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Ametallographic study of advanced low-carbon martensitic steel with high resistance of the supercooled austenite to the transformations by pearlitic and bainitic mechanisms under continuous cooling is performed. The critical temperatures of the transformations occurring under heating and cooling are determined. The kinetics of the isothermal transformation of the austenite in the temperature range of 300 – 400°C is analyzed. The possibility of successful implementation of isothermal quenching of the steel after carburizing is demonstrated.

The authors have studied grade R260 rail steel after austenization in the temperature range from 800 to 930°C and isothermal soaking at 500, 550, and 600°C followed by quenching in water. The steel structure and abrasive wear resistance were characterized. Recommendations concerning an increase in steel wear resistance were provided.

The principal possibility of formation of a structure of lower carbide-free bainite under long-term holds is shown for steels of two alloying systems Kh3G3MFS and KhN3MFS containing 0.25 – 0.30 and 0.40 – 0.45% carbon. Dilatometric analysis of the phase transformations is performed and the bainitic range of the diagrams of decomposition of supercooled austenite is plotted. The structure is studied by the methods of light, scanning electron, and transmission microscopy.

A dilatometric study of steel 45Kh5MF used for making rolling mill rolls is performed. The results of the dilatometric and metallographic analyses are used to determine the kinetics of the perlitic transformation in the steel under continuous cooling. The developed method of numerical simulation of structural transformations under continuous cooling is used to solve the inverse problem and to plot a computational isothermal diagram of decomposition of supercooled austenite in steel 45Kh5MF in the range of perlitic transformation. To check the solution obtained an experimental isothermal diagram of decomposition of austenite in steel 45Kh5MF in the perlitic range is plotted. The data of both diagrams match accurately enough.

The main results obtained by metal scientists of the Chelyabinsk Polytechnic Institute (the South-Ural State University) under the guidance of M. M. Shteinberg in the field of the kinetics of martensitic transformations in iron alloys and its variation under the action of alloying, precipitation hardening and phase hardening are considered.

Isothermal and thermokinetic diagrams are worked out for the decomposition of supersaturated solid solution of for heat treated alloy V91 of the Al – Zn – Mg – Cu system. Curves for these diagrams correspond to the start of forming GP zones, metastable T′ and stable phases (T, η). An isothermal diagram for the alloy is plotted from the results of hardness and corrosion rate determination. Critical cooling rates are determined for alloy V91 from the diagrams obtained.

A method for plotting thermokinetic and isothermal diagrams of decomposition of supercooled austenite in powder porous steels with the help of a high-accuracy magnetometer is suggested. Thermokinetic and isothermal diagrams of decomposition of supercooled austenite in hot-formed powder steel PK40N2M are analyzed.

The effect of constant magnetic fields on isothermal martensitic transformation and structure of alloy Fe-14%Ni-4%Mn is studied. A C-diagram of the isothermal martensitic transformation and volume kinetic diagrams in a stationary magnetic field are plotted.

A physical model for determining the relative amount of phase components and the size of ferrite grains after decomposition of austenite in the process of cooling of double-phase steels is suggested. The main products of austenite transformation, i.e., polygonal ferrite, pearlite, bainite, and martensite, are considered. The driving forces of the transformation and the concentration of carbon on the phase surface are determined with the use of methods of computational thermodynamics. The model is based on equations of the classical theory of nucleation and growth. It allows for the structural features of the occurrence of γ → α transformation and contain some empirical parameters. The latter are determined using data of dilatometric measurements of the kinetics of austenite transformation and metallographic measurements of the size of ferrite grains. The model is used for predicting the kinetics of the transformation under the complex cooling conditions implemented by the VOEST-ALPINE STAHL LINZ GmbH rolling mill within the computer system for control of mechanical properties of hot-rolled strip.

A magnetometric complex for studying the decomposition of supercooled austenite is described. The kinetics of phase transformations is studied by the developed method. The kinetic parameters are determined and kinetic diagrams of phase transformations for conditions of stage hardening are plotted. A mathematical model of decomposition of supercooled austenite under conditions of isothermal and stage hardening is developed on the basis of these kinetic parameters and diagrams. An experimental check of adequacy of the models obtained is performed for the whole temperature and time range of the transformation.

A brief review of works devoted to phase transformations in titanium alloys and published by the authors during a long period and of recent publications on the topic is presented. When plotting phase diagrams of hardened titanium alloys (metastable diagrams) the authors used as a basis their relationship to the equilibrium diagrams and the transformations of the β-phase of various compositions during hardening from the β-range temperatures. After studying the structure of hardened alloys they investigated their transformations during aging, isothermal treatment, and continuous cooling from the β-range temperatures. Using the interrelation between the phase diagrams of hardened alloys and the phase transformations of β-phase under various kinds of heat treatment the authors developed theoretical diagrams of isothermal and anisothermal transformations, which were later confirmed by experimental diagrams for pilot and commercial alloys. The suggested classifications of diagrams of isothermal and anisothermal transformations are applicable to titanium alloys of various structural classes with allowance for the coefficient Kβ of stabilization of the β-phase.

The results of research in the field of martensite transformations reflect the complexity and inexhaustible nature of this many-sided process, which shows the importance and expediency of continuation of the work in this field with the aim of finding new, still unknown possibilities for varying this process and discovering new aspects of its practical use. Concluding the review, I would like to list some fundamental notions formulated in the theoretical work ofG. V. Kurdyumov andA. L. Roitburd “On the nature of martensite transformations” [92], that reflect the large role of the martensite problem in the general theory of phase transformations: “The martensite problem is a cornerstone of the general problem of the structure and properties of crystalline bodies… The formation of substructure in martensite tranformation is a result of relaxation of stresses and a decrease in the elastic energy. This occurs with “pumping” of the elastic energy into the energy of interphase boundaries… There are grounds to assume that isothermal transformation occurs with formation of a more relaxed phase… This is proved indirectly by calorimetric measurements of the stored energy, which turns out to be substantially lower in an isothermal transformation than in an adiathermal one [157]… The main features of martensite kinetics consist in manifestation of the transformation proper rather than reflection of the specific features of the prepared centers of heterogeneous nucleation,”… they are “connected with inherent features of the transformation and, are not brought from outside together with the prepared center.”… “The significance of lattice defects in martensite transformations is similar to the role of impurities in crystallization of liquids. They affect the kinetic parameters substantially but do not change the general laws. In conclusion it should be stressed that many characteristics of martensite transformations are typical of all phase transformations. This is one of the causes of the long-term and fundamental nature of the problem.”

The structure of powder phosphorus steels is investigated. New models of isothermal decomposition of super-cooled austenite are suggested which make it possible to predict the kinetics of the γ→α transformation in the pearlite and bainite temperature ranges. It is shown that mechanical alloying has a favorable effect on the formation of structure in steels. Grain disintegration accelerates the decomposition of austenite mainly due to the growth of the specific surface. At the same time, the contribution of 0.65–1.1% phosphorus additives to the process is low.

1. Formation of e-phase under isothermal conditions leads to a change in the ratio of phases and correspondingly the mechanical properties of steel G22. The latter develops particularly after thermal cycling treatment. 2. The possibility is demonstrated of repeated restoration of the capacity of austenite in steel G22 to undergo isothermal γ → e-transformation after hardening or stabilizing thermal cycling treatment by the scheme 400 ⇆ 20°C. 3. The combination of cold plastic deformation with intermediate heating which provides e → γ-transformation make it possible to obtain high total deformation which may be used in shaping operations.

1. Isothermal holding of graphitized steels, and also of CIGG, leads to a better (in comparison to other forms of heat treatment) combination of strength and ductility, and also of the impact strength. 2. The damping ability of graphitized steels after isothermal hardening is an order of magnitude less that after tempering, and is 2–3 times less than the damping ability of bainite CIGG. 3. A better combination of strength and ductivity is obtained by the graphitized steels than by CIGG after isothermal holding.

Retardation of the austenite transformation in the interval “end of the pearlite—austenite transformation to start of the ferrite—austenite transformation” in low-carbon steels, which coincides with the magnetic transformation in ferrite, is connected with the fact that the latter inhibits diffusional transport of the interstitial atoms.

The growth of pearlite in the process of isothermal transformation of austenite decreases mainly in sections between branches rich in manganese. At 560°C the growth rate of pearlite within branches and between branches is almost identical. With an increase of the manganese concentration from 1.5 to 3.6% the difference in the growth rates of pearlite in branches and between branches increases 7–64 times at 560°C, which is due to the change in the segregation coefficient from 0.82 to 0.56.

1. With increasing carbon concentrations the rate and extent of the isothermal transformation of austenite increase in Cr-Mo and Cr-Mn cast irons. 2. The temperature range in which isothermal transformation of austenite occurs becomes broader with increasing carbon concentrations, but the temperature of minimal stability of austenite decreases. 3. With increasing carbon concentrations the initial martensitic transformation temperature decreases greatly in Cr-Mo cast irons.

1. For the Cr−Ni−Co−Mo steels investigated the character of the martensitic transformation changes with decreasing MS (induced by slight changes in chemical composition) — the isothermal transformation is superposed on the athermal transformation. 2. The rate of the isothermal transformation increases with decreasing MS and evidently for this reason the susceptibility of austenite to thermal stabilization increases in the process of cooling and during repeated heating after quenching. 3. For steels with>15% retained austenite after the original quenching the cooling conditions after austenitizing must be strictly controlled (cooling with compressed air, for example), since the amount of retained austenite may be affected by stabilization and destabilization processes and also stresses that occur in the process of cooling during quenching. 4. Steels of the transition class investigated cannot be heated in the temperature range of 200–500° (welding, cutting, et al.) before the cold treatment.

1. On heating low-carbon steels containing 0.025–0.35% C at rates of 7.5–200 deg/min in the intercritical temperature range Ac1–Ac3 it is possible to isolate three basic temperature ranges: transformation of pearlite into austenite Ac3S–Ac3f, absence of transformation, and transformation of ferrite into austenite Ac3S–Ac3f. 2. The range for absence of transformations between temperatures for the finish of P→A-transformation and the start of F→A-transformation is greater as carbon content in the steel decreases. With a carbon content in the steel of 0.35% or more Ac1f and Ac3S almost coincide.

1. It was found that the structure of steels 22F1B, 27F1B, and 35F2B resulting from isothermal transformation of supercooled austenite at 680–720°C consists of evenly distributed high-strength, oriented, dispersed vanadium carbide in a ferrite matrix. Such vanadium carbides impart high strength at normal and elevated temperatures, while ferrite provides fairly high ductility. 2. The structure of the steels after the double heat treatment (normalization at 1130°C for 5 min, tempering at 720°C for 3 h) consists of ferrite and unevenly distributed vanadium carbides. The strength characteristics at normal and elevated temperatures are lower than after the isothermal treatment at 680–720°C for 30 min. 3. The original structure of the steel after quenching or after isothermal treatment has different effects on the hardness during subsequent tempering. Tempering at 400–450°C after quenching or incomplete decomposition of supercooled austenite leads to reduction of the hardness due to coalescence of cementite, while complete decomposition of austenite, with formation of a ferrite-carbide structure, leads to an increase in hardness beginning at very low temperatures.

1. Measurements of the lattice constant of M6C and MC carbides, and α- and γ-phases, and the width of lines of α- and γ-phases in the process of heating at 20–1200°C make it possible to calculate the partial coefficients of linear and volume expansion of these phases and demonstrate the relationship between the type of solid solution and the partial coefficients of expansion of M6C carbides. 2. Measurements of the lattice constants and line width of α- and γ-phases and the intensity of the lines (quantity) of M6C carbides in the process of isothermal holding for as long as 3 h at 1200–250°C after austenitizing at 1220 made it possible to substantiate the processes and the temperature-time range for such technological operations as HTTMT, LTTMT, austempering, etc.

1. Plastic deformation of austenite during HTTMT shifts the isothermal transformation curves to the left and reduces the hardenability of steel. 2. A method was proposed for plotting a calibration curve in order to determine the relationship between the hardenability of steel after HTTMT and after standard quenching. 3. A calibration curve was plotted and an empirical equation was obtained making it possible to determine the standard end-quench hardenability of alloy structural steels with a high carbon content (0.50–0.70% C) after HTTMT under the following conditions: TAust.=900–1000°, TDeform.=850–950°, e=20–30%.

1. The development of the processes of dispersion hardening of supercooled austenite leads to improved strength properties of steels 30Kh3N3MFB and 30KhNSMFB (with a slight reduction of ductility and real fracture stress) after hardening with isothermal holding in the range of subcritical temperatures and low tempering compared with the properties of these steels obtained by conventional hardening and the same tempering. 2. The uniform distribution of regions enriched with vanadium and carbon forming in the process of isothermal transformation of supercooled austenite and inherited by martensite, the possibility of their being cut by dislocations cause the uniform development of plastic deformation throughout the bulk of the steel without stress concentrations arising. Such a structural state ensures the possibility of attaining good strength properties. Then failure of steels 30Kh3N3MFB and 30KhNSMFB proceeds by the ductile micromechanism, but when the isothermal holding before quenching is extended to more than 10 min, the tendency of steel 30KhNSMFB to brittle failure increases; this may be due to the effect of silicon. 3. The good strength properties combined with satisfactory ductility and real fracture stress enable us to recommend the newly devised steels, especially 30Kh3N3MFB, as material for improtant machine parts.

1. Isothermal treatment of steel 12Kh1MF at 350–500°C leads to the formation of a structure consisting of upper and lower bainite. At 500° the structure consists primarily of upper bainite, and at 350° of lower bainite. 2. With tempering of the steel with a structure of upper and lower bainite at 730° for 3 h the dislocations undergo redistribution of the polygonization type within ferrite needles, with development of a cellular substructure. The acicular structure of the matrix is retained in this case. The density and evenness of the distribution of carbides is higher in upper bainite than in lower bainite. 3. Steel 12Kh1MF with a structure of upper bainite is more susceptible to recrystallization as compared with a structure of lower bainite, which is responsible for the higher heat resistance of the latter.

1. The pearlitic transformation of austenite is a two-stage process: the first stage is associated with the formation of the phase that controls the growth of the pearlite column — cementite; the formation of ferrite and a thickening of the cementite plates occur in the second stage. Both stages are separate in time. 2. The formation of upper bainite consisting of a cementite skeleton and ferrite envelope, the latter forming by diffusionless means from decarbonized austenite, occurs within a narrow temperature interval (500–400°C). 3. Acicular (lower) bainite forms below 400°C via segregation of disperse cementite inclusions with their subsequent envelopment by martensite. The martensite is decarbonized by means of diffusion, transforming into ferrite. 4. At temperatures below 240°C, the occurrence of the diffusionless transformation of austenite of initial composition is possible in eutectoid steel. The martensite that forms during sufficient isothermal holding may decompose in accordance with two mechanisms: with the segregation of cementite (the formation of lower bainite), and with the segregation of e-carbides inaccordance with the mechanism of spinoidal stratification (the formation of temper martensite). 5. All of the transformations that we have enumerated are reflected in diagrams of the isothermal decomposition of austenite, which have been constructed schematically for eutectoid, subeutectoid, and transeutectoid steels. A second region of pearlitic transformation, which expands with increasing carbon content in the steel, is located on these diagrams between the curves of the upper and lower bainite.

1. Ferromagnetic carbides of the cementite type are present in the steel after quenching (without tempering) and after tempering at temperatures below 620° for 1 h. Carbides of the cementite type have the shape of thin platelets with a length up to 100 A (tempering at 200°), up to 1000 A (tempering at 350°), and up to 2000 A (tempering at 500° and higher). 2. MC secondary carbides (VC) are precipitated in the process of tempering at temperatures ≥540°. The electrochemical properties of these dispersed carbides and MC carbides differ greatly after quenching or annealing. The properties become similar with increasing tempering temperatures. 3. M2C secondary carbides and chromium carbides are precipitated during tempering at 560–580° and higher, and M6C carbides with tempering above 640–650°. No preferential precipitation of special carbides on cementite particles or the boundaries of martensite crystals, twins, or other lattice defects was observed. 4. Precipitation of special carbides is accompanied by a reduction of the austenite and martensite lattice constant. Formation of M6C carbides is observed after almost complete isothermal transformation of retained austenite and reduction of the lattice constant of α phase to a value close to that obtained after annealing.

1. From the temperature dependence of Ms of isothermal and athermal components of the martensitic transformation in relation to the carbon and nickel concentrations a relationship was established between the isothermal martensitic and isothermal bainitic transformations respectively in Fe−Ni alloys and nickel steels and also between the athermal and isothermal martensitic transformations. 2. Steels with 10% Ni treated to isothermal temper martensite and isothermal bainitic temper martensite have a higher impact strength at 20 and −196° than steels with the same strength treated to athermal temper martensite.

1. Ultrasonic vibrations applied to the sample during isothermal transformation have different effects on the phase transformations in the pearlitic and bainitic regions. 2. Ultrasonic vibration has opposite effects on steels with and without nickel. 3. In steels 30KhGSNA and 30KhGSA ultrasonic vibration induces more complete transformation of austenite in the bainitic region.

1. We confirmed the data published in [2] concerning the increase in the temperature of the secondary transformation of austenite under the effect of cobalt. 2. In the steel investigated the cobalt decreases the stability of residual austenite to such an extent that after quenching from 1240°C and tempering 45 min at 600°C the bainitic transformation occurs at 340–290°C during continuous slow cooling. 3. The increase of the quenching temperature to 1260°C changes the transformation rate during subsequent tempering: the temperature of the intermediate transformation of austenite increases and the martensitic transformation begins at lower temperatures (140–150°C). Thus, the decrease of the temperature of martensitic transformation is not equivalent to the increase of the stability of residual austenite in steels containing high concentrations of cobalt. 4. An increase of the tempering temperature to 600–650°C leads to a more complete transformation of residual austenite in the intermediate region. 5. We confirmed the results obtained in [9] concerning the intensification of the transformation of austenite as the result of rapid cooling after tempering: rapid cooling partially or completely eliminates the intermediate transformation of austenite or increases the effect of the martensitic transformation. 6. The martensitic transformation occurring during tempering of the cobalt steel investigated results from the preceding bainitic transformation. The rate of the secondary martensitic transformation can be changed by changing the bainitic transformation.

1. The differences in the properties of different batches of 40Kh steel are due to the presence of surface-ac-tive SiO. The concentration of SiO in the metal can be predicted from the physicochemical conditions of the smelting process. 2. The presence of surface-active impurities considerably slows down the rate of isothermal transformation of supercooled austenite, which leads to an increase of the depth of the effect of heat treatment. 3. In the presence of a surface-active impurity in the steel the martensite and other products of isothermal decomposition of supercooled austenite become oriented. Furthermore, the intensity of the process of coagulation of the carbide phase decreases when the metal is kept for long periods of time near the critical temperature. 4. The effect of surface-active impurities is particuarly great in the α-phase and during the γ»α transformation at certain temperatures. The intensity of the effect is determined by the ratio between the amount of the impurity and the rate of its displacement, which is characterized by a curve with a peak between 450 and 650°C.

1. We plotted isothermal transformation diagrams of an Al−Zn−Mg−Mn−Zr alloy, which show that decomposition of the solid solution is more rapid in unrecrystallized than in recrystallized plates of the alloy. Transmission electron microscopic examination indicates that the high rate of decomposition of the supercooled solid solution in unrecrystallized plates is due to the well-developed network of subgrain boundaries, promoting nucleation and growth of secondary phase during aging. 2. Unrecrystallized plates age more rapidly than recrystallized plates at temperatures of 100–200o. On aging at temperatures below ∼140o the strengthening effect is higher in unrecrystallized than in recrystallized plates, while at higher temperatures it is lower. 3. The strength characteristics of artificially aged plates with an unrecrystallized structure are more sensitive to the time interval between quenching and artificial aging than in the case of phates with a recrystallized structure.

The investigations carried out show that isothermal formation of the e-phase can be controlled not only by alloying, but also by means of repeated phase transitions. Reduction of the temperature at which the formation of the e-phase commences intensifies the isothermal transformation near the martensite point and at room temperature. The latter is valid if the γ→e-transformation during continuous cooling begins not lower than the temperature at which the isothermal soaking is carried out, otherwise the intensity of the isothermal process decreases, and the lower the commencement of the γ→e-transformation with respect to the temperature isothermal line, the more the intensity drops. It was found that if the e-phase forms at −60, −70°C, there is a complex dependence between the rate of the isothermal process and the temperature.

1. Omnilateral pressure, which reduces the martensitic point, does not alter the “explosive” character of the martensite transformation in type N31 Fe−Ni alloy. Omnilateral pressure greatly reduces the transformation rate for an isothermal martensite transformation (type N23G3 alloy) and causes a reduction in the temperature of the kinetic maximum. 2. Preliminary plastic deformation does not reduce the transformation temperature range for N31 alloy but causes a large decrease in the extent of the first and subsequent explosions. Preliminary plastic deformation of Fe−Ni−Mn alloys, which suppresses the isothermal transformation, causes almost no reduction in the extent of the athermal transformation in a magnetic field. Tempering of the austenite inhibits the isothermal transformation but has no effect on the athermal transformation excited by pulsed fields.

1. On the basis of the diagram of the isothermal transformation of the β2-solution in the YuNDK35T5 alloy we have shown the possibility of quenching this alloy from two temperature ranges: 1260–1240 and 900–850°C.

From the diagrams of isothermal transformation of austenite one can determine which region of isothermal transformation characterizes the hardenability of the steel and determine the numerical value of the hardenability (the value of the critical diameter, the depth of the quenched layer in the end quenched samples of pipes, the thickness of the quenched layer in steel articles of different diameters).

1. Plastic deformation of stable austenite with limitation of the development of recrystallization, provides strengthening of the final structure in all forms of transformation-martensitic, intermediate, and pearlitic. 2. The strengthening effect is different with different forms of transformation in eutectic carbon steel. Strengthening as the result of deformation is a maximum in the martensitic transformation, with the tensile strength increasing 30% and the ductility dropping by half; a minimum in the pearlitic transformation, with an increase in ductility of 20% and in strength of 15–20%; and with an intermediate value for intermediate transformation, an increase in strength of 20% and in ductility also of 20%. 3. The maximum strengthening is obtained with an austenite deformation of 25–35% for the martensitic transformation; of 15–20% for the intermediate transformation; and of 6–8% for the pearlitic transformation. 4. As a result of changes in the fine structure of the strengthened steel the addition of a small quantity of titanium provides high ductility. 5. The desirability of using particular methods of treating eutectic carbon steel, such as high temperature thermomechanical working, high temperature thermomechanical working with isothermal transformation, and high temperature thermomechanical working with diffusion transformation, is determined by the dimensions of the starting piece, service conditions, and production conditions.

1. In the process of soaking at 1200–850°C austenites cooled from austenitizing temperatures are considerably impoverished and M6C carbides are precipitated. Therefore high-speed steels should not be subjected to HTTMT at above-critical temperatures. 2. In the upper (pearlite) transformation region carbides precipitate in austenite during the γ→α transformation and in the ferrite being formed. 3. In the region of stable austenite (600–400°C) there are no changes of any kind in the supercooled austenite which can lead to any reduction of the performance characteristics, and therefore soaking at such temperatures can be used in heat treatment and TMT. 4. In the intermediate region (550–200°C) small amounts of M6C carbides are formed only during γ→α transformation. No carbides are precipitated from austenite or ferrite, and the intermediate transformation products have a high hardness and red hardness and satisfactory structure after tempering. 5. Only carbides of the M6C type (Fe3W3C) of different compositions were found in the austenite, ferrite, and in the γ→α transformation in high-speed R18 steel during soaking at all temperatures investigated. No carbides of other types or, specifically, cementite were found.

1. For hypereutectoid steels the duration of the α→γ transformation in parts and the intensive isothermal solution time of excess carbides during heating to quenching temperature are linear functions of the dispersity of the initial structure as characterized by the hardness HRA. 2. The temperature dependence of the intensive isothermal solution time of excess carbides in austenite has an exponential character for hypereutectoid steels. 3. The data obtained make it possible to calculate the effect of the initial structure and the temperature on the transformation during heating of high-carbon steels and more precise calculation of the soaking time. 4. The monographs facilitate the technical calculations of the intensive isothermal solution time of carbides in U12A steel with different initial structures.

1. From the change in the hardness and the results of metallographic and x-ray analysis we plotted isothermal transformation diagrams for alloys of the critical composition of the VT22 type in which molybdenum and vanadium were partially or completely replaced with iron and chromium. 2. Increasing the amount of iron and chromium, with a corresponding reduction in the amount of molybdenum and vanadium, accelerates the beginning of decomposition of the β phase. 3. Alloying with over 4.2% iron and chromium lowers the upper limit at which ω phase exists (as compared with alloy VT22). 4. No intermetallic compounds were observed in any of the alloys investigated with isothermal holding for 100 h.

1. The data presented confirm the fact that the formation of "combination structures," consisting of bainite and troostite, is the result of decomposition of austenite in a relatively narrow temperature range where the pearlitic and bainitic ranges overlap. 2. The data obtained on the mechanism by which combination structures consisting of high-carbon troostite and low-carbon bainite are formed lead us to assume that the redistribution of carbon occurs only in the beginning of the isothermal transformation in regions commensurate in size with fragments of grains.

1. The kinetics of the isothermal transformation of austenite in cast nickel steel is similar to that in wrought steel. 2. We obtained additional data on the possibility of graphitizing carbides in nickel steel and plotted curves of the beginning and end of graphitization of carbides in the cast and quenched conditions. The addition of silicon and quenching accelerate graphitization.

1. Hot deformation accelerates precipitation of excess ferrite and the pearlitic transformation in steels 40 and U8. With decreasing transformation temperatures this accelerating effect of deformation gradually decreases. Hot deformation of steel U8 slows down the transformation in the bainitic region. 2. The difference in the rate of decomposition of deformed and undeformed austenite increases with the austenitizing temperature. 3. Hot deformation prevents the formation of Widmanstatten ferrite and refines the ferrite and pearlite grains; the dispersity of pearlite remains unchanged in this case.