Metal Science and Heat Treatment

Vol. 50, Nos. 9 – 10, 2008

STEELS AND ALLOYS UDC 669.14.018.295

HIGH-STRENGTH AUSTENITIC STEELS WITH DIFFERENT ALLOYING SYSTEMS AND CARBIDE HARDENING I. I. Kositsyna1 Translated from Metallovedenie i Termicheskaya Obrabotka Metallov, No. 10, pp. 6 – 13, October, 2008.

The effect of the composition of the austenite matrix on the level of strengthening ensured by precipitation hardening and on the characteristics of high-temperature strength is considered. The relation between the alloying, the observed mechanisms of decomposition, and the morphology of carbide segregations in different-base austenitic steels and the level of their strength, ductility, and heat resistance is discussed.

heat treatment of precipitation-hardening steels is relatively simple. The mechanical properties of the steels can be controlled in a wide range by changing the temperature and time parameters of the aging and choosing the required combination of strength and ductility. The present work2 was devoted to studying the regular features of formation of structure and properties of precipitation-hardening austenitic steels with different matrices using the achievements in the field of aging of austenitic steels and the data of the known Moscow scientists O. A. Bannykh and V. M. Blinov [2], Ural scientists I. N. Bogachev and M. I. Gol’dshtein [5], M. A. Filippov, A. I. Uvarov, V. V. Sagaradze [4], and others. It is assumed that the most effective carbide-forming element is vanadium [2, 4, 5]. In contrast to stronger carbide-formers (Nb, Zr, Ti) vanadium forms VC carbides with carbon. These carbides are easily soluble in heating for hardening and form ultrafine and uniformly distributed segregations isomorphic to the g-matrix; these particles serve as effective barriers for dislocation slip and strongly harden the matrix without embrittling the material. Introduction of 12 – 18 wt.%3 Cr into the austenite matrix makes the material corrosion resistant and capable of self-induced passivation under natural conditions transforming it to the class of stainless or refractory steels. However, the austenite matrix of steels bearing 18% Cr cannot be stabilized with the help of only manganese. For this reason medium-

INTRODUCTION In addition to the main property required from austenitic steels, i.e., the nonmagnetic nature, they should often meet additional requirements of high strength, corrosion resistance under atmospheric conditions and in various aggressive media, high heat resistance and refractoriness, and good ductility at various temperatures including cryogenic ones. Austenitic steels based on the Fe – Mn, Fe – Cr – Mn, Fe – Cr – Ni, and Fe – Mn – Ni systems have always been and remain materials meeting these requirements to this or that degree [1]. A substantial drawback of hardened austenitic steels is a relatively low yield strength, which limits their use for heavily loaded parts and units of high-duty equipment. Ways for raising the yield strength of steels of the class in question are quite few [2 – 4], namely, solid solution hardening of austenite by additional alloying, cold and warm plastic deformation (strain hardening), phase hardening due to cyclic g ® a ® g transformations, and precipitation hardening (carbide, nitride, and intermetallic aging). Precipitation hardening accompanied by segregation of ultrafine particles of a second phase (carbides, nitrides, intermetallics) from supersaturated g-solid solution is undoubtedly a more universal and effective method for strengthening austenitic steels than the methods of solid-solution hardening or phase hardening. Precipitation hardening strengthens austenitic steels with virtually any combination of the main alloying elements, and 1

2 3

Institute for Metals Physics of the Ural Branch of the Russian Academy of Sciences, Ekaterinburg, Russia.

The author is grateful to V. V. Sagaradze, Yu. I. Filippov, and T. N. Kochetkova for their participation in the study. Here and below in the text the content of elements is given in mass percent.

459 0026-0673/08/0910-0459 © 2008 Springer Science + Business Media, Inc.

460

I. I. Kositsyna sr ; s0.2 , ÌPà

TABLE 1. Chemical Composition of Studied Steels Content of elements, wt.%

1800

Grade of steel

26G20M4F2 73G20M4F2 45G20M2F2 50Kh16G15N6M2F2 45Kh18N10G10M2F2 40Kh18N18M2F2 45N26M2F2

C

Mn

Cr

Ni

V

Mo

0.26 0.73 0.45 0.50 0.45 0.40 0.45

20.0 19.7 19.4 15.0 10.7 – –

– – – 15.8 17.7 18.0 –

– – – 5.9 10.0 18.0 25.7

2.05 2.20 1.71 1.74 1.60 1.66 1.79

3.9 3.9 1.64 1.71 1.61 1.51 1.78

d; y, % HRC sr 70

1400

s0.2

1200

HRC

1000

40

600

30 20

y

200

10

d

0 0

0.2

0.4

0.6

0.8

1.0

1.2

C, wt.%

METHODS OF STUDY We studied manganese, chromium-nickel-manganese, chromium-nickel, and nickel austenitic steels bearing 0.2 – 0.5% C, 16 – 18% Cr, at least 18% Mo + Ni, and alloyed with vanadium and molybdenum in different proportions. The chemical composition of the steels is presented in Table 1. These steels remain stable with respect to the g ® a and g ® e transformations at room temperature; a-martensite does not form after deformation ensured by tensile mechanical tests. We melted the metals in an open induction furnace. The obtained ingots with a mass of 20 kg were forged for preforms 8 ´ 8 mm in size at a temperature of 1200 – 1000°C. The mechanical properties after uniaxial stretching at room temperature and at 600°C were determined for fivefold specimens 3 mm in diameter. The tests for high-temperature strength were performed using standard AIMA-5-2 machines at a temperature of 500, 600, 650, and 700°C at a stress of from 150 to 600 MPa. The duration of the tests was at least 20,000 h. For each tested steel we plotted primary three-stage creep curves and used them to compute the creep rate in the steady stage (&e), the time before failure (t), and the long-term ductility (d). Then we plotted experimental dependences of the rate of steady creep (&e) and of the time before failure on the stress (s). All the data on high-temperature strength were processed with the help of the least squares statistical method [6]. The structure was analyzed using transmission electron microscopy with the help of a JEM-200CX microscope. The mean size of the carbides and the number of carbide particles per unit volume were computed from the results of statistical processing of the sizes of carbide particles measured on dark-background electron microscope images. The accuracy of evaluation of the sizes of the particles was no worse than ± 0.5%. Quantitative x-ray spectrum microanalysis of the composition of primary carbides was performed using a SuperprobeJCXA-733 microanalyzer accurate to 1% at a voltage of 5 and 25 kV. The heat treatment consisted in water quenching

50

800

400

carbon austenitic Cr – Mn steels are commonly alloyed with 6 – 8% Ni [1].

60

Fig. 1. Effect of carbon content on the properties of steels of type G20M4F2 after aging at 650°C for 10 h.

of the steels from 1150 – 1175°C and aging at 600, 650, 700, 750, and 800°C for 1, 5, 10, and 20 h. RESULTS When developing the high-strength austenitic steels and choosing the carbide-forming elements we first studied manganese austenitic steels of type G20 alloyed with various carbide-forming elements, i.e., from 1.5 to 4% vanadium, from 2 to 9% molybdenum, from 0 to 7% chromium, and from 0.1 to 1.2% carbon. Preliminary durometer measurements (after quenching and aging) showed that molybdenum alloying of the steels of type 80G20 hardened them during aging inconsiderably (the hardness was < 18 HRC ); the hardening due to chromium carbides (40 HRC ) was less than that due to vanadium carbides (47 HRC ). Only joint alloying with 2% V and 2 – 4% Mo ensured maximum hardening (52 HRC ) in high-carbon austenitic manganese steels. The positive effect of molybdenum manifests itself not only in elevation of the hardness but also in elevation of the heat resistance of austenitic steels; in the metals alloyed with molybdenum and vanadium noticeable decrease in the hardness due to overaging begins at an aging temperature 50 – 100°C higher than in the steels bearing only vanadium [7]. Figure 1 presents the dependences of mechanical properties of an aged steel of type G2M4F2 on the concentration of carbon (0.1 – 1.2%). It can be seen that the contribution of precipitation hardening into the elevation of the strength characteristics of the steels depends considerably on the carbon concentration. Austenitic manganese steel 26G20M4F2 after quenching and aging has the following mechanical properties: sr » 1300 MPa, s0.2 » 1150 MPa, d » 18 %;

High-Strength Austenitic Steels with Different Alloying Systems and Carbide Hardening y » 32 %; KCU + 20 ³ 1.4; KCU – 196 ³ 0.4 MJ/m2. It should be noted that it is possible to obtain s0.2 » 1600 MPa by raising the content of carbon in such steels (Fig. 1). Development of stainless austenitic steels with carbide hardening was started with alloying the base steel 20G20M2F2 with chromium in an amount of 12%. However, steels 19G20Kh12M2F2 and 29G20Kh12M2F2 acquired a (g + a) double-phase structure after the hardening, which led to formation of s-phase after aging. The austenite was stabilized by introducing nickel or copper in an amount of 3%. However, steels 20G20Kh12NM2F2 and 20G20Kh12D3M2F2 did not have high strength properties after hardening and aging, and the ductility of the metal with copper was lower than that of the metal with nickel. A high-strength stable austenitic steel was obtained after raising the carbon content to 0.4%. In further study we tested precipitation-hardening austenitic steels 45G20M2F2, 50Kh16G15N6M2F2, 45Kh18N10G10M2F2, 40Kh18N18M2F2, and 45N26M2F2 alloyed jointly with vanadium and molybdenum (Table 1). In order to ensure active precipitation hardening in aging it is necessary that carbon and the carbide-forming elements pass into the solid solution as fully as possible. With allowance for the necessity for the fullest dissolution of primary carbides and for prevention of undesirable grain growth, we hardened the studied steels from 1150 – 1175°C. A certain amount of undissolved primary carbides remained in the metals after the hardening. It was especially high in the steels alloyed with 16 – 18% Cr, i.e., about 3 – 4 vol.%, whereas in the nickel and manganese steels it was 1 vol.%. Quantitative x-ray spectrum microanalysis showed that the coarse primary carbides had a complex variable composition; the Me23C6 carbides contained 40 – 55 at.% Cr, 9 – 13 at.% Mn, 22 – 27 at.% Fe, 6 at.% V, and 2 – 3 at.% Mo and the MeC carbides contained up to 8 at.% V and Mo [8]. The mechanical properties of the steels after hardening are presented in Table 2. Manganese austenite can be distinguished among the various matrix compositions; it possesses a higher ultimate rupture strength (940 MPa) and a higher elongation (41%), which seems to be connected with the low energy of packing defects (EPD) of the manganese austenite, formation of packing defects due to stretching, and develop-

461

sr , ÌPà

s0.2 , ÌPà

1600

1350 1250 1150 1050 950 850 750 650 550 450

1500

1

1400 1300 1200

2

1100

3

1000 900

4 5

800 700 600

650

700

750

1

3 2 5 4 600

d, %

y, %

40

70

35

60

650

30

5

20

4

15

3 1

10 5

40

2

30

3

20

1

10

0 600

650

700

750

600

650

tag , °Ñ

s0.2 , MPa

sr , MPa

d, %

360 580 420 410 500 470 440

940 1040 940 850 850 830 780

52 26 41 39 36 33 30

750

Fig. 2. Mechanical properties of studied steels after 10-g aging at various temperatures: 1 ) 45G20M2F2; 2 ) 50Kh16G15N6M2F2; 3 ) 45Kh18N10G10M2F2; 4 ) 40Kh18N18M2F2; 5 ) 45N26M2F2.

ment of the process of strain-induced twinning [4]. The nickel and chromium-nickel austenites are superior with respect to the contraction (67 and 61%, respectively). The presence of a high amount of undissolved carbides affects the yield strength of hardened steels 40Kh18N18M2F2 and 45Kh18N10G10M2F2, which is quite high, i.e., 470 and 500 MPa, respectively). Figure 2 presents the mechanical properties of steels 50Kh16G15N6M2F2, 45G20M2F2, 45Kh18N10G10M2F2, 40Kh18N18M2F2, and 45N26M2F2 as a function of the aging mode. On the whole, the strength of the steels increases and the ductility decreases in the process of isothermal hold

y, %

Hardening from 1150°C

26G20M4F2 73G20M4F2 45G20M2F2 50Kh16G15N6M2F2 45Kh18G10N10M2F2 40Kh18N18M2F2 45N26M2F2

700

tag , °Ñ

TABLE 2. Mechanical Properties of Studied Steels Grade of steel

750

5 4

50

2

25

700

s0.2 , MPa

sr , MPa

d, %

y, %

Quenching + aging at 600°C for 10 h

60 36 54 51 49 61 67

1140 1580 1290 850 900 680 795

1310 1750 1440 1260 1080 1060 940

18 4 17 28 18 26 19

32 5 31 42 46 53 61

462

I. I. Kositsyna log e, 10 – 5 [%/h] 5

5

4

4

1

3

2 3

2 1

à 0 150 200 250 300 350 400 450 500 550 600 650

s, ÌPà

log t, h 5 4 3 2

1

4 3

1

5 2 b

0 150 200 250 300 350 400 450 500 550 600 650

s, ÌPà

Fig. 3. Dependence of the creep rate (e&) in the steady stage (a) and of the time before failure (t) (b ) on the applied stress (s) at 600°C for steels: 1 ) 45G20M2F2; 2 ) 50Kh16G15N6M2F2; 3 ) 45Kh18N10G10M2F2; 4 ) 40Kh18N18M2F2; 5 ) 45N26M2F2.

in the temperature range of segregation of secondary carbides, which is typical for precipitation hardening. When developing high-strength austenitic steels we carefully chose the modes of heat treatment (hardening and aging) and studied their effect on the mechanical properties. In the case of aging at 600°C the processes of precipitation hardening occurred too slowly. In steel 26G20M4F2 the yield strength increased less after 20 h of aging at 600°C than after 1 h of aging at 650°C, but the high ductility of the metal was preserved (d = 25% and y = 45%). At 700°C the phenomena of overaging became noticeable and the strength decreased upon growth in the aging time. For all the studied steels the peak of precipitation hardening corresponded to aging at 650°C (t = 10 h). It should be noted that the level of strength attained in the studied steels after aging differed (for example, after 10-h aging at 650°C, see Fig. 2 and Table 2). The highest strength was observed in the manganese austenite (s0.2 = 1290 MPa) and the lowest was observed in the chromium-nickel austenite (s0.2 = 600 MPa) at about the same content of carbon and of carbide-forming elements V and Mo. The stable manganese and chromium-manganese steels possessed a lower margin of ductility than the chromium-nickel and nickel

steels; for example, y = 31 and 42% and 53 and 61%, respectively (Table 2). The lowering of the strength characteristics of austenitic steels upon transition to g-matrices with chromium is primarily connected with the presence of primary Me23C6 carbides in the latter, which have not dissolved during heating for hardening and do not participate in precipitation hardening but bond the elements, especially carbon that is necessary for formation of the main hardening phase, i.e., the VC carbide. Computations show that when primary carbides form in steel 40Kh18N18M2F2 in an amount of 4.5 vol.%, about 0.26 wt.% C leaves the g-matrix and the solid solution preserves only 0.14 wt.% C, which lowers substantially the intensity of carbide formation in further aging and affects the final level of strengthening. It can be seen from Fig. 1 that when the carbon content decreases from 0.75% to 0.1%, the yield strength of the steels type G20M4F2 after aging decreases from 1550 to 500 MPa. Despite the fact that the studied chromium-manganese-nickel austenitic steels 50Kh16G15N6M2F2 and 45Kh18N10G10M2F2 are noticeably inferior to the manganese steel 45G20M2F2 with respect to the level of maximum strengthening due to carbide aging, they can be recommended as promising high-temperature corrosion-resistant austenitic steels sparingly alloyed with nickel. The presence of chromium in the steels makes them high-temperature carbide-hardened materials, the service characteristics of which depend on the composition of the g-solid solution, on the size and content of the carbide particles, and on the thermal stability of the structure. Before testing the test pieces of all the steels for high-temperature strength we subjected the former to quenching and aging at 700°C for 10 h. This treatment ensured formation of thermally stable hardened structure in the steels. The results of short-term high-temperature tests at 600°C show that manganese and chromium-manganese austenites having s0.2 = 625 and 570 MPa and d = 10 and 12% and y = 44 and 45%, respectively, are superior to nickel and chromium-nickel austenites for which s0.2 = 414 and 412 MPa, d = 28 and 26%, and y = 30 and 35%. Figure 3 presents the obtained dependences of the creep rate in the steady stage (&e) and of the time before failure (t) on the applied stresses (s) in semilogarithmic coordinates after testing for creep at 600°C. Similar dependences have been obtained for the test temperature of 700°C. For all the steels the creep rate in the steady stage increases upon growth in the test temperature and in the load applied. The creep rate of the steels at a temperature of 600°C and a stress of 200 – 300 MPa was &e ³ 1 ´ 10 – 4 %/h; at a temperature of 700°C and a stress of 200 MPa we obtained &e ³ 1 ´ 10 – 3 %/h; at a temperature of 700°C and a stress of 400 – 500 MPa we obtained &e ³ 1 ´ 10 – 1 %/h. It can be seen from Fig. 3 that the creep rate in the steady stage depends on the composition of the austenite matrix; for example, at a stress of 320 MPa the

High-Strength Austenitic Steels with Different Alloying Systems and Carbide Hardening

463

TABLE 3. Long-Term Strength of Studied Steels Grade of steel

45G20M2F2 50Kh 16G15N6M2F2 45Kh18G10NI0M2F2 40Kh18N18M2F2 45N26M2F2

s t1000 , MPa, at a temperature, °C

s t10,000 , MPa, at a temperature, °C

500

600

700

500

600

700

343 490 444 417 543

240 310 340 290 364

143 124 182 130 189

290 399 381 349 457

190 200 256 211 265

95 114 114 44 104

maximum creep rate is observed in steel 45G20M2F2, i.e., &e = 3.8 ´ 10 – 2 %/h. In steel 50Kh16G15N6M2F2 &e = 3.2 ´ 10 – 4 %/h, in steel 45Kh18N10G10M2F2 &e = 5.7 ´ 10 – 4 %/h, in steel 40Kh18N18M2F2 &e = 6.4 ´ 10 – 4 %/h, and in steel 45N26M2F2 &e = 9 ´ 10 – 4 %/h. Note that the known correlation [9] between the creep rate in the steady stage and the yield strength is not observed, i.e., the manganese steel with the highest strength has the highest creep rate and the lowest endurance. The deformation (elongation) before failure of the test pieces of the manganese, chromium-manganese, and chromium-nickel steels does not exceed 1% at a temperature of 600°C and 2.5% at 700°C in the used range of loads; in the nickel steel d = 4 – 5% at 600°C and 4 – 8% at 700°C, which is typical for high-strength steels. In order to estimate the working capacity of the studied austenitic steels with different matrices we used the parametric equation of Larsen – Miller [10], plotted parametric diagrams of the long-term strength, and predicted (by computation) the long-term strength in a temperature range of 500 – 700°C for the service time of 1000 and 10,000 h (Table 3). The results of the prediction and of the experiments on active stretching for 1000 h turned out to be close. It can be seen from the data of Table 3 that the ultimate long-term strength depends on the alloying system of the austenitic steel. The highest strengthening of manganese austenite in precipitation hardening is characterized by a high yield strength (at test temperatures of 20 and 600°C) but the long-term strength at 600°C in this case is the lowest (s1000 = 240 MPa and s10,000 = 190 MPa). The highest longterm strength at 600°C is exhibited by steel 45N26M2F2 (s1000 = 364 MPa and s10,000 = 265 MPa) with the lowest strength characteristics but with the best ductility. The studied chromium-manganese and chromium-nickel steels 50Kh16G15N6M2F2 and 45Kh18N10G10M2F2 are superior to the known steels ÉI388 (40Kh15N7G7FMS) and ÉI481 (37Kh12N8G8MFB) with respect to the long-term strength at 500 – 700°C for test bases of 1000 and 10,000 h [11]. RESULTS AND DISCUSSION In order to determine the causes of the observed difference in the mechanical properties of precipitation-hardening

austenitic steels with different matrices we studied the processes of nucleation and growth of carbide particles and the special features of their morphology and determined quantitative parameters of the hardening phases (the size of the carbides and the number of particles per unit volume). The mechanisms of carbide nucleation determine not only the morphology (the size, the shape, and the kind of distribution in the matrix) of the hardening particles but also the density of their segregation and hence the entire set of properties of the steel after aging. It is known [5] that when supersaturated austenite decomposes, carbide particles can nucleate by one of the following mechanisms (according to the prevalence upon decrease in the degree of supersaturation of the solid solution or growth in the aging temperature): nucleation in the matrix, on sweeping partial Frank dislocations (stacking fault dislocations), on sweeping perfect dislocations a/2á110ñ, on initial hardening dislocations, and, finally, on twin, subgrain, and grain boundaries. Manifestation of a specific mechanism of decomposition is primarily connected with the action of two competing factors, i.e., the degree of supersaturation of the solid solution (determined by the composition of the steel, by the hardening temperature, and, chiefly, by the aging temperature) and the diffusion mobility of atoms in the aging process (also controlled by the aging temperature and the composition of the austenite matrix). The maximum density of the segregating VC particles and hence the maximum strengthening of the precipitationhardening austenite can be attained due to implementation of a low-temperature mechanism of nucleation of particles, i.e., due to matrix homogeneous nucleation [2, 4, 5]. For all the studied steels maximum strengthening is ensured by 10 – 20-h aging at 650°C. According to an electron microscope study (see Fig. 4) homogeneous matrix segregation of VC carbides occurs after this treatment in all the studied steels (Fig. 4a ). However, only in the case of manganese austenite (steel 45G20M2F2) does a matrix mechanism of segregation of VC carbides develop at all the tested aging temperatures (600 – 800°C) [12]. It has been shown in [7] on the basis of measurement of the lattice parameter of vanadium carbide in x-ray diffraction analysis of carbide precipitate that molybdenum passes from g-solid solution into vanadium carbide during aging, which is accompanied by growth in its lattice parameter from

464

I. I. Kositsyna

0.15 mm

a

0.1 mm

b

c

0.2 mm

d

e

0.4 mm

f

Tr. of plane 111 111 002 (110)g

0.2 mm

111 002 (110) Tr. of plane

(100) (010)

0.1 mm

Fig. 4. Structure of austenitic steels after hardening and aging: a) 45G20M2F2 (hardening from 1175°C in water, aging at 650°C for 10 h), dark-background image in reflection (111)VC ; b ) 45N26M2F2 (hardening from 1175°C in water, aging at 650°C for 20 h), dark-background image in reflection (111)VC ; c) 45N26M2F2 (hardening from 1175°C in water, aging at 700°C for 5 h), the traces of planes in the plane of the foil (110)g are indicated); d ) 50Kh16G15N6M2F2 (hardening from 1175°C in water, aging at 650°C for 10 h); e) 40Kh18N18M2F2 (hardening from 1175°C in water, aging at 650°C for 20 h), the traces of planes in the plane of the foil (010)g are indicated); f ) 40Kh18N18M2F2 (hardening from 1175°C in water, aging at 700°C for 10 h), dark-background image in reflection (111)VC .

a = 0.4153 nm in the steel without molybdenum (this corresponds to VC0.82 [13]) to a = 0.4175 nm in the steels with vanadium and molybdenum. The lattice with parameter a = 0.4175 nm does not suit the range of homogeneity on unalloyed vanadium carbide. The anomalously large lattice parameter is explainable by formation of a composite carbide, i.e., alloying of the vanadium carbide with molybdenum, when up to 25% Mo can be dissolved in the VC and the lat-

tice parameter of the carbide can grow to a = 0.4220 nm [13]. However, the size of the particles and their amount at the same matrix mechanism of nucleation are affected substantially by the time and temperature parameters of the aging (Table 4). When the duration of isothermal hold at all the studied aging temperatures is increased, the VC particles become larger and their number per unit volume increases. This

High-Strength Austenitic Steels with Different Alloying Systems and Carbide Hardening

465

TABLE 4. Mean Size of VC Carbides (d, nm) and Number of Particles per Unit Volume (n ´ 1016, cm – 3 ) in the Steels d

n

d

n

d

n

d

n

d

n

Aging mode 45G20M2F2

600°C, 20 h 650°C, 10 h 650°C, 20 h

3.7 2.7 4.7

700°C, 5 h 700°C, 10 h 700°C, 20 h

4.2 4.5 6.2

2 9 10 6 8 5.4

50Kh16G16N6M2F2

7.9 6.1 25 ´ 5 10.8 12.5 30 ´ 5

0.6 1.0 0.6 0.02 0.4 0.2

is accompanied by a noticeable increase in the strength, which corresponds to growth in the volume fraction of segregated carbide particles. For example, the yield strength of steel 45G20M2F2 after aging at 650°C for 5, 10, and 20 h amounts to 980, 1290, and 1360 MPa, respectively. The finest carbides with high distribution density (Table 4) segregate in the manganese austenite; the mean size of the particles d = 3 – 6 nm and the number of particles per unit volume n = (2 – 10) ´ 1016 cm – 3. In the nickel austenite (steel 45N26M2F2) aged at 600 and 650°C the VC carbide also segregates homogeneously (Fig. 4b ). The number of particles per unit volume is an order of magnitude lower than in the manganese steel 45G20M2F2 and the mean size of the particles is 1.5 – 2 times larger (Table 4). It seems that nickel increases the mobility of vacancies in the austenite but prevents fixation of enough of it during hardening, which leads to a considerable decrease in the number of nucleation centers of the hardening VC particles and, on the contrary, accelerates the diffusion growth of particles. In addition, the solid solution of nickel austenite is more depleted of carbon than that of manganese austenite. In the early stages of aging at 700°C a new mechanism of nucleation of hardening particles appears in nickel austenite, i.e., heterogeneous aging (Fig. 4c ). Given that the orientation of the crystal is favorable, clusters of particles form in planes {110}g and {111}g , which indicates nucleation and growth of particles on perfect sweeping dislocations a/2á110ñ. Upon increase in the time of aging at 700°C, when there are not enough places for nucleation of particles on dislocations, homogeneous nucleation begins again and VC carbides grow in the matrix. The number of particles per unit volume increases by an order of magnitude (Table 4). Transition to chromium-bearing steels causes participation of secondary chromium carbides in the decomposition of the supersaturated solid solution of austenite, but the Me23C6 carbides do not become the main hardening phase, because they segregate over grain boundaries during aging (Fig. 4d ). When the duration of isothermal holds in the aging is increased, the grain-boundary Me23C6 carbides grow from 25 – 50 nm (650°C, 5 h) to 400 nm (700°C, 20 h). Grain boundaries are favorable places for formation of Me23C6 carbides due to grain boundary distortions of the lattice (this

45Kh18N10G10M2F2

8.0 7.3 10 ´ 4 7.0 8.0 15 ´ 7

40Kh18N18M2F2

45N26M2F2

0.8 2.1 1.0

8.0 8.7 11.0

0.1 0.5 0.8

6.8 6.4 7.1

0.3 0.8 0.8

1.4 2.0 1.3

15.0 16.0

0.02 0.04 0.04

15.0 9.0 10.0

0.04 0.2 0.1

35 ´ 8

promotes carbide nucleation) and a high rate of diffusion over grain boundaries (this ensures growth of the carbides). It should be noted that grain boundary carbides in the manganese austenite are not large; they are VC particles 10 – 20 nm in size when segregated at aging temperatures of 600 – 650°C and 50 – 80 nm in size when segregated at 700 – 750°C. However, they produce no less an effect on the decrease in the elongation, contraction, and the characteristics of hightemperature strength. In Cr – Mn – Ni and Cr – Ni austenitic steels the main hardening phase due to precipitation hardening is also represented by VC carbides, but the structural mechanism of their nucleation and their shape change upon growth in the aging temperature. At aging temperatures of 600 and 650°C matrix segregation of VC carbides is observed. However, when the temperature and the duration of the isothermal hold are increased, the particles coarsen to 30 – 35 nm and acquire a disc-like shape (Fig. 4e ). Disc-like shape of VC particles is observed only in Cr – Mn – Ni and Cr – Ni steels in late stages of aging (after 20 h at 650, 700, and 750°C). Discs form when the particles lose their coherence to the matrix in two directions and continue to grow as plates lying in planes {100}g . It is possible that in chromium-bearing steels chromium alloys the VC carbide, which affects the symmetry of the lattice and changes the shape of the carbide in late aging stages. In chromium-alloyed steels aging at 650°C for 10 h yields less VC carbides; the number of their particles per unit volume decreases by an order of magnitude, i.e., n = (0.5 – 2) ´ 1016 cm – 3 (Table 4) because the other Me23C6 carbides attract a considerable part of carbon from the solid solution. This leads to a substantial decrease in the level of strength attained in aging of Cr – Mn – Ni and Cr – Ni austenitic steels as compared to manganese austenitic steels (Fig. 2). The lowest strengthening (s0.2 does not exceed 680 MPa) due to matrix nucleation of VC carbides in steel 40Kh18N18M2F2 is connected with the low number of VC particles per unit volume (aging at 650°C for 10 h), i.e., n = 5 ´ 1015 cm – 3 (Table 4). Replacement of a part of nickel by manganese makes it possible to raise the content of VC particles per unit volume. Aging of steel 45Kh18G10N10M2F2 in the same mode (650°C, 10 h) pro-

466 duces four times more particles (n = 2.1 ´ 1016 cm – 3 ) and the yield strength amounts now to 900 MPa. In steels 50Kh16G15N6M2F2 and 40Kh18N18M2F2 (where the supersaturation is especially low due to the low content of carbon in the g-solid solution) 5- and 10-h holds at 700°C produce heterogeneous segregation of VC carbides on the initial dislocations (Fig. 4e ). The content of the VC particles per unit volume falls by an order of magnitude (as compared to homogeneous decomposition) and this affects the level of strengthening. For example, in steel 40Kh18N18M2F2 aged for 5 h at 700°C the yield strength s0.2 is virtually equal to that in hardened state (470 MPa) and increases by only 60 MPa after 10-h aging at 700°C. The ductility is lowered substantially primarily due to grain boundary segregations of carbides, i.e., d decreases from 34% (after hardening) to 24 and 19% and y decreases from 60% (after hardening) to 49 and 34%, respectively. Vanadium-bearing austenitic aging steels are characterized by the presence of zones free of fine segregations (SFZ) near grain boundaries. When the aging temperature is increased, the width of these zones increases. For example, for steel 45N26M2F2 aged at 600, 650, and 700°C it is equal to 160, 250, and 340 nm, respectively. The width of the SFZ also increases upon growth in the nickel content in the austenite matrix, which promotes diffusion of vacancies and atoms to grain boundaries. For example, after aging at 650°C the width of the SFZ in the manganese austenite is equal to 80 nm; in the Cr – Mn – Ni austenite it is equal to 95 nm, in the Cr – Ni austenite it is 200 nm, and in the nickel austenite it is 250 nm. It seems that the width of the SFZ influences the ductility of aged steels; the contraction grows upon increase in the nickel content in the composition of the austenite matrix and, respectively, upon growth in the width of SFZ in the steels (see Table 2). In the SFZ not reinforced by vanadium carbides the stresses concentrated in the places of blocking of slip bands by grain boundaries relax more fully. It seems that this effect of the SFZ hinders the nucleation and propagation of cracks that cause early failure due to creep and lower the ductility characteristics. Since the failure of materials under conditions of high-temperature creep can be controlled by generation and growth of micropores and microcracks on grain boundaries, it is obvious that steel 45N26M2F2, which has no coarse grain boundary segregations, is obviously more advantageous. CONCLUSIONS 1. Joint alloying with carbide-forming elements, i.e., carbon, vanadium, and molybdenum in a mass proportion of C : V : Mo » 0.5 : 2 : 2, makes it possible to obtain highstrength precipitation-hardening austenitic steels. 2. We have determined the temperature-and-time range of maximum effect of precipitation hardening. After the hardening treatment (water quenching from 1150°C + aging at 650°C for 10 h) austenitic steel of type 26G20M4F2 has

I. I. Kositsyna the following properties: s0.2 = 1150 MPa, sr = 1300 MPa, d = 21%, y = 32%, and KCU + 20 =1.4 MJ/m2. 3. We have established that the composition of the austenite matrix has an important effect on the formation of structure and on mechanical properties of carbide-hardening austenitic steels. At the same content of the main carbideforming elements manganese austenite is reinforced more intensely in the aging process (steel 45G20M2F2: s0.2 = 1290 MPa, d = 17%, y = 31%) than nickel austenite (steel 45N26M2F2: s0.2 = 795 MPa, d = 19%, y = 61%) and chromium-nickel or chromium-manganese-nickel austenite (steel 45Kh18N10G10M1F2: s0.2 = 900 MPa, d = 18%, y = 46%). The ductility of all the steels decreases in the aging process upon growth in the hold time; the elongation is virtually independent of the composition of the g-matrix and the contraction increases upon growth in the content of nickel in the steel. 4. Despite the high short-term mechanical properties at room and elevated (600°C) temperatures manganese steel 45G20M2F2 possesses a lower ultimate long-term strength than steel 45N26M2F2 having a low short-term strength. A high level of high-temperature properties at 500 – 700°C has been attained in chromium-manganese-nickel steels 50Kh16G15N6M2F2 and 45Kh18N10G10M2F2, namely, 600 600 = 310 and 340 MPa, s10 s1000 , 000 = 200 and 260 MPa, respectively. 5. The effect of the composition of the austenite matrix on the level of mechanical properties primarily manifests itself through changes in the morphology, sizes, and density of segregation of particles of the main hardening phase, i.e., vanadium carbide alloyed with molybdenum. In the manganese steel these quantities have extreme values of d = 3 – 5 nm, n = (8 – 10) ´ 1016 cm – 3; in the other steels the particle size is 2 – 3 times larger and the segregation density is 5 – 10 times lower. Alloying of the austenite with a high amount of chromium lowers the intensity of precipitation hardening due to bonding of a part of carbon in primary complexly alloyed Me23C6 carbides. REFERENCES 1. M. V. Pridantsev, N. P. Talov, and F. L. Levin, High-Strength Austenitic Steels [in Russian], Metallurgiya, Moscow (1969). 2. O. A. Bannykh and V. M. Blinov, Precipitation-Hardening Nonmagnetic Vanadium-Bearing Steels [in Russian], Nauka, Moscow (1980). 3. K. A. Malyshev, V. V. Sagaradze, I. P. Sorokin, et al., Phase Hardening of Iron-Nickel-Base Austenitic Alloys [in Russian], Nauka, Moscow (1982). 4. V. V. Sagaradze and A. I. Uvarov, Strengthening of Austenitic Steels [in Russian], Nauka, Moscow (1989). 5. M. I. Gol’dshtein and V. M. Farber, Precipitation-Hardening Steels [in Russian], Metallurgiya, Moscow (1979). 6. M. N. Stepnov, Statistical Processing of Results of Mechanical Tests [in Russian], Mashinostroenie, Moscow (1972). 7. I. I. Kositsyna, V. V. Sagaradze, and E. N. Frizen, “Structure and properties of aging Mn – V – Mo austenitic steels with different

High-Strength Austenitic Steels with Different Alloying Systems and Carbide Hardening

content of carbon,” Fiz. Met. Metalloved., 62(3), 556 – 565 (1986). 8. I. I. Kositsyna, V. V. Sagaradze, and O. N. Khakimova, “Special features of carbide aging of austenitic steels with different bases. 1. Mechanical properties and influence of primary carbides,” Fiz. Met. Metalloved., 82(1), 112 – 120 (1997). 9. V. A. Pavlov, Physical Foundations of Plastic Deformation of Metals [in Russian], Nauka, Moscow (1962). 10. V. M. Rozenberg, Fundamentals of High-Temperature Strength of Metallic Bodies [in Russian], Metallurgiya, Moscow (1973).

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11. S. B. Maslenkov and E. A. Maslenkova, Steels and Alloys for High Temperatures [in Russian], Metallurgiya, Moscow (1991), Vol. 1. 12. I. I. Kositsyna, V. V. Sagaradze, and O. N. Khakimova, “Special features of carbide aging of austenitic steels with different bases. 2. Kinetics and mechanisms of segregation of carbides,” Fiz. Met. Metalloved., 82(1), 121 – 130 (1997). 13. J. Goldschmidt, Interstitial Alloys [Russian translation], Mir, Moscow (1971), Vol. 1.