523-529

UDC 669.14.018.298

DISPERSION OF THE STRUCTURE OF STEELS

UNDER INTENSE THERMAL EFFECT.

PART 1. CHOICE OF OPTIMUM ALLOYING SYSTEM

L. Ts. Zayats,1 D. O. Panov,1 and Yu. N. Simonov1

Translated from Metallovedenie i Termicheskaya Obrabotka Metallov, No. 11, pp. 13 19, November, 2010.

Results of a complex study of the kinetics of transformations, structure, and characteristics of mechanical

properties of low-carbon alloy steels are used to formulate principles for designing chemical composition. The

principles make it possible to lower the factor of carbon activity and to raise the resistance of the system to dif-

fusion relaxation, which promotes dispersion of the structure in heat treatment.

Key words: low-carbon steels, alloying system, diffusion relaxation, activity of carbon, harden-

ability.

INTRODUCTION

Optimum combination of mechanical properties in ge-

neral-purpose structural steels is obtained when their struc-

ture meets two conditions, i.e., (1 ) enough density of struc-

ture defects (imperfections), especially of dislocations and

boundaries and (2 ) minimum scale of structural and chemi-

cal inhomogeneity of the material.

All modern methods for strengthening metallic materi-

als, i.e., for raising the density of imperfections, are based on

providing a most nonequilibrium state in them. This state is

characterized by the value of the excess free enthalpy (amor-

phous state, mechanical stresses due to severe plastic defor-

mation, heat of phase hardening, supercooled austenite after

hardening) due to intense mechanical and thermal actions

[1]. Maximum possible level of nonequilibrium widens the

spectrum of the mechanisms of subsequent lowering (relax-

ation) of the excess energy of the system (closeness to equi-

librium) and thus widens the possibility of control of the pro-

cess of structure formation. Growth in the dislocation density

and formation of new boundaries as a result of evolution of

the dislocation structure is a result of shear relaxation. How-

ever, a more effective mechanism of relaxation without frac-

ture, which is most close to equilibrium, is diffusion. Diffu-

sion controls the processes of recrystallization (primary and

secondary), oxidation, and decarburization at elevated tem-

peratures, segregation of excess ferrite, development of

pearlitic and bainitic reactions in supercooled austenite, and

decomposition of martensite due to tempering. It is obvious

that diffusion relaxation in multicomponent systems not only

lowers the strength but also increases the scale of inhomo-

geneity of the material, i.e., causes layering of the system

with respect to the chemical composition in two-phase do-

mains.

Thus, in order to obtain an optimum structure in a struc-

tural steel, which should ensure a high level of strength char-

acteristics and reliability parameters, we should implement a

series of states in the system (stages of dispersion), namely,

(1) a state with maximum deviation from the equilibrium

condition, i.e., a state with the highest driving force of subse-

quent relaxation;

(2) a state with high dislocation density formed as a re-

sult of shear relaxation (the attained dislocation density is the

higher the greater the driving force of the relaxation);

(3) a state characterized by evolution of the dislocation

structure including controllable mechanisms of diffusion re-

laxation with minimum scale, primary recrystallization in the

first turn.

Primary recrystallization in a crystal with high density of

uniformly distributed dislocations of one sign can ensure

maximum turning of neighbor volumes of the crystal, i.e.,

can raise as much as possible the density of high-angle

boundaries and reduce the size of the characteristic structural

component [2].

In the suggested model of structure formation in a steel

states (1 ) and (2 ) can be implemented as a result of a low-

Metal Science and Heat Treatment, Vol. 52, Nos. 11 12, March, 2011 (Russian Original Nos. 11 12, November December, 2010)

523

0026-0673/11/1112-0523 2011 Springer Science + Business Media, Inc.

1Perm State Engineering University, Perm, Russia (e-mail:

[email protected]).

temperature intense mechanical action, by hardening of the

melt (amorphization), and by hardening with polymorphic

transformation. All the three states can be created succes-

sively by controlled thermomechanical treatment, heating of

an amorphous material, and rapid and short-term austeni-

zation of steels with initial martensitic structure under the

condition that the austenite inherits high dislocation density.

In heat treatment of steel the stages mentioned can be imple-

mented both in the cycle of austenization and in the cycle of

hardening cooling.

The aim of the present work consisted in formulating

principles for designing chemical composition for low-car-

bon alloy steels in order to create an optimum structure that

would provide high strength characteristics and reliability

parameters and in studying the effect of the alloying system

meeting these principles on the stability of supercooled aus-

tenite both under isothermal conditions and in continuous

cooling.

METHODS OF STUDY

We studied two low-carbon martensitic steels (LCMS),

i.e., a promising LCMS of grade 12Kh2G2NMFT with ele-

vated strength level and a sparingly alloyed grade 08Kh2FB.

For comparison, we chose a bainitic steel of grade

15Kh2GMF. The chemical compositions of the steels are

presented in Table 1.

Test pieces were cut from hot-rolled preforms in the axial

direction. The pieces were subjected to austenization at 940C

for 40 min and cooled at various rates. The average cooling

rate in the range of 650 350C was computed from the

change in the temperature, which was detected with the help

of a thermocouple calked into one of the test pieces. The ten-

sile tests were performed using an UM-10T testing machine

for short five-fold test pieces in accordance with the require-

ments of the GOST 149773 Standard. The impact tough-

ness KCU and KCT was determined in accordance with

GOST 945478 in an MK-30 impact testing machine (speci-

mens of type 3 and 17, respectively). The kinetics of the

transformations was studied in an updated Akulov aniso-

meter that permitted detection of low amounts of -phase.

The structure of the steels was studied using a Neophot-32

light microscope and a JEM-200CX electron microscope.

RESULTS AND DISCUSSION

Principles of Dispersion of Structure by Heat Treatment

V. D. Sadovskii has shown [3] that the recrystallization

due to austenization of steels with crystallographically or-

dered structure occurs in two stages, i.e., a phase transforma-

tion proper, which yields phase-hardened austenite orienta-

tion-bound to the initial structure, and recrystallization,

which causes refinement of grains, removal of phase harden-

ing, and elimination of intragrain texture. It is obvious that

the effect of grain refinement in austenization is the more

considerable the higher the dislocation density by the mo-

ment of the start of primary recrystallization of phase-hard-

ened austenite and the lower the rate of secondary recrys-

tallization, i.e., the lower the susceptibility of the metal to

diffusion relaxation.

In its turn, the size of the austenite grains affects sub-

stantially the dispersion behavior in hardening. It has been

shown in [4] that when the austenite grain diameter in steel

45KhNMFA exceeds 20 m, the size of a martensite lath is

virtually independent of the size of the austenite grains. This

is connected with the effective fragmentation of the structure

during formation of lath martensite. However, in the case

when the austenite grains are less than 10 m in size, the lath

size decreases markedly. Starting with 5 m, only one

martensite lath with a size equal to that of the austenite grain

forms in a grain due to hardening.

The authors of [5] report a nonmonotonic effect of the

size of austenite grains on the location of martensite point

Ms, which begins to decrease abruptly upon formation of

austenite grains less than 5 m in size.

The authors of [6, 7] show that it is possible to obtain a

principally new type of martensite (block martensite) that

consists of equiaxed crystals-blocks with a high density of

uniformly distributed dislocations separated by high-angle

boundaries and having a size of 100 200 nm. In order to

obtain such a structure they used thermocycling (3 7 cy-

cles), every cycle of which included induction heating to the

austenitic range and subsequent hardening, These data show

the possibility of effective dispersion of structure in a steel

both in the process of austenization and in subsequent hard-

ening for martensite.

524 L. Ts. Zayats et al.

TABLE 1. Chemical Composition of Studied Steels

Steel

Content of elements, wt.%

C Cr Mn Mo V Other

12Kh2G2NMFT 0.120 2.38 2.23 0.43 0.09 1.38 Ni; 0.03 Ti

08Kh2G2FB 0.085 2.20 2.25 0.08 0.03 Nb

15Kh2GMF 0.165 1.95 1.00 0.20 0.13

Note. In addition to the listed elements all the metals contained 0.26 0.30% Si and at least 0.02% S and 0.02% P.

Repeated cyclic dispersion of structure is also possible. It

will be effective only in the case of attainment of a higher

level of nonequilibrium in the first stage of each subsequent

cycle.

It is obvious that two tasks should be solved for succes-

sive implementation of stages of dispersion of structure,

namely, we should

(1) choose a material (alloying system) that would permit

successive formation of the required intermediate states and

minimization of the possibility of diffusion relaxation and of

risks of fracture;

(2) choose a method for implementing the intermediate

states, i.e., a specific dispersion process based on an intense

action on the system.

Study of the Effect of Alloying System on Stability

of Supercooled Austenite

Let us consider the role of alloying system in solution of

one of the most important tasks of todays physical metal-

lurgy of steels, i.e., provision of stability of austenite in har-

dening.

It cannot be doubted today that formation of a high-

strength state in iron-carbon alloys is closely related to cre-

ation of martensitic structure in them. However, it is by far

not always possible to implement a martensitic trans-

formation. In the first turn, this concerns large-size and mas-

sive preforms and articles the cooling rate of the core of

which is considerably lower than that on the surface.

The problem of hardenability is especially urgent for

steels with a low content of carbon, which is known to be a

-stabilizer. For a long period, heat treatment specialists have

assumed that low-carbon steels were not hardenable at all.

At the same time, commercial use of low-carbon steels with

elevated and high strength promises much, because the high

level of their structural strength is accompanied by a whole

set of technological advantages with respect to medium-car-

bon steels just due to the low content of carbon. These ad-

vantages include good weldability, low deformation in heat

treatment, and low susceptibility to decarburization and

crack formation. In the case of high hardenability of steels

with low carbon content the production cycle of many parts

and structures will be simplified considerably.

The problem of formation of highly stable supercooled

austenite can be solved is we manage to satisfy at least two

conditions, i.e., ensure high stability of the austenite in the

range of pearlitic transformation and suppress the bai-

nitic transformation. The temperature range of the marten-

sitic transformation should be regulated in order to minimize

the possibility of accommodation by twinning (inevitable at

low temperatures Mi M

f) and to guarantee formation of a

structure of lath martensite in the process of slow hardening

cooling and simultaneously to eliminate the undesirable pro-

cesses of tempering of the freshly obtained lath martensite,

which are inevitable at high temperatures Mi M

f.

At first, the possibility of implementation of martensitic

transformation in slow cooling of low-carbon alloyed

austenite was discovered in a study of the Fe Ni system. At

optimum alloying of iron with nickel the temperature of the

thermodynamic equilibrium decreases substantially,

and a thermodynamic stimulus for transformation

arises only in a temperature range below 450C. At these

temperatures a transformation can develop only by a

shear mechanism. The task of suppression of bainitic trans-

formation in these alloys has been solved radically, i.e.,

bainitic transformation is simply absent due to the exception-

ally low carbon content.

Another way for obtaining a martensitic structure in slow

cooling of low-carbon austenite is to raise the stability of

supercooled austenite at a specific proportion of carbon and

alloying elements. It has been shown in [8, 9] that at 0.04

0.12% C combined with additives of Cr, Mn, Ni, and Mo

(5 6% in total) the incubation period of pearlitic transfor-

mation at the temperature of the lowest stability of austenite

(650C) lasts from tens of minutes to several hours. Such al-

loys have been called low-carbon martensitic steels (LCMS).

If the carbon content in such steels exceeds a certain level, a

bainitic reaction is activated. The diagrams of isothermal de-

composition of the austenite exhibit only a range of bainitic

transformation [9].

It is suggested in [10, 11] that massive parts with elevated

strength level should be produced from steel 15Kh3G3MF

(0.10 0.20% C, 3.2 4.0% Cr, 3.2 4.0% Mn, 0.5 1.0%

Mo, 0.3 0.5% V). In this steel supercooled austenite is

characterized by high stability; in the range of 700 400C

the transformation does not occur for 4 h. In the range of

400 200C bainite appears in an amount of 5% after a hold

for 5 sec and in an amount of 55% after a hold for 700 sec.

Introduction of an additional low content of carbon (over

0.15%) into this steel results in substantial activation of the

intermediate transformation. Industrial testing of this steel

has shown that a martensite structure forms in massive parts

(with a length of at least 1800 mm and a diameter of

185 280 mm) only after oil hardening [11]. It should be

noted that the steel melted for industrial testing contained

equal amounts of molybdenum and vanadium (0.5%).

The author of [12] writes that steel 12Kh2MFB (I531)

is a modification of the Croloy 2.25 grade widely known in

foreign countries, which contains 0.08 0.12% carbon,

2.25% chromium, and 1% molybdenum. In the USSR this

steel has been used widely due to its high susceptibility to air

hardening (additional hardening, especially in the near-weld

zone, due to arc welding). Addition of 0.5 0.8% niobium

to this composition has reduced the air hardenability substan-

tially.

Steel 12Kh2MFB contains 0.08 0.12% C, 0.4 0.7% Si,

0.4 0.7% Mn, 2.1 2.6% Cr, 0.5 0.7% Mo, 0.20 0.35%

V, and 0.5 0.8% Nb, i.e., is characterized by high concen-

tration of carbide-forming elements. The minimum stability

of austenite in this steel in the range of pearlitic (normal)

transformation is about 100 sec at 700 720C; active deve-

Dispersion of the Structure of Steels under Intense Thermal Effect. Part 1 525

lopment of bainitic transformation occurs in 50 100 sec at

450 350C [8].

The object of study of [13, 14] is steel 55Kh2MFB

(0.55% C, 0.6% Mn, 1.86% Cr, 0.51% Mo, 0.26% V, 0.10%

Nb). The minimum stability of its austenite in the range of

pearlitic (normal) transformation is about 1000 sec at 650C;

bainitic transformation begins in 200 300 sec at 350

280C. Thus, the proportion of carbide-forming elements in

steel 55Kh2MFB decreases upon decrease in their suscepti-

bility to carbide formation. Despite the high carbon content

steel 55Kh2MFB possesses higher stability of supercooled

austenite than steel 12Kh2MFB with less carbon.

In the 1980s L. M. Kleiner suggested a low-carbon mar-

tensitic steel 12Kh2G2NMFT. It was discovered later that

this steel exhibits exceptionally high stability of supercooled

austenite [15 17]. A special feature of this grade of steel is

the presence of a virtually full set of carbide-forming ele-

ments in its composition. The content of the carbide-forming

elements decreases progressively upon growth in the suscep-

tibility to carbide formation.

Thus, by the 1990s researchers possessed ample data on

a favorable effect of multicomponent alloying with carbide-

forming elements on growth of stability of supercooled aus-

tenite. The empirical approach was used for obtaining steels

possessing high stability of supercooled austenite.

It becomes clear from the facts mentioned above that

systematic multicomponent alloying is a necessary condition

for high stability of supercooled low-carbon austenite in the

process of slow cooling.

Theoretical Substantiation of the Choice of Alloying

System with Elevated Resistance to Diffusion Relaxation

In heat hardenable Fe C alloys carbon is a poison and

a drug. On the one hand, its presence in the solid solution is

a necessary condition for hardening to occur, because carbon

lowers the temperature range of martensitic transformation

and thus the possibility of diffusion relaxation; it also intensi-

fies the effect of phase hardening. On the other hand, the

high thermodynamic activity of carbon (thermodynamic ac-

tivity is understood as the probability of presence of a carbon

atom in a position of crystal lattice with arbitrary coordi-

nates) provokes development of diffusion relaxation due to

decrease in the activation energy of self-diffusion of iron and

appearance of diffusion flows of point defects upon forma-

tion of carbides or segregations. For this reason, it is neces-

sary to provide an enough carbon concentration in the Fe C

system at minimum carbon activity.

The thermodynamic activity of carbon in an alloy (aC

all)

can be lowered without changing its concentration by alloy-

ing with carbide-forming elements [18] in accordance with

the expression

aC

all= (

CfC

A1fC

A2 ... f

C

Ai) N

C, (1)

where fC

A1, f

C

A2, f

C

Aiare relative factors of activity in the

Fe C A1

... Ai

system, which allow for the effect of

the alloying element (Ai) on the activity of carbon;

Cis the

factor of activity of carbon in the Fe C alloy; and NC

is the

atomic fraction of carbon in the alloy. For ternary alloys with

a carbide-forming element fC

A< 1, and hence the activity of

carbon decreases.

In its turn, the relative factor of activity is related to the

atomic concentration of the alloying element through an ex-

pression

ln fC

Ai= K

AiN

Ai, (2)

where NAi

is the atomic fraction of the alloying element in

the alloy; KAi

is a factor allowing for the atomic interaction

between carbon atoms and the alloying element in the solid

solution of the ternary system (Fe C Ai).

In fact, the activity of carbon in an alloy reflects the

probability of residence of carbon atom at a point with arbi-

trary parameters; the carbide-forming elements limit the dif-

fusion mobility of carbon as a result of the Me C binary

atomic interaction and decrease this probability. The activity

is decreased even upon the introduction of an enough content

of one strong carbide-forming element. However, this inevi-

tably results in formation of a carbide, appearance of diffu-

sion flows, and lowering of the content of carbon in the solid

solution. Therefore, the decrease in the activity of carbon due

to alloying with a carbide-forming element should not be ac-

companied by formation of a special carbide.

If the carbide-forming elements do not form isomorphic

carbides and their influence on the activity of carbon is the

same, i.e., fC

A1= f

C

A2= ... = f

C

Ai, the interaction between

these elements and carbon as a result of competition is li-

mited by formation of binary bonds in the solution, and a car-

bide does not form. It is obvious that the interaction between

the alloying elements and carbon can also be estimated in

terms of the value of the energy of formation of the corre-

sponding carbides. However, for strongly diluted solutions it

seems more expedient to use the coefficient KA

.

We arrive at a principle of formation of an alloying sys-

tem with minimum thermodynamic activity of carbon and

probability of carbide formation, i.e., the content of carbide-

forming elements should be inversely proportional to the

value of the coefficient or, in a formalized form,

NA1

: NA2

: ... : NAi

= (1KA1

) : (1KA2

) : ... : (1KAi

). (3)

Basing ourselves on the reasoning presented above we

should use a low-carbon steel (< 0.2% C) alloyed with five

six carbide-forming elements, the concentration of which

should decrease in an inverse proportion to the value of KA

for these elements, and the effective and rational total con-

centration of the alloying elements should not exceed

4 6 wt.%.

In order to check this concept we performed a complex

study including an investigation of the effect of the alloying


system on the stability of supercooled austenite both under

isothermal conditions and in continuous cooling.

A Study of Stability of Supercooled Austenite

in Low-Carbon Martensitic Steels

We studied LCMS of grades 12Kh2G2NMFT and

08Kh2G2FB and a bainitic steel 15Kh2GMF for comparison.

Analysis of the chemical composition of steel

12Kh2G2NMFT has shown that the principle of system al-

loying is obeyed in it to a maximum degree. As a conse-

quence, the study of the kinetics of the transformation

has not shown regions of pearlitic and bainitic transforma-

tions in this steel (Fig. 1a ). The study of the microstructure

and the electron microscopy showed the presence of only

lath and packet -phase in the structure. After austenization

at 940C and cooling at various rates regular segregations of

carbides and retained austenite have not been detected [15].

The strength of steel 12Kh2G2NMFT does not change

both upon growth in the length of the isothermal hold in the

range of pearlitic transformation and upon decrease in the

cooling rate from 150 to 0.0044 Ksec (Fig. 2a ). The de-

crease in the impact toughness observed for the cooling rate

of 0.0044 Ksec is a result of the occurrence of tempering of

the martensite, as it has been shown in [16].

The exceptionally high stability of supercooled austenite

in steel 12Kh2G2NMFT is explainable only by the complex

alloying with carbide-forming elements with progressively

increasing susceptibility to carbide formation, which are

taken in a specific proportion. It is important that the chain of

the carbide-formers should be continuous, i.e., the activity of

chromium and manganese (with respect to that of carbon)

should be suppressed by a lower content of molybdenum (or

tungsten), and the activity of these elements should be ba-

lanced by a still lower content of vanadium, niobium or tita-

nium.

The results of the study of steel 08Kh2G2FB illustrate

the important role played by the presence of a continuous se-


700

650

600

550

500

450

400

350

300

250

700

650

600

550

500

450

400

350

300

250

1 10 100 1000

1 10 100 1000

, sec

, sec

t, C

t, C

3%

3%

3%

3%

50%

50%

50%

95%

95%

95%

b

Fig. 1. Diagrams of isothermal decomposition of austenite for

steels 12Kh2G2NMFT (a) and 15Kh2GMF (b ) after austenization

at 930C for 40 min.

r 0.2; , P

r

r

r

0.2

0.2

0.2

vcool , K sec

KCU KCT; , J m 2

KCU

KCU

KCU

KCT

KCT

KCT

1200

1000

800

600

400

200

0

1200

1000

800

600

400

200

0

1200

1000

800

600

400

200

0

2.4

2.0

1.6

1.2

0.8

0.4

0

2.4

2.0

1.6

1.2

0.8

0.4

0

2.4

2.0

1.6

1.2

0.8

0.4

0

1000 100 10 1 0.1 0.01 0.001

1000 100 10 1 0.1 0.01 0.001

1000 100 10 1 0.1 0.01 0.001

c

b

r 0.2; , P

r 0.2; , P

vcool , K sec

vcool , K sec

KCU KCT; , J m 2

KCU KCT; , J m 2

Fig. 2. Dependence of the characteristics of mechanical properties

of steels 12Kh2G2NMFT (a), 08Kh2G2FB (b ), and 15Kh2GMF (c)

on the rate of cooling after 40-min austenization at 930C.

ries of carbide-formers. The only substantial difference of

steel 08Kh2G2FB from 12Kh2G2NMFT is the absence of

molybdenum (or tungsten). As a result, the chain of car-

bide-forming elements is broken and the joint effect of

chromium and manganese is insufficient for suppressing the

activity of vanadium. This causes decrease in the content of

carbon in the austenite due to segregation of special vana-

dium (or niobium) carbides. This is reflected by the 25

30C growth in the temperature of the start of martensitic

transformation at cooling rates less than 1 Ksec as com-

pared to the values of Midetermined under isothermal condi-

tions. It should also be noted that rapid transfer of a specimen

heated to the austenization temperature into a furnace heated

a temperature 10 15C higher than Mi

isoth, which makes it

possible to suppress the formation of special carbides, does

not shift Mi. Segregation of special carbides results in

destabilization of the austenite. For this reason, we observed

lower bainite in the structure of steel 08Kh2G2FB after

cooling at a rate of 0.15 Ksec; at a still lower rate of

0.0044 Ksec we observed upper bainite.The strength of steel 08Kh2G2FB decreases after cool-

ing at a rate of 0.035 Ksec, and the impact toughness de-

creases after cooling at a rate of 0.15 Ksec. The dynamic

crack resistance turns out to be the most sensitive to embrit-

tlement due to the appearance of products of a nonmar-

tensitic transformation and falls upon cooling at a rate of

3.5 Ksec (Fig. 2b ).Steel 15Kh2GMF has an almost similar content of alloy-

ing elements but differs substantially from steel

12Kh2G2NMFT in their quantitative proportion (Table 1).

For example, steel 15Kh2GMF contains more carbon and va-

nadium, whereas the concentration of manganese and molyb-

denum is lower. This causes general destabilization of aus-

tenite, which is manifested in the occurrence of the transfor-

mation in both pearlitic and bainitic regions (Fig. 1b ). A

structure of lath martensite in steel 15Kh2GMF is detected

only after cooling in water; upper bainite forms even after

cooling in air (vcool

= 3.5 Ksec). The strength characteristics

of steel 15Kh2GMF decrease continuously upon decrease in

the cooling rate, and the impact toughness and the dynamic

crack resistance are at a low level and do not rise upon de-

crease in the strength up to a level of 0.2

600 MPa

(Fig. 2c ).

CONCLUSIONS

1. We have formulated principles for designing chemical

composition, which provide high stability of supercooled

austenite and, as a consequence, ensure a martensitic struc-

ture in low-carbon alloy steels, namely,

base alloying with Cr, Mn, and Ni with total content of

at least 4%, which is required for providing enough stability

of supercooled austenite in the range normal (pearlitic) trans-

formation;

balanced alloying with strong carbide-forming ele-

ments (Mo, W, V, Ti, Nb, ...) in order to prevent bainitic

transformation; the number of these elements should be ma-

ximum (at least 3 4 elements) and their content should be

minimum; the elements may be arranged in a series accord-

ing to the decrease in their susceptibility to carbide for-

mation;

the principle of balancing implies that the content of

each carbide-forming element should be inversely propor-

tional to the value of the coefficient of interatomic interac-

tion between this element and carbon (KA

);

the used carbide-forming elements should not form

isomorphic carbides; otherwise, the temperature ranges of

formation of carbides should not be overlapped (for example,

for Ti and V that form isomorphic carbides the possibility of

segregation of VC due to cooling of the austenite at their low

enough concentrations in the solution is realized only when it

has been exhausted for TiC).

2. In fact, these principles of design of steels make it pos-

sible to minimize the activity of carbon in the solid solution

and to lower the probability of formation of carbides and

promote growth in the resistance to diffusion relaxation in

the system.

3. The use of an alloying system with elevated resistance

to diffusion relaxation is a condition for effective dispersion

of the structure of the steel as a result of austenization and

subsequent hardening.

REFERENCES

1. D. E. Kaputkin, Nonequilibrium states of the structure of har-

dened multicomponent iron alloys and their closeness to equi-

librium, Fundament. Prob. Sovr. Metalloved., 4(1), 58 65

(2007).

2. S. S. Gorelik, Recrystallization of Metals and Alloys [in Rus-

sian], 2nd ed., Metallurgiya, Moscow (1978), 568 p.

3. V. D. Sadovskii, Origin of inheritance of structure in steel,

Fiz. Met. Metalloved., 57(2), 213 223 (1984).

4. V. I. Izotov, V. V. Voznesenskii, and A. P. Bashchenko, Effect

of the size of initial grains on the structure and yield strength of

a steel hardened for martensite, in: Problems of Physical Me-

tallurgy and the Physics of Metals, Coll. Works of TsNIIChermet

[in Russian], Metallurgiya, Moscow (1976), No. 3, pp. 192 199.

5. A. Garcia-Junceda, C. Capdevila, F. G. Caballero, and C. Garcia

de Andres, Dependence of martensite start temperature on fine

austenite grain size, Scr. Mater., 58, 134 137 (2008).

6. L. F. Porter and D. S. Dabkowski, Regulation of grain size by

thermocycling, in: Superfine Grains in Metals [Russian trans-

lation], Metallurgiya, Moscow (1973), pp. 135 164.

7. R. A. Grange, Strengthening steel by austenite grain refine-

ment, Trans. Quart. ASM, 59, 26 47 (1966).

8. L. M. Kleiner, L. I. Kogan, and R. I. ntin, Properties of al-

loyed low-carbon martensite, Fiz. Met. Metalloved., 33(4),

824 830 (1972).

9. L. I. Kogan, L. M. Kleiner, and R. I. ntin, Special features of

transformation of austenite in low-carbon alloy steels, Fiz.

Met. Metalloved., 41(1), 118 124 (1976).

10. Yu. A. Bashnin, V. K. Galkin, and V. M. Korovina, Low-car-

bon high-strength alloy steels of martensitic class, Metalloved.

Term. Obrab. Met., No. 10, 50 52 (1980).


11. Yu. A. Bashnin, V. K. Galkin, and Yu. V. Vasilev, Computer-

aided design of low-carbon sparingly alloyed high-strength

steels, Metalloved. Term. Obrab. Met., No. 8, 34 37 (1989).

12. K. A. Lanskaya, Refractory Steels [in Russian], Metallurgiya,

Moscow (1969), 247 p.

13. L. E. Popova and A. A. Popov, Diagrams of Transformation of

Austenite in Steels and of Beta-Solution in Titanium Alloys, A

Handbook of Heat Treatment Specialist [in Russian], 3rd ed.,

Metallurgiya, Moscow (1991), 503 p.

14. V. G. Sorokin, Yu. A. Krasyuk, S. P. Gubova, et al., Inventors

Certif. No. 956604 of 07.05.82, Byull. Izobr. Polezn. Modeli,

No. 33 (1982).

15. A. P. Kamenskaya et al, Special features of transforma-

tion in steel 12Kh2G2NMFT, Fiz. Met. Metalloved., 93(1),

90 93 (2002).

16. A. P. Kamenskaya, L. Ts. Zayats, L. M. Kleiner, and Yu. N. Si-

monov, Special features of formation of structure and proper-

ties in low-carbon martensitic steel 12Kh2G2NMFT, Metallo-

ved. Term. Obrab. Met., No. 3, 10 12 (2003).

17. Yu. N. Simonov, Conditions of formation of a structure of lath

martensite due to slow cooling of low-carbon austenite, Fiz.

Met. Metalloved., 97(5), 77 81 (2004).

18. B. M. Mogutnov, I. A. Tomilin, and L. A. Shvartsman, The

Thermodynamics of Iron-Carbon Alloys [in Russian], Metallur-

giya, Moscow (1972).


AbstractKey wordsINTRODUCTIONMETHODS OF STUDYRESULTS AND DISCUSSIONPrinciples of Dispersion of Structure by Heat TreatmentStudy of the Effect of Alloying System on Stability of Supercooled AusteniteTheoretical Substantiation of the Choice of Alloying System with Elevated Resistance to Diffusion Relaxation

CONCLUSIONSREFERENCES

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