9781898326434 9780585237466 English Physical metallurgy, Steel--Metallurgy. 1997 TN693P49 1997eb 669/.96142 Physical metallurgy, Steel--Metallurgy.

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Physical Metallurgy of Thermomechanical Treatment of Structural Steels Karel Mazanec and Eva Mazancová

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Published by Cambridge International Science Publishing 7 Meadow Walk, Great Abington, Cambridge CB1 6AZ, UK http://www.demon.co.uk/cambsci/homepage.htm First published December 1997 ©K Mazanec and E Mazancová ©Cambridge International Science Publishing Conditions of sale All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopy, recording, or any information storage and retrieval system, without permission in writing from the publisher British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library ISBN 1 898326436 Translated by V E Riecansky Production Irina Stupak Printed by Transtech Printers Ltd, London, England

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About the Authors Prof Ing Karel Mazanec, DrSc graduated from the Technical University in Ostrava in 1949. Since then, he has worked in a number of leading posts, including the Vitkovice Iron Works Research Institute in the area of the development of new types of steel for power engineering, research into the thermomechanical treatment of structural and high-strength steels, their weldability and brittle fracture resistance. In 1962, he defended his DrSc dissertation on 'Delayed fracture in martensite'. Prof Mazanec has published more than 350 studies in many leading journals and conference proceedings. In 1965 he became a professor at the Faculty of Metallurgy of the Technical University in Ostrava and later Dean of the Faculty. In 1983 he was elected a member of the Czechoslovak Academy of Sciences. He is also a honorary member of the French Metallurgical Society. At present, is a Professor Emeritus at the Technical University of Ostrava where he lectures and works in the research of unconventional methods of heat treatment of steel and physical-metallurgical problems of strengthening and increase of fracture resistance. Recently, he has been paying attention to physical-metallurgical problems of thermoelastic martensite in shape memory TiNi alloys. Ing Eva Mazancová, CSc, graduated from the Technical University of Ostrava in 1975 and started her career at the Research and Testing Institute of the New Iron Works in Ostrava where she now works in the structural and microfractographic analysis of structural steels, microanalysis of nonmetallic inclusions and research into segregation phenomena, including their effect of the properties of steel. In 1981, Eva Mazancová defended her PhD dissertation on 'Thermomechanical treatment of structural steels' in which she examined the application of thermomechanical treatment of microalloy steels in seamless tube production. She has also worked on problems of hydrogen and sulphide brittleness of materials for oil industry, the effect of chemical constitution and microstructure on corrosion cracking; problems associated with optimization of the heat treatment of oil pipes with higher strength and their physical-metallurgical characteristics also attracted her attention. Her current interests include problems of microfractographic analysis of structural steels and evaluation of relationships between microstructure, fracture parameters and cleanness of steel. She is also involved in evaluating the microcleanness of continually cast steels and its effect on the properties of structural steels after various heat treatments. During her scientific and research activity, Eva Mazancová has published 105 studies in leading journals and conferences, both home and abroad.

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Contents

1 Introduction

1

2 Main Physical Metallurgy Characteristics of Thermomechanical Treatment

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2.1 Sources of Hardening the Metallic Matrix

2.2 Analysis of Physical Metallurgy Characteristics of the Initial Austenitic Matrix Taking into Account the Subsequent Heat Treatment

2.3 Influence of the State of the Initial Austenitic Matrix on Phase Transitions in Thermomechanical Treatment

2.4 Physical Metallurgy of the Ferritic Phase Transformation in Work-Hardened Austenite 3 Thermomechanical Treatment of High-Strength Martensitic Steels

3.1 Effect of Thermomechanical Treatment on Properties

3.2 Effect of Thermomechanical Treatment on the Modification of Microstructural Characteristics

3.3 Evaluation of the Mechanical and Metallugical Properties of High-Strength Martensitic Steels after HTMT

3.4 Summary of the Achieved Results and Their Physical Metallurgy Analysis 4 Thermomechanical Treatment and Controlled Rolling of Carbon and Low-Alloy Steels

4.1 Application of HTMT in Strengthening the Matrix of Steels for Wider Technical Application and Determination of the Relationships between the Initial Structure and the Resultant Physical Metallurgy Properties of Products of Austenite Decomposition

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4.2 Technical and Technological Characteristics of Controlled Rolling

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4.2.1 Deformation in the Austenite Range Associated with Recrystallization

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4.2.2 Deformation in the Region of Restricted Recrystallization

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4.2.3 Deformation in the AusteniteFerrite (Two-Phase) Region

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4.3 Application of Various Variants of Accelerated Cooling after Controlled Rolling

4.4 A Model of Predicting Changes in the Microstructure in Controlled Rolling and Accelerated Cooling of Structural Steels for Wider Technical Application

4.4.1 Perspective Technical and Technological Solutions

4.5 Effect of Controlled Rolling and Accelerated Cooling on Reducing the Susceptibility to Hydrogen Embrittlement

4.6 Technical and Technological Comments Regarding Application of Controlled Rolling and Accelerated Cooling

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5 Conclusions

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References

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Index

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1 Introduction The development of science and technology imposes new requirements on obtaining higher applied properties of metallic materials, especially higher strength and ductility properties. The main methods based on complex or higher alloying or on the development of metallic materials with new constitutions usually enable higher material and technical properties to be obtained, but this is achieved with higher or more extensive requirements on alloying elements. This increases the production cost. Despite the fact that in many cases this is the only possible approach that can be used, on a wider scale higher applied material and technical properties can be achieved in various types of metallic systems by other methods usually based on the application of new, unconventional heat treatment technologies. These solutions form suitable conditions for the maximum utilization of the processing properties of appropriate types of metallic materials. This has resulted in the development of a large number of unconventional heat treatment methods of metallic materials, some of which are based on using highly complicated technical and processing variants, as discussed in, for example, Ref.1. According to the currently available experience, different variants of thermomechanical treatment (TMT) represent one of the most efficient methods of unconventional heat treatment. This method of unconventional heat treatment can be used widely not only for steels but also in heat treatment of, for example, alloys of aluminium, titanium, nickel, etc.2 In this book, the authors pay special attention to analysis of the physical metallurgy parameters and structuralmaterial characteristics obtained in structural steels or iron-based alloys using various types of TMT. This method makes it possible to obtain higher applied properties of these materials, without any large changes in their chemical constitution. From the viewpoint of developing this heat treatment method, the proposed solution will include not only an analysis of basic variants of TMT in the region of stable and unstable austenite in high-strength martensitic steels and alloys but also the problems of application of TMT in optimizing the properties of low-alloy or carbon structural steels in which the austenite

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to martensite transition does not take place during the phase transformation. Consequently, for these steel types we shall also discuss selective problems of the physical metallurgy of controlled rolling or the characteristics of accelerated cooling, which are closely linked with the problem of the processes of TMT of structural steels. The main feature of all these variants of TMT is that the plastic deformation of the initial austenitic matrix and the formation of its controlled structural and mechanical state with subsequent heat treatment, phase transformation, are linked. This link makes it possible to utilize the hardening or refining of the grains of the initial austenitic matrix and also the structural modification of the products of the phase transformation of austenite associated with the modification of structural and substructural characteristics of the austenitic matrix leading to a higher level of the set of the mechanical and metallurgical properties of various types of structural steel processed by this treatment. The book summarizes the results obtained in extensive research of the hardening characteristics of metallic materials (structural steels), with special attention given to the problems of the physical metallurgy nature of hardening of martensitic steels and control-rolled structural steels or structural steels subjected to subsequent accelerated cooling and used for wider technical applications. Taking this analysis into account, selected problems, whose solution leads to obtaining optimized technical and technological conditions of processing these types of steel by the selected unconventional heat treatment method, will also be discussed. The aim of this book is to present a detailed analysis of the main parameters of unconventional treatment of structural steels resulting in higher strength and ductility properties of these materials, including higher utility properties. This makes it possible to create suitable conditions for obtaining a higher level of safety and reliability of structures or structural components produced from materials treated in this manner. The main technical and technological characteristics of TMT of structural steels of different chemical composition and the conditions of their optimum use will also be determined.

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2 Main Physical Metallurgy Characteristics of Thermomechanical Treatment Information on the controlling role of the structure and the conditions of influencing the so-called structure-sensitive characteristics of metallic materials has resulted in conclusions according to which the considered structure-sensitive properties are strongly affected by the defectiveness of the matrix. It is well known that lattice defects, present in the initial matrix, also strongly influence by both the mechanism and kinetics of phase transformations in heat treatment, i.e. the characteristics of the resultant structure and the final mechanicalmetallurgical properties of this type of metallic material. This shows that the aim of the solution must be to obtain both the favourable density of structural defects, namely dislocations, and their arrangement in the metallic matrix in various stages of TMT in order to obtain, in controlling the appropriate processes, the required structure and, consequently, the optimum level of the final mechanical and metallurgical characteristics. Plastic deformation and subsequent recovery processes in the parent metallic matrix are the main metallurgical processes enabling control of the dislocation density in the metallic matrix. These considerations then lead logically to a simple conclusion, according to which the most efficient method is to combine efficiently the appropriate physical metallurgy characteristics of phase transformations and plastic deformation (including control of the subsequent or simultaneous development of recovery and recrystallization processes) and the formation of a single technological unit. TMT is then a process resulting in the detailed coupling of controlled modification of the structure in deformation of the initial austenitic matrix and its decomposition in heat treatment. As indicated by previous studies concerned with the analysis of TMT conditions of various types of steels, this process must be regarded as a set of partial processes taking place during heating, deformation and cooling of metallic materials. This results in the formation of a structure enabling properties to be obtained similar to those which can be achieved in, for example, materials at a higher density and specific arrangement of lattice defects gen-

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erated in plastic deformation in the metallic matrix prior to the final phase transformation.1,2 Initial studies in the area of research into the physical metallurgy and technical and technological characteristics of TMT in, for example, structural steels, were aimed at determining the conditions for obtaining the maximum level of strength whilst maintaining the required toughness level.1,3 Therefore, special attention was given to two main TMT variants: evaluation of the technical and technological parameters of TMT in the stable austenite range (high-temperature thermomechanical treatment, HTMT), and thermomechanical treatment in the unstable austenite range (low-temperature thermomechanical treatment, LTMT). In the former case, plastic deformation is carried out in stable austenite, whereas in the latter case plastic deformation takes place in unstable austenite, usually connected with subsequent quenching to martensite1 leading to a large increase of the mechanical and metallurgical properties of alloyed high-strength martensitic steels. During the subsequent development of TMT technologies, the coupling of plastic deformation with subsequent heat treatment was also extended to carbon and low-alloy steels, where the product of final heat treatment is usually the formation of a ferriticbainitic, bainitic or bainiticmartensitic structure.4 In all these cases, this is achieved by the rapid cooling of appropriate types of structural steels with the defined parameters of the initial austenitic structure, e. g. in a quenching press.1 In addition, in the further development of unconventional methods of heat treatment of structural steels, as a part of a wider concept of technical and technological variants of TMT, work is being carried out on technical and technological variants of controlled rolling or the superposition effect of accelerated cooling. This method is used mainly in lowcarbon or microalloy steels for wider technical applications alloyed with, for example, niobium, vanadium, titanium, or with several microalloying additions. Usually, the phase transformation of the austenitic matrix is accompanied by the formation of non-martensitic products, with a special position in optimizing the resultant properties played by the optimum form of the resultant ferrite grains. The addition of microalloying elements, together with the modification of the reduction schedule in forming and the selection of holding time between passes (control of the temperature of completion of forming and the development of recovery processes in the parent metallic matrix austenite) has a beneficial effect on the resultant austenite grain size and the development of recrystallization processes (static and dynamic recrystallization) and the extent of strain-induced precipitation of carbides or carbonitrides of microalloying additions.5

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These approaches will be used in subsequent parts of this work to discuss the main physical metallurgy sources of hardening the metallic matrix under the conditions of TMT of high-strength martensitic steels and low-alloy or carbon steels. The strengthening conditions and the conditions of increasing the level of the toughness parameters in controlled rolling or subsequent accelerated cooling of carbon and microalloy structural steels for wide technical applications will be analyzed separately. 2.1 Sources of Hardening the Metallic Matrix In this section, we shall analyze the main physical metallurgy sources of hardening the metallic matrix of structural steels which are linked closely with the application of the discussed methods of unconventional heat treatment, in highstrength (martensitic) steels, low-alloy or carbon steels processed using conventional TMT techniques and the physical metallurgy conditions leading to strengthening of the metallic matrix under controlled rolling or accelerated cooling after rolling structural steels for wide use (carbon or microalloy structural steels). The shear modulus G of plain iron reaches around 84 000 MPa, but it must be considered that in iron solid solutions the presence of the impurity elements usually reduces this value. On the basis of a simplified consideration of the so-called homogeneous shear in a perfect crystal, the theoretical shear stress, required for inducing plastic deformation in alpha iron, can be expressed by the following simple equation:

where the value of b for the atom spacing in the α direction is 0.24 nm, and the value of d for the distance of the {110}α planes reaches 0.202 nm. Equation (1) shows that the theoretical value of the shear stress in this case corresponds to approximately (G/10). This theoretical level of the strength characteristics can be obtained approximately in a thin, highly perfect iron whisker. In the presence (generation) of dislocations in the iron whisker the strength level rapidly decreases and plastic deformation can take place under considerably lower stresses. Nevertheless, in accordance with equation (1), it is possible to estimate the upper attainable level of the strength properties of materials based on alpha iron under the condition that the dislocations are not mobile and anchored. In polycrystalline materials with various grain orientation, this level reaches 0.050.10 G (this corresponds to approximately 40008000 MPa). This level has not as yet

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Fig. 1 Stress-strain dependence. a) thin fibres (whiskers), b) alpha iron single crystal, c) steel with ideal strength characteristics. been reached, although in certain cases the resultant strength is approaching this level. On the other hand, in pure iron containing mobile dislocations or areas in which dislocations can generate preferentially, plastic deformation at room temperature may already take place at a shear stress of around 10 MPa. In polycrystalline pure iron, the yield strength is not usually greater than 40 MPa, i.e. its value is around 10-4 G. This shows that the strength of steel at room temperature could reach values close to theoretical ones, around 40008000 MPa, whereas the lower limit is ~200 MPa. Because of various influences exerted on the structure, the actual strength characteristics should fit inside this wide range. The resultant strength level is controlled by the level of stresses required for generating dislocations and for their motion, e.g. through the ferritic matrix. This means that to obtain a high strength level, the main metallic matrix must contain strongly anchored dislocations and sufficiently 'strong' obstacles to dislocation motion. The previously mentioned (although highly perfect) thin crystals (whiskers) are not suitable because the stress required for their motion through the matrix rapidly decreases as soon as dislocations form in them, as shown in Fig.1.6 This graph also compares schematically the stressstrain curves for single crystals and for an idealized case of a highstrength steel. To provide further information, Fig.2 shows the schematic dependence of the resultant strength properties of iron or steels on the density of defects in the matrix, indicating that TMT utilizing the strain hard-

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Fig. 2 Dependence of strength on the density of defects in the metallic matrix. ening of the metallic matrix is effective. However, it may be concluded that the following mechanisms contribute to the hardening of structural steels to an extent which depends on the treatment method: a) dislocation hardening, b) grain boundary hardening, c) solid solution hardening hardening of the parent metallic matrix by substitutional or interstitial atoms, d) dispersion hardening by precipitates. The aim of analyzing these main hardening mechanisms is not to discuss their physical principles, but analysis will be carried out taking into account the possible connections with the expected set of the physical metallurgy properties of the strengthened metallic matrix of structural steels. As mentioned previously, the effect of strain hardening, caused by an increase of the dislocation density or by interaction of dislocations, is an important parameter as part of unconventional heat treatment. The corresponding increase of yield strength ∆σd, resulting from a higher dislocation density ρ (cm-2) in the matrix, may be expressed by the following relationship:

where σ0 is the value of the corresponding yield strength of the matrix formed in this case by pure polycrystalline alpha iron. The value of constant α depends on dislocation distribution, e.g. Keh7 states that α = 1. Equation (2) indicates that when the dislocation density greatly increases, the effect of strain hardening should make it possible to obtain values close to the theoretical strength level. However, the dislo-

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cation density does not simply increases with increasing strain. The hardening effect of dislocations depends not only on their spacing but also on the interaction characteristics. For example, after cold deformation, the dislocation density is around 1012 cm-2 which, according to equation (2), corresponds to high strain hardening reaching values of approximately 3000 MPa. In this case, the mean dislocation spacing is around 10 nm which corresponds to approximately 30 interatomic distances. A shortcoming of this strengthening mechanism is that the material in the work-hardened condition has very low plastic properties. Like the dislocations in the parent matrix, a hardening effect is also exerted by the grain boundaries which form obstacles to dislocation motion. The yield strength increases with an increasing number of grain boundaries, i.e. with a reduction of grain diameter D. In accordance with the HallPetch relationship, the corresponding contribution to the increase of the yield strength ∆σD can be expressed by the following equation:

The value of constant ky (in MPa m1/2) in structural steels with a ferritic matrix depends strongly on the 'degree' of segregation of impurity elements at the grain boundaries and on their effect on pinning of dislocations located in the region of the grain boundaries which represent potential places for the development of plastic deformation. From a general viewpoint, the effect of grain refining may be regarded as highly significant, from both the viewpoint of increasing the strength characteristics and that of exerting a beneficial effect on the resistance to brittle fracture. The total effect of this mechanism is restricted by the possibilities of refining ferrite grains to the level corresponding to approximately 10-3 mm. In this very 'fine recrystallization' of the austenitic matrix it is usually not possible to obtain finer ferrite grains than the previously mentioned level. The effect of segregation of interstitial atoms on the value of constant ky in a ferritic steel (in relation to cooling rate and, consequently, different degrees of development of segregation of interstitial atoms), is shown in Fig.3. The grain boundaries are highly effective obstacles to dislocation motion in cases of extensive segregation of interstitial atoms to the grain boundaries. The atoms dissolved in the matrix of alpha iron form also obstacles to dislocation motion and contribute by an appropriate amount, denoted ∆σM, to the increase of yield strength. However, it is quite difficult to specify separately the effect of hardening of the matrix by dissolved at-

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Fig. 3 Dependence of ky on test temperature (ferritic steels). oms with the simultaneous effect of other hardening mechanisms. In addition, the effects of substitutional and interstitial atoms in the ferritic matrix overlap. Therefore, for general evaluation, it may be assumed that, in accordance with Fleischer's theory,8 the contribution to hardening the matrix caused by dissolved atoms ∆σM can be expressed as follows:6

where c is the concentration of dissolved atoms (in at%), the hardening parameter a depends on the amount and type of dissolved atoms in the matrix and on their effect on the lattice parameter; shear modulus G also depends on the content of dissolved atoms in the matrix. 9 This type of hardening is the subject of a number of discussions. Above all, it should be mentioned that, in accordance with the conclusions drawn in Labusch's studies, in the case of a diluted solution of dissolved atoms in the parent matrix the modification of the strength properties is expressed by the value proportional to c2/3.10 At a higher concentration of impurity elements where it is not possible to approximate the conditions corresponding to the formation of a diluted solution in the matrix and where ordering processes in the matrix may occur to various degrees, the hardening effect is far more complicated. In this case, the effect can be expressed approximately, in relation to the concentration of dissolved atoms, by the following relationship: ∆σ≈ c(1 c2).9

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However, in accordance with previous considerations (equation (4)), the possibilities of hardening are greatly restricted and determined by maximum solubility cmax in the solid state, as indicated by appropriate equilibrium diagrams. However, this solubility usually decreases with an increase of the value of hardening parameter a in equation (4). The hardening effect of interstitial atoms in the body-centred cubic lattice is far stronger than the effect of substitutional atoms. The source of this hardening is the high local tetragonal deformation of the lattice which reaches on average around 36% in the case of interstitially dissolved carbon. Since the solubility of carbon atoms in the ferritic matrix is very low, the high hardening effect can be utilized only if it is possible to achieve a high degree of supersaturation of the metallic matrix, as in the case of martensitic phase transformations. Figure 4 shows that solid solution hardening is an important component of the strength of martensite, as observed previously by, for example, Winchell and Cohen in FeNiC alloys.11 As indicated by this graph, in the supersaturated solid solution it is possible to retain a high content of interstitial (carbon) atoms in the case of martensitic phase transformation. However, lattice defects (dislocations, phase boundaries of martensite plates or laths) form during this phase transformation and also contribute to matrix hardening. As regards the examined FeNiC alloys, it should be mentioned that the ratio of the nickel and carbon content was also varied in such a manner that the Ms temperature of the alloys in all cases was around 30°C.11 Precipitation hardening is another hardening mechanism contributing significantly to the hardening of structural steels. This type of hardening

Fig. 4 Example of hardening of the matrix by dissolved atoms.

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is caused by precipitated particles in the parent matrix; its contribution to the yield strength is denoted ∆σp. The effect of precipitation hardening depends on both the size of particles (can be approximated by their diameter d) and their mean spacing D* and their volume fraction f. The hardening effect also depends on whether the precipitates are coherent or not. For precipitation hardening resulting from the effect of particles precipitated in the matrix, the value of ∆σp, expressed by the stress required for dislocation motion, is given by the relationship:

Constant α' is equal to approximately 1 in the case of spherical precipitates uniformly distributed in the matrix. Equation (5) shows that in the case of a relatively high volume fraction of precipitates f, or for precipitation of very fine particles (precipitates with a small diameter d), it is possible to obtain values close to the theoretical strength of the iron matrix, mentioned previously. Figure 5 shows the calculated values of the yield strength of alpha iron containing precipitated (hard) particles

Fig. 5 Dependence of yield strength on the size of precipitates (for different volume fraction of precipitates f).

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with different volume fractions f. These dependences also hold for very small particles, provided they are not intersected by dislocations. If the precipitates are intersected by dislocations, the yield strength values, resulting from the hardening effect of the precipitates, are considerably lower. For example, at a precipitate content of f = 0.1, it is possible to obtain a high yield strength level, almost 4000 MPa (if intersection of the particles by the dislocations is suppressed). The dependences given for various volume fractions of the precipitated particles represent the upper attainable values of dispersion hardening of alpha iron. If very fine precipitates form (usually coherent precipitates), it is quite possible that the precipitates will partly 'contribute' to the deformation process, i.e. they will be intersected by the dislocations. In this case, the strength characteristics of structural steels are lower than those corresponding to equation (5). The resultant maximum strength level will depend on the parameters of the induced stress field around the precipitates, on the physical metallurgy properties of the precipitates,6 and on the ratio of the intersected and the so-called hard precipitates which the dislocations must bypass. However, it may be assumed that the hardening effect of the precipitated coherent particles becomes stronger with an increase of the difference of their lattice parameters in comparison with the lattice parameter of the matrix.12 A suitable example of efficient dispersion hardening is the hardening effect of very fine niobium carbonitrides in a low-carbon steel microalloyed with a niobium addition. The increase of the yield strength of this steel is in excellent agreement with equation (5). For example, dispersion hardening is utilized widely in maraging steels where the strength of the martensitic matrix is increased by precipitates of various types of intermetallic phases. The development of plastic deformation depends on the stress required for generating dislocations (σVD) and on the stress (σMD) required for the dislocations to move through the metallic lattice. The differences in the effect of these parameters on the development of plastic deformation are reflected in differences in the form of the stressstrain curve when the yield strength is exceeded. When σVD < σMD, the resultant stressstrain curve is continuous, without the formation of an instability, whereas when σVD > σMD, the stressstrain curve usually indicates the formation of a discontinuity, characterized by the so-called pronounced yield point. The first of these cases is usually observed in the majority of metallic materials and their alloys, including austenitic steels. The second case is found mainly in tensile deformation of ferritic steels. It may be concluded, in a simplified manner, that in ferritic steels the nature of the discontinuity observed at the yield stress is similar to that of the nature of the discontinuity observed in whiskers.

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The reasons for this phenomenon are obviously identical and are associated with the creation of suitable conditions for generation of mobile dislocations. Preferential areas for generation of dislocations at the start of plastic deformation are mainly grain boundaries and/or the release of matrix dislocations. These dislocations are usually pinned by interstitial atoms or precipitates formed on them. In the presence of a pronounced yield point, the stress required for the formation of dislocations σVD is higher than the stress required for dislocation movement σMD. As the difference between the values of σVD and σMD increases, the probability of mobile dislocations forming simultaneously in different places of the volume of the material decreases. Consequently, plastic deformation takes place preferentially in places of the structure with the minimum stress required for dislocation generation and extends from these places with the formation of Lüders bands. The deformed regions of the material behind the front of the Lüders bands have such a high density of mobile dislocations that no new sources of dislocations are activated. Under these conditions, it may be written that σ0 + ∆σd < σVD. As soon as the Lüders bands fill up the entire volume of the matrix (when σ0 + ∆σd > σVD), new dislocation sources are activated and strain hardening of the matrix takes place. In metals characterized by the continuous stressstrain curve without any pronounced yield point, the strain hardening process starts in the matrix immediately in accordance with an increase of the dislocation density. This difference in the behavior of the ferritic matrix is caused by the high susceptibility of interstitial (carbon, nitrogen) atoms to segregation to defects of the matrix (to both grain boundaries and dislocations in the grain volume). This susceptibility to segregation can be defined by the energy difference ∆U of the interstitial atoms dissolved in the matrix lattice and the atoms segregated into areas at the grain boundaries or to dislocations. This value of ∆U is considerably higher in the ferritic matrix than in the matrix formed by the face-centred cubic lattice (for example, austenite). The value of ∆U in alpha iron is considerably higher than in other bcc metals, for example, chromium, molybdenum, tungsten, etc. This high value of ∆U is caused by high strains generated by the interstitial atoms during dissolution in the alpha iron matrix.13 These considerations indicate that the concentration of the atoms, segregated to the grain boundaries and characterized by the degree of segregation Cs, increases with an increase of the value of ∆U. According to Mc Lean,14 the value of Cs can be expressed by the following equation

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where c expresses the carbon concentration in the parent lattice, ∆U ≈1 eV for the FeC system, and k is the Boltzmann constant. The value Cs also determines the magnitude of the stress required for generation of dislocations σVD and is associated with the dependence of ky in equation (3) on the degree of segregation of the interstitial atoms at the grain boundaries. This dependence is also shown in Fig.3. As already mentioned, the segregation of interstitial atoms does not take place only towards the grain boundaries but also to the dislocations in the volume of the matrix resulting in pinning of these dislocations. These dislocations are then immobile and plastic deformation of the matrix starts only when new sources are activated, or the dislocations are released. A case of simultaneous segregation of the interstitial atoms to the moving dislocations may also occur. This strenghtening mechanism operates at elevated temperatures, as also indicated by the schematic stressstrain curve determined at different test temperatures, Fig.6. As indicated by Fig.6, at room temperature (for example, in low-carbon structural steels) a significant discontinuity is detected at the yield stress. At elevated temperatures between 80 and 250°C, the stressstrain curve shows a jerky flow usually associated with the development of Snoek's interaction of dislocations with interstitial (carbon) atoms at lower test temperatures and the so-called serrated flow at test temperatures situated in this case above approximately 110°C. This is caused by the develop-

Fig. 6 Effect of test temperature on the stressstrain curve.

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ment of Cottrell's interaction.15,16 The higher hardening effect is obtained in testing at temperatures between 150 and 250°C. At test temperatures higher than 250°C, the stressstrain curve is continuous (smooth) over the entire plastic deformation range. These effects of test temperature on the form of the stressstrain curve are valid for a constant strain rate. At lower strain rates the formation of these instabilities in the stressstrain curve is observed at lower test temperatures. At higher strain rates the characteristics of the stressstrain curve are displaced in the opposite direction. The data on the displacement of the yield point in relation to the strain rate are summarized in Fig.7 which shows three most important regions, according to the response of the given steel to plastic deformation. This solution is similar to that used previously in defining the individual regions of the controlling mechanisms of interaction of dislocations with interstitial atoms in as-quenched martensite.13,17,18 These three regions are defined by the following conditions: νd > νI, νd≈νI and νd < νI, where νd defines the velocity of dislocation motion and νI the diffusion velocity of interstitial atoms in the matrix. This analysis is based on the wellknown relationship between strain rate

and the mean velocity of dislocation motion

:

where b is the Burgers vector, and ρm expresses the density of mobile dislocations. In region III (Fig.7), the mobility of dislocations is higher than that of the interstitial (carbon) atoms and the mechanism of solid solution hardening operates efficiently in this region, as indicated by equation (4). On the other hand, at the higher temperatures considered (region I), the mobility of carbon atoms is sufficiently high, so that after their segregation to the dislocations these atoms are also capable of moving together with the dislocations. These two regions are separated by region II, in which plastic deformation is accompanied by the formation of, for example, a serrated flow in the stressstrain curve. In this case, the dislocations after travelling over a specific distance are pinned by interstitial atoms or interact with dislocations. This interaction restricts dislocation mobility. In a limiting case, dislocations may be released from pinning or new dislocations may be generated during further plastic deformation. As reported in Ref.19, the second of these mechanisms is more likely to operate. For example, in pinning of dislocations and restricting their mobility, characterized by a reduction of the mean velocity of dislocations , to retain the original strain rate

new dislocations must be generated in accordance with equation (7) and this increases

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Fig. 7 Deformation characteristics of steel in relation to strain rate. their density. This process increases the level of strain hardening (equation (2)), as follows from the schematic dependences shown in Fig.6 for the given temperature range. The discussed hardening characteristics are also consistent with the activation energy values determined for the transition from one region to another in the strain rate dependence, as shown in Fig.7. The results obtained for lowcarbon ferritic steel show that the activation energy corresponding to the start of region II is around 0.9 eV, i.e. similar to the activation energy of carbon diffusion in alpha iron; the activation energy for the transition from the second to first stage was higher, around 1.40 eV. This value consists of the sum of the activation energy for carbon diffusion in ferrite and the interaction energy of carbon atoms with a dislocation whose value for the Cottrell interaction is around 0.50 eV. Although the previously mentioned hardening mechanisms are generally valid, their importance from the viewpoint of overall hardening of the parent matrix changes in relation to the type of their crystallo-graphic lattice. All the previously mentioned data on the hardening conditions were considered for the body-centred cubic lattice of alpha iron; however, in the face-centred cubic metallic matrix (for example, in an austenitic matrix) the importance of the individual hardening mechanisms greatly changes. The following characteristic properties of the austenitic matrix are the reason why the mechanical properties of austenitic steels

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often differ greatly from the properties obtained for ferritic steels: considerably higher solubility of interstitial atoms in the matrix; lower mutual interaction energy between the interstitial atoms and lattice defects (dislocations); considerably lower stacking fault energy of the austenitic matrix; instability of the face-centred cubic lattice of a large number of technical systems; reduced diffusion coefficient of dissolved atoms (reduced by as much as two orders of magnitude). The higher solubility and reduced diffusion coefficient cause that unstable precipitates of 'interstitial' phases are not detected in the austenitic matrix. Because of the low interaction energy, the austenitic matrix does not contain instabilities on the stressstrain curve which could be caused by the mutual influence of the dislocations and the atoms interstitially dissolved in the lattice. This is also linked with a reduction of constant ky (equation (3)) and leads to a considerably weaker dependence of the resultant yield strength values on grain size. Low stacking fault energies of the austenitic matrix are the reason for the detected high level of the strain hardening exponent of austenitic steels. Due to their relatively extensive splitting, the dislocations cannot deviate from their slip plane during plastic deformation and, consequently, their density increases more markedly at a comparable plastic strain in comparison with the ferritic matrix. The strain-induced martensitic phase transformation, which takes place during plastic deformation of unstable or metastable austenite, further increases the strain hardening exponent. It must also be remembered that due to the reduced diffusion coefficient the austenitic matrix is characterized by coalescence processes which are considerably slower than in ferritic steels so that the effect of dispersion hardening is also retained to a large extent at elevated temperatures and long exposure times. In most cases, all discussed hardening mechanisms operate in the majority of structural steels, although the effect of the individual mechanisms can vary, depending on the processing method. The overall effect can be expressed schematically using the following equation:6

This relationship, based on the simplified additive coupling of the individual elementary hardening effects (as indicated by equations (2) to (5)) describes only formally the overall hardening effect, because under practical conditions the effects of the individual considered hardening

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mechanisms may condition or enhance each other (for example, the combined effect of dispersion and dislocation hardening). These hardening mechanisms operate mainly in the modification of the strength properties of low-carbon structural steels, for example, produced by controlled rolling and used widely in industry. The microstructure of these steels consists mainly of ferrite or ferrite with low-carbon bainite, or martensite or a mixture of ferrite and pearlite. These considerations may also be applied to the conditions under which the so-called transformation hardening operates.6,13 Other transformation hardening mechanisms can operate during a phase transformation, for example, during simultaneous transformation of austenite to martensite. Under a certain degree of approximation (under certain simplifying assumptions), they may be connected with the generation of a large number of dislocations. This corresponds to the well-known fact, i.e the martensitic transformation can be regarded to a first approximation as a specific form of the deformation process.13,20,21,22 These questions will be discussed in greater detail in analyzing the physical metallurgy characteristics of the thermomechanical treatment of high-strength martensitic steels. In addition, the properties of the grain boundaries of original austenite are also modified, for example, by the induced interaction of growing martensite plates with this barrier to their propagation. Analysis of the individual contributions to the hardening of the matrix of martensitic steels, including the effect of dispersion hardening, for example, after low-temperature tempering of a high-strength martensitic steel, is shown in Fig.8.6,23 As indicated by the scheme, the region 12 expresses the

Fig. 8 Analysis of strengthening contributions in a high-strength martensitic steel.

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contribution to the increase of yield strength caused by strengthening by dissolved atoms in the matrix (hardening of the solid solution). The region 23 represents the effect generated during phase transformation (transformation hardening) of austenite to martensite, for example, under the effect of an increase of the dislocation density and depending on the degree of refining of martensite. This expresses, to a certain degree of approximation, an effect similar to that attributed to the effect of the grain boundaries. Consequently, region 34 corresponds to the hardening effect caused by, for example, the effect of precipitation of carbides or intermetallic phases (in maraging steels). The mutual ratios of these mechanisms in the individual hardening regions may vary partially, depending on the phase transformation conditions or the properties of the initial austenitic matrix, although it is evident that this diagram is generally valid for high-strength martensitic steels. For example, in maraging steels, the effect of dispersion hardening will be stronger than in at low-temperature tempered martensite after conventional heat treatment or TMT. These considerations can also be used in discussing the conditions for hardening the metallic matrix of TRIP steels formed in the initial state by unstable austenite and precipitated carbides. During the strain-induced martensitic transformation, these steels greatly harden and this high hardening level also leads to satisfactory ductility, as indicated by a very wide range of uniform plastic deformation prior to necking (localised plastic deformation). In TMT of high-strength steels it may be assumed that, taking into account the discussed strenghtening mechanisms of the matrix, there will also be other mechanisms causing higher hardening of martensite, favourable structural and metallurgical characteristics of martensite (refining of martensite plates, modification of their structure) and the change of the parameters of dynamics of the martensitic phase transformation, namely the restriction of the so-called dynamic effects in this transformation.13 These problems will be analyzed in greater detail when discussing the physical metallurgy conditions of TMT of the discussed types of high-strength structural steels and the conditions leading to the optimized properties of these steel types in the following section of this book. However, this treatment, based on general analysis of the main hardening mechanisms, includes only one part of the required increase of the mechanical and metallurgical characteristics of structural steels. The solution of the hardening problem without considering the associated views on the brittle fracture resistance is not complete with respect to the detailed analysis of the conditions leading to higher applied properties of these steel types. In fact, in certain cases unconventional heat treatment is applied in order to obtain higher brittle fracture resistance whilst maintaining a sufficiently high strength level, which in conven-

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tional treatment (CT) leads to low brittle fracture resistance (this is often required in high-strength martensitic steels).13 This solution in its consequences results in increasing the safety and service reliability of the materials processed in this manner in technical service. These problems will be discussed in greater detail in analyzing the relationships between the hardening mechanisms of the examined types of structural steels and the corresponding brittle fracture resistance parameters. 2.2 Analysis of Physical Metallurgy Characteristics of the Initial Austenitic Matrix Taking into Account the Subsequent Heat Treatment The properties of the initial deformed austenitic matrix prior to the subsequent phase transformation are very important with respect to the resultant TMT effect. Depending on the degree of development of its strain hardening, the mechanism and kinetics of the subsequent phase transformation are modified (either diffusionless or diffusion type) or the resultant level of the mechanical and metallurgical properties of structural steels after TMT is strongly affected.1,2 The requirements on the modification of structural and metallurgical characteristics of the initial austenitic matrix differ depending on whether unconventional heat treatment of the high-strength martensitic steel is applied or whether the controlled rolling conditions (or even accelerated cooling conditions) of low-carbon or microalloy steels for wide technical applications are used. However, the controlling parameter in both TMT cases mentioned is the degree of development of recrystallization processes or the extent of strain hardening of the initial austenitic matrix.4,24 For this reason, the main aim of this part of the book will be an analysis of the effect of the individual physical metallurgy parameters of deformation of the initial austenitic matrix and determination of the conditions leading to obtaining the most favorable structural state from the viewpoint of the maximum efficiency of the TMT effect in optimizing the properties of these steel types.1 This means that it is required to carry out a detailed analysis of the physical metallurgy factors affecting the development of recrystallization processes or ensuring strain hardening of the initial austenitic matrix. It is therefore necessary to determine the conditions leading to the optimum size of initial austenite grains and modified substructure characteristics of the austenitic matrix from the viewpoint of the final requirements on obtaining satisfactory properties of structural steels when applying the examined variants of unconventional treatments.5,25 If the austenitic matrix is deformed in the appropriate temperature range at a higher strain (reduction in pass) than the socalled critical value (this value increases with decreasing deformation temperature),

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recrystallization takes place both during the deformation process development of dynamic recrystallization or after completing forming, or in handling periods between the individual reductions in pass or prior to final heat treatment austenite breakdown. This corresponds to the development of static recrystallization. The size of recrystallized austenite grains depends on a large number of factors, e.g. deformation temperature, strain and strain rate, the initial austenite grain size prior to deformation and the chemical composition of the austenitic matrix. The size of the recrystallized austenite grains decreases with a reduction of forming temperature and an increase of strain, Fig.9. The reduction of the initial austenite grain size (it is also determined by the temperature prior to forming) also leads to a reduction of the size of recrystallized austenite grains whilst other parameters of the deformation process remain unchanged, Fig.10.26 The addition of microalloying elements (for example, niobium and titanium) also influences efficiently both refining of the recrystallized austenite grains and reduces the rate of the actual recrystallization process. Here, it must be taken into account that the degree of refining of the austenite grains in recrystallization is limited. According to the currently available data, it may be assumed that the attainable minimum size of the recrystallized austenite grains is on average 1015 µm.26,27 However, this austenite grain size can ensure the formation of a very fine ferritic structure in controlled rolling of steels for wide technical applications with a mean ferrite grain size of 45 µm only in an extreme case. This size is essential for obtaining the peak properties of the given type of structural steel, especially from the viewpoint of the resultant brittle fracture resistance.27 Therefore, the final strain in this case must be selected to retain to the maximum extent the elongation of the austenite grains and strain hardening of the austenitic matrix. In both cases, this increases the frequency of occurrence of potential areas for the nucleation of fine ferrite grains. As already mentioned, the selection of a suitable temperature prior to deformation (forming) is also important in affecting the austenite grain size. As indicated by Fig.11, which shows the relationship between the heating temperature and the austenite grain size of three types of low-carbon steels (including steels microalloyed with the addition of niobium or vanadium), the examined dependence contains two regions in which the rate of grain growth decreases, i.e. at temperatures around 950°C and around 1100°C.26,27 In heating in these temperature ranges the austenite grains are usually uniform with a minimum scatter. On the other hand, in heating at temperatures between these two temperature ranges, the austenite grains are partially nonuniform. As indicated by Fig.11, the low-carbon structural steel, microalloyed with niobium, has very fine austenite grains under these conditions. The exceptionally rapid increase

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Fig. 9 Dependence of the size of recrystallized austenite grains on deformation temperature and level.

Fig. 10 Effect of total strain and deformation temperature on the austenite grains size. of the size of the austenite grains is found in these microalloyed steels only at temperatures higher than 1200°C (Fig.12).27 These data can also be applied when optimizing the rolling conditions, for example, in controlled rolling. During heating to lower tem-

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Fig. 11 Effect of heating temperature on the austenite grain size (holding time at temperature 30 min).

Fig. 12 Coarsening of austenite grains in relation to heating temperature (holding time at temperature 30 min). peratures (around 1100°C), it is possible to select shorter holding periods between rolling in the individual forming stages, e.g. holding between rolling completed in the area of rapid development of recrystallization processes and subsequent rolling (in the second stage) in the temperature range in which there is no rapid development of austenite recrystallization, or its development is suppressed.

However, heating to lower temperatures also leads to a number of other problems of the physical metallurgy nature. Heating to lower tem-

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peratures reduces the content of dissolved microalloying elements in the matrix (mainly when microalloying with niobium and titanium), because the initial structure contains a higher content of carbonitrides of these microalloying elements. The undissolved carbonitrides of the microalloying elements (niobium, titanium) have a beneficial effect on refining the austenite grains. However, a reduction of their content in the matrix restricts the 'braking' effect on the development of austenite recrystallization during or after forming and reduces the efficiency of utilizing the strengthening effect of the microalloying elements in possible subsequent dispersion hardening.26,29 This assumption may also be linked with the conclusions obtained in analyzing the conditions of optimizing the chemical composition of low-carbon steels for wider technical application microalloyed with various elements. When reaching the optimum content of, for example, niobium in the matrix at the usual conditions of heating prior to forming at lower temperatures, the content of 'bonded' carbon should not exceed 0.05% to minimize the niobium content of precipitates. These losses in hardening can be compensated by, for example, modified and complexly balanced alloying, for example, by adding vanadium, molybdenum, manganese and titanium. A higher manganese content (1.51.8%) improves the hardenability of the matrix. A molybdenum contributes to restricting the development of recrystallization processes and to refining the austenite grains, like the niobium addition. It also contributes significantly to hardenability. From the viewpoint of the optimum chemical composition of these steel types, it can be concluded that, to a first approximation, the molybdenum addition should have approximately the same effect as the total effect of the niobium and manganese addition. The vanadium addition is efficient mainly in dispersion hardening of the matrix. Attention should also be given to a titanium addition which reduces the rate of recrystallization processes, mainly after heating to lower temperatures. On the basis of these assumptions, Tamura26 recommends the optimum chemical composition of a low-carbon microalloy steel with higher strength characteristics based on manganese, molybdenum, niobium, with an addition of approximately 0.01% titanium. According to Tamura's data, in controlled rolling of this steel with a balanced chemical constitution it is possible to obtain both favorable strength properties and a sufficiently high toughness level, whilst maintaining satisfactory weldability. Although this problem is not of primary importance from the viewpoint of the examined problem of development of recrystallization processes in the austenitic matrix and the resultant mechanical and metallurgical properties, it represents a very important parameter from the viewpoint of the resultant applied properties of these steels.

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Taking into account the modification of the chemical composition of these steel types, heating to temperatures of ~1050° C can also be selected. At these temperatures, the deformation process will take place without accompanying rapid development of recrystallization processes in the austenitic matrix. A practical technological consequence of selecting these heating conditions is that it is possible to ensure that the forming process takes place in a single deformation cycle without holding; usually, between the deformation conducted under the conditions of development of recrystallization processes and deformation rolling carried out at lower temperatures suppressing the development of recrystallization in the austenitic matrix. As indicated by the information obtained on the conditions of development of recrystallization processes in controlled rolling of these types of structural steels, and taking into account the subject of further detailed discussion, the fine austenite grains also lead to the development of fine ferrite grains in the subsequent austeniteferrite phase transformation. The limiting size of the austenite grains, obtained in recrystallization of the austenitic matrix, is ~1520 µm,26,27 which, however, would make it difficult to ensure efficient refining of the ferrite grains defined by the ferrite grain size below 5 µm, as discussed previously. Taking into account the fact that in this initial recrystallised austenite matrix it is not possible to ensure efficient refining, it is important to select final deformation in such a manner has to prevent the development of recrystallization of the austenite grains. In this case, the austenite grains are elongated in the forming direction, including the formation of deformation bands in the matrix.30 Both these effects increase the frequency of potential nucleation areas for the formation of ferrite. For example, Kosazu et al.29 have introduced a concept of effective phase boundary areas SV of austenite given by the sum of the grain boundaries of recrystallised and elongated austenite grains and the phase boundary, corresponding to the resultant deformation bands in the unit volume of the matrix. This value is linked with the final ferrite grain size. The value of SV increases with a reduction of the initial austenite grain size and with increasing strain. An example of the relationship between the size of the austenite grains and the resultant ferrite grains in a steel microalloyed with niobium is shown in Fig.13. The initial austenite grain size varies depending on the final deformation (forming) conditions with 1 to 3 passes at temperatures above 1050°C (after heating the experimental material to 1250°C). Final deformation was carried out at 850°C.29 Sellars31 proposed an empirical relationship between the initial size diameter of the austenite grain d0 and its final size d, determined on the basis of evaluation of the development of static recrystallization after a single pass deformation. For example, for a steel microalloyed with

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Fig. 13 Effect of forming in the region in which austenite does not recrystallize and the effect of the initial austenite grain size on the ferrite grain size (different initial austenite grain sizes were obtained by forming in one or three passes at temperatures higher than 1053°C and after heating to 1250°C. Final deformation was carried out at 850°C). a niobium addition at strain ε he obtained the following relatively simple empirical relationship:

However, this relationship converges quite rapidly to an almost constant value d for initial different austenite grain sizes d0 if the deformation process is repeated several times. The rate of this convergence and of obtaining the expected limiting value of d depends on strain ε in individual passes. An example of the effect of different values of strain in a single pass (0.33 or 0.06) on the resultant austenite grain sizes for two initial austenite grain sizes is shown in Fig.14a, b. As indicated by the diagrams, at a higher strain the refining of austenite grains is more rapid, but at a lower strain and at an initial austenite grain size of 100 µm the austenite grain size even slightly increased. These considerations do not take into account further deformation parameters, such as the effect of temperature and strain rate, as considered in, for example, Ref. 31. The

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Fig. 14 Effect of repeated recrystallization on the austenite grain size (two different austenite grain sizes and deformation levels in individual passes). a) strain 35% in a single pass, b) strain 6% in a single pass.

Fig. 15 Effect of strain rate on the size of recrystallized austenite grains (deformation temperature 1050°C, strain in a single pass 70%).

effect of strain rate in a steel microalloyed with a niobium addition (0.05% Nb) can be seen in Fig.15 showing the dependence of the resultant size of the recrystallised austenite grains on the mean strain rate (in a single pass with a strain of 70% at 1050°C in the strain rate range 0.03526 s-1). As indicated by this diagram, the size of the recrystallized austenite grains increases with a reduction of strain rate. Finer recrystallized austenite grains can be obtained using higher reductions in the individual rolling passes, i.e. with increasing strain rate. The relationship between the austenite

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Fig. 16 Effect of strain in a pass on the size of recrystallized austenite grains (total deformation 65 and 75% in a temperature range 1060980°C). grain size and strain in the individual passes was determined for the previously mentioned microalloyed steel and is shown in Fig.16.27 This graph shows that when selecting the maximum values of strain, the size of the recrystallised austenite grains is even smaller than 20 µm, which represents the optimum value from the viewpoint of the current technical and technological possibilities. An important technical problem in selecting the optimum forming conditions is the partial development of recrystallization processes. If in the case with complete development of recrystallization the structure consists of equiaxed austenite grains, then in partial development of recrystallization the resultant structure is of the duplex type. It usually consists of a mixture of large and fine grains. In this case, it is a mixture of: a) fine recrystallised austenite grains b) deformed initial grains c) grains coarsened under the effect of strain-induced migration of the grain boundary in different areas of the structure during deformation.32 In this case, structural and phase analysis shows that in grains whose size is greater than the size in the initial condition, the grains grew as a result of migration of the grain boundaries. This means that in this 'transition' region the density and distribution of precipitates is not sufficient or not sufficiently uniform to ensure efficient pinning of the grain boundaries in the volume of the metallic matrix. An important problem in the formation of this duplex type of structure is that it is very difficult to remove in further forming.27 From the practical viewpoint it is convenient to define, for the individual steel types, the minimum temperatures leading

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to complete recrystallization at the selected strain. The data obtained in Ref.27 are presented in Fig.17. The steel types examined were heated to different temperatures resulting in different degrees of enrichment of the matrix with dissolved atoms of the microalloying additions and in their subsequent reprecipitation. The specimens were then formed by different numbers of passes, until only partial recrystallization of the austenitic matrix was observed. As indicated by Fig.17, the temperature for the start of partial recrystallization uniformly increases with increasing content of dissolved microalloying additions in the matrix at the start of the deformation process. On the other hand, it was observed that the initial austenite grain size has a very small or even negligible effect on the process under examination. Taking into account the conclusions obtained in the basic discussion of the physical metallurgy characteristics of the deformed (non-recrystallised) austenitic matrix, we shall carry out an additional analysis of selected microstructural characteristics. An important structural parameter (in elongation of the austenite grains in deformation) for the subsequent phase transformation is the 'height' of these austenite grains characterized by the dimension parallel with the compression axis in forming. Figure 18 shows that the height of the austenite grains decreases with increasing total strain. However, the height is greater than the value corresponding to the dependence resulting from fulfilling the theoretical conditions in cold deformation. This dependence is also shown in this figure. The difference between the measured and the so-called ideal (theoretical) con-

Fig. 17 Change of the temperature of the end of recrystallization in relation to the content of dissolved microalloying additions.

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Fig. 18 The effect of total deformation of austenite (with recrystallization suppressed) on the change of size (height h) of austenite grains. Deformation temperature 880-800°C. ditions indicates that partial migration of the grain boundaries took place probably during deformation. The discussed relationship between the values of h, d0 and ε can be expressed approximately using the following empirical relationship:27

where d0 is the initial size (diameter) of the austenite grain.

From the viewpoint of obtaining the optimum strength properties of control-rolled structural steels, an important role is also played by the possibility of utilizing the physical metallurgy characteristics of work

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hardening the metallic matrix induced in forming steels of this type in the two-phase (austeniticferritic) region. In comparison with the properties obtained after deformation in the region of development of recrystallization processes or after deformation in the region in which the development of recrystallization is suppressed, forming in the two-phase region further improves both the strength properties and leads, under certain conditions, to higher brittle fracture resistance. This beneficial effect is caused both by additional strain hardening of the existing ferrite and by its partial polygonisation and texture development. A general rule in obtaining the optimum results in forming in this range is that the strength of the matrix increases but the level of absorbed energy in evaluating fracture toughness in the upper transition temperature region (during ductile failure) decreases. However, these dependences are modified if higher strains are applied. With increasing strain in forming the resultant toughness of the matrix initially decreases but then slightly increases. This tendency corresponds to changes of the physical metallurgy properties of work-hardened ferrite. The dislocation density in the ferrite grains increases with the strain increasing to 1520%, thus resulting in extensive strain hardening. During a further increase of strain the distribution of dislocations is modified with the formation of subgrains during dynamic recovery which takes place in the ferritic matrix.26 The addition of alloying elements, such as molybdenum, niobium, titanium or vanadium reduces the rate of the recovery process so that the addition of these elements also contributes in this manner to increasing the level of strain hardening of the matrix. Forming under the conditions leading to suppression of austenite recrystallization results in the development of a texture in the matrix which is then 'inherited' in the subsequent austeniteferrite phase transformation. Usually, the following texture variants were detected: {332}~{554} and {112}~{113}. In rolling in the two-phase (austeniticferritic) region, the ferrite texture develops further and the orientations of the {100} type are also detected under these conditions. This type of texture rapidly develops at higher strains or with decreasing deformation temperature.33 In deformation leading to the formation of recrystallised grains and in cases in which sulphide inclusions (MnS) are elongated after forming, the examined materials show the so-called delamination on the fracture surface. This type of failure is not found if the sulphide inclusions have the form of small spheres (for example, after adding calcium, or if the sulphur content of steel is very low). If a structural steel is formed in the two-phase region, delamination of the material is already detected even at a very low sulphur content because the texture of the {100} type rapidly develops in ferrite, as mentioned previously. It is well known

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that {100} planes are the cleavage planes of ferrite supporting the development of delamination. However, as stated by Tamura,26 the occurrence of delamination can not be only attributed to the rapid development of a texture of the {100} type. However, it may be assumed that the matrix also contains bands formed by ferrite grains with the same orientation {100} or {111}, parallel with the rolling plane. Taking into account the anisotropy of the plastic properties of the {100} and {111} type bands, the phase boundary of these bands of different quality are the regions of initiation of brittle fracture in tests at lowered temperatures. If these boundaries are normal to the fracture surface, they reduce the notch toughness corresponding to the formation of ductile failure, i.e. they lead to low-energy ductile failure. However, it should be mentioned that in delamination the transition temperature is usually displaced to lower values.34 These considerations can be utilized efficiently in defining the physical metallurgy principles of optimizing the technological conditions of controlled rolling low-carbon and microalloyed structural steels, with certain restrictions associated with the consequences of forming in the area of complete austenite recrystallization. In forming with suppression of austenite recrystallization a favourable dislocation structure forms in the matrix and restricts the action of dynamic effects in subsequent martensitic phase transformation,1,13,14 as observed in TMT of high-strength martensitic steels1,3 or steels with a mixed bainiticmartensitic microstructure.4 Controlled rolling can also be applied in heat treatment of steels with a higher carbon content (i.e. usually in steels with more than 0.40% carbon) where pearlite is an important structural constituent. In controlled rolling of this type of steel, it is efficient to select lower forming temperatures than in conventional heat treatment because the interlamellar spacing can slightly increase due to an increase of the temperature of the subsequent austenite to pearlite transformation (obviously associated with the strain-induced effect). This may then result in a small reduction of the strength properties, as also indicated by the results shown in Fig.19. On the other hand, the brittle fracture resistance, determined after controlled rolling, is higher than in conventional heat treatment because the pearlite nodules are smaller, Fig.20.26 These problems can be partially solved by accelerated cooling after rolling. More extensive application of 'rolling during treatment' of high-carbon (pearlitic) steels is prevented at the moment by a shortage of information on the physical metallurgy principles of this process in the given type of steel (mainly efficient combination of controlled rolling with accelerated cooling).35 An integral part of the detailed evaluation of the effect of the physical metallurgy characteristics of the initial matrix from the viewpoint of properties obtained after subsequent heat treatment is the availability

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Fig. 19 The effect of carbon content on the yield and tensile strength obtained after controlled rolling and conventional treatment (CT).

Fig. 20 The effect of carbon content on elongation, reduction in area and transition temperature (T(KCV 50%)) after controlled rolling and conventional treatment (CT).

of values of the deformation resistance of austenite in forming, including determination of the level of pressures acting on rolling equipment.1 However, the deformation resistance of austenite is not affected only by temperature, strain and strain rate but also by the austenite grain size

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and the content of alloying additions. Generally, deformation resistance increases with refining of the austenite grains and with increasing content of the alloying additions. It must also be taken into account that in forming which is accompanied by the suppression of recrystallization of the austenitic matrix there is the 'gradual' effect of strain hardening (accumulation of strain hardening). This (accumulation) effect of strain hardening must be supplemented by data on the effect of structural and metallurgical characteristics in order to determine with the maximum accuracy the deformation resistance of the matrix and, consequently, prepare theoretically substantiated data for obtaining the optimum conditions of the deformation process and its control.36 2.3 Influence of the State of the Initial Austenitic Matrix on Phase Transitions in Thermomechanical Treatment The initial state of the austenitic matrix strongly affects the parameters of subsequent phase transformations, both from the viewpoint of modifying its microstructure during development or suppression of recrystallization processes and from the viewpoint of modifying the substructural characteristics of initial austenite. Despite the fact that in the general sense of the word there are different TMT variants, in this book this method of unconventional heat treatment will be divided to the actual process of thermomechanical treatment (HTMT or LTMT) consisting of the process of plastic deformation of austenite and subsequent quenching in, for example, a quenching press, and the process of controlled rolling (determined by the technological possibilities), supplemented by the superposition effect of accelerated cooling. TMT is then used in high-strength martensitic steels (where a martensitic structure forms in subsequent phase transformation)1,2,3 and in treatment of low-carbon, low-alloy steels where subsequent final quenching, with various degrees of work hardening of the initial austenitic matrix, usually leads to the formation of a structure formed by low-carbon bainite or its mixture with low-carbon martensite.4 In this case, in TMT of high-strength steels, the material is usually tempered at a low temperature, i.e. temperature below the temperature of possible occurrence of irreversible, low-temperature temper brittleness, although the development of this type of brittleness is, according to some authors, partially suppressed in TMT,3 and in low-carbon, low-alloy steels it is tempering at higher temperatures.4 On the other hand, in controlled rolling or accelerated cooling the structural steels processed in this manner are usually used in the as-received condition (after forming and phase transformation without additional tempering). From the viewpoint of subsequent heat treatment after quenching and the resultant mechanical and metallurgical properties, an important role

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Fig. 21 Dependence of the tensile strength of martensite on the austenite grain size for steel CSN 14160. in TMT of high-strength martensitic steels is played by the size of the initial austenite grains.37 As indicated by Fig.21, showing the dependence of tensile strength on the mean size of the initial austenite grains,1 exceeding a specific limiting size of the austenite grain results in a rapid reduction of the strength of the resultant martensitic matrix. The dislocation substructure of austenite has also a strong effect it influences the dynamics of development (growth) of martensite plates and changes the nucleation characteristics of the martensitic phase transformation. In the former case, the dislocation substructure, formed in the deformed recrystallised austenitic matrix by dislocation cells, has a damping or restricting effect on plate growth. In a number of cases, the faces of martensite plates were even arrested on the walls of these dislocation cells or their growth deviated during passage through the walls of the dislocation cells. This obviously shows that they have a braking effect and restrict the development of the burst process usually accompanying the martensitic phase transformation.1,13 An example of this is the observed change of the dependence of electrical resistance on the stage of development of martensite formation according to the level of undercooling in the conventionally processed austenitic matrix and after its preliminary plastic deformation, Fig.22. After preliminary deformation of austenite, the burst phenomenon was completely suppressed in a number of cases and this restricted the consequences of the dynamic effects of growing martensite plates, for example during their interaction with the grain boundaries of initial austenite or with other structural obstacles.38,39 The higher dislocation density in the deformed austenitic matrix also increases the probability of initiation (nucleation) of martensite. This leads to refining of the martensitic structure and shortening of the so-

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Fig. 22 Results of resistance measurement of the kinetics of martensitic phase transformation. called free path for the growth of martensite plates, thus restricting the dynamic effects in the martensitic transformation, as shown schematically in Fig.23. This assumption, applied previously in Ref. 1, is based on the original KnappDehlinger concept of the growth of martensite plates. These considerations can also be applied in analyzing the favorable effect of the initial fine-grained austenitic microstructure where the free path for the growth of martensite plates is also restricted with further consequences in restricting the dynamic effects of martensite plates on obstacles when arresting their growth, as shown in Fig.21.13 It can be concluded that the main factor in all previously discussed

Fig. 23 Growth of martensite plates.40

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cases is the restriction of the consequences of dynamic effects in the martensitic transformation and, consequently, the restriction of the formation of defects in the martensitic structure. Defects are found mainly in areas of arrest of rapidly growing martensite plates, as frequently observed under the conditions of conventional heat treatment.1,13,20,22 This assumption should also be linked with the effect of the austenite grain boundaries on the development of defects in the martensitic matrix, mainly the higher susceptibility to intercrystalline failure taking place at the initial austenite grain boundaries. In TMT, the susceptibility to intercrystalline failure of martensitic steels is suppressed. This susceptibility is manifested by the formation of brittle intercrystalline fracture surfaces, for example, in formation of quench cracks, and premature and delayed fractures in as-quenched martensite.1,41-44 As demonstrated earlier, the physical metallurgy properties of austenite grain boundaries have a very important influence on the initiation of this type of failure,41,42 especially their modification due to the development of microsegregation processes at the austenite grain boundaries.45,46 Grain boundary microsegregation processes, leading to the enrichment of grain boundaries with, for example, harmful elements, are selective, as given in Ref.47. A high degree of enrichment on the high-angle austenite grain boundaries of CrNi martensitic steels with, for example, phosphorus was found. Enrichment of the individual grain boundaries was different and, therefore, the authors47 concluded that the grain boundary energy of the original austenite is not evidently the only parameter influencing the increased susceptibility of the martensitic structure to intercrystalline failure (see the considered superimposition of grain boundary 'weakening' due to dynamic effects accompanying the martensitic phase transformation).1,13,41-43 One of the methods of increasing the mechanical and metallurgical properties of martensite (in addition to restricting the degree of development of microsegregation processes) is, for example, restriction of the dynamics of the martensitic phase transformation in TMT induced by refining of austenite grains, reduction of the free path for the growth of martensite plates after preliminary deformation of the austenitic matrix prior to the phase transformation of austenite to martensite. It is thus possible to reduce the probability of initiation of defects in the area of arrest of the growth of martensite plates and reduce the level of localised stresses induced in these areas48-50 in cases in which the microsegregation of harmful elements at the austenite grain boundaries is not modified by other metallurgical measures.1,41,42 As a result of these changes of the initial microstructure of austenite and their consequences on the modification of the kinetic parameters of the subsequent martensitic transformation and martensite morphology

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the detrimental effect of the microsegregation of harmful elements at the original austenite grain boundaries and the restricted superimposed influence of the dynamic effects of martensite plates at the obstacles no longer cause the previously mentioned 'weakening' of the martensitic structure in the areas of arrest of growth of the martensite plates. Consequently, all processes causing or supporting the development of intercrystalline failure are suppressed and this fact, together with refining the structure of martensite, increase the mechanical and metallurgical properties of the martensitic matrix. The final effect is the higher degree of utilizing of the attainable level of the mechanical properties and an increase of the resistance to brittle fracture and/or suppression of the formation of premature and delayed fractures in the martensitic matrix. When applying TMT to low-carbon or low-alloy steels, the effect of the initial state of the austenitic matrix is reflected in its effect on the formation of fine-grained low-carbon martensite or a mixture consisting of low-carbon bainite and low-carbon martensite, and in its effect on hardening the matrix due to higher dislocation density in the basic (initial) austenitic matrix in comparison with the conventional heat treatment conditions. The assumption on the possible modification of dynamic effects in a phase transformation is less important under these conditions because these effects do not play any significant role in the formation of low-carbon lath martensite or in bainitic transformation in the total hardening characteristics of this type of structural steels subjected to unconventional heat treatment. In the final form, the modification of the initial structure or substructure of the austenitic matrix is reflected not only in strengthening effects but also in the resultant level of brittle fracture resistance in a different way. The controlling common parameter is, however, the refining of the initial austenitic structure or the development of a dislocation substructure which leads to considerable refining of the resultant products of transformation of austenite to martensite or bainite or a mixture of bainite and low-carbon lath martensite. However, different requirements are imposed on the structural and metallurgical properties of the initial austenitic matrix in the subsequent austenite to ferrite transformation under the conditions of controlled rolling or subsequent accelerated cooling. Deformation of the austenitic matrix, at the lowest possible temperature, has a beneficial effect reflected in the maximum refining of resultant ferrite. The frequency of occurrence of deformation bands in elongated austenitic grains which, in addition to the austenite grain boundaries, contribute significantly to the nucleation of ferrite grains, increases with increasing strain of the austenitic matrix.

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An important problem in optimizing the microstructural parameters of the resultant ferrite is the determination of physical metallurgically substantiated relationships between the microstructure state of initial austenite and the resultant ferrite, as a basis for controlling the microstructure of the resultant ferrite and predicting these microstructural characteristics. Therefore, special attention will be paid to predicting the size of ferrite grains formed during an isothermal phase transformation in a work-hardened austenitic matrix. For example, according to Scheil, these data can be applied to the continuous cooling transformation conditions, as carried out in Ref. 51. 2.4 Physical Metallurgy of the Ferritic Phase Transformation in Work-Hardened Austenite. The control of the size of ferrite grains in austenite decomposition is a very important parameter because the resultant mechanical properties and the brittle fracture resistance, especially at subzero service temperatures, are linked directly with the microstructural characteristics of the appropriate type of steel. In close connection with the development of technology of controlled rolling or accelerated cooling,52 it is necessary, to obtain a high level of toughness at lowered temperatures, to pay special attention to the conditions leading to extensive refining of the ferrite grains. Since the previously mentioned methods of unconventional treatment of the steel are usually accompanied by austenite decomposition under the work hardening conditions of this constituent, the proposed solution will be concerned with the analysis of the physical-metallurgical conditions leading to the formation of fine ferrite grains in the work-hardened austenitic matrix. To obtain basic and accurately defined conditions of controlling the ferrite grain size, we shall analyze the conditions of isothermal decomposition of work-hardened austenite to ferrite. Taking this solution into account, the parameters leading to the formation of fine ferrite grains using subsequent accelerated cooling of low-carbon or microalloy steels, will be processed. The high nucleation rate of ferrite grains is a very important characteristic for obtaining a fine-grained ferritic microstructure. This rate is associated with utilizing the higher concentration of potential sites for the nucleation of these grains in the initial austenitic matrix. The parameters of nucleation of ferrite grains in work-hardened austenite and the associated refining of the ferrite grains depend on the level of final deformation of the matrix prior to the austeniteferrite phase transformation. A very important result is the one which shows that the resultant ferrite grains are considerably finer when they formed in the initial

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Fig. 24 Dependence of the nominal ferrite grain size on the surface area of austenite grains for work hardened and recrystallized austenite. deformed austenitic matrix than in recrystallised austenite, even if the effective surface of the austenite grains is the same, Fig.24. In this connection, it has been concluded that the high level of 'localised energy' of the grain boundaries and deformation bands intensifies the nucleation process of ferrite grains in comparison with the nucleation of ferrite in the recrystallised austenitic matrix.53 The effect of work hardening of the initial austenitic matrix on the refining of the resultant ferrite grains is evaluated using the so-called transformation ratio (DF/DA), where DF or DA expresses the nominal diameter of the ferrite or austenite grains.53,54 In the solution obtained in this work, the following physical metallurgy principles of decomposition of the austenitic matrix were considered: a) the austenite grain surfaces are potential sites for the nucleation of ferrite grains. After plastic deformation, the initial austenite grains, being spherical, change to ellipsoids, where one axis is elongated in the deformation direction. In the solution, it is assumed that no recovery processes take place in the austenitic matrix prior to the subsequent phase transformation. In addition, deformation bands nucleate in the work hardened austenitic matrix. These bands also represent additional potential sites for the nucleation of ferrite grains. b) the surface of the deformation bands in the unit volume of the austenitic matrix Sdp is the quadratic function of strain ε = [ln(1r)]:

where A*is a constant which for the given type of low-carbon steel is

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on average equal to 30 mm-1,55 as verified by experiments on low-carbon microalloy steels, and r is the strain or the reduction in forming. , is the same as the rate of ferrite c) the rate of ferrite nucleation, related to the unit area of the deformation band nucleation on the surface of austenite grains in the work-hardened austenitic matrix (on the surface of the deformed austenite grains)

.

d) assuming the unit radius of the spherical austenite grains in the initial condition, the surface of the austenite grains can be expressed by the following simple relationship:

If one half-axis of the ellipsoid formed in plastic deformation is a = [1/(1 r)], the second half-axis does not change, b = 1, and the third half-axis c = (1 r), it is possible, if the volume is constant, to determine the surface of the considered ellipsoid using the following relationship

where

The elliptical integrals tabulated values for E(α, ϕ) and F(α ,ϕ), since k = sin α.56

can be determined from the

The ratio of the surface of the austenite grains prior to and after forming in relation to the strain is shown in Fig.25. As indicated by the graph, the ratio rapidly increases at strains higher than 40%. In agreement with Umemoto et al,57 the controlling parameter of the grain size of ferrite is the ratio of the rate of nucleation and growth of ferrite (the parabolic constant of the growth rate g according to Zener is used).58 For the workhardened austenitic matrix, this ratio can be expressed by means of the following empirical relationship:55,59

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Fig. 25 Relationship between the relative increase of the surface work-hardened austenite and its deformation. and g0 are the rate of nucleation of ferrite at the austenite grain boundaries or the constant of the growth rate where of ferrite grains in the recrystallised (undeformed) austenitic matrix, B is a constant equal to approximately 100,55 is the rate of ferrite nucleation at the boundaries of deformed austenite grains and on the surface of deformation bands, and g(r) is the constant of the growth rate of ferrite grains in the deformed austenitic matrix. The values

or

are related to the unit area of the austenite grain boundary or the unit area of the grain boundary and the strain band. The value of the ratio is equal to approximately 3000 mm-1, as reported in Ref.57. This corresponds to the conditions of austenite transformation at temperatures between 650 and 700°C.53 Assuming that the potential nucleation sites for ferrite grains in the recrystallised matrix are only the austenite grain boundaries and the ferrite grains are ellipsoidal with the axis ratio 3:1, as shown schematically in Fig.26, the 'radius' of the ferrite grain r, formed on the surface of the austenite grain at time t (if transformation starts at τ0), can be expressed as follows:60

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Fig. 26 Expected form of a ferrite grain nucleated at austenite grain boundary. The parabolic constant of the growth rate of ferrite g0 can, in agreement with Zener,58 be expressed as follows:

where Dγ is the coefficient of carbon diffusion in austenite, C0 is the initial carbon content, equilibrium carbon contents in austenite and ferrite.57

and

are the

The increase of the fraction of the area of the ferrite grain boundaries Yfr nucleated on the unit surface of the grains of undeformed austenite at the nucleation rate

at time t is given by the relationship:61

After substituting the expression in equation (15) into equation (17) and after rearranging the terms, the final expression for determining Yfr has the form:

The fraction of the area of the austenite grain boundary occupied by ferrite grains at time t, Ys, can be calculated from the equation:

The total number of ferrite grains nucleated on the unit area of the austenite grain boundaries can be determined using the following equation:55

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After substituting the expression in equation (19) into equation (20) and after rearranging the terms, we obtain the following equation:

The nominal value of the diameter of ferrite grains DF can be calculated as follows:

where is the surface of the grain boundaries of recrystallised ferrite in the unit volume of the matrix. From equation (22) and after substituting the expression for ns (see equation (21)), it is possible to determine DF and, after appropriate rearranging, to obtain the final relationship for expressing the transformation ratio (DA/DF) for the recrystallised (undeformed) austenitic matrix:

For the deformed austenitic matrix, equation (23) can be written in the following form:53,54

where Sv(r) expresses the sum of the potential nucleation surfaces (austenite grain boundaries and deformation band surfaces) for the formation of ferrite grains in the unit volume of austenite and expression included in equation (14).

is given by the

The effective surface of austenite grains Sv(r) is given by the equation:

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where the values of S(r) and Sdp (r) can be substituted by the expressions given in equations (13) and (11). Figure 25 shows relative surface changes not only for the case in which only the deformation of the original spherical form of the initial austenite grain to ellipsoid is considered (equation (13)) but also for a case in which the effect of deformation bands (equation (25)) in relation to the initial condition

(equation (12)) is taken into account.

The dependence shown in Fig.27 indicates that the ferrite grains, formed in the work-hardened austenitic matrix, are greatly refined at higher strains. Analysis of the results also shows that the detected refining of the ferrite grains nucleated in the work-hardened austenite is caused not only by the increase of the effective surface of the austenite grains defined by the value Sv(r) (equation (25)), which can here be regarded as the 'geometrical factor', but also by the higher nucleation intensity of ferrite associated with the higher deformation energy localised in the area of the grain boundaries and the surface of the deformation bands. The effect of this higher level of nucleation intensity of the ferrite grains is included in the second term on the right hand side of equation (14).

Fig. 27 Dependence of the nominal ferrite grain size on the effective austenite grain surface for different degrees of work hardening (r).

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The results also show that the nucleation intensity of ferrite in the deformed austenitic matrix during its transformation is also greatly influenced by the development of recovery or recrystallization processes.52,62 Taking into account this evaluation, the additional assessment the effect of the cooling rate on the modification of the ferrite grain size was carried out. For subsequent accelerated cooling, the empirical dependence for estimating the nominal ferrite grain size has the following form:57

where q is the cooling rate (in K s-1). However, equation (26) shows that accelerated cooling results in further refining of the ferrite grains, although the value of the exponent at q is relatively low.52,61 For the work-hardened initial austenitic matrix, the value of DA in equation (26) can be replaced by an effective value which also includes the effect of deformation bands induced in austenite grains and forming potential sides for the nucleation of ferrite, as discussed previously. The nominal value DA can be expressed as follows:

where the value Sv (r) is given by the expression in equation (25). The values of DA can be used to determine, for the appropriate level of the cooling rate q, the minimum value of DF in accelerated cooling of work-hardened austenite, following the process of controlled rolling, as discussed in detail in Ref.52. Consequently, the proposed evaluation method makes it possible to predict quite reliably (at the given level of strain hardening of the austenitic matrix) the resultant microstructural characteristics in order to create suitable conditions for predicting the final mechanical and metallurgical properties of this type of structural steel. Despite the fact that this solution is related to evaluating the microstructural parameters of ferrite formed in work-hardened austenite under isothermal conditions of its breakdown, using suitable methods it is possible to carry out conversion to continuous cooling transformation conditions of austenite, as mentioned previously.51 Taking these results into account, it may be concluded that the modification of the nucleation intensity of ferrite in the deformed austenitic matrix is also greatly influenced by the development of recovery or recrystallization processes in austenite, as discussed later in connection with examining the effect of development of these restoration processes

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on the nucleation characteristics of ferrite under accelerated cooling.63 Controlled rolling can be combined with accelerated cooling from the finish rolling temperature (after completing forming) which leads to further refining of the resultant ferrite grains. This is caused by an increase of the rate of ferrite nucleation in the austenitic matrix during the phase transformation and by suppression of the growth characteristics of ferrite. However, the transformation characteristics can also partly change in cases in which the austenitic matrix undergoes recovery or recrystallization processes. This is associated to a certain extent with the duration of the time period between completed forming (work-hardened condition of the austenitic matrix) and the start of accelerated cooling.64 More detailed analysis was carried out using the diagram of forming and application of accelerated cooling, Fig.28. The application of accelerated cooling with a simultaneous reduction of transformation temperature increases the nucleation rate of ferrite grains, Figs.29 and 30. The tendency to the increase of the nucleation rate of ferrite depends on the strain at temperatures around 800°C; if the strain is zero, the number of ferrite nuclei formed under different isothermal breakdown conditions changes only slightly (intracrystalline nucleation does not take place). On the other hand, higher deformation of the austenitic matrix (in deformation at temperatures around 800°C) results in a rapid increase of the nucleation rate of ferrite, especially at the lowest temperatures considered, due to the superposition effect of intracrystalline nucleation. However, the rate of intracrystalline nucleation in transfor-

Fig. 28 Diagram of selected forming conditions.

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Fig. 29 Dependence of the nucleation frequency of ferrite on the strain and transformation temperature.

Fig. 30 The dependence of intracrystalline nucleation ferrite on strain and transformation temperature.

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mation at temperatures higher than 680°C is relatively low, although the final deformation, carried out at 800°C, was relatively high (even at ε = 0.69), as indicated by the dependence shown in Fig.30. However, intercrystalline nucleation of ferrite was also observed at lower deformations (around 10%) which do not cause deformation bands to form in the austenitic matrix. It may be assumed that some other potential nucleation sites operate under the accelerated cooling conditions and this increases the intracrystalline nucleation intensity of ferrite. After deformation in which recrystallization is restricted (ε = 0.36, 800°C) and during subsequent isothermal holding for approximately 100 s, the austenitic matrix undergoes almost complete static recrystallization; after isothermal holding for approximately 1015 s partial static recrystallization of the work-hardened austenitic matrix is already recorded. The development of the processes of recovery or (static) recrystallization reduces the number of ferrite nuclei, especially the formation of nuclei in the volume of austenite grains (intracrystalline nucleation), as indicated by the results presented in Fig.31a, b. These data show that accelerated cooling applied immediately after deformation (suitable conditions for recovery or recrystallization processes are not created) supports additional refining of the ferrite grains. The intracrystalline nucleation of ferrite is an important parameter and a source of these fine ferrite grains. It can be concluded that by a) suitable combination of the degree of deformation, in suppressing the development of dynamic recrystallization during forming, b) suppressing subsequent static recovery and recrystallization, and c) higher rate of accelerated cooling after forming (deformation), required for higher rates of intracrystalline nucleation of ferrite, it is possible to obtain a favourable and balanced level of the mechanical and metallurgical properties of structural steels for wide technical applications (high-strength low-alloy steels). It is useful to mention here that the effects under b and c will have the controlling effect in accelerated cooling. To increase hardenability, it would be possible to increase the carbon and manganese content of this type of steel, but this modification of chemical composition, especially the increase of the carbon content, would lead unavoidably to a reduction of the achieved toughness and weldability characteristics of the structural steel. However, the manganese content should not be increased above 2% because of processing reasons. The susceptibility of the manganese addition to the formation of segregation bands results in a risk of the local manganese content increasing above 2.5%. Consequently, in these regions the subsequent phase transformation may cause even martensite or lower bainite to form, with subsequent consequences resulting in a reduction of the brittle fracture resistance. The hardenability, especially ferritic hardenability, is effi-

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Fig. 31 Effect of development of recovery processes taking place in the austenitic matrix on the nucleation characteristics of ferrite. a) changes in the value of the softening parameter, b) changes in the number of nucleated ferrite grains and the development of intracrystalline nucleation of ferrite. ciently increased by the addition of molybdenum, chromium or nickel. However, with a large increase of hardenability, the decomposition of austenite leads to the formation of a bainitic structure followed by the reduction of the toughness and brittle fracture resistance. It is therefore necessary to select a balanced modification of the chemical composition and cooling rate, in order to ensure maximum refining of the ferrite grains and controlled precipitation of carbides or carbonitrides of microalloying additions (for example, niobium, titanium or vanadium) without reducing brittle fracture resistance.

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The optimum properties of these steels can be obtained efficiently by utilizing the phase transformation in the workhardened austenitic matrix to fine-grained ferrite. These conditions of decomposition of austenite result in the formation of a fine ferritic structure which has, in addition to satisfactory strength properties, also high toughness. Figure 32 shows the dependence of the ferrite grain size on the selected cooling rate (at temperatures around 600°C) of a low-carbon steel microalloyed with niobium at different initial austenite grain sizes.26 The graph shows that the effect of the initial austenite grain size is exceptionally strong and that the application of accelerated cooling is evidently an effective supplementing factor. The results show that to determine the optimum controlled rolling conditions and obtain balanced strength properties with brittle fracture resistance, it is essential to know the relationships between the structural and metallurgical characteristics of the initial austenitic matrix and the resultant ferrite grain size or the distribution of the ferrite grain size. An exceptionally large increase of the final strain does reduce the 'height' (h) of elongated austenite grains in the cross section of the formed material, thus exerting a favourable effect on the number of potential sites for ferrite nucleation but, at the same time, it may increase partially the temperature of the start of the austenite to ferrite transformation. This effect may also slightly increase the size of the resultant ferrite grains and this restricts the expected beneficial effect of the higher nucleation rate after the selected high final plastic deformation of austenite. The relationships between these mutually competing processes are

Fig. 32 Effect of cooling rate on the ferrite grain size.

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shown in Fig.33,27 which shows the relationship between the resultant austenite and ferrite grain sizes and the final deformation carried out in the temperature range 880 to 820°C in a steel microalloyed with niobium. In the examined case, after final deformation (last pass) of the specimen, the latter was cooled in a salt bath at a temperature of 650°C at which isothermal transformation of austenite to ferrite took place (over a period of 120 s) followed by air cooling. As shown by the graph, the mean ferrite grain size with a diameter of 15 µm, was obtained in cases in which the initial austenitic structure consisted of recrystallized grains with a mean height (h) of 30 µm. If the material was formed at a total deformation of austenite (with suppressed recrystallization) of around 90%, the mean size of the ferrite grains was 5 µm. Figure 33 shows that the height h of the austenite grains, determined in the cross section of the material, decreases with increasing final deformation under the conditions of suppressing the development of recrystallization in the austenitic matrix. Similarly, the mean size of the ferrite grains decreases. The graph shows that the ratio between the height h of the austenite grains and the mean diameter of the resultant ferrite grains DF is approximately 2:1, so that:27

Fig. 33 Changes of the height (h) of austenite grains and ferrite grain size in relation to final strain of austenite (46 passes, deformation temperature in the range 880-820°C).

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This relationship also shows that under the conditions of exceptionally high final deformation of the austenitic matrix, the ferrite grains nucleate preferentially at the elongated austenite grain boundaries and grow into the 'centre' of the elongated austenite grains, up to contact with the grains growing from the other side of the austenitic boundary. This shows that the ferrite grain size DF is approximately equal to (h/2). This means that areas other than the elongated austenite grain boundaries are not longer, under the given conditions, characterized by the very fine initial austenite grains, representing potential nucleation areas for ferrite formation. However, examination of larger work-hardened austenite grains shows unambiguously that deformation bands exert a strong effect in ferrite nucleation.65 Figure 34 shows27 the relationship between the austenite and ferrite grain sizes, together with the data for the austenite to ferrite phase transformation, for both the recrystallised matrix and the work-hardened austenitic matrix. The dependence shows that for larger recrystallised austenite grains potential areas also include areas other than the surface of the austenite grains (constant ferrite grain size). It is likely that this is caused by the effect of, for example, annealing twins and inclusions, etc. This graph also shows that the relationship between the austenite grain size and the size of the resultant ferrite grains does not hold for larger austenite grains. When the height of these grains is greater than approximately

Fig. 34 Dependence of the ferrite grain size on the austenite grain size(height).

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35 µm, the relationship between the sizes of the two compared grain types is no longer linear. In the examined type of steel (low-carbon steel microalloyed with 0.05% niobium) the growth of the ferrite grains is 'arrested' with increasing austenite grain size, if the height of the austenite grains is greater than approximately 80 µm (see Fig.34).27,28 Taking into account the analysis of the physical metallurgy characteristics of controlled rolling, especially the evaluation of the nucleation parameters of the ferrite grains, it may be concluded that the examined process can be divided into three main stages:66,67 a) deformation takes place in the region in which austenite recrystallization occurs; b) deformation takes place in the region in which no recrystallization occurs; c) deformation occurs in the two-phase austeniticferritic region. These conditions are summarised in Fig.35 which shows schematically the appropriate changes of the microstructure of both the initial austenitic matrix and the resultant products of the phase transformation during austenite decomposition. As indicated by the discussion, the important parameters in controlled rolling are the deformation temperature and the degree of strain of the austenitic matrix. The technical and technological consequences of this process will be discussed separately in a special section of this work. When discussing the effect of TMT of high-strength and low-carbon steels and controlled rolling of low-carbon (or microalloyed) steels, special attention will be given to accurate determination of both the individual

Fig. 35 Schematic representation of three stages of controlled rolling.

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physical metallurgy parameters which contribute significantly to optimizing the achieved mechanical and metallurgical properties and the technical and technological parameters leading to higher applied properties of these types of structural steel.

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3 Thermomechanical Treatment of High-Strength Martensitic Steels As indicated by previous analysis, TMT is a heat treatment technique which includes plastic deformation of, for example, steels (but it can also be applied in processing aluminium or titanium alloys, to increase their mechanical and metallurgical properties) prior to or during a phase transformation. This increases the strength properties or ductility and brittle fracture resistance. However, this characteristic covers a very wide range of various types of TMT. The two main TMT variants used for high-strength martensitic steels are as follows:13 a) TMT in the stable austenite range which is usually referred to as high-temperature TMT, HTMT; b) TMT in the metastable austenite range referred to as low-temperature thermomechanical treatment, LTMT. Figure 36a, b shows schematically both types of TMT: HTMT and LTMT. From the technical viewpoint, HTMT can be regarded as highly promising because it requires a considerably lower investment. This method of unconventional heat treatment can be used in the existing industrial

Fig. 36 Diagram of low-temperature (LTMT) and high-temperature (HTMT) thermomechanical treatment.

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equipment, provided a suitable quenching system (quenching press) is available. Its main characteristics in sheet production will be described in a separate section,68 although these principles can be generally utilized in HTMT of other types of metallurgical products. It should be mentioned here that in applying TMT, not only in processing highstrength martensitic steels that also from the viewpoint of wider technical application, this heat treatment method not only increases the main mechanical and metallurgical characteristics but also results in a considerable energy saving because final treatment (quenching) is carried out directly from the finish rolling temperature to obtain the 'controlled' state of the microstructure of initial austenite.1,4,24,69 In comparison with HTMT, the possibilities of applying LTMT in the treatment of steels with very high strength have been studied far less extensively because relatively low temperatures at which LTMT can be efficiently applied (in the region of increased 'stability' of austenite, see Fig.36b) result in a rapid increase of the deformation resistance of austenite which greatly restricts the technological possibilities of this heat treatment method. In HTMT technology, the steel is formed in the stable austenite range, i.e. at higher temperatures. After forming, the steel is immediately, or after a certain controlled period of time, quenched in such a manner that its final microstructure is martensitic. This method can greatly increase the strength properties of the resultant martensitic structure whilst maintaining sufficiently high or even increasing the ductility properties and brittle fracture resistance. The state of the initial austenite grains formed in this process differs depending of the type of processed steel, forming temperature, strain or the sequence of individual deformation cycles, and the time elapsed since the completion of forming to the start of quenching. Starting with the optimum condition, i.e. the selected type of steel contains elements preventing the development of recrystallization and the forming temperature is slightly higher than temperature Ac3, with a sufficient deformation level (the time between completion of forming and the start of quenching is minimum), it can be expected that the structure of austenite will contain no recrystallised grains. Of course, these considerations do not take into account the possible effect of dynamic recrystallization during the deformation process and special attention is given to the final deformation cycles. In addition, as shown previously by Grange,70 satisfactory characteristics can also be obtained with the fine-grained recrystallization of the initial austenitic structure where according to previous views,1,13,24,69 the free path for, for example, the growth of martensite plates, decreases thus restricting the dynamic effects in the martensitic phase transformation.38,39,71

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Fig. 37a Effect of isothermal holding at 900°C after forming in HTMT on the strength and ductility properties of type 2Cr2NiMnSiMo steel with 0.40%C (tempering at 100°C/4h/air).

Fig. 37b Effect of strain of 2Cr2NiMnSiMo steel at 0.27%C on the resultant strength (tempering at 100, 150, 200°C/4h/air).

These considerations on the effect of technological parameters show the results of evaluating the effect of the isothermal holding time prior to quenching and of the degree of deformation of the austenitic matrix. A complexly alloyed 2Cr2NiMnSiMo steel with a carbon content of 0.40% C was formed at 900°C with a total deformation of 85%. The steel was used to examine the effect of isothermal holding at 900°C for 1 min to 1

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hour on the strength and ductility characteristics of martensite. After isothermal holding, specimens 3 mm thick were quenched in oil and tempered at 100°C for four hours. For comparison, these characteristics for the case of instantaneous quenching were also examined. The handling time between forming and quenching was 4 s. This tempering temperature was selected on the basis of previous experience because it makes it possible to examine, in the optimum form, the mechanical and metallurgical characteristics of martensite without the risk of occurrence of heterogeneities detected in evaluating as-quenched martensite because both relaxation processes and processes of ageing the matrix take place in this constituent to various degrees. After tempering at 100°C the properties of the matrix become more uniform but the number of defects in the martensitic matrix, directly associated with the initial state of the austenitic matrix and the nature of the austenite to martensite transformation, does not drop greatly.72-75 Figure 37a shows the dependence of tensile strength and elongation on isothermal holding time. As indicated by these dependences, after isothermal holding for around 60 s these properties decrease. The ductility decreases appreciably: after isothermal holding for 1 hour the elongation decreases by approximately 1% from the initial level of 1011%. This indicates premature failure of the tensile specimen. This type of steel (with a lower carbon content of 0.27%) was also used to evaluate the effect of strain on the strength properties. Strains of 30 and 50% were applied in a single pass at a rolling speed of 0.76 mm s-1 at 900°C. A strain of 85% was applied in three passes. In all cases, quenching was carried out 4 s after rolling. Figure 37b shows the dependences determined for the specimens quenched in oil and tempered at 100, 150 and 200°C for four hours. As indicated by the graph, at strains lower than 50% the effect of HTMT is minimum, the strength increase is very small, and the ductility level is also influenced only slightly. In evaluating the effect of strain in a steel with similar chemical composition but with a carbon content of approximately 0.40%, HTMT at a strain of 50 or 30% should have a stronger effect because in this steel with higher carbon content conventional heat treatment is usually followed by the drastic occurrence of premature failures. After forming with this strain, the susceptibility to the occurrence of 'a premature failure' in the stressstrain curve is partially suppressed.13 The main technological principle of LTMT is to obtain the maximum deformation level in the region of metastable austenite below the recrystallization temperature of austenite with subsequent quenching to produce a martensitic structure. This treatment should result in a higher strength level than HTMT, but a reduction of the ductility characteristics must be expected. However, effort should be made to ensure that this reduction does not exceed some limit. As mentioned previously,1,3

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this method of unconventional heat treatment can be efficiently utilized only in cases in which the region of metastable austenite is sufficiently wide to carry out all required operations in this region. Experience shows that LTMT can be applied efficiently in chromium-alloyed steels; in unalloyed (carbon) steels this method can not be used. In addition to the discussed parameters of the forming process such as the forming temperature and deformation, in LTMT it is also important to take into account the potential risk of the breakdown of austenite into non-martensitic products. The possibility of isothermal transformation of austenite during forming in the course of LTMT cannot be completely ignored because deformation of steel in the metastable austenite region greatly shortens the incubation period for the decomposition of the austenite matrix and accelerates the transformation kinetics. This undesirable austenite decomposition may reduce the final strength properties in comparison with the expected level for the purely martensitic structure. In addition, the structure formed in this manner, which contains a certain proportion of bainite, has considerably lower brittle fracture resistance. The positive point is that the strength of this microstructure is higher than that of the structure of a similar type formed under conventional heat treatment.1 Another very important parameter which must be considered is the requirement to avoid exceeding a certain limiting state of loading, for example, rolls, to avoid failure or damage to rolling equipment when the pressure on the rolls rapidly increases. Figure 38 shows the dependence of the mean specific pressure on forming temperature in a structural low-alloy steel of the 2Cr2NiMnSiMo type with 0.50% C. Simultaneously, the values of the applied deformation at the given forming temperature are presented in this figure. In all cases, the given deformation was obtained in a single pass. The deformation in a single pass was usually on the level corresponding to 40% (see forming at 500 to 700°C). At the lowest forming temperatures, the deformation in a single pass was limited taking into account the strength of the rolls and the stiffness of rolling equipment (experimental rolling mills). At deformations exceeding 20% the mean specific pressure rapidly increased to approximately 20002300 MPa (400450°C). Using the experimental method described previously,1 the mean specific pressure on the roll pmean was calculated from the measured rolling pressure P (N) using the following relationship (N mm-2):

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Fig. 38 Dependence of mean specific pressure on forming temperature of 2Cr2NiMnSiMo structural steel with 0.50%C (deformation in a single pass). where b0 is the initial width of the rolled specimen in mm; b1 is the final width of the rolled specimen in mm, and R is the roll radius in mm. In equation (29) h0 is the initial height of the rolled specimen and h1 is the final height of the rolled specimen (in mm). It can be concluded that the strength properties in LTMT (for the given type of steel) are controlled by the total strain and deformation temperature. A very important parameter is the need to prevent 'premature' transformation of metastable austenite during deformation. The strain rate and deformation mode are also important, probably because of the possibility of adiabatic heating with increasing strain rate. As indicated by the data for the characteristics of the deformation processes taking place during HTMT and LTMT, the extent of application of LTMT is restricted and it is important to take a number of essential technological measurements to prevent the phase transformation of metastable austenite into non-martensitic products and avoid an excessive increase of the rolling pressure or even failure of the rolls or rolling mills, usually of conventional design. On the other hand, the application of HTMT imposes considerably smaller requirements on ensuring the quality of the forming process. 3.1 Effect of Thermomechanical Treatment on Properties In examining the possibilities of obtaining the optimum properties of steels with very high strength, an important effect is played by the carbon content, in addition to the characteristics of the martensitic phase

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transformation and high hardenability. It is well known that higher carbon content greatly increases the susceptibility to premature failure.42,69,71 This strongly affects the extent of application of the high-strength steels of this type. The significance of this effect increases as the tempering temperature decreases, i.e. with increasing effect of the characteristic properties of the austenite to martensite transformation, especially 'dynamic consequences' of the formation of plate martensite.13 Basic data on the modification of the mechanical properties of low-temperature tempered martensite (after tempering at 100°C/4 h/air), in relation to carbon content after conventional treatment and HTMT, are presented in Figs.39 and 40 obtained for the 2Cr2NiMnSiMo type steel. For comparison, the graphs also show the dependence of tensile strength and elongation obtained after LTMT. Deformation was carried out in the metastable austenite range at 550°C. As indicated by these graphs, conventional treatment at a carbon content of around 0.40% is followed by a relatively rapid reduction of strength, whereas the maximum strength obtained after HTMT and LTMT is displaced to a content of approximately 0.52% (after HTMT) and 0.50% C (LTMT). The values of tensile strength obtained after LTMT are only slightly higher than those after HTMT, Fig.39. Similar dependences were recorded in evaluating elongation (A5). If conventional treatment is followed by a reduction of elongation at 0.38% C, then in HTMT a reduction is observed only at a carbon content higher than 0.50%. An elongation of 1011% is still obtained at a carbon content of approximately 0.50%. In LTMT, although the initial strength is only slightly higher than after HTMT, the ductility properties are lower, Fig.40. Detailed analysis of the properties of low-temperature tempered martensite shows that after conventional treatment, starting from a carbon content of 0.40% C, tensile specimens showed premature failures, usually accompanied by extensive intercrystalline failure at the initial austenite grain boundaries.41,42,69,71,74 After tempering at 200°C where the expected defects in martensite are already partially eliminated,13 as observed in the so-called as-quenched state or in a slightly modified form after tempering at 100°C, the dependence of elongation and strength on carbon content radically changes. The data are summarised in Figs.41 and 42. For comparison, the graphs again show the dependences obtained after HTMT and conventional treatment. The changes of the strength and ductility properties after LTMT are not included here because they were in satisfactory agreement with those obtained after HTMT. As indicated by the graphs, the susceptibility to premature failures of the tensile specimens is almost completely suppressed up to a carbon content of approximately 0.50% C. In the examined carbon content range, 0.250.55%, the increase of strength is constant and its average value is 250 to 350 MPa. Similar dependences were obtained

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Fig. 39 Dependence of the strength of 2Cr2NiMnSiMo structural steel on carbon content (HTMT and CT; tempering at 100°C/4h/air).

Fig. 40 Dependence of values of elongation of 2Cr2NiMnSiMo structural steel on carbon content (HTMT and CT; tempering at 100°C/4h/air).

in determining elongation (A5). Comparison of Fig.41 with Fig.42 shows that after tempering at 200°C/4h/air no premature failures were recorded at a carbon content of up to 0.50%. The elongation in this carbon content range is virtually the same after both HTMT and conventional treatment and its value is between 11 and 12%.69 For the examined type of structural steel, investigations were carried

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out to examine the dependence of strength on tempering temperature for two different carbon contents (0.32 and 0.54%) after HTMT and conventional treatment. Perfect martensitic hardenability was obtained in both cases so that the compared structural states are completely identical. Figure 43 shows the appropriate dependence, i.e. for a carbon content which after tempering at 100 °C in the case of conventional treatment was situated below the maximum strength, and for a carbon content which was located, after tempering the steel, in the region of premature failure. As indicated by this graph, in tempering at temperatures above 200°C the compared dependences are parallel and the increase of tensile strength after HTMT continues after tempering at 400°C.1,42 These conclusions are also supported by the results obtained for two types of steel, MnCr and CrMnSiMo, after HTMT and conventional treatment. These steels were quenched and then tempered at 200°C/4 h/air. Again, a constant increase of the values of Rm of approximately 200 MPa after HTMT can be seen over a very wide range of carbon content (to approximately 0.65%), Fig.44. As indicated by the initial results obtained in TMT,76 a number of metallurgical effects may influence strongly the resultant level of the strength and ductility properties after this unconventional heat treatment method. Similar conclusions were obtained using HTMT in processing low-alloy 30CrNiMo steel and in evaluating the susceptibility of this steel to delayed failure.1 The results show that the effect of production technology of high-strength (lowtemperature tempered) martensitic steels is important and that it can greatly modify the results obtained after HTMT in comparison with conventional heat treatment. Figures 45 and 46 show dependences similar to those in, for example, Figs.39 and 40, which were obtained after applying HTMT to 2Cr2NiMnSiMo steel produced in vacuum from highly clean initial materials. The phosphorus content of these melts was 0.0010.003% and the average sulphur content was around 0.0050.007%. As shown by Fig.45, in relation to the carbon content, in the examined range (0.300.56%) after conventional treatment (quenching in water and tempering at 100°C/4 h/air) there was no rapid reduction of tensile strength values at a carbon content higher than 0.40%. This was observed in melts produced by conventional metallurgical technology. Similarly, the elongation values were higher in vacuum-melted melts, Fig.46. After HTMT, the determined dependences Rp %C and A5%C, like in the commercially produced melts, Figs.39 and 40, showed no large reductions at contents of up to 0.50 or 0.52%. It may be concluded that the Rm values are almost parallel to the Rm %C dependence (CT) with the values increasing by approximately 200220 MPa in the examined carbon content range. The elongation values (A5) in this carbon content

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Fig. 41 Dependence of the strength of 2Cr2NiMnSiMo structural steel on carbon content (HTMT and CT; tempering at 200°C/4h/air).

Fig. 42 Dependence of elongation of 2Cr2NiMnSiMo structural steel on carbon content (HTMT and CT; tempering at 200°C/4h/air).

range in steels produced by vacuum technology and with high metallurgical cleanness are almost identical after HTMT and conventional treatment. This means that, in this case, the strength after HTMT increases on average by 200 MPa without any reduction of the ductility properties. For comparison, Figs.45 and 46 also show the dependences recorded for this

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Fig. 43 Dependence on the strength determined after CT and HTMT on tempering temperature for two melts of 2Cr2NiMnSiMo steel with 0.32% and 0.54%C.

Fig. 44 Dependence of strength on carbon content in type 2Cr2NiMnSiMo and CrMo steels after various treatments (HTMT, CT) and tempering at 200°C/4h/air). type of steel in production by commercial technology, as indicated by Figs.39 and 40.1,69,71

After tempering at 200°C/4 h/air, the Rm %C and A5 %C dependences of the vacuum melted 2Cr2NiMnSiMo steel have almost the identical course as with those obtained after tempering at 100°C. Figure 47 shows, for information, the dependence of the strength properties after tempering at 200°C for both compared variants (HTMT and CT). The graph shows that the maximum strength decreases by approximately 200250 MPa (at a carbon content of 0.450.50%), whilst the difference of

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Fig. 45 Effect of different production methods of 2Cr2NiMnSiMo steel on the strength obtained after HTMT and CT and tempering at 100°C/4h/air.

Fig. 46 Effect of various methods of production of 2Cr2NiMnSiMo steel on the resultant elongation values after HTMT and CT and tempering at 100°C/4h/air. the values of Rm between HTMT and CT remains almost constant in the entire examined carbon content range in the high-strength steel of the 2Cr2NiMnSiMo type.

The evaluation of the strength and ductility properties of the melts of the 2Cr2NiMnSiMo steel shows that in the melts produced using highly clean initial materials and vacuum technology conventional treatment

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never resulted in any large reduction of the values of Rm and A5 of the steels processed by CT at a higher carbon content and after tempering at 100°C. It may therefore be assumed that testing in the non-tempered condition will not show any reduction of these parameters, as in the case of commercially produced melts of these steels. In addition, in the melts processed by CT and tempered at 100°C did not lead to intercrystalline fracture, as observed in the commercial melts. Of course, this conclusion is valid fully for the examined carbon content range. Taking into account the previous experience,13,41,42 it may be expected that a certain degree of susceptibility to occurrence (although greatly restricted) of intercrystalline failure will also be observed in processing steels with a higher metallurgical cleanness, as indicated by further discussion of the physical metallurgy nature of this phenomenon. These results show that the effect of specific (vacuum) metalurgy will be stronger with increasing strength properties, although as indicated by the dependence discussed previously, the use of HTMT will 'weaken' this beneficial effect. Of course, it may also be assumed that HTMT can to a certain extent replace the negative consequences of the influence of the metallugical characteristics in commercially processed structural steels with very high strength. This dependence was also noted by May who stated that the use of vacuum technology increases the strength properties by approximately 57%. However, in tempering in 200°C these differences are reduced to a large degree,1,13,24,49,69 as indicated by Fig.47.

Fig. 47 Dependence of strength on the carbon content of 2Cr2NiMnSiMo steel after HTMT and Ct and tempering at 200°C/4h/air in melts produced by various technologies.

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These results are in complete agreement with the usual assumptions1,13,41,42,69 according to which the martensitic phase transformation results in the formation of defects at the austenite grain boundaries due to the dynamic effect of this transformation because the boundaries represent potential barriers to the growth of martensite plates.77 In accordance with the discussed results, this susceptibility is considerably stronger in melts produced by commercial metallurgy which are characterized by lower metallurgical cleanliness. Similar conclusions were also drawn from the results obtained in examining the susceptibility to intercrystalline failure of martensite in relation to the carbon content and the stage of tempering;1,41,42 in this conection, it has also been shown that the deviations in the 'extent of intercrystallinity' on the fracture surfaces may be caused by the effects of various metallurgical parameters in steelmaking.69 For example, in high-strength martensitic steels it was found that differences in the susceptibility of melts of the same type of steel to the occurrence of intercrystalline failure of martensite are obviously associated with changes in the level of the surface energy of the austenite grain boundaries. Melts characterized by higher susceptibility to intercrystalline failure on tensile specimens (in the as-quenched condition or after tempering at 100°C) showed higher values of the dihedral angle of etched grooves after vacuum etching.41 Assuming the surface energy in the transcrystalline region is constant, the increase of the angle θ in accordance with the equation:

expresses the reduction of the surface energy of the austenite grain boundaries γb. In this equation, γtr is the transcrystalline surface energy (energy of the so-called free surface). The melts with a higher susceptibility to intercrystalline failure show higher mean values of the dihedral angle (by 1012°), which represents a reduction of the relative surface energy of the grain boundaries by almost 40%. From the viewpoint of the simple energy balance, we can write the following dependence for the formation of intercrystalline failure in a clean matrix (without microsegregation processes at the grain boundaries)78

where γ expresses the energy required for the formation of brittle intercrystalline failure. If brittle fracture develops in a metallic matrix in which the microsegregation of harmful elements reducing the surface energy

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of the grain boundaries (γb) took place, the value of γ greatly decreases when intercrystalline failure takes place because the value of γtr, owing to the conserved microsegregation effect on the formed fracture surfaces, decreases very strongly. In addition, because of the development of microsegregation at the grain boundaries, the possibility of development of plastic deformation, usually characterized by the value γpl, may be also affected. According to approximate evaluation, it may be assumed that the following relationship expresses the reduction of γpl, in dependence of the reduction of γb:78

Consequently, at, for example, the previously mentioned 40% reduction of the surface energy at the grain boundaries, the level of the plastic surface energy which can be reached decreases by more than an order of magnitude. An increase of the surface energy at the austenite grain boundary results in a reduction of the probability of formation of intercrystalline defects after the martensitic phase transformation. This is obviously associated in the first instance with the high 'resistance' of the austenite grain boundaries to the dynamic effect of growing martensite plates whose propagation is usually arrested at this obstacle. In addition, another factor, characterized as a static effect and associated with the geometry (width) of martensite plates, can operate in this case.30,45 These considerations can be linked with the strong effect of the 'strength' of the austenite grain boundaries on the possible formation of microcracks (or microdefects) which represent nuclei for subsequent intercrystalline failure of martensite. These nuclei of intercrystalline failure, initiated in the area of arrest of growth of martensite plates, are directly linked with the occurrence of so-called premature failures in martensite and with the formation of delayed failures. These types of failure are accompanied by a developed intercrystalline failure. The resultant fracture surface follows to a large extent the original austenite grain boundaries. The results show that the strength of the grain boundaries or its reduction is strongly affected by the intercrystalline microsegregation of harmful impurities. Since in the melt produced by the commercial method this harmful effect was reduced after applying TMT (HTMT and LTMT), it may be assumed that in these unconventional methods of heat treatment the superimposition effect of martensite plates on the 'weakened' austenite grain boundaries is reduced. As already mentioned, Fig.22, in HTMT the dynamic effect of martensite plates on the austenite grain boundaries is reduced,1,13,31,39,41,42,44,45,69 in addition to a further restriction of the

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static effect because TMT is accompanied by refining of the martensite plates.13 As indicated by the dependences shown in, for example, Figs.39 and 40, these special characteristics are observed at higher carbon contents at which the microstructure of the martensite of the examined steels consists mainly of plate martensite with a large number of internal twins. The formation of this type of martensite is usually accompanied by the development of the so-called 'burst' phenomenon. In TMT, this phenomenon is suppressed or restricted so that one of the main sources of weakening of the austenite grain boundaries is eliminated (in most cases, martensite plates stop to grow at the austenite grain boundaries). The limiting case of the positive effect of TMT is the restriction of the discussed dynamic and/or also static effects in the growth of martensite and suppression of defect formation or its restriction resulting in the situation in which nucleated the microcracks are 'smaller', internal induced stresses are lower and, the level of threshold stress required for the formation of delayed fracture increases and the resistance to the formation of premature fractures is also higher. These results indicate that to restrict the formation of premature fractures in martensite and to obtain the required level of the ductility properties, it is necessary to prevent or restrict the formation of nuclei of microcracks at the austenite grain boundaries. This means that the initial austenitic matrix must contain a large number of obstacles for the growth of martensite plates. It is necessary to restrict the length of the free path for the growth of martensite plates and reduce the impact effects of these plates on obstacles in their development.13,69 These considerations must also be supplemented by the effect of carbon content because at lower carbon contents where lath martensite is the main component, the dynamic effect of the martensitic phase transformation is far weaker. These considerations are also fully supported by the results of examination of the effect of recrystallization of austenite in type 55MnCr steel in HTMT on the resultant mechanical properties. Taking into account the relatively high carbon content, this steel shows, because of its chemical constitution, a very high suscepibility to premature failure indicated by the dependence shown in Figs.39 and 40 obtained after tempering at 100°C.1,13,69,70 Figure 48 shows the dependence of the strength values on the isothermal holding time at 800°C after completing forming at a total deformation of 85%. The effect of static recrystallization is taken into account in this connection. The diagram shows the dependence for the case in which the specimens were tempered for four hours at 100 and 200°C. As indicated by these dependences, after approximately 3040 s of isothermal holding the strength of the specimens tempered at 100°C rapidly decreased. After tempering at 200°C, there was only a

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Fig. 48 Effect of the course of recrystallization of type MnCr steel on the strength obtained after HTMT and tempering at 100°C/4 h/air and/or 200°C/4 h/air. small change in the values of Rm in comparison with the initial condition. During isothermal holding for approximately 80 s both compared dependences 'intersect'. After tempering at 200°C, the tensile strength values are considerably higher than those recorded for the specimens after tempering at 100°C. The strength values show that tempering at 200°C does not cause premature failure. This evaluation was also supplemented by the determination of the austenite grain size carried out on specimens selectively etched in a picric acid solution with an addition of a surface-active agent. Figure 49 shows the dependence of the

Fig. 49 The effect of recrystallization of MnCr steel on the austenite grain size prior to the phase transformation to martensite.

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austenite grain surface area S (in mm2/mm3) on isothermal holding time at 800°C. These values were determined both in the longitudinal section in relation to the forming direction and in the transverse direction, to obtain maximally reliable results. As indicated by Fig.49, after isothermal holding for 90100 s the differences between the values in the longitudinal and transverse directions was eliminated. Comparison of the dependence of the values of Rm after tempering at 100°C on the isothermal holding temperature (Fig.48) and of the values of S on this temperature (Fig.49) shows that the start of the rapid change of the values of Rm and S is almost identical. Taking into account the previous considerations, it is evident that the rapid reduction of the values of Rm after isothermal holding for 3040 s is associated with a rapid increase in the austenite grain size (a reduction of the surface in the unit volume) and with the formation of a long free path for the growth of martensite plates and for the development of dynamic effects with the appropriate consequences on the initiation of defects in the area of arrest of growth of the martensite plates. Figure 50 shows the dependences of the values of A5 for the selected isothermal holding time at 800°C and also the frequency of occurrence of premature failures in the tensile test (in both cases, the tests were carried out on specimens tempered at 100°C). This evaluation also confirms that the start of the reduction of ductility corresponds to the start of the reduction of the strength values and that the susceptibility to premature failure increases at the same time.

Fig. 50 Effect of recrystallization of MrCr steel on elongation values after HTMT and tempering at 100°C/4h/air. Susceptibility to premature fracture is also given.

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Similtaneously, Figs.48 and 50 give the values of Rm and A5 determined after CT. In this case, the specimens were initially normalised (temperature 900°C) and then quenched from 840°C. The strength after this treatment and tempering at 100°C was only 1780 MPa which is lower than the value obtained after isothermal holding for 200 s at 800°C. Similar results were obtained in evaluating elongation. In the specimens treated by CT, the values of A5 were very low and did not exceed 1%, Fig.50. Even the longest isothermal holding time did not increase the austenite grain size to the level obtained after CT (S = 100 mm2/mm3).13,69 To supplement these results, a similar evaluation was carried out for the effect of the development of recrystallization on the resultant strength and ductility properties of 55MnCr steel modified by adding zirconium to restrict the recrystallization processes in austenite during HTMT. Figure 51 showed the dependence of the values of Rm on isothermal holding time at 800°C in a modified steel of type 55MnCrZr (0.08%) and, for comparison, in 55MnCr steel (Fig.48). The applied experimental procedure was the same as in the previously discussed case. As indicated by the compared dependences of Rm on the isothermal holding time (specimens were tempered at 100°C), the start of reduction of Rm in the modified steel of 55MnCrZr type was shifted by an average of half an order of magnitude to longer times (the start of reduction of the values of Rm was observed after isothermal holding at 800°C for 350400 s).

Fig. 51 Effect of recrystallization of MnCr steel after adding 0.08% Zr on the strength achieved after HTMT and tempering at 100°C/4h/air.

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Fig. 52 Effect of recrystallization of MnCr steel after adding 0.08% Zr on the elongation values after HTMT and tempering at 100°C/4h/air. The susceptibility to premature failure is also shown. The graph also shows the dependences of the values of Rp, which show that at the moment of rapid premature failure in the tensile specimens the values of Rp and Rm 'merge' together. Figure 52 shows similar dependences for the values of A5 and also the frequency of occurrence of premature failures. Again, the start of the rapid reduction of the values of Rm and A5 coincides. 3.2 Effect of Thermomechanical Treatment on the Modification of Microstructural Characteristics The beneficial effect of TMT on the mechanical and metallurgical characteristics1,41,42,69 has often been linked with the modification of the microstructure of the resultant martensite. It is obvious that the morphology of martensite plays a significant role in this case. Generally, it may be assumed that the free path length in austenite for the growth of martensite plates decreases, for example, fine martensite will form in HTMT. However, formation of lath martensite is accompanied by changes which are difficult to detect. The formation of this type of martensite is attributed to a lower carbon content. In this case, the packet dimensions, being an important structural component of lath martensite, are smaller, although this reduction in the dimension is not proportional to the refining of initial austenite. For example, in refining the austenite matrix grains by approximately 60% the packet dimensions are only approximately 40% smaller. On the other hand, no changes were detected in the mutual arrangement and geometry of the martensite laths.13 However,

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after preliminary deformation, the dynamics of formation of lath martensite are also reduced.38,39 It may be assumed that the laths form athermally by gradual nucleation, and the growth under the CT conditions takes place by means of a burst process with the lath being formed by a system of parallel, fine 'plates'.79 In both compared cases, it was found that the individual laths are usually separated by low-angle boundaries. High-angle disorientation between the laths was detected in only a very small number of cases and in certain cases there was even mutual twin disorientation.80 The substructure of lath martensite is characterized by a high dislocation density, so that the quantification of the contributions of HTMT from the viewpoint of maintaining the high dislocation density in the work-hardened austenitic matrix is very difficult.13 The effect of HTMT in steels whose microstructure consists mainly of lath martensite can be linked primarily with the reduction of the size of the lath packet and, consequently with a high dislocation density. In this case, the modification by the dynamic effect will not be strong.1,41,42,69 This is indicated by suppression of the susceptibility to the formation of premature or delayed failures in processing by HTMT or CT. If the martensite formed in the given type of steel is of the plate type, depending on the chemical constitution of the steel, TMT suppresses the formation of premature and delayed fractures which are directly linked with the formation of defects in the areas of interaction (arrest) of martensite plates with obstacles to their growth usually represented (especially in the initial stages of transformation) by the austenite grain boundaries. These questions were discussed in detail in the previous chapter. The quantitative evaluation of the change of the length and width of the martensite plates after HTMT in commercial steels of the 2Cr2NiMnSiMo or MnCr types could not be caried out because of the relatively low degree of resolution of the microstructure, in comparison with the FeNiC steels. However, on the basis of the results of examination of the substructure by TEM of thin foils it can be reliably assessed that not only the width but also the length of the martensite plates changes to the same degree as in the FeNiC steels.4 We shall now pay attention to evaluating the effect of work hardening of the austenitic matrix in TMT or restricting the dynamics of the martensitic transformation associated with shortening the free path for the growth of martensite plates on the microstructural and substructural characteristics of the martensite plates. A typical feature of this type of martensite in proccessing by CT is the formation of the internal so-called transformation twins which in various stages of transformation occupy larger or smaller parts of the resultant plate. The highest frequency of occurrence of the transformation twins was always observed in the central part of the plate and in the vicinity of the so-called midrib in the volume of the martensite plate.13

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After CT, the type FeNiC steels showed a high susceptibility to the formation of transformation twins, but after preliminary cold deformation at 2030% the occurrence of transformation twins in martensite was greatly restricted and in the majority of cases localised in a relatively narrow central region. The peripheral regions of the resultant plates, of course considerably smaller than after CT, showed a very high dislocation density. However, the measurements did not make it possible to state unambiguously whether they are only the 'surviving' dislocations from the work-hardened austenitic matrix or whether it is the superimposition process of generation of dislocations in martensite plates which is associated with the 'deformation' process in the plate in generation of transformation twins. The cellular dislocation substructure of austenite (formed in the work-hardened austenite of FeNiC steel) forms in the subseqent phase transformation a semipermeable obstacle to the growth of martensite plates (leads to gradual arrest of the growth of martensite plates on the wall of the dislocation cell),1,42 and this cellular substructure is 'inherited' by the martensite plates.13 An example of this inheritance of the cellular dislocation substructure in a martensite plate in the FeNiC steel is shown in Fig.53. Comparison of the dimensions of the dislocation cells in the initial work-hardened austenite and in the 'inherited' cellular substructure in the martensite plate shows that the size of the dislocation cells is similar in the two cases.81 Examination of the nature of transformation twins in type 2Cr2NiMnSiMo steel with 0.40% carbon shows that the transformation twins are present in the martensite plates after processing by both CT and HTMT, in contrast to the type FeNiC steels where the occurrence of transformation twins after deformation at high strains (around 50%) is almost completely

Fig. 53 An example of an inherited cellular dislocation substructure of austenite in a martensite plate of FeNiC aloy.

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Fig. 54 Substructure of 2Cr2NiMnSiMo steel after HTMT and tempering at 100°C/4h/air.

Fig. 55 (right) Substructure of 2Cr2NiMnSiMo steel after HTMT and tempering at 100°C/4h/air. suppressed. In the examined 2Cr2NiMnSiMo steel, the probability of occurrence of transformation twins after HTMT (total deformation 75%) decreased only slightly. For example, the probability of formation of transformation twins in martensite plates of 53%, measured after CT, decreased after HTMT to 48%. These results show that in this type of complexly alloyed martensitic steel, in contrast to the type FeNiC steel, the frequency of occurrence of transformation twins after HTMT is only slightly changed taking into account the results obtained after CT. Typical photographs of the substructure of the martensite of 2Cr2NiMnSiMo steel after HTMT are shown in Fig.54 and 55. The differences in the behaviour of the FeNiC and 2Cr2NiMnSiMo steels under the TMT conditions may also be associated with the fact that in 2Cr2NiMnSiMo steel partial recrystallization may take place during

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or after forming. This is due mainly to the possibility of development of static recrystallization because in the examined cases the period between the completion of forming (rolling) and the start of quenching was an average 15 s at temperatures of around 830850°C.42,73 Similar conclusions were also drawn in HTMT of a steel with 0.08% C and 5% Cr, as stated in, for example, Ref.13. In this steel, HTMT reduced the probability of formation of transformation twins after 60% deformation from 67% after CT to 59%, or after 30% deformation to 45%, which are values which fully correspond to the results obtained for the 2Cr2NiMnSiMo steel. Taking into account this 'direction' of the effect of the strain, it is not possible to reject completely the possibility of the 'competing' effect of the partial development of recrystallization on the probability of occurrence of transformation twins in martensite after HTMT. In a number of cases, the transformation twins were regarded as having a strong effect on increasing the strength of plate martensite, and it was assumed that the width or distance between the twins plays an important role, for example, analogously as it is in the case of evaluation of the pearlite interlamellar spacing. Figure 5658 show polyhedrons of the frequency of the relative width of the transformation twins in the martensite plates in relation to various types of heat treatment (CT, HTMT, TMT) in the type 2Cr2NiMnSiMo steel. As indicated by these diagrams, the maximum number of the determined values of the relative width of the transformation twins was in the range 811 nm in all examined cases. Figure 59 shows the mean values of the relative width and the relative

Fig. 56 A polyhedron of the frequency of the relative width of transformation twins after CT of steel 2Cr2NiMnhSiMo.

Fig. 57 A polyhedron of the frequency of the relative width of transformation twins after HTMT of steel 2Cr2NiMnSiMo.

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Fig. 58 A polyhedron of the frequency of the relative width of transformation twins after LTMT of steel 2Cr2NiMnSiMo.

Fig. 59 Mean width and mean spacing of transformation twins after CT, HTMT and LTMT, steel 2Cr2NiMnSiMo. distance of the transformation twins, as determined for the two melts of the previously mentioned steel. On the whole, on the basis of these results it may be concluded that both after CT and TMT the dimensions of the transformation twins do not change, as observed in the FeNiC steel. It is evident that the transformation twins do not play any significant role in increasing the strength properties of the steels after TMT and that, from this viewpoint, they are only a secondary parameter.1,42

It was also observed that the density of precipitates after low-temperature tempering of martensite is the same in both examined cases of heat treatment (CT and HTMT). This is in complete agreement with re-

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sults obtained by the majority of other authors. In none of the cases examined HTMT resulted in any increase of the tendency to the growth of carbides in comparison with CT.82 3.3 Evaluation of the Mechanical and Metallugical Properties of High-Strength Martensitic Steels after HTMT In addition to an increase of the strength properties without affecting or with a slight increase of the ductility properties, HTMT results in a considerably higher resistance to the formation of delayed1,41 and premature fracture42 or, in certain cases, this type of failure was completely eliminated. In comparison with CT, HTMT also results in favourable characteristics from the viewpoint of the resistance to stress corrosion cracking (for example, in an aqueous solution of NaCl) and to hydrogen embrittlement in cathodic hydrogen charging of specimens.13,73 In this section, we shall compare the values of fracture toughness after HTMT and CT of selected melts of 2Cr2NiMnSiMo high-strength martensitic steel (1.101.28% Mn, 1.101.40% Si, 1.401.08% Ni, 1.802.5% Cr, 0.230.30% Mo) with a carbon content of 0.2548%, and type MnCrSi steel (0.52% C, 0.65% Mn, 1.35% Si and 0.57% Cr) produced in an electric arc furnace by a commercial technology. The first of the steels was cast into 3 t ingots, the other one to 7 t ingots. Initially, the ingots were rolled by the conventional methods to plates (slabs) 40 mm thick which represent the initial material for final treatment. Prior to final rolling, the plates (slabs) were annealed at 680°C/20 h/air. The HTMT technology applied in this case consisted of heating to 950°C and rolling of plates in a three-high mill to a thickness of 1216 mm. The schedule of rolling was selected to ensure that the strain between the individual passes is constant. The total rolling time did not exceed 30 s and the finish rolling temperature was between 850 and 870°C. Immediately after rolling, the plates were quenched in water in a special quenching system whose main part was a quenching press.68 The handling time between the completion of rolling and quenching did not exceed 15 s. After quenching, the plates were immediately tempered for four hours at 200°C, followed by cooling in air. The comparison material, processed by CT, was initially normalised and tempered at 650°C for four hours. Final treatment of the steel consisted of quenching from 880°C in water and tempering at 200°C/4 h/air. The mechanical and metallurgical properties after conventional heat treatment of the MnSiCr steel were not examined because after quenching and tempering at 200°C premature fractures already occurred at a very low stress thus preventing the use of this steel in the condition after this type of heat treatment.83 The microstructure of both steel types after HTMT consisted of very fine martensite plates.

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The fracture toughness values KIC were determined in accordance with the appropriate ASTM standards in three-point bending of single-edge notch (SEN) flat specimens. The susceptibility to stress corrosion cracking in a 3.5% aqueous solution of NaCl was also determined on SEN flat specimens subjected to tensile loading for no more than 104 min.83 The specimen of this form were also used in evaluating the susceptibility of the examined steels to hydrogen embrittlement: the specimens were subjected to electrolytic hydrogen charging in 1 MHCl +0.1 MN2H2 solution at a current density of 2×102 A m-2 with simultaneous static tensile loading.42 The threshold values of the stress intensity factor in the corrosive medium (KISSC) and under electrolytic hydrogen charging (KISH) were defined as the highest values of the stress intensity factor at which failure or stable crack growth does not take place after exposure for 104 min. Examination of the effect of environment on the resistance to fracture was accompanied by evaluation of the dependence of the stable crack growth rate on the instantaneous value of stress intensity factor (K). The loading timecrack length dependence in stable crack growth was determined from deformation of the specimen using a compliance technique. 3.4 Summary of the Achieved Results and Their Physical Metallurgy Analysis Figure 60 shows the dependence of fracture toughness (KIC) on tensile strength (Rm), obtained in room temperature tests of 2Cr2NiMnSiMo steel melts with a graded carbon content. The arrow on the graph indicates the displacement of the examined dependence in HTMT, in comparison with CT, to the range of higher values of KIC and higher strength. The carbon content of the two compared melts was the same (0.33%).1 Figure 60 shows that at the given value of Rm the fracture toughness after HTMT is ~1720 MPa m1/2 higher than the values of KIC determined for the specimens processed by CT. At the same time, it may be seen that at the selected fracture toughness level the strength obtained after HTMT in on average 200230 MPa higher than after CT. Figure 61 shows the temperature dependences of the fracture toughness values determined for the melts A and B after HTMT andCT in the temperature range 196 to 20°C. These dependences also confirm that the fracture toughness after HTMT is higher not only in the room temperature test but also over a wide range of subzero temperatures. As indicated by the graph, at the same strength level of the compared melts, the fracture toughness in the examined temperature range increases by 1520 MPa m1/2. As already mentioned in the previous section, this beneficial

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Fig. 60 Dependence of fracture toughness (KIC) on strength (Rm) of two melts of 2Cr2NiMnSiMo steel after HTMT and/or CT and tempering at 200°C/4h/air.

Fig. 61 Temperature dependence of fracture toughness (KIC) for two selected melts of 2Cr2NiMnSiMo steel (melts denoted A,B).

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effect can be attributed to the refining of the martensitic structure after applying HTMT. The beneficial effect of HTMT was also reflected in the dependence of the load on the crack tip opening displacement (δ) determined in the fracture toughness test. Typical examples of the examined dependences for the specimens processed by the CT and HTMT methods are presented in Fig. 62. The dependences show that the compared treatments greatly differ in the nature of stable crack growth during loading. The R-curve was determined from the dependence of the load on the crack tip opening displacement (δ). Its form shows that the R-curve of the material processed by HTMT is considerably steeper than after CT. This form of the R-curve corresponds to the conditions under which brittle fracture resistance was higher.83 Figure 63 shows the relationship between the initial (initiation) value of stress intensity factor KIi and the time to fracture, determined for three selected melts of the examined 2Cr2NiMnSiMo steel with a carbon content of 0.350.39% after applying HTMT in testing in a 3.5% aqueous solution of NaCl, on the stress corrosion cracking resistance. The graph shows that the dependences have the same form and that the threshold values KISCC are distributed over a very narrow range. However, it is surprising to note that in the melt with the lowest carbon content, the threshold value is situated between the values determined for a carbon content of

Fig. 62 Load- crack tip opening displacement (δ) dependence of 2Cr2NiMnSiMo steel after HTMT and CT (melt B).

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Fig. 63 Dependence of the initial value of the stress intensity factor (Kli) and time to fracture (steel 2Cr2NiMnSiMo-HTMT, tempering at 200°C/4h/air).

Fig. 64 Dependence of the initial value of the stress intensity factor (Kli) and time to fracture after HTMT and CT of steel 2Cr2NiMnSiMo, melt B after tempering at 200°C/4h/air.

0.38 and 0.39%. However, as shown by experience,1,13,69,70 the scatter of this type in the high-strength steels is understandable because of the significant effect of the metallurgical characteristics.84 Figure 64 shows a similar dependence of the values of KIi on the time to fracture for the case of evaluation of stress corrosion cracking resistance and the resistance to hydrogen embrittlement in electrolytic hydrogen

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charging for the selected melt of 2Cr2NiMnSiMo steel with 0.33% carbon (melt B, see above) after CT and HTMT. The threshold value KISCC of this melt after HTMT was 28 MPa m1/2. The threshold value of KISH after HTMT was 10 MPa m1/2 and for CT is was 40% lower (6 MPa m1/2). Figure 65 shows the dependence of KIi on the time to fracture obtained for MnSiCr steel after HTMT, for both the state after electrolytic hydrogen charging and for testing in an aqueous solution of NaCl. The specimens treated by the CT methods were not tested because of the exceptionally high susceptibility of this steel to premature fracture after applied tempering.83 In this steel (in testing at a temperature of + 20°C) the fracture toughness was KIC = 38 MPa m1/2; after exposure in a 3.5% aqueous solution of NaCl the threshold value of KISCC was only 14 MPa m1/2, and after electrolytic hydrogen charging the threshold value KISH decreased to 50% of KISCC (KISH = 7 MPa m1/2). Examination of the effect of HTMT on the resultant mechanical and metallurgical properties was supplemented by a detailed analysis of the fracture surfaces. The fracture surfaces of the fracture toughness specimens tested at room temperature showed that the specimens treated by HTMT fracture mainly by the transcrystalline ductile failure mechanism accompanied by the occurrence of dimples. As the test temperature was

Fig. 65 Dependence of the initial value of the stress intensity factor (Kli) and time to fracture after HTMT of MnSiCr steel (after tempering at 200°C/4h/air).

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Fig. 66a Fracture surface of a fracture toughness specimen of 2Cr1NiMnSiMo steel (melt A) after HTMT and tempering at 200°C/4h/air. Test temperature 20°C.

Fig. 66b Fracture surface of a fracture toughness specimen of 2Cr2NiMnSiMo steel (melt A) after HTMT and tempering at 200°C/4h/air. Test temperature -196°C. reduced (Fig.61), the proportion of quasicleavage fracture, usually found on the quenched and tempered high-strength steel fracture surfaces, increased. Figure 66a shows the photograph of the fracture surface of melt A of 2Cr2NiMnSiMo steel after HTMT obtained in the room temperature fracture toughness test. Figure 66b shows a similar fracture surface of a specimen tested at 196°C. In the specimens of the steel processed by CT the fracture surfaces contained intercrystalline fracture areas at the austenite grain boundaries; these area formed not only in the room temperature test but also in testing at reduced temperatures. An example of the fracture surface determined on a specimen taken from melt A in testing at 40°C is shown in Fig. 67. The fracture area consists in this case of areas of intercrystalline fracture and regions of ductile transcrystalline fracture.74,75

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Fig. 67 Fracture surface of a fracture toughness specimen of 2Cr2NiMnSiMo steel (melt A) after CT and tempering at 200°C/4h/air. Test temperature -40°C. The fracture surfaces of the specimens which were tested for the resistance to stress corrosion cracking showed in all cases mostly intercrystalline brittle failure, as indicated by Fig.68a (fracture of the specimen of melt B Processed by HTMT in loading corresponding to K = 50 MPa m1/2). Figure 68b is the photograph of the fracture surface of MnSiCr steel after testing in a 3.5% aqueous solution of NaCl at a stress intensity factor of K = 20 MPa m1/2. In electrolytic hydrogen charging, the high frequency of occurrence of resultant quasicleavage fracture surfaces was found. Areas of intercrystalline brittle or transcrystalline ductile failure were observed in dependence on the applied method of heat treatment. Figure 69 shows the fracture surface of a specimen after HTMT (melt B) corresponding to the region situated close to the fatigue crack. The specimen was fractured at K = 12 MPa m1/2. In addition to transcrystalline cleavage failure, the fracture surface also contained localized areas of intercrystalline failure. In the specimens taken from the steel processed by the CT method, the number of areas of intercrystalline failure was always considerably greater than after HTMT. However, quantitative evaluation is very difficult because parts of the fracture area are attacked by corrosion thus preventing more detailed evaluation. The results show that in the 2Cr2NiMnSiMo steel HTMT not only increases the strength by 1215% in comparison with the strength obtained after CT, but the fracture toughness values are also higher. The increase of fracture toughness was recorded over a wide test temperature range from room temperature to 196°C (Fig.61). In comparison with CT, after HTMT the frequency of areas of intercrystalline brittle failure on the fracture surfaces of the fracture toughness specimens is smaller and the susceptibility to stress corrosion cracking and hydrogen

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Fig. 68a Fracture surface of a specimen for determining stress corrosions susceptibility of 2Cr2NiMnSiMo steel (melt B) after HTMT and tempering at 200°C/4h/air.

Fig. 68b Fracture surface of a specimen for determing the corrosion crack susceptibility of MnSiCr steel after HTMT and tempering at 200°C/4h/air.

Fig. 69 Fracture surface of a specimen for determining the susceptibility to hydrogen embrittlement of 2Cr2NiMnSiMo steel after HTMT and tempering at 200°C/4h/air.

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embrittlement is lower. Consequently, after CT the energy required for fracture formation is lower than after HTMT, as indicated, for example, by the fracture toughness value KIC (Fig.60). The beneficial effect of HTMT on the resultant threshold values of KISCC and KISH is confirmed by the results shown in Fig.64. In the type MnSiCr steel which has a high carbon content, HTMT results in a considerably higher level of the values than after CT after which the steel shows exceptionally high susceptibility to premature fracture. After HTMT, fracture toughness KIC is very high and the threshold values KISCC and KISH are satisfactory. This is a very important result because processing by HTMT creates suitable conditions for enlarging the range of application of this type of steel. It is evident that the resultant mechanical and metallurgical properties after HTMT are greatly affected by the refining of martensite plates, due to shortening of the free path available for their growth.14 However, it is also important to take into account the effect of preliminary deformation in HTMT. Shortening of the free path for the growth of martensite plates restricts the dynamic characteristics of the martensitic phase transformation and their consequences on the possible initiation of microcrack nuclei in the areas of arrest of martensite plate growth, usually frequently at the austenite grain boundaries or during mutual interaction of the largest martensite plates.13,69 In processing by CT in the areas of this interaction high localised internal stresses appear in the material and microcrack nuclei form preferentially in these areas. In the case of HTMT this effect is greatly reduced and this is confirmed not only by the reduction of the number of areas of intercrystalline brittle failure on the fracture surface but also by the suppression of the previously discussed premature and delayed fractures.1,41,42,85 3.5 Main Technological Principles of Application of HTMT in High-Strength Steels Taking into account the results of HTMT and LTMT investigations carried out in a laboratory rolling mill, we shall discuss the technological principles of application of HTMT on the production scale.1,42,70 Processing by HTMT the application of LTMT will not be discussed any further because of the considerable demands on rolling (forming) equipment enables the efficient utilisation of the existing rolling systems, but because of technical reasons, it is essential to ensure almost instanteous quenching in a suitable quenching press located behind the rolling mill.68 In HTMT in plate production, it is necessary to install straightening equipment in front of the quenching press in order to obtain the required flatness of the plates processed in this manner with minimum requirements on additional straightening. In HTMT of high-

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Fig. 70 Diagram of quenching equipment for HTMT plates. strength martensitic steels, the equipment should include a tempering furnace or tempering bath in order to carry out tempering as quickly as possible after quenching thus preventing the formation of defects on plates (products) in the asquenched condition.13,14,42,69 At present, the equipment for HTMT of plates is installed behind the three-high rolling mill; it consists of straightening and quenching equipment (quenching press). Straightening equipment removes unevenness of the surface of rolled plate in the hot state in such a manner that the actual quenching press 'fulfills the function' of the compressive force compensating quenching stresses induced in the plate. This equipment enables volume quenching to be carried out because the plate enters the water bath and the system of the water jets ensures the uniform flow of the quenching bath. This method results in the maximum uniformity of hardening along the width and length of the plates. As indicated by the results of hardness HB measurements, the scatter of the diameter of the indentation of a ball with a diameter of 10 mm at a pressure of 30 kN applied for 30 s did not exceed 0.05 mm.68 Uniform quenching throughout the length of the plate is ensured by the design of the pressing wheels in the water bath because they are in contact with the surface of the quenched plate only 'in a straight line' in sections 60 mm long. In this system, the contact of the wheels with the surface of the plate is not stable because the plate travels through the quenching equipment at a constant speed. The overall schematic of straightening equipment and of a quenching press is shown in Fig.70.68,69 The straightening equipment is of the roller design with five bottom and top pressure rollers which can be independently adjusted. The bottom rollers are driven by two electric motors with a power of 30 kW. Through an epicyclic gear and a drive housing, the force is transferred to rollers with a diameter of 220 mm. Their cir-

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cumferential speed is adjusted to the speed of the rolling mill and is approximately 0.7 m s-1. In the existing rolling system, the effective width of plates processed by HTMT is not higher than 1100 mm, the maximum thickness is 25 mm. Load on the equipment was calculated for the hot strength of the material of 120 MPa. As shown by Fig.70, straightening equipment inserts the HTMT treated plate into two pairs of feed rollers in front of the quenching bath which then moves the plate through the entry mechanism into the operating space of the quenching press. The press consists of two independent containers, a lower collection container with a volume of 6 m3 and an upper container with a volume of 3.5 m3 of water. The upper container contains holding wheels whereas in the lower container the wheels are driven by a 30 kW motor, as in straightening equipment through the epicyclic gear housing. The upper wheels can be centrally adjusted by means of two rack bars. Quenching water is pumped from the lower container with a pump with a capacity of 2 m3 per minute at a pressure of 3.5 × 105 Pa into the upper container and is distributed into two sections of the upper and lower jets. Jets with a diameter of 4.5 mm are directed against the movement of quenched plate. The upper quenching container contains a total 140 jets. The injector effect of the water stream and the jets stirs the surface of the water in the container so that the steam formed at the lower plate area is immediately removed. The height of the level in the upper container is determined by a overflow orifice which transfers water back to the lower reserve container where it is mixed with the supplied cold water. To maintain the water temperature below 25°C, it is necessary to supply 320 litres of cold water per minute. After quenching, the plate leaves through the output gate to a storage grid where the plates are assembled and then tempered in batches. This quenching equipment can be used for HTMT of plates with a thickness of 425 mm and a maximum width of 1100 mm and a length of up to 3300 mm, and ensures the required straightness of the plates thicker than 6 mm, i.e. 10 mm per 1 m of length. In thinner plates, the resultant straightness is slightly worse. These dimensions, the width and length of the plate, processed by HTMT, are not linked with a restriction which results from the capacity of the described equipment (straightening equipment and quenching press), but is associated with the capacity of the rolling mill and the maximum weight of the applied slabs which can be heated in a walking beam furnace which is available in the given rolling mill train.68,69 This type of equipment can also be used to process other types of metallurgical products (sections,tubes). In addition, another type of quenching equipment was developed in this area which enables highly promising results to be obtained for thinner plate.1,42

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4 Thermomechanical Treatment and Controlled Rolling of Carbon and Low-Alloy Steels As already mentioned, another area of technical application of TMT is the use of this heat treatment method in processing steels for wider technical applications where the transformation usually also results in the formation of nonmartensitic microstructural components. For example, when using conventional TMT technology followed by quenching in a quenching press, low-carbon or low-alloy steels have a structure of either low-carbon bainite or a mixture of low-carbon bainite and martensite. In controlled rolling, possibly supplemented by accelerated cooling, the steel usually contains a mixture of fine-grained ferrite and bainite (or pearlite) or bainite and low-carbon martensite, depending on the final cooling from finish rolling temperature after controlled rolling.63,64 Therefore, in the following part of this work attention will be given to the physical metallurgy characteristics of using HTMT in optimising processing of the selected types of steels for wide technical application and technical and technological characteristics of controlled rolling (together with subsequent accelerated cooling) will be discussed. The conditions of modifying the structural and mechanicalmetallurgical properties of the examined structural steels after the given methods of unconventional heat treatment will be analyzed. In the final part of this section, detailed analysis will be made of the possibilities of further development of controlled rolling or other methods of heat treatment and the possibilities of predicting the resultant physical metallurgy characteristics of structural steels for wider technical applications. 4.1 Application of HTMT in Strengthening the Matrix of Steels for Wider Technical Application and Determination of the Relationships between the Initial Structure and the Resultant Physical Metallurgy Properties of Products of Austenite Decomposition As already indicated in the previous part of this work, the resultant strength or mechanicalmetallurgical properties of these types of steel can be greatly improved by applying HTMT. This treatment results in higher strength properties without reducing the ductility characteristics or brittle fracture resistance.1,4,13 In this section, we shall summarise the results ob-

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tained after applying HTMT in unconventional heat treatment of plates of low-carbon or low-alloy structural steel. We shall discuss the relationships between the technical and technological parameters used in this treatment and the resultant microstructural characteristics and the mechanical properties of the examined types of structural steels. The main bulk of research was carried out on two types of commercially produced steel for wider technical application with the chemical composition given in Table 1. Table 1 Chemical composition of steels for wider application (wt.%) Steel

C

Mn

Si

P

S

A

0.18

1.38

0.40

0.022

0.016

B

0.16

1.13

0.34

0.016

0.014

Cr

Ni

Mo

B

0.63

0.20

0.52

0.0032

The effect of HTMT was examined under the conditions obtained in rolling plates 11 and 18 mm thick under the optimum technical and technological conditions of this method of unconventional heat treatment obtained in its application to high-strength martensitic steels.1,4 The main technological parameters in this case are: heating temperature, rolling and finish rolling temperature, rolling time and the applied deformation level (possibility of development of dynamic recrystallization), the holding time between completed forming and the start of quenching in a quenching press (possibility of development of static recrystallization). The experimental material was heated to 930950°C. Plates 11 mm thick were rolled in 7 passes and 18 mm thick in 5 passes. During forming, the experimental material was rotated by 90° in order to minimise the anisotropy of the properties in transverse rolling. The average rolling time was 100 s and the holding time after completing rolling and the start of quenching was 10, 30 and 50 s. This processing method makes it possible to examine the effect of development of static recrystallization of the material prior to its quenching. Attention was also given to the effect of a possible reduction of temperature prior to quenching in the quenching press to avoid the undesirable phase transformation of austenite to ferrite. According to the currently available experience, the temperature at the start of quenching must not decrease below 820°C.4 In addition, the temperature at the start of quenching should not be higher than 900°C to minimise the risk of development of static recrystallization. For comparison, the same melts were subjected to conventional heat treatment (CT). Steel A was, after preliminary normalisation at 920°C, quenched from 880°C. Steel B was initially normalised at 920°C and then quenched from 870°C. Several of the experimental specimens (af-

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ter HTMT and CT) were examined in the untempered condition to obtain, although indirectly, basic information on the structural and metallurgical parameters of initial austenite (especially from the viewpoint of evaluating the development of recrystallization). The remaining part of the material was then tempered for four hours between 550 and 625°C for steel A and between 550675°C for steel B, followed by cooling in water. The development of recrystallization processes when applying HTMT to the steel A and B in the period between completed rolling and the start of quenching in a quenching press was examined on untempered specimens by evaluating changes in the size of the surfaces of selectively etched austenite grains by the intercept method as the size of the surface of austenite grains in the unit volume S (mm2/mm3).86,87 In this evaluation, the effect of the anisotropy of the properties on the measurement results was taken into account.4 Table 2 summarises the results of measuring the surface of austenite grains (in the unit volume) after CT and after the selected variants of isothermal holding in HTMT. As indicated by Table 2, in the 18 mm thick plates, produced from both steels examined, the degree of development of recrystallization is higher than in the plates 11 mm thick. The difference in the degree of development of recrystallization is higher in steel B. In the 11 mm thick plates, quenched after isothermal holding for only 10 s the surface of the austenite grains is on average 2.22 times greater after HTMT than after CT. In the 18 mm thick plates (steel B) after isothermal holding for the same period of time this ratio is 1.48. After isothermal holding for 50 s this ratio decreases to 1.1. In steel A this ratio is (isothermal holding time 10 s) 1.30 for the 18 mm thick plates and 1.45 for the 11 mm thick plates. A surprising result was obtained in measuring the surfaces of austenite grains after CT in steel A in the 18 and 11 mm thick plates where a difference of up to 60% was found between these values (see Table 2). Table 2 Surface of austenite grains after HTMT and CT S (mm2/mm3) Steel

18 mm plate CT

11 mm plate

HTMT

CT

10 s

30 s

50 s

HTMT 10 s

30 s

50 s

A

70.5

92.2

82.8

75.0

116

169

158

127

B

256

376

317

292

263

570

506

475

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However, the reason for this could not be determined in the presented solution. The higher degree of recrystallization in the 18 mm thick plates can be evidently attributed to the fact that the temperature at which forming was completed was in these plates on average 30°C higher than in the 11 mm thick plates. However, a certain disadvantage of this solution is the fact that the possible extent of dynamic recrystallization cannot be defined more accurately. The compared values of the surfaces of the austenite grains can be used to determine reliably only the effect of static recrystallization taking place in the matrix of austenite during isothermal holding after completing rolling up to the start of quenching in the quenching press. As further indicated by Table 2, in steel A after isothermal holding 50 s the difference between the austenite grain surface areas S after HTMT and the compared value of CT is minimum and can be characterized by a ratio 1.07 and 1.09. Figures 71 and 72 summarised the basic mechanical properties (determined at a test temperature of 20°C on material taken from 11 mm

Fig. 71 Mechanical properties of steel A.

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Fig. 72 Mechanical properties of steel B. thick plates) in relation to tempering temperature for both structural steels examined. These tests were accompanied by evaluation of the hydrogen embrittlement susceptibility (determined in electrolytic hydrogen charging) of the thermomechanically processed structural steels A and B on the basis of the changes of the threshold value for the formation of delayed fractures determined in testing under the conditions corresponding to the optimised heat treatment conditions of these steels.88,89 In the material taken from 11 mm thick plate (steel A), quenched after isothermal holding for 10 s, after forming (HTMT) increased the threshold value from 400 MPa (after CT) to 450 MPa. In steel B (plates of the same thickness, tempered at 675°C, the threshold value for the formation of delayed fractures increased from 400 MPa after CT to 500 MPa after HTMT. On average, in steel A HTMT resulted in a 1012% increase of the threshold value, whereas in steel B the threshold value increased by 2025%. Figure 73 shows the dependence of transition temperature (defined by the level of notch toughness 50 J cm-2 on KCU specimens) on yield

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Fig. 73 Dependence of transition temperature (50J cm-2) on yield strength. strength (Rp). After HTMT, the transition temperature (defined in this manner) of steel A decreased by 20°C; in steel B, the transition temperature decreased by as much as 4050°C in comparison with CT. A similar tendency was also recorded in evaluating the critical crack opening displacement δ (COD).4 In steel A, processed by HTMT and tempered after quenching at 625°C, the yield strength increased by 60 MPa in comparison with yield strength values after CT. In steel B after HTMT the yield strength increased on average by 120 MPa. The temperature dependences of the values of Rp were used to determine the athermal component σµ. and thermally activated component σ*. This analysis shows that the detected increase of the yield strength is given mainly by the increase of the value of σµ. The temperature dependence of the thermally activated component of yield strength σ* is almost identical for both compared heat treatment variants of steels A and B. In the as-quenched condition, the microstructure of the examined steels consisted of a mixture of low-carbon bainite and low-carbon lath martensite. These products of transformation of the austenitic matrix are closely linked with the initial dimensions of the austenite grains, namely the size of packets of resultant bainite and lath martensite is directly given by the initial austenite grain size.90 This means that the obtained strength parameters91 are closely linked with the values of the austenite grain surfaces S, see Table 2, and, therefore, the final mechanical and metallurgical properties, determined after tempering, depend to a certain extent on the morphology and/or the structural and metallurgical state of initial austenite.4

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According to Refs.92 and 94, the resultant level of the yield strength can be expressed as follows:

The second term on the right-hand side of this equation expresses the effect of dispersion hardening according to Orowan, the third term gives the effect of dislocation hardening and the fourth term includes the effect of the size of packets of bainite or low-carbon martensite (with the appropriate relationship with the initial austenite grain size).94 The effect of dispersion hardening, determined on the basis of statistical analysis of the distance between the precipitates (carbides), as determined by measurements carried out on carbon extraction replicas, can, in accordance with equation (33), be expressed as follows:93

where G is the shear modulus, b is the Burgers vector, d is the diameter of precipitated particles, and D* is the distance between the precipitates (interparticle spacing) on the potential slip plane. This value of D* can be determined from the following relationship:

where λ is the mean free path between the precipitates determined on the surface of the carbon extraction replica, f is the volume fraction of the carbide phase.93 Table 3 gives the appropriate values of σD (in MPa) see equation (34) determined for steel A after tempering at 625°C and for steel B after tempering at 675°C (in both cases, evaluation was carried out for isothermal holding of 10 s between the end of rolling and the start of quenching). Table 3 also gives the differences between the values of Rp and σD. The yield strength were determined at a strain rate of Assuming that after tempering at high temperatures of 625 or 675°C the effect of high dislocation density, obtained after HTMT, is almost negligible, it is possible to plot the relationship between the value of (RpσD) and the size of the surface of the austenite grains (S), characterising the main 'strength' parameter of bainite or lath martensite, given by the size of the lath packets.90,91

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Fig. 74 Dependence of the value (Rp 2σD) on the austenite grain surface S (mm2/mm3). Table 3 Evaluation of the effect of dispersion hardening Steel

CT

HTMT

CT

HTMT

Plate thickness [mm]

σD [MPa]

σD [MPa]

(RpσD) [MPa]

(RpσD) [MPa]

A

230

280

255

265

11

B

321

395

335

385

11

A

215

260

240

250

18

B

320

380

325

370

18

Figure 74 shows that there is a linear relationship between the expression (RpσD) and the austenite grain surface. This also confirms the original assumption according to which the effect of dislocation hardening in analysing the steels processes in this manner (hightemperature tempering) can be ignored. Differences in the stage of development of recrystallization of the parent matrix, with the exception of the effects of precipitates, greatly affect the austenite size and, after austenite decomposition, the dimensions of the packets of lath martensite and bainite and, consequently, the resultant strength level as indicated by the equation (33).4,94 Comparison of the mechanical and metallurgical properties of the examined steels shows that the beneficial effect of HTMT depends on

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the development of recrystallization in the work-hardened austenitic matrix. It is evident that the high dislocation density in austenite after HTMT also affects the conditions for precipitation of the carbide particles and the favourable effect of dispersion hardening. In addition, for the examined types of structural steels we determined the main technical and technological parameters for determining optimum technology of producing plates 11 and 18 mm thick using HTMT technology. The results show that there is a very close relationship between the stage of development of recrystallization of the austenitic matrix in HTMT and the obtained mechanical and metallurgy properties. 4.2 Technical and Technological Characteristics of Controlled Rolling. Controlled rolling of steels for wide technical applications is one of the most advanced methods of TMT because it makes it possible to obtain, with a minimum energy requirement, higher mechanical and metallurgical properties of structural steels. Controlled rolling is a technological process and it makes it possible, directly after finish rolling or using additional accelerated cooling to obtain considerable refining of the resultant ferritic structure, accompanied by high strength properties and higher brittle fracture resistance, without any need for additional energy-demanding heat treatment. The physical metallurgy parameters of controlled rolling have been studied from several viewpoints; for example, attention was given to the problems of development of dynamic recrystallization, evaluation of the state of the austenitic matrix prior to a subsequent phase transformation, the conditions leading to the development of 'abnormal' growth of austenite (formation of an austenitic structure with nonuniform grains), the effect of partial development of recrystallization prior to a phase transformation, or evaluation of the effect of forming under the conditions of a heterogeneous matrix formed by austenite and ferrite in rolling in the intercritical temperature range. Taking these problems into account, attention was also given to the effect of the addition of microalloying elements, especially from the viewpoint of the effect of restricting the development of recrystallization. In this research, special attention must be given to evaluating the modification of the nucleation parameters of the ferritic phase transformation in relation to various efficiency of the austenite grain boundaries and intragranular potential areas for ferrite nucleation and, in the final phase, the possibility of mutual combination of controlled rolling and accelerated cooling which is in fact a superposition technological variant of the controlled rolling process. Therefore, a model of the entire technological cycle of this treatment was developed (detailed regime of water supply and its control), thus creating suitable conditions for development of new (economical) types of steel with high

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Fig. 75 Effect of various heat treatment methods, including accelerated cooling after controlled rolling, on the microstructure of steels for wide technical applications. strength parameters and a very low carbon equivalent. These steels have excellent weldability parameters.96-98 The processes, taking place in controlled rolling and subsequent accelerated cooling can be divided into the following stages, as already carried out in the chapter concerned with the analysis of physical metallurgy of TMT:52,54,62 a) forming in the region in which the austenite matrix recrystallises, b) forming in the region in which the development of recrystallization is suppressed, c) forming is carried in a heterogeneous austeniteferrite region after partial decomposition of austenite to ferrite, d) subsequent process of accelerated cooling after controlled rolling which then has a beneficial effect on refining the ferritic structure and the resultant mechanical and metallurgy properties.57,64 Taking into account Fig.35, which gives the main three stages of controlled rolling, Fig.75 shows schematically the effects of various variants of heat treatment, including accelerated cooling, on the resultant microstructures.98 Figures 76ad show the microstructure for the case of cooling in accordance with diagrams shown in Fig.75, denoted F1F4. For example, if rolling is completed at relatively high temperatures (in the high-temperature austenitic range) and if it is followed by cooling in air, these steels usually contain a mixed microstructure of ferrite and upper bainite. If accelerated cooling after rolling is applied under the conditions of limited (suppressed) recrystallization of austenite or from the region of existence of the two-phase ferriticaustenitic microstructure, the resultant microstructure of the steel (lowcarbon steel) consists

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Fig. 76a-d Microstructures obtained after cooling methods shown in Fig.75. of fine-grained ferrite with dispersed islands of bainite or martensite. The most important contribution of this method of heat treatment is that after accelerated cooling the microstructure is finer and more uniform than after only controlled rolling (even if final rolling is carried out in the intercritical ferriteaustenite region). It should be also mentioned that after accelerated cooling the formation of pearlite bands in the microstructure is almost completely suppressed. On the other hand, after air cooling the microstructure contains a 'banded' structure. The steel processed by accelerated cooling is also less susceptible to delamination because pearlite bands, i.e. potential areas of this weakening, do not form under these conditions. Controlled rolling and accelerated cooling play a different role in modification of the resultant microstructure and it must be taken into account that these treatment processes mutually 'supplement' each other to a certain extent. The main result in the first case is the increase of the yield strength (or generally the strength properties) and toughness level (brittle fracture resistance) at subzero temperatures. This is attributed mainly to the resultant fine-grained microstructure. In the second case (accelerated cooling), the contribution to the increase of the level of the resultant mechanical and metallurgical characteristics is caused not only by refining of ferrite but also by a change in the morphology of the second phase precipitated in the ferritic matrix (islands of bainite or low-carbon martensite).63,98 Taking into account the four main stages of TMT of steels for wide technical applications, we shall now discuss in detail the effect of the individual technical and technological parameters on the modification of the resultant microstructure and/or physical metallurgy properties of structural steels of this type processed by this method. 4.2.1 Deformation in the Austenite Range Associated with Recrystallization The aim of this treatment is to ensure refining of the initial austenitic structure by repeated deformation and recrystallization processes. However, the steels produced by controlled rolling may contain in certain cases a

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'nonuniform' ferritic microstructure (with different ferrite grain sizes) which then reduces the efficiency of controlled rolling and results in degradation of the resultant notch toughness values and a reduction of the brittle fracture resistance. The non-uniform ferritic microstructure is the consequence of the formation of the initial nonuniform austenitic microstructure which is in certain cases observed in, for example, steels microalloyed with niobium. Therefore, it is important to suppress the formation of a nonuniform austenitic structure and ensure its high homogeneity and fine grain size as the initial microstructural condition prior to the subsequent phase transformation. The nonuniform structure of austenite is the result of coarsening which is supported by the development of straininduced migration of the grain boundaries at low reductions (low deformations) in the initial rolling state.99 In addition to grain coalescence, it is also important to consider the effect of these processes: a) abnormal grain growth during relatively long holding after deformation, b) effect of the initial austenitic microstructure leading to the formation of nonuniformly recrystallised austenite, c) partial recrystallization either prior to subsequent deformation (under repeated deformation cycles) or prior to the actual phase transformation.100 4.2.2 Deformation in the Region of Restricted Recrystallization The final effect of deformation in the region in which the development of recrystallization is suppressed is the increase in the number of potential sites for ferrite nucleation associated with the formation of deformation bands inside elongated (deformed) austenite grains. Figure 77 shows schematically the transformation characteristics and parameters of the microstructure of the hot-deformed steel in relation to the treatment used. The main difference between the steel produced by conventional treatment and controlled rolling is that, in the former case, ferrite nucleation takes place exclusively at the grain boundaries of initial austenite, whereas in the latter case ferrite nucleates both by the intercrystalline and intracrystalline mechanism. However, the rate of nucleation of ferrite grains is considerably higher at the boundaries of the deformed austenite grains in comparison with the boundaries of the recrystallised grains. At the same time, in the case of non-recrystallised (work-hardened) austenite grains ferrite also nucleates by the intracrystalline mechanism.101 The result of these different nucleation characteristics is the formation of a different type of ferrite microstructure. The fact that the deformation bands are equivalent by their nucleation effect to the austenite grain boundaries makes it possible to conclude that the austenite grains can be regarded as 'divided' by deformation bands to individual blocks.53 These results also indicate that the deformation bands play the same role in nucleation of ferrite as the grain boundaries because the disori-

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Fig. 77 Transformation characteristics and the parameters of the resultant microstructure after various heat treatment variants. entation between the adjacent regions of the deformation bands is relatively high so that this 'phase' boundary can be regarded as a certain type of high-angle boundary.102 However, it has been concluded that these deformation bands can act under favourable conditions as potential sites for the preferential nucleation of recrystallised austenite grains. Figure 78 shows schematically the relationships between the ferrite grain size and the effective surface of all austenite grains (including both the austenite grain boundaries and the deformation bands formed in the work-hardened matrix of austenite). The resultant ferrite grains are considerably finer in the case of the initial deformed austenitic matrix than when they form in recrystallised austenite, although the effective size of the surface of the austenite grains is the same in both cases. This result can be explained by the fact that the rate nucleation of ferrite at the austenite grain boundaries is high in the deformed austenitic matrix than in recrystallised austenite.52,56 In contrast to the recrystallised material, in the deformed material the individual ferrite grains also form in the volume of the deformed austenite grains by intracrystalline nucleation.103 The beneficial effect of controlled rolling on refining the ferrite grains is basically associated with the higher nucleation efficiency of the grain boundaries (related to the unit surface area of the austenite grain boundaries) which are areas with strongly localised deformation. The differences in the efficiency of the grain boundaries of recrystallised and workhardened austenite become clear in this sense. Consequently, the level of the

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Fig. 78 Dependence of the ferrite grain size on the effective value of the austenite grain surface (microalloy steels). nucleation intensity of ferrite in the work-hardened austenitic matrix can be arranged into the following set of the individual effects:104 1. The austenite grain boundaries are potential areas of ferrite nucleation with the edges of the austenite grains being preferred sites.51,52 Higher nucleation intensity is recorded at the grain boundaries of the deformed austenitic matrix. 2. The boundaries of the coherent twins formed in austenite grains cannot be regarded as potential nucleation sites because of the low energy of this phase boundary. In the case of intensive deformation in the temperature range characterized by suppression of recrystallization of the austenitic matrix strongly deformed areas can form in the vicinity of the coherent twins where ferrite then nucleates. 3. Deformation bands are areas with a high intensity of nucleation of ferrite grains. The intensity of nucleation of these deformation bands differs. This is probably associated with the fact that the energy of the appropriate interface greatly varies. The occurrence of deformation bands with high energy and, consequently, the formation of other potential sites for ferrite nucleation are evidently associated with a higher degree of deformation of the austenitic matrix. On the other hand, it should be taken into account that the development of recovery processes (the recovery of the dislocation substructure) in areas of the austenitic matrix between the deformation bands greatly reduces the intensity of nucleation processes of ferrite in the austenitic matrix.53,54

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4. Nucleation of ferrite at the phase boundary of undissolved carbides and nitrides is observed in a relatively large number of cases. 5. However, it is very difficult to obtain convincing material on the nucleation of ferrite at the boundaries of some subgrains formed in the austenitic substructure during matrix recovery. The formation of ferrite grains in the volume of austenite grains observed in a number of cases without the presence of particles of a second phase or deformation bands can be regarded as a manifestation of the occurrence of this nucleation process (nucleation of ferrite grains at the subgrain boundaries). The efficiency of ferrite nucleation in the work-hardened matrix of austenite depends on the final deformation level prior to the austeniteferrite phase transformation. Figure 79 shows the dependence of the ferrite grain size on the initial austenite size, as determined for a steel microalloyed with a niobium addition at different final strain levels (cooling rate after forming was 0.2°C S-1). As indicated by the diagram, at a given austenite grain size the ferrite grains become finer with increasing final strain. These dependences again indicate that a high level of localised energy at the grain boundary, like the grain boundary itself, plays an important role in controlling the ferrite grain size. In previous analysis it was shown that ferrite nucleates preferentially at the grain boundaries and at deformation bands in the work-hardened austenite matrix. Secondary nucleation of ferrite grains is detected at the phase boundary of undissolved phases in the austenite matrix and at

Fig. 79 Effect of different austenite grains sizes and strain prior to phase transformation εt on the resultant ferrite grain size (Nb microalloyed steel, cooling rate 0.2°C/s).

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the boundaries of coherent twins in the austenite grains. Of these examples of possible potential sites, the austenite grain boundaries represent the effective areas for ferrite nucleation for all considered states of the austenite matrix. The high nucleation intensity of the grain boundary is associated with the high energy of the austenite grain boundaries. This is also linked with the requirement according to which the initial austenite grains (prior to a subsequent phase transformation) should be as small as possible and, combined with the frequent occurrence of deformation bands, it is then possible to obtain a very fine ferritic structure. The mean distance between 'effective' boundaries which include both the actual austenite grain boundaries and the deformation bands is considerably smaller in the fine-grained initial austenite than in the structure consisting of large grains formed by, for example, coalescence processes. In order to obtain a uniform fine-grained ferritic microstructure, it is necessary to produce a fine-grained and uniform (recrystallised) austenitic microstructure which is then deformed in the region of suppressed recrystallization. 4.2.3 Deformation in the AusteniteFerrite (Two-Phase) Region Deformation in the two-phase austeniteferrite region increases the strength properties, especially tensile strength. Steels which were deformed in the two-phase region have a microstructure consisting of a equiaxed (soft) ferrite grains and 'hard' (strain hardened) grains with a specific substructure. However, the nature of the so-called hard grains differs (depending on the degree of recovery) from the recrystallised grains up to the ferrite grains with a high degree of strain hardening. The behaviour of this structure is then similar to that of the structure of the two-phase steels in which the strength level depends on the volume fraction of 'hard' grains and the ratio of the hardness of the 'hard' and 'soft' grains. Figure 80 shows the relationship between the resultant strength level and the deformation of ferrite determined for a steel microalloyed with niobium after deformation in the two-phase (austeniteferrite) region. As indicated by the graph, the resultant strength linearly increases with increasing deformation of the ferrite matrix. It is difficult to determine the increase of the yield strength by means of a simple relationship because this value is affected by a number of factors, such as the volume fraction or 'hard' ferrite grains, the ratio of the hardness of the 'soft' and 'hard' ferrite, the density of mobile dislocations, etc. Deformation in the two-phase region affects the resultant toughness level in two ways: one is the formation of a texture and the other is represented by the resultant changes of the microstructure. It is well known that hot forming in the twophase region is followed by the formation of a crystallographic texture which becomes more distinctive with

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increasing deformation and with an increase of the volume fraction of deformed ferrite. Two types of texture form: in the direction α which is parallel to the rolling direction, and in the direction α which is normal to the rolling direction. The texture of the first type causes the formation of anisotropy of the mechanical and metallurgical characteristics, whereas the second type of texture contributes to 'weakening' the material (increasing the susceptibility to brittle fracture) in the thickness of rolled product. This then increases the susceptibility to the so-called delamination. Delamination is reflected in a reduction of the transition temperature of notch toughness specimens taken in the longitudinal direction or in transverse direction, but it also reduces the toughness level in the direction of material thickness. It may therefore be concluded that delamination has a beneficial effect on the resultant notch and fracture toughness determined in both the longitudinal and transverse direction in tests at reduced temperatures. On the other hand, it has a highly detrimental affect resulting in a very large reduction of toughness when the mechanical properties are evaluated in the direction of thickness in the given material. As already mentioned, in deformation in the heterogeneous austenitic-ferritic region deformed ferrite transfers to a recrystallised, fine-grained or recovered ferritic structure or remains as work-hardened ferrite, depending on the degree of development of recovery processes, in the parent matrix. Toughness at subzero temperatures increases with the formation of a recrystallised and fine-grained ferritic structure. A toughness

Fig. 80 Relationship between the strength and strain of ferrite (steel microalloyed with Nb addition, deformed in the intercritical region).

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remains almost unchanged when some subgrains form in ferrite as a result of its recovery. However, a large reduction of the brittle fracture resistance is observed when non-recovered ferrite grains are found in the matrix. If we summarise the effect of the individual factors affecting the resultant notch toughness level, we can predict the level of transition temperature Tt in the longitudinal or transverse direction using the following equation:104

where d is the mean ferrite grain diameter, Ttext is the contribution to the transition temperature resulting from the effect of formation of a texture, ns is the number of areas of delamination determined in the thickness of evaluated material, σp is the precipitation hardening, σdh is the dislocational hardening, σsub represents the substructural hardening resulting from the contribution of subgrain formation; β, γ, k1, k2 and k3 are appropriate constants. In the steel formed in the two-phase region, the transition temperature is strongly affected by the fourth term on the right-hand side of the equation (36), γ ns, and by the last term of this equation. If the contribution to the transition temperature due to the effect of the last term in the equation (36), namely dislocational hardening σdh, is stronger than the effect of the delamination (γ ns), the transition temperature increases. On the other hand, if the contribution of delamination is greater than the effect of dislocational hardening, transition temperature Tt decreases. 4.3 Application of Various Variants of Accelerated Cooling after Controlled Rolling In the previous section it was shown that the ferritic phase transformation also nucleates by the intracrystalline mechanism at deformation bands and that the efficiency of ferrite nucleation at the grain boundaries of deformed (unrecrystallized) austenite is higher and that in the case of intercrystalline nucleation of ferrite in the recrystallized matrix.54 However, at the same time, it must be taken into account that in the case of the nucleation effect of the deformation bands on the formation of ferrite there may be cases in which there is not a single unambiguously defined structural parameter because in addition to the actual strong nucleation effect of 'actual' deformation bands a certain effect may also be exerted by the coherent twin boundaries. Extension of the conditions of ferrite nucleation from intercrystalline to intracrystalline and an increase of

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Fig. 81 Effect of deformation of austenite in the region of suppressed recrystallization on the position of temperature Ar3 (steel microalloyed with Nb). the nucleation capacity of the grain boundaries of deformed austenite increases the intensity of the ferritic phase transformation. Figure 81 shows the effect of deformation of austenite on temperature Ar3 for three types of niobiummicroalloyed steels. The increase of the nucleation capacity for the formation of ferrite grains has been utilized in developing new steel types. Figure 82 shows schematically the main concept of utilising the increase of the nucleation potential of ferrite. If plastic deformation of austenite takes place in the region of suppression of recrystallization, the subsequent phase transformation is also accompanied by the intracrystalline nucleation of ferrite grains leading to its more intense development and also to higher stability of untransformed austenite. This effect leads to the formation of islands of martensite or bainite in subsequent cooling. Therefore, from this viewpoint, it may be concluded that the intracrystalline nucleation of ferrite leads to two opposing effects: a) to a reduction of the ferritic hardenability in connection with a higher density of potential nucleation sites for ferrite grains; b) to an increase of the 'hardenability' of the matrix because of the higher stability of austenite.

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Fig. 82 Diagram of three types of phase transformation taking place in steels with different hardenability. Depending on the chemical composition and the cooling rate, three types of phase transformation and the associated formation of three types of microstructures take place during the subsequent phase transformation after forming. 1. The ferritic phase transformation takes place in the non-hardenable steel or at the conventional cooling rates; this results in the formation of a fine-grained ferritic structure. 2. If the steel with medium hardenability is cooled at a slightly higher rate, two types of phase transformation take place: a diffusion phase transformation and a phase transformation of the shear nature leading to the formation of acicular ferrite. 3. In hardenable steels, or when using accelerated cooling, two separate phase transformation processes take place: ferritic and martensitic (or bainitic) with the formation of a two-phase structure. The acicular ferritic structure usually forms during continuous cooling and both the diffusion and shear mechanism of the ferritic phase transformation, taking place at a slightly higher temperature than that corre-

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sponding to the formation of upper bainite, operate at a specific mutual 'ratio'. In fact, a mixed phase transformation takes place under these conditions, and the formation of a ferritic network at the austenite grain boundaries is suppressed. The resultant fine-grained ferrite, fine cementite particles and martensite islands are scattered in the parent acicular ferritic microstructure. Taking into account the fact that during the formation of this type of microstructure the original austenite grain boundaries are eliminated, the resultant facets of cleavage fracture are finer than the initial austenite grains and, consequently, the transition temperature is lower. Consequently, the ferriticmartensitic (two-phase) steels are the result of the effect of separate phase transformations: formation of ferrite and stabilisation of untransformed austenite in its initial transformation stage. In a subsequent stage, the untransformed austenite transforms to martensite as a secondary process. The development of these 'separate' phase transformations is supported by a suitable combination of a) chemical composition, b) controlled rolling, c) conditions of accelerated cooling.104 In the following section, we shall analyse in detail the conditions for the formation of a two-phase steel when applying controlled rolling and accelerated cooling. The resultant optimisation of the microstructure is based on: a) considerable refining of the ferrite grains, b) obtaining a fine dispersion of the precipitated phases, c) formation of bainite or martensite areas instead of pearlite. The microstructure formed in this manner corresponds to the two-phase microstructure in a broader sense of the word. The formation of the two-phase microstructure enables the carbon equivalent of these unconventionally processed structural steels to be reduced, including improvement of the level of the fracture toughness values at subzero temperatures. Figure 76ad98 shows examples of the effect of various methods of accelerated cooling on the resultant microstructure where the method of accelerated cooling is characterized by rapid cooling in a relatively narrow range of transformation temperatures followed by air cooling. Controlled rolling forms the main technical and technological condition for the efficient utilisation of subsequent accelerated cooling in order to obtain a fine-grained ferritic structure. Controlled rolling and subsequent accelerated cooling lead to the formation of a finer ferritic microstructure than only controlled rolling. Since precipitation takes place at reduced temperatures, the precipitated secondary particles are finer and their volume fraction is smaller. It is important to note that, after accelerated cooling, the formation of pearlite bands is almost completely suppressed, as already discussed previously in the section concerned with the restriction of the development of delamination of the matrix.

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Figure 83 shows the effect of the finish rolling temperature on the resultant mechanical properties of a steel microalloyed with an addition of niobium, control-rolled and then subjected to accelerated cooling. The graph shows that: a) the finish rolling temperature must be close to temperature Ar3 to obtain a low transition temperature, b) conventional hot rolling conditions (characterized by a relatively high finish rolling temperature) lead, even in subsequent accelerated cooling, to a relatively high transition temperature, c) accelerated cooling does increase the mechanical properties but its effect on the resultant level of yield stress and transition temperature is minimum, d) it can be concluded that the combination of controlled rolling and accelerated cooling increases the strength properties without 'impairing' the toughness level.52,61

Fig. 83 The effect of finish rolling temperature on the mechanical properties of steel microalloyed with Nb addition (cooling in air and accelerated cooling).

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Fig. 84 Relationship between the increase of strength and the volume fracture of the secondary phase (martensite or bainite). The contribution of controlled rolling and accelerated cooling can be summarised as follows: controlled rolling is used to ensure the formation of a fine-grained ferritic microstructure and, consequently, increase the yield strength and reduce transition temperature. The aim of subsequent accelerated cooling is mainly to improve the morphology of secondary precipitated phases thus increasing the strength. From this viewpoint, the technical and technological combination of controlled rolling and accelerated cooling can be regarded as a highly suitable solution. Figure 84 shows the relationship between the increase of the strength values obtained after accelerated cooling, in relation to the strength determined after air cooling, and the volume fraction of the secondary phase (fraction of martensite or bainite in the parent matrix). The observed linear dependence is steeper when evaluating the microstructure formed by ferrite and martensite than in the case of the microstructure consisting of ferrite and bainite. The microstructure formed by three components: ferrite, bainite and martensite, is characterized by an intermediate gradient positioned between the two previously mentioned types of mixed microstructures. This leads to a logical conclusion according to which the increase of the strength (hardness) of the particles of the secondary coexisting phase (martensite or bainite) increases the overall level of tensile strength.52 The relationship between the transition temperature and the ferrite grain size (value d-1/2) for SiMn and low-carbon MnNb steels is shown

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in Fig.85 (evaluation was carried out for 50% ductile fracture on the fracture surfaces of Charpy V-notch toughness specimens). The proposed dependence shows that: a) the transition temperature is considerably lower in the steel with the ferriticpearlitic structure than in the two-phase ferriticbainitic or ferriticmartensitic steel, b) the difference in the transition temperature between both structures decreases with refining of the ferrite grains and is almost zero when the ferrite grains are smaller than 5 µm, c) low-carbon bainite or martensite have more favourable properties than bainite and martensite with a high carbon content. The concept of production of a fine-grained two-phase steel is shown in Fig.86. Controlled rolling causes separation of phase transformations which is further supported by a niobium addition. Subsequent accelerated cooling leads to a fine dispersion of bainite or martensite into suppressing the start of the formation of a pearlitic banded structure. The increase of strength after accelerated cooling makes it possible to reduce the carbon equivalent of the steel. A reduction of this equivalent improves the toughness of reduced temperatures and weldability parameters. The relationship between the preheat temperature required for preventing cracking in welding and the carbon equivalent is shown in Fig.87. The preheat temperature is considerably lower in subsequent accelerated cooling of the steel after controlled rolling than in only control-rolled or normalised steel.98 Different conditions in comparison with processing of sheets with a two-phase structure occur in the production of strips with a two-phase

Fig. 85 Relationship between the ferrite grain size and 50% occurrence of brittle fracture on the fracture surface of notch toughness specimens of a steel with a dual-phase structure.

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Fig. 86 Main concepts of processing dual-phase steel using thermomechanical treatment (structural steel with higher strength).

Fig. 87 The effect of the carbon equivalent on selection of the preheat temperature in welding of high-strength structural steels (normalizing, controlled rolling and accelerated cooling). structure. The optimised conditions of cooling for producing strips with this type of microstructure in low-carbon steel (0.05% C, 1.5% Mn, 1% Si and 1% Cr) after finish rolling at temperatures around 800°C are shown

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schematically in Fig.88. The control-rolled material is cooled to the optimum temperature which leads, together with the selected holding time at this temperature, to the formation of a sufficient amount of ferrite and this is followed by rapid cooling to coiling temperature.104 As a result of this sequence of the processes of controlled rolling and subsequent accelerated cooling in coiling of strips, the austenite is enriched with carbon and nitrogen (from ferrite) which then increases the stability of austenite which did not transform to ferrite. This cooling methods corresponds to the dependence denoted by S in Fig.88. The most suitable temperature at the start of the phase transformation at which the maximum amount of ferrite forms corresponds, according to experimental results, to a temperature of approximately 700°C. The finish rolling temperature also plays an important role in optimising the microstructure characteristics of this type of product. If the final deformation temperature is high, the effect of strain hardening is partially eliminated prior to the start of the subsequent phase transformation. This results in 'the loss' of part of the driving force (energy) for the ferritic phase transformation. The cooling methods denoted A and B (Fig.88) do not correspond to the favourable conditions for obtaining the two-phase structure. However, in this connection,

Fig. 88 Schematic representation of optimised cooling (denoted S) for obtaining the dual-phase structure in rolling strip steel (heat treatments A, B does not lead to the formation of a dual-phase structure). it should be mentioned that the optimised processing conditions shown in Fig.80 (processing method S) are not rigidly determined but depends to a large degree on the chemical composition of the steel subjected to this mechanical treatment and on the selected production - technical or technological methods.

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4.4 A Model of Predicting Changes in the Microstructure in Controlled Rolling and Accelerated Cooling of Structural Steels for Wider Technical Application To construct a model of predicting changes in the microstructure taking place during controlled rolling and accelerated cooling, it is essential to pay attention to systematic investigation of changes in the microstructure taking place in the individual stages of the entire technological cycle from heating the initial material up to the stage of controlled rolling and accelerated cooling.94,104 If models of predicting the changes in the microstructure taking place in the individual stages are known, it is then possible to control both the hot forming conditions and the level of the resultant mechanical and metallurgical properties. As shown by previous experience, control and prediction of the microstructure and the resultant mechanical properties of hot-formed steels will constitute a main problem in the further development of unconventional heat treatment methods. In this connection, the authors of, for example, Refs.31 and 105 proposed a model of predicting the microstructure and the size of recrystallised austenite grains in rolling. This model takes into account both the development of recrystallization processes and subsequent austenite grain growth. Saito et al106 summarised the results obtained in this area into the following main physical metallurgy principles operating in the examined types of conventional TMT of steels for wider technical application. They proposed a model of predicting the processes taking place in this treatment: a) strain induced precipitation, for example, niobium carbide or carbonitride, b) localisation of plastic deformation during forming of austenite at reduced temperatures, c) retention of the effect of strain hardening in the ferritic matrix, d) producing fine ferrite grains in the relation to the austenite grain size obtained after austenite forming, e) retaining the effect of plastic deformation in the matrix prior to phase transformation, f) selection of the suitable cooling rate. It was also shown that changes in the microstructure, taking place in hot forming, are also reflected in the achieved level of the deformation characteristics. However, to improve the efficiency of prediction, it will be necessary obtain more accurate physical metallurgy data on the relationships between different forming parameters, the resultant microstructures and the mechanical-metallurgy properties of these types of steel. In this connection, it should be mentioned that in contrast to controlled rolling which can be applied with minimum investment in the existing rolling trains,107 the introduction of the method of accelerated cooling requires construction of a special system of 'controlled' (accelerated) cooling which takes into account the requirement on processing a

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wide range of thicknesses, for example, processing sheets with higher mechanical properties (Fig.89).104 The figure shows the diagram of complex technological 'facilities' of the cooling system enabling various final treatment to be selected, both accelerated cooling and direct quenching or air cooling. This system is being installed at Kawasaki Steel Corporation98 and represents one of the optimum variants of the so-called cooling 'in line' in hot rolling. 4.4.1 Perspective Technical and Technological Solutions When developing this type of processing of structural steels for wider technical application, it is possible to develop various technical and technological variants which have also been applied in production or should be used widely.52,98,108,109 They can be included in the following groups: 1. Interrupted accelerated cooling in which the cooling regime is interrupted at a temperature above temperature Ms (cooling rate is usually selected in the range 520°C s-1); 2. Continuous accelerated cooling, in which cooling is not interrupted and is often extended until normal temperature is reached. In this case, the cooling rate is the same as at 1, i.e. between 5 and 20°C s-1; 3. Interrupted direct quenching, in which the steel is cooled to a temperature situated only slightly below temperature Ms or Bs. In this case, the cooling rate is between 10 and 100°C s-1; 4. Direct quenching, in which cooling is not interrupted (at a high cooling rate reaching up to 100°C s-1) until normal temperature is reached. Accelerated cooling does not require additional tempering and, consequently, widens the possibilities of technical application of thick plates or other types of rolled products in the as-rolled condition. However, when using direct quenching, it may be necessary to carry out tempering. This method of heat treatment is used for quenched and tempered structural steels. Another rule is that the cooling rate in accelerated cooling is lower than in direct quenching. Both these cases of final heat treatment cannot be separated only on the basis of the cooling rate used. Direct quenching requires obtaining the maximum cooling capacity in the given quenching medium in all plate thickness produced, although the cooling rate depends strongly on plate thickness. On the other hand, accelerated cooling as a variant of the 'controlled cooling' process is aimed at obtaining and/or maintaining the constant cooling rate of plates after completed forming over a relatively wide range of plate thickness. Here we pay attention mainly to unconventional heat treatment of plates because this processing method is used mainly in this area, for example in comparison with processing other types of rolled products (bars, sections, tubes, etc.). A general rule is that the cooling in accelerated cooling should be similar to the cooling rate obtained in the thickest plates in direct

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Fig. 89 Diagram of multi-purpose equipment for applying accelerated cooling by processing plates.98

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Fig. 90 Dependence of cooling rate on plate thickness in accelerated cooling and direct quenching. quenching. Changes in the resultant cooling rate with a change of the plate thickness (see equipment for accelerated cooling in Fig.89) result from the dependence shown in Fig.90.98 As indicated by this Figure, a cooling rate of approximately 10°C s-1 in accelerated cooling corresponds on average to the quenching rate obtained in direct quenching plates approximately 50 mm thick.98 In interrupted accelerated cooling, the actual cooling process starts almost immediately after completed rolling and accelerated cooling is usually arrested at a temperature above temperature Ms and is then followed by a cooling in air. The temperature at which accelerated cooling is arrested, like the cooling rate, has a strong effect on the resultant microstructure and mechanical-metallurgy properties of produced thick plates (thick plates produced by application of TMT). In interrupted direct quenching108 cooling is carried out down to temperatures below Ms or Bs, to obtain a complete or mixed martensiticbainitic structure. When using continuous accelerated cooling or direct quenching,104 it is necessary to ensure rapid cooling until normal temperature is obtained. 4.5 Effect of Controlled Rolling and Accelerated Cooling on Reducing the Susceptibility to Hydrogen Embrittlement Since steels used for many technical applications are often used under the superimposition effect of hydrogen, it is very important to examine the effect of controlled rolling and accelerated cooling on the development of this type of embrittlement.110 It is has already been mentioned that the steels for wide technical applications, processed by this method,

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show low susceptibility to delamination and formation of localised pearlite bands in comparison with the properties of steels processed by controlled rolling only.52,104 The susceptibility to hydrogen development was examined in three types of low-carbon steels microalloyed with Nb, Nb+V and Nb+B whose carbon equivalent was on a comparable level (on average 0.30). In all cases, the steel was produced using CaSi in order to modify sulphide inclusions, although the effect in the individual cases differed, as indicated by comparing the values of the effective sulphide shape controlling parameter: (ESSP = [Ca] (1124[0])/1.25[S]). In the steels microalloyed with Nb and Nb+B this value was 2.4 or 1.15, in the steel microalloyed with Nb+V it was only 0.74, which means that the occurrence of elongated sulphide inclusions of the MnS type was not completely eliminated. The minimum value of the ESSP, ensuring elimination of elongated sulphide inclusions, is 1.0. The experimental material was finish-rolled at temperatures no more than 50°C above Ar3 and the total deformation at temperatures below 950°C was on average 7580%. Accelerated cooling of 16 mm thick plates was carried out at a rate of 3035°C s-1.111 Attention was given mainly to determining the effect of different temperatures of arresting accelerated cooling on the hydrogen embrittlement resistance, including the evaluation of the relationship of this resistance and the microstructure characteristics obtained in the selected types of microalloyed steels.112,113 In addition to evaluating the main mechanical properties Rp, Rm, DWTT and KCV-Charpy V-notch impact toughness specimens (40°C) for the examined steels, the susceptibility to hydrogen embrittlement (expressed as hydrogen induced cracking, HIC) was determined in a test solution of an H2S-saturated solution containing 5% NaCl and 0.5% acetic acid with pH 3.03.5 on the basis of the extent of the failure area in the specimen (using the ultrasonic method developed by NACE and expressing the crack area ratio CAR).114 Figures 9193 show the resultant values in relation to the temperature at which cooling was arrested and, for comparison, for the case of cooling in air.111 As indicated by Fig.91, in the steel microalloyed with niobium the susceptibility to hydrogen embrittlement is completely suppressed in a relatively wide range of the cooling arrest temperature. In cooling in air from finish rolling temperature, the susceptibility to hydrogen embrittlement is very high. In the steel microalloyed with Nb+V addition it was not possible to ensure complete suppression of embrittlement, although for cooling arrest temperatures in the range 400500°C susceptibility is greatly reduced (Fig.92). In comparison with the case shown in Fig.91, in this case after air cooling the susceptibility to hydrogen embrittlement is considerably higher (~1.5 times). It is evident that, in this type of steel, the segregated areas show the presence of a large number of the so-called harder phases (due to a higher

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Fig. 91 Effect of the temperature at which accelerated cooling is stopped on the susceptibility to hydrogen embrittlement and mechanical properties. Steel microalloyed with 0.04%Nb. manganese content) which precipitate in the ferritic matrix and, consequently, increase susceptibility.111 In the third examined case, Fig.93, the susceptibility to hydrogen embrittlement is completely suppressed only in a narrow range of the cooling arrest temperature (around 490°C). This is probably caused by the less favourable form of sulphide inclusions, as indicated by the low value of the ESSP. Another important parameter which affects the susceptibility to hydrogen embrittlement is the selection of the suitable initial temperature of accelerated cooling. As indicated by Fig.94, the susceptibility to hydrogen embrittlement is suppressed if the initial cooling temperature is higher than temperature Ar3. If cooling starts at temperature below Ar3, the occurrence of bands of 'hard' microstructural components in the ferritic

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Fig. 92 Effect of temperature at which accelerated cooling is stopped on the susceptibility of hydrogen embrittlement and mechanical properties. Steel microalloyed with 0.04% Nb and 0.09%V. matrix is not suppressed. Therefore, it may be concluded that the susceptibility to hydrogen embrittlement is suppressed after accelerated cooling. The optimum temperature for completing accelerated cooling of the examined types of steel is the range between 400500°C at cooling rates higher than 15°C s-1 and at the start of cooling at temperatures slightly above temperature Ar3.

The microhardness of the evaluated steel types can also be regarded as a supplementary criterion of the resistance to hydrogen embrittlement. As indicated by Fig.95, the critical microhardness value is 250 Hm. The susceptibility to hydrogen embrittlement of the examined types of microalloyed steels is suppressed if Hm does not exceed 250 which approxi-

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Fig. 93 Effect of temperature at which accelerated cooling is stopped on the hydrogen embrittlement susceptibility and the resultant mechanical properties. Steel microallyed with 0.04%Nb and 0.001%B. mately corresponds to 22 HRc. This therefore shows that the susceptibility to hydrogen embrittlement increases with increasing strength or hardness of the matrix.111 As mentioned previously, with accelerated cooling the formation of a banded microstructure and the occurrence of hard phases (martensite or bainite) in these bands are also suppressed. This is also evidently associated with the higher resistance to hydrogen embrittlement. However, if in cooling narrow bands of bainite or martensite 'decorate' elongated austenite grains, then after accelerated cooling the occurrence of these phases in the grain boundary region is discontinuous and their total volume fraction is greatly reduced there. The data are summarised

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Fig. 94 Effect of the difference between the temperature of the start of accelerated cooling and temperature Ar3 on the susceptibility to hydrogen embrittlement. Steel microallyed with an addition of 0.04%Nb. in Fig.96 (type Nb+B steel) in relation to the cooling arrest temperature. The most suitable conditions were recorded when cooling was arrested at a temperature of ~500°C. Comparison of Fig.97a, b shows that in accelerated cooling carbon is distributed more uniformly and a difference between the regions without detected segregation and with extensive segregation is minimum (Fig.97a). In air cooling, the concentration differences are highly significant (Fig.97b). In this case, there is a marked change in the distribution of carbon caused by its diffusion to regions in which the phase transformation of austenite takes place at a later stage (in this case, this leads to a localized increase of the carbon content). However, favourable conditions exist only if the temperature of the start of accelerated cooling is above temperature Ar3 (Fig.94). If the temperature of the start of accelerated cooling is reduced below temperature Ar3, the resistance to hydrogen embrittlement does not increase. In this case, as in final air cooling after controlled rolling, the large change of the distribution of the carbon content creates favourable conditions for the formation of a large amount bainite and martensite in the segregated zone. The duration of transformation of austenite to ferrite in the non-segregated and segregated region differs greatly and enables intense carbon diffusion and an increase of its content in the segregated areas characterized by a higher level of hardenability (Fig.98). The width of the region (X) from which carbon diffuses to the segre-

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Fig. 95 Effect of cooling rate in accelerated cooling on the susceptibility to hydrogen embrittlement. Steel microalloyed with an addition of 0.04%Nb.

Fig. 96 Fraction of microstructural components precipitated at the austenite grain boundaries. Steel microallyed with 0.04%Nb and 0.001%B.

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Fig. 97 Distribution of carbon content in the segregated region in the central part of plate. Steel microalloyed with an addition of 0.04%Nb and 0.008%V. gated area can be expressed by a means of the following equation:115

where is the coefficient of carbon diffusion in austenite, Cv is the volume concentration of carbon, is the is the carbon concentration in ferrite carbon concentration in austenite at the austeniteferrite phase boundary, and at the austeniteferrite phase boundary.

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Fig. 98 Difference in the duration of phase transformation of austenite to ferrite in the segregation zone and in the zone without detected segregation. Effect of the temperature at the start of phase transformation in accelerated cooling. This evaluation of the susceptibility to embrittlement is very important both from the viewpoint of basic resistance and hydrogen embrittlement in welding these steels. In addition, deoxidation with CaSi suppresses or restricts in the parent material the occurrence of elongated sulphide inclusions (usually MnS type) which contributes significantly to the development of hydrogen embrittlement89 (see the previously given values of ESSP of the examined steel).111 From this viewpoint, the proposed analysis defines preferentially the conditions of the state of the microstructure on the development of hydrogen embrittlement after accelerated cooling. Evaluation of embrittlement under the effect of hydrogen is closely linked with the evaluation of stress corrosion cracking resistance. In the steel microalloyed with the Nb+B addition (with two carbon contents, 0.035 and 0.08%) it was found that in testing in the same solution as in the previous case (susceptibility to hydrogen embrittlement) accelerated cooling greatly increases the stress corrosion cracking resistance. As indicated by Fig.99 which gives, for comparison, also the data obtained for cooling in air after controlled rolling, at a comparable time to fracture the applied stress level increases on average by 2030%.112

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Fig. 99 Effect of accelerated cooling on the resistance to stress corrosion cracking. The results obtained for controlled rolling followed by air cooling are given for comparison. 4.6 Technical and Technological Comments Regarding Application of Controlled Rolling and Accelerated Cooling As already mentioned, these methods are used on a wide scale in the final TMT of flat products. In this connection, special attention is given to selecting the suitable cooling system, for example, systems used in multipurpose equipment, with diagram as shown in Fig.89. It is mainly the need to ensure uniform cooling of the appropriate products and the required cooling rate in the given temperature range.116 Figure 100 shows the dependences of the mean cooling rate obtained in the centre of 24 mm thick plate on the flow rate of cooling water (temperature 38°C), as determined for upper and lower jets of the cooling system developed at the Kawasaki plant.117 The cooling conditions at the lower jet are indicated by the diagram in Fig.101. The high intensity in this cooling method results from mutual comparison for the conventional design of the jet and the more efficient (modified) cooling methods (Fig.100). Since the edges of the plates are cooled at a higher rate, in the previously mentioned multipurpose device these edges (faces) are protected with a special masking device.117 Currently, special attention is also given to using the so-called laminar jets, especially with the need to obtain a higher cooling intensity also for large products in both accelerated cooling and direct quenching.118 In this method of cooling from finished rolling temperature it is possible, in large products, in addition to the previously mentioned high and reliably controlled cooling intensity, to minimise the distortion of prod-

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Fig. 100 Effect of the water flow rate on the cooling rate of 24 mm thick plate.

Fig. 101 Design of the lower jet in multi-purpose equipment for TMT of plate (Kawasaki). ucts due to a suitably 'dispersed' aqueous medium.119 It is also promising to use controlled cooling or accelerated cooling in final processing of seamless tubes. In Timicroalloyed steels, controlled rolling of seamless tubes with a diameter of 168 mm and a wall thickness of 20 mm resulted in a large increase of the mechanical properties, including brittle fracture resistance in rolling in conventional equipment.120,121 Recently, accelerated cooling in the form of direct quenching has also been utilized. For example, Fig.102 shows a diagram of such a system which uses laminar jets and central jetting of the cooling medium into

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Fig. 102 Diagram of equipment for direct quenching of tubes after controlled rolling (laminar cooling). the internal space of the tube (at a speed of approximately 1013 m s-1) with simultaneous rotation of the tube to obtain a high and controlled quenching intensity with minimum distortion of seamless tubes treated in this unconventional thermomechanical method.122 As shown by the results, in rotation at a minimum rate of 60 rpm it is possible to ensure a very high final straightness of the tubes. However, at a lower rotation rate, a relatively high distortion was recorded during cooling. Final deformation (for example at 30 rpm) was considerably higher (approximately double) than when the tube at approximately 60 rpm.

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5 Conclusions. We have discussed the main physical-metallurgy principles of the processes of TMT which makes it possible to obtain considerably higher both mechanical and ductility properties as well as higher resistance to brittle fracture, stress corrosion cracking and hydrogen embrittlement. Since TMT leads mainly to high mechanical and metallurgical properties, both in high-strength martensitic steels and low-carbon or microalloyed or low-alloyed structural steels for wider technical applications, in the introduction we carried out a detailed physical metallurgy analysis of the sources of hardening of these steels. In addition to the problems of TMT of high-strength martensitic steels, we have also discussed the physical metallurgy problems of strengthening processes in controlled rolling and in applying subsequent accelerated cooling. An integral part of this study is also the discussion of the technical and technological parameters of the investigated processes of unconventional heat treatment. It is mainly a detailed description of quenching equipment (quenching press) and the possibilities of controlled rolling with accelerated cooling, as indicated by the diagram of special equipment. This systems enables not only normalising but also accelerated cooling at the selected constant rate in the given temperature range or the so-called direct uninterrupted cooling. In the section concerned with the analysis of the physical metallurgy characteristics and the nature of strength obtained after applying TMT of high-strength martensitic steels, it was shown that the suppression of the dynamic effects of the martensitic phase transformation is a significant parameter of the beneficial effect of TMT in comparison with conventional heat treatment. Shortening of the free paths for the growth of martensite plates in work-hardened austenite or even after its recrystallization leading to refining of martensite, restricts or suppresses the formation of localised regions with high internal stresses and the possibility of initiation of microcracks in the areas of arrest of growth of martensite plates. In addition this, a 'braking' effect is also exerted by a high dislocation density in work-hardened austenite in interaction of dislocations with a growing face of the martensite plates. This effect on increase in the strength including suppression of the formation of premature fractures in tensile specimens, was observed mainly in steels with a high carbon content where the controlling process was

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the formation of plate martensite. In the region where formation of lathe martensite was the main process, the strength after applying TMT increased by approximately 200 MPa. In martensitic steels after tempering at 100°C the region of technical application after TMT was shifted on average by 0.12%C, i.e. from 0.40% to 0.500.55%C. After tempering at 200°C the strength increased by a constant amount in the entire carbon content range examined, i.e. by approximately 200 MPa. The important effect of the metallurgical characteristics has also been confirmed in this work. In the melts of the examined steel of the 2Cr2NiMnSiMo type, produced from high-purity initial materials using vacuum metallurgy, the strength increase after TMT in the entire carbon content range was approximately 200250 MPa even after tempering at 100°C. This represents an average increase of approximately 1015% in comparison with conventional treatment. These melts did not show any large reduction of the strength and ductility properties at a higher carbon content (above 0.40%) as observed in commercial melts of this type of steel in conventional treatment. The results also show that TMT in the stable austenite region is an efficient variant. The application of this technology in the region of metastable austenite (at lower deformation temperatures, around 500580°C, for example) leads to an unacceptable increase in the rolling pressures and to the risk of damage in rolling equipment without any large increase of the strength of the steel. After TMT, the martensitic steels showed an increase of fracture toughness and threshold values of KISCC' KISH and suppression of the susceptibility to delayed and premature fracture formation. The results also showed a lower probability of intercrystalline failure which is in agreement with the previously mentioned conclusions on obtaining the higher austenite grain boundary strength and/or lower probability of the initiation of microcracks at the austenite grain boundaries where the growth of martensite plates is usually arrested. In research of TMT of steels for wider technical applications, we developed a physical metallurgy model of four stages of this process, including accelerated cooling. The characteristics of the deformation process in the austenite range, associated with austenite recrystallization, the characteristics of subsequent decomposition of austenite after deformation in the region of suppression of the recrystallization process, the effect of deformation in the heterogeneous austeniteferrite region and, finally, the effect of subsequent accelerated cooling were also studied in detail. Special attention was given to examining the conditions of formation of ferrite grains in work-hardened austenite and determining the relationship between the initial austenite grain size, the degree of deformation of the austenite matrix and refining of ferrite grains characterized by the so-called transformation ratio DA/DF (the ratio of the size of austen-

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ite and ferrite grains under the conditions of isothermal transformation of austenite). The effect of subsequent accelerated cooling on the refining of the resultant ferrite grains was also defined in this sense. Other mechanical and metallurgical characteristics of the steels processed by this treatment for wider technical applications have also been determined, including the beneficial affect on the brittle fracture resistance. The results also show that the high degree of refining of ferrite increases the resistance to hydrogen embrittlement and the level of further technological parameters. The application of TMT (controlled rolling and accelerated cooling) makes it possible to reduce the temperature or eliminate preheating in welding this type of structural steel without any risk of cracking.

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Index A abnormal growth of austenite 101 accelerated cooling 115 austenite breakdown 21 B Boltzmann constant 14 Burgers vector 15 burst phenomenon 71 burst process 35 C coefficient of carbon diffusion in austenite 43 coiling temperature 118 constant of the growth rate of ferrite 42 continuous accelerated cooling 120 controlled rolling 5, 22, 101, 115 cooling 'in line' 120 Cottrell's interaction 15 crack area ratio 123 crack tip opening displacement 84 D deformation bands 25 delamination 109 delamination 31, 110 diffusion velocity 15

direct quenching 120 dislocation density 3 dislocation hardening 7 dispersion hardening 7 dynamic effects 19 dynamic recrystallization 4, 21, 101 E effective boundaries 108 effective phase boundary area 25 effective sulphide shape controlling parameter 123 effective surface of austenite grains 44 F ferrite nucleation 107 ferritic hardenability 49

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ferritic phase transformation 112 Fleischer's theory 9 free surface 69 G. grain boundary hardening 7 H HallPetch relationship 8 hard grains 108 hard precipitates 12 hardenability 49 hardening parameter 9 harder phases 123 high-temperature thermomechanical treatment 4 high-temperature TMT 56 homogeneous shear 5 hydrogen embrittlement 81 hydrogen induced cracking 123 I intercrystalline nucleation of ferrite 49 interrupted accelerated cooling 120 interrupted direct quenching 120 intracrystalline nucleation 49 J jerky flow 14 K Kawasaki Steel Corporation 120

KnappDehlinger concept 36 L laminar jets 131 lath martensite 75 localised energy 40 low-carbon lath martensite 38 low-temperature thermomechanical treatment 4, 56 Lüders bands 13 M martensitic hardenability 64 mean dislocation spacing 8 mean specific pressure on the roll 60

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mean velocity of dislocations 15 midrib 76 N nominal ferrite grain size 45 P parabolic constant of the growth rate of ferrite 43 plate martensite 62 premature failures in martensite 70 premature fracture 81 pronounced yield point 12 R. rate of ferrite nucleation 42 recrystallised austenite grains 27 rolling pressure 60 S SEN flat specimens 82 serrated flow 14 shear modulus 5 Snoek's interaction 14 soft grains 108 solid solution hardening 7 strain hardening 34 strain hardening exponent 17 stress corrosion cracking 81 stress intensity factor 84 stressstrain curve 6, 13

structure-sensitive characteristics 3 T theoretical shear stress 5 thermomechanical treatment 1 transcrystalline surface energy 69 transformation hardening 19 transformation ratio 40, 135 transformation twins 76 transition temperature 97 W work hardening 76 work-hardened ferrite 31

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