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High-strength precipitation-hardening austenitic Fe–Mn–V–Mo–C steels with shape memory effect
Materials Science and Engineering A 481–482 (2008) 747–751
High-strength precipitation-hardening austenitic Fe–Mn–V–Mo–C steels with shape memory effect V.V. Sagaradze, I.I. Kositsyna ∗ , M.L. Mukhin, Y.V. Belozerov, Yu.R. Zaynutdinov Institute of Metal Physics, Ural Branch RAS, 18 S. Kovalevskaya St., 620041 Ekaterinburg GSP-170, Russia Received 19 May 2006; received in revised form 20 November 2006; accepted 15 February 2007
Abstract Austenitic Fe–20% Mn alloys with the shape memory effect (SME) were studied. The alloys contained different carbon concentrations and were doped with V and Mo so as to achieve the maximum strengthening. It was shown that the contribution of the precipitation hardening depended strongly on the carbon concentration, the number of carbide-forming elements, and the temperature-and-time parameters of aging. The SME was controlled by stabilization and destabilization of the austenite with respect to the formation of the -martensite. For this purpose, conditions of formation of V(Mo)C carbides were adjusted. © 2007 Elsevier B.V. All rights reserved. Keywords: High-strength austenitic steel; Carbide strengthening; -Martensite; Shape memory effect
1. Introduction
2. Experimental
The shape memory effect (SME), which arises from the ␥ ↔ martensitic transformation in austenitic manganese steels, is discussed in Refs. [1–4]. A drawback of known austenitic Mn27–Si6 steels consists in their poor strength characteristics [2]. The precipitation hardening, which is realized due to the precipitation of dispersed MeC carbide phases, ensures considerable strengthening of austenitic steels. By changing the temperature-and-time parameters of aging of austenitic manganese steels, which are alloyed with carbide-forming elements, it is possible to adjust mechanical properties of the steels over broad limits and select the proper combination of the strength and plasticity [5]. Moreover, processes of stabilization and destabilization of the austenite with respect to the formation of the -martensite during carbide aging can be used to control the shape memory effect (SME) [3–4]. The controlled SME in Fe–Mn–V steels after carbide aging may be as large as ∼3% and the yield stress may be equal to 900 MPa [3,4]. The objective of the present study was to achieve the maximum possible strengthening by means of carbide aging of austenitic manganese steels alloyed with two carbide-forming elements (V and Mo).
The subjects of study were austenitic Fe–20% Mn steels, which contained 0.2–0.8% carbon and were doped with 2% V, 2% Mo and 1% Si: 0.20Mn20V2Mo2, 0.26Mn20V2Mo4Si, 0.31Mn20V2Mo3Si, 0.45Mn20V2Mo2, 0.78Mn20V2Mo2Si, and 0.29Mn20V2Si. The steels were made in an open inductionarc furnace. Ingots of a mass of 20 kg were forged to produce blanks 6 mm × 6 mm in size at temperatures of 1200–1000 ◦ C. The thermal treatment of the steels as studied included water quenching from 1150 to 1175 ◦ C and aging at temperatures of 600, 650, 700 and 750 ◦ C for 1, 5, 10 and 20 h. The mechanical properties were determined using samples of diameter 3 mm, which were subject to the uniaxial tension at room temperature. The structure was analyzed in a JEM-200CX transmission electron microscope. The average size of carbides precipitated during steel aging and their number per unit volume were calculated from statistically processed measurements of carbide particles in dark-field electron microscopy images by standard methods. The particle size was estimated to within ±0.5 nm and the number of particles per unit volume was determined to within ±10%. The quantitative electron probe microanalysis (EPMA) of the carbide composition was performed in a Superprobe-JCXA733 microanalyses (the accuracy of 1 mass%). The carbide precipitates as separated in the electrolyte (7.5% KCl, 10% glycerin, 82.5% HCl) at the current density of 0.01 A/cm2 and the
∗
Corresponding author. Tel.: +7 343 3783879; fax: +7 343 3745244. E-mail address: [email protected] (I.I. Kositsyna).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.02.158
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V.V. Sagaradze et al. / Materials Science and Engineering A 481–482 (2008) 747–751
temperature of −5 ◦ C were analyzed using a DRON-3.0 X-ray diffractometer by means of Cr K␣ radiation. The lattice spacing was evaluated by graphical extrapolation to the first five reflections. The influence of the bending strain on the recovery of their original shape was analyzed by determination of deformation characteristics due to the SME in these compounds. For the purpose of the quantitative estimation, the reversible deformation (eSME ) of plane samples having the cross-sectional area of 1.5 mm × 10 mm was measured when these samples were heated to the austenitic range over Af = 300 ◦ C. The samples were predeformed by bending on a mandrel of diameter D = 30 mm at room temperature (Md < Af ). 3. Results 3.1. Strengthening of Mn–V–Mo steels during aging The precipitation hardening is followed by considerable changes in strength and plasticity properties. The strength is improved during aging and plasticity is impaired under isothermal treatment conditions. The analysis of the mechanical properties, which were measured using different temperatureand-time parameters of aging, demonstrated that the maximum strengthening was achieved at 650 ◦ C after 10–20 h. At the aging temperature of 600 ◦ C, processes of the precipitation hardening were very slow. For example, the yield stress of the 0.31Mn20V2Mo2Si steel increased less after aging at 600 ◦ C for 20 h than after aging at 650 ◦ C for 3 h. However, the mate-
rial preserved a high plasticity (δ = 25%, ψ = 45%) after aging at 600 ◦ C. Overaging phenomena became apparent at 700 ◦ C. The ultimate strength and the yield stress were a maximum after aging for 1–3 h and then decreased slightly. The strength and plasticity characteristics were lower at 700 ◦ C than those at 650 ◦ C at all aging times (Table 1). The change of the carbon concentration had a considerable effect on the precipitation-hardening kinetics and the achieved properties of the steels supersaturated with strong carbideforming elements. For example, when the carbon concentration of the Mn20–V2–Mo2 steels changed from 0.2 to 0.78%, their yield stress increased from 1000 to 1560 MPa (after treatment at 650 ◦ C for 10 h) and the plasticity properties were impaired by a factor of 10 (Table 1). The main factors that determine the strength of an aged material are the number, the size and the shape of carbide particles and the level of elastic stresses around them. The steels in the quenched state had austenite grains with rare coarse carbides 160–340 nm in size, which did not dissolve during quenching heating. The carbides were present both in the bulk of the grains and on the grain boundaries. The quantitative EPMA of the coarse primary carbides suggested their complex variable chemical composition (V92 Mo8 )C. The structure of the quenched steels contained stacking faults and annealing twins, which were indicative of small stacking-fault energy of the iron–manganese austenite. Dispersed particles of the vanadium carbide precipitated during aging in all steels as investigated. This conclusion was confirmed by results of the X-ray diffraction analysis of the carbide precipitates and by
Table 1 Mechanical properties of the steels, the carbide size (n), the number of carbides per unit volume (d), and the reversible deformation (e) Thermal treatment
HRC
σ B (MPa)
σ 0.2 (MPa)
δ (%)
ϕ (%)
1150 ◦ C,
water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
17 37 28 35
810 1210 1160 1160
280 1040 840 810
65 26 29 29
64 37 42 39
1150 ◦ C,
water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
17 40 29 35
990 1310 1240 1160
350 1140 930 870
52 17 21 12
43 22 26 15
1150 ◦ C, water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
17 39 33 28
1010 1220 1180 1150
280 970 850 670
44 22 30 21
34 23 29 23
0.31Mn20V2Mo3Si
1150 ◦ C, water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
18 42 41 41
980 1400 1300 1310
440 1190 1030 980
54 17 14 10
56 14 22 16
0.45Mn20V2Mo2
1150 ◦ C, water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
20 45 43 42
940 1440 1480 1145
420 1290 1200 620
41 17 13 10
54 31 26 16
0.78Mn20V2Mo2Si
1150 ◦ C, water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
25 52 48 50
1030 1750 1670 1580
640 1580 1500 1510
26 3 3 1
36 5 4 0.5
Steel
0.20Mn20V2Mo2
0.26Mn20V2Mo2Si
0.29Mn20V2Si
d (nm)
n (cm−3 )
eSME (%)
4.2
9 × 1015
6.1
7 × 1015
0.75 0.83 1.33
4.3
3.6 × 1016
5.9
1.2 × 1016
5.3
1.8 × 1016
18
0.87 0.90 1.61
3.9 × 1013 0.98 1.08 1.26
3.7 4.2 6.2
9 × 1016 6 × 1016 5.4 × 1016
0.91 0.52 1.32
4.1
9.2 × 1016
7.8
4.5 × 1016
0.78 0.98 1.69
V.V. Sagaradze et al. / Materials Science and Engineering A 481–482 (2008) 747–751
transmission electron microscopy (TEM) examination. At early stages of aging (temperatures of 600–650 ◦ C and holding time of 1 h), the structure had a matrix deformation “tweed” contrast with two-dimensional bands along 1 0 0 directions. This contrast was probably due to the formation of a quasi-periodic (modulated) structure of a system of clusters or matrix-coherent particles in the presence of a periodic field of elastic distortions caused by these particles. The electron diffraction patterns did not contain reflections of a new phase and diffuse effects, which were symmetric to the main reflections in 1 0 0 directions of the reciprocal lattice, appearing near reflections of the ␥-matrix. When the aging time of the steels with 0.2–0.3 mass% C increased to 3 h, the matrix contrast from precipitates was enhanced in the electron microscopic images, but the diffraction patterns still contained reflections of the ␥ solid solution only. The strength of the steels was considerably improved already at this stage of aging. The selected area diffraction patterns changed after aging at 650 ◦ C for 5 h or longer. Additional reflections appeared near matrix reflections. The former were indexed as reflections of dispersed particles of the vanadium carbide with the matrix-isomorphous FCC lattice. The “cube–cube” orientation relations (0 0 1)VC || (0 0 1)␥ and [0 0 1]VC || [0 0 1]␥ of the lattices of VC and the ␥ solid solution were noted. The qualitative changes in the fine structure of the steels under study were insignificant after the time of aging at 600–700 ◦ C increased to 10–20 h and the carbon concentration increased from 0.2 to 0.7 mass%. A homogeneous matrix precipitation of MeC carbides was observed in all the steels at hand at the aging temperatures of 600–800 ◦ C. The maximum strengthening of the steels after aging at 650 ◦ C for 10 h was due to a high precipitation density of carbide particles and their dispersion (Table 1). A characteristic deformation contrast (in the form of a pair of dark segments separated by the zero-contrast line) was seen around the VC particles. This contrast, which usually is observed near coherent particles, suggested the spherical symmetry of the matrix distortion fields. The deformation contrast around particles, which indicated to the coherence of the VC particles in the austenite matrix, was preserved after aging at 700 ◦ C for 3 h (Fig. 1a). The aging treatment at 750 ◦ C for 12 h was followed by the attenuation of the deformation contrast around the particles, the loss of the coherence between the lattices of the vanadium carbide and the matrix, and the growth of the particle size (Fig. 1b). The precipitation of molybdenum carbides in the steels investigated was confirmed by the X-ray diffraction analysis of the carbide precipitates. However, fine Mo2 C carbides could not be detected during the electron microscopy examination because of a large density of VC precipitates. Separate coarse Me6 C carbides were observed on grain boundaries in the high-carbon Mn–V–Mo steel. The strengthening effect of molybdenum in the austenitic manganese steel was due not so much to the formation of intrinsic carbides as to the influence of molybdenum on the quantity, the size, the stability and the composition of the main strengthening phase, namely the vanadium carbide.
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Measurements of the lattice spacing in the vanadium carbide during the X-ray spectrum analysis of the carbide precipitate showed that molybdenum passed from the ␥ solid solution to the vanadium carbide during aging. The lattice spacing of the vanadium carbide in the molybdenum-free steel was a = 0.4153 ± 0.0001 nm, which corresponded to VC0.82 [5]. The lattice spacing of the vanadium carbide in the Fe–Mn–Mo–V steel increased to a = 0.4175 ± 0.0001 nm. The lattice with this spacing did not fit the homogeneity region of the unalloyed vanadium carbide. The anomalously large lattice spacing was explained by alloying of the vanadium carbide with molybdenum, because, in accordance with [5], up to 25% molybdenum could dissolve in VC. As a result, the lattice spacing of the carbide increased up to a = 0.4220 nm. Alloying of the vanadium carbide with molybdenum led to improvement of the strength, the hardness and the heat resistance of the austenitic manganese steels, which were alloyed simultaneously with vanadium and molybdenum. The strength of the molybdenum-free 0.29Mn20V2Si steel was less after the steel was treated under analogous thermal conditions (Table 1). The quantitative analysis demonstrated that the precipitation density of vanadium carbide particles depended on the carbon concentration of the steels, the presence of molybdenum, and the temperature-and-time parameters of aging. For example, the number of carbides in the aged Mn20V2Mo2 steel with 0.20 mass% C was smaller (9.0 × 1015 cm−3 ) than in the steel with 0.26 mass% C (3.6 × 1016 cm−3 ), and all the smaller than in the steel with 0.73 mass% C (9 × 1016 cm−3 ). Strengthening of the steels changed correspondingly (Table 1). If the carbon concentration was the same, the effect of molybdenum was seen. The number of MeC carbides increased in the molybdenumcontaining steels. For example, the number of particles per unit volume (after aging at 650 ◦ C for 10 h) was 1.8 × 1016 cm−3 in the 0.29Mn20V2Si steel and 3.6 × 1016 cm−3 in the 0.26Mn20V2Mo2Si steel or 3.8 × 1016 cm−3 in the 0.75Mn20V2Si steel and 9.2 × 1016 cm−3 cm−3 in the 0.78Mn20V2Mo2Si steel. Alloying of the vanadium carbide with molybdenum especially retarded the growth of particles at high aging temperatures. For example, in the case of overaging, the size of the (V, Mo)C carbide increased insignificantly in the Mn–V–Mo steels, while the size of some VC particles increased four to five times in the 0.75Mn20V2Si steel. It is seen from Table 1 that the maximum strengthening of the austenitic manganese steels with 0.2–0.4 mass% C was achieved after aging at 650 ◦ C for 10 h and was equal to 1040–1290 MPa, which was much larger than strengthening of analogous molybdenum-free steels. 3.2. The change of the SME magnitude in aged Mn–V–Mo steels Martensite phases were not detected in all the heat treatment conditions. However, tion of a few percent led to the formation of Fig. 1c and d presents the steel structure
the steels under plastic deformathe -martensite. after the ␥ →
750
V.V. Sagaradze et al. / Materials Science and Engineering A 481–482 (2008) 747–751
Fig. 1. Microstructure (TEM) of the 0.18Mn20V2Mo2Si (a and b) and 0.78Mn20V2Mo2Si (c and d) steels after aging in different conditions: (a) aging at 700 ◦ C, 3 h; (b) aging at 750 ◦ C, 12 h; (c) aging at 750 ◦ C, 12 h, deformation, dark-field image in the reflection of (V, Mo)C; (d) aging at 750 ◦ C, 12 h, deformation, dark-field image in the reflection of the -phase.
transformation, which was caused by 2–4% plastic deformation. The SME shows up in austenitic Mn–V–Mo steels as a partial restoration of the shape of a bend-strained sample under heating. The restoration is caused by the forward martensitic ␥ → transformation during cold deformation and the reverse shear → ␥ transformation during subsequent heating to a temperature above Af . According to the resistometric analysis, Af was 300 ◦ C in the alloys studied. It was found (Table 1) that the SME was enhanced with the steel aging temperature and its magnitude depended on the quantity of the -martensite. The absolute values of the reversible deformation under heating were as large as 1.7% and 0.9% in the steels, which were aged at 750 ◦ C for 12 h and 700 ◦ C for 3 h, respectively. The concentration of the -martensite in the deformed steel after destabilization aging at 750 ◦ C for 12 h was much larger than its concentration after stabilization aging at 700 ◦ C for 3 h. This
fact was explained by different stability (which was determined by changes in the coherence and the morphology of vanadium carbide precipitates) of the aged austenite with respect to the ␥ ↔ transformation and the saturation of the ␥-matrix with austenite-forming elements. The elastic field of the fine coherent (V, Mo)C carbides, which precipitated during aging, hampered the formation and the growth of the -martensite and stabilized the austenite phase. The loss of the coherence, the growth of (V, Mo)C particles and the increase in their volume fraction during aging at 750 ◦ C for 12 h decreased the stabilization effect of the carbides. Moreover, the martensite point was elevated due to the depletion of the matrix in carbon, vanadium and molybdenum after precipitation of carbides and the austenite was destabilized. As a result, the SME was enhanced after aging at 750 ◦ C for 12 h. It should be noted that 2% Mo added to the austenitic Mn–V steel stabilized it with respect to the formation of the strain-induced -martensite.
V.V. Sagaradze et al. / Materials Science and Engineering A 481–482 (2008) 747–751
In this case, the SME diminished in general. The thermal or thermomechanical treatment conditions need be optimized further so as to increase the SME in high-strength Mn–V–Mo steels. 4. Conclusion High-strength austenitic manganese steels showing SME and reaching the yield stress of 1100–1500 MPa were proposed. The proposed steels can be strengthened and their SME may be controlled by changing the temperature-and-time parameters of aging. In this case, the austenite is stabilized or destabilized with respect to the formation of the -martensite.
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Acknowledgment This study was supported by the RFBR (Project No. 06-03 32715). References [1] K. Enami, F. Nagasavo, S. Nenno, Scripta Metall. 9 (9) (1975) 941– 948. [2] K. Murakami, H. Otsuka, H.G. Suzuki, S. Matsuda, ICOMAT-96. Nara Imao Tamura Eds. (1987) 985. [3] V.V. Sagaradze, Ye.V. Beloserov, N.L. Pecherkina, et al., ICOMAT-2005, Abstract Book, Shanghai, 2005, p. 153. [4] V.V. Sagaradze, Ye.V. Beloserov, M.L. Mukhin, et al., Phys. Met. Metallogr. 101 (5) (2006) 506–512. [5] I.I. Kositsyna, V.V. Sagaradze, Metals 6 (2001) 65–74.
High-strength precipitation-hardening austenitic Fe–Mn–V–Mo–C steels with shape memory effect V.V. Sagaradze, I.I. Kositsyna ∗ , M.L. Mukhin, Y.V. Belozerov, Yu.R. Zaynutdinov Institute of Metal Physics, Ural Branch RAS, 18 S. Kovalevskaya St., 620041 Ekaterinburg GSP-170, Russia Received 19 May 2006; received in revised form 20 November 2006; accepted 15 February 2007
Abstract Austenitic Fe–20% Mn alloys with the shape memory effect (SME) were studied. The alloys contained different carbon concentrations and were doped with V and Mo so as to achieve the maximum strengthening. It was shown that the contribution of the precipitation hardening depended strongly on the carbon concentration, the number of carbide-forming elements, and the temperature-and-time parameters of aging. The SME was controlled by stabilization and destabilization of the austenite with respect to the formation of the -martensite. For this purpose, conditions of formation of V(Mo)C carbides were adjusted. © 2007 Elsevier B.V. All rights reserved. Keywords: High-strength austenitic steel; Carbide strengthening; -Martensite; Shape memory effect
1. Introduction
2. Experimental
The shape memory effect (SME), which arises from the ␥ ↔ martensitic transformation in austenitic manganese steels, is discussed in Refs. [1–4]. A drawback of known austenitic Mn27–Si6 steels consists in their poor strength characteristics [2]. The precipitation hardening, which is realized due to the precipitation of dispersed MeC carbide phases, ensures considerable strengthening of austenitic steels. By changing the temperature-and-time parameters of aging of austenitic manganese steels, which are alloyed with carbide-forming elements, it is possible to adjust mechanical properties of the steels over broad limits and select the proper combination of the strength and plasticity [5]. Moreover, processes of stabilization and destabilization of the austenite with respect to the formation of the -martensite during carbide aging can be used to control the shape memory effect (SME) [3–4]. The controlled SME in Fe–Mn–V steels after carbide aging may be as large as ∼3% and the yield stress may be equal to 900 MPa [3,4]. The objective of the present study was to achieve the maximum possible strengthening by means of carbide aging of austenitic manganese steels alloyed with two carbide-forming elements (V and Mo).
The subjects of study were austenitic Fe–20% Mn steels, which contained 0.2–0.8% carbon and were doped with 2% V, 2% Mo and 1% Si: 0.20Mn20V2Mo2, 0.26Mn20V2Mo4Si, 0.31Mn20V2Mo3Si, 0.45Mn20V2Mo2, 0.78Mn20V2Mo2Si, and 0.29Mn20V2Si. The steels were made in an open inductionarc furnace. Ingots of a mass of 20 kg were forged to produce blanks 6 mm × 6 mm in size at temperatures of 1200–1000 ◦ C. The thermal treatment of the steels as studied included water quenching from 1150 to 1175 ◦ C and aging at temperatures of 600, 650, 700 and 750 ◦ C for 1, 5, 10 and 20 h. The mechanical properties were determined using samples of diameter 3 mm, which were subject to the uniaxial tension at room temperature. The structure was analyzed in a JEM-200CX transmission electron microscope. The average size of carbides precipitated during steel aging and their number per unit volume were calculated from statistically processed measurements of carbide particles in dark-field electron microscopy images by standard methods. The particle size was estimated to within ±0.5 nm and the number of particles per unit volume was determined to within ±10%. The quantitative electron probe microanalysis (EPMA) of the carbide composition was performed in a Superprobe-JCXA733 microanalyses (the accuracy of 1 mass%). The carbide precipitates as separated in the electrolyte (7.5% KCl, 10% glycerin, 82.5% HCl) at the current density of 0.01 A/cm2 and the
∗
Corresponding author. Tel.: +7 343 3783879; fax: +7 343 3745244. E-mail address: [email protected] (I.I. Kositsyna).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.02.158
748
V.V. Sagaradze et al. / Materials Science and Engineering A 481–482 (2008) 747–751
temperature of −5 ◦ C were analyzed using a DRON-3.0 X-ray diffractometer by means of Cr K␣ radiation. The lattice spacing was evaluated by graphical extrapolation to the first five reflections. The influence of the bending strain on the recovery of their original shape was analyzed by determination of deformation characteristics due to the SME in these compounds. For the purpose of the quantitative estimation, the reversible deformation (eSME ) of plane samples having the cross-sectional area of 1.5 mm × 10 mm was measured when these samples were heated to the austenitic range over Af = 300 ◦ C. The samples were predeformed by bending on a mandrel of diameter D = 30 mm at room temperature (Md < Af ). 3. Results 3.1. Strengthening of Mn–V–Mo steels during aging The precipitation hardening is followed by considerable changes in strength and plasticity properties. The strength is improved during aging and plasticity is impaired under isothermal treatment conditions. The analysis of the mechanical properties, which were measured using different temperatureand-time parameters of aging, demonstrated that the maximum strengthening was achieved at 650 ◦ C after 10–20 h. At the aging temperature of 600 ◦ C, processes of the precipitation hardening were very slow. For example, the yield stress of the 0.31Mn20V2Mo2Si steel increased less after aging at 600 ◦ C for 20 h than after aging at 650 ◦ C for 3 h. However, the mate-
rial preserved a high plasticity (δ = 25%, ψ = 45%) after aging at 600 ◦ C. Overaging phenomena became apparent at 700 ◦ C. The ultimate strength and the yield stress were a maximum after aging for 1–3 h and then decreased slightly. The strength and plasticity characteristics were lower at 700 ◦ C than those at 650 ◦ C at all aging times (Table 1). The change of the carbon concentration had a considerable effect on the precipitation-hardening kinetics and the achieved properties of the steels supersaturated with strong carbideforming elements. For example, when the carbon concentration of the Mn20–V2–Mo2 steels changed from 0.2 to 0.78%, their yield stress increased from 1000 to 1560 MPa (after treatment at 650 ◦ C for 10 h) and the plasticity properties were impaired by a factor of 10 (Table 1). The main factors that determine the strength of an aged material are the number, the size and the shape of carbide particles and the level of elastic stresses around them. The steels in the quenched state had austenite grains with rare coarse carbides 160–340 nm in size, which did not dissolve during quenching heating. The carbides were present both in the bulk of the grains and on the grain boundaries. The quantitative EPMA of the coarse primary carbides suggested their complex variable chemical composition (V92 Mo8 )C. The structure of the quenched steels contained stacking faults and annealing twins, which were indicative of small stacking-fault energy of the iron–manganese austenite. Dispersed particles of the vanadium carbide precipitated during aging in all steels as investigated. This conclusion was confirmed by results of the X-ray diffraction analysis of the carbide precipitates and by
Table 1 Mechanical properties of the steels, the carbide size (n), the number of carbides per unit volume (d), and the reversible deformation (e) Thermal treatment
HRC
σ B (MPa)
σ 0.2 (MPa)
δ (%)
ϕ (%)
1150 ◦ C,
water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
17 37 28 35
810 1210 1160 1160
280 1040 840 810
65 26 29 29
64 37 42 39
1150 ◦ C,
water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
17 40 29 35
990 1310 1240 1160
350 1140 930 870
52 17 21 12
43 22 26 15
1150 ◦ C, water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
17 39 33 28
1010 1220 1180 1150
280 970 850 670
44 22 30 21
34 23 29 23
0.31Mn20V2Mo3Si
1150 ◦ C, water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
18 42 41 41
980 1400 1300 1310
440 1190 1030 980
54 17 14 10
56 14 22 16
0.45Mn20V2Mo2
1150 ◦ C, water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
20 45 43 42
940 1440 1480 1145
420 1290 1200 620
41 17 13 10
54 31 26 16
0.78Mn20V2Mo2Si
1150 ◦ C, water 650 ◦ C, 10 h 700 ◦ C, 3 h 750 ◦ C, 12 h
25 52 48 50
1030 1750 1670 1580
640 1580 1500 1510
26 3 3 1
36 5 4 0.5
Steel
0.20Mn20V2Mo2
0.26Mn20V2Mo2Si
0.29Mn20V2Si
d (nm)
n (cm−3 )
eSME (%)
4.2
9 × 1015
6.1
7 × 1015
0.75 0.83 1.33
4.3
3.6 × 1016
5.9
1.2 × 1016
5.3
1.8 × 1016
18
0.87 0.90 1.61
3.9 × 1013 0.98 1.08 1.26
3.7 4.2 6.2
9 × 1016 6 × 1016 5.4 × 1016
0.91 0.52 1.32
4.1
9.2 × 1016
7.8
4.5 × 1016
0.78 0.98 1.69
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transmission electron microscopy (TEM) examination. At early stages of aging (temperatures of 600–650 ◦ C and holding time of 1 h), the structure had a matrix deformation “tweed” contrast with two-dimensional bands along 1 0 0 directions. This contrast was probably due to the formation of a quasi-periodic (modulated) structure of a system of clusters or matrix-coherent particles in the presence of a periodic field of elastic distortions caused by these particles. The electron diffraction patterns did not contain reflections of a new phase and diffuse effects, which were symmetric to the main reflections in 1 0 0 directions of the reciprocal lattice, appearing near reflections of the ␥-matrix. When the aging time of the steels with 0.2–0.3 mass% C increased to 3 h, the matrix contrast from precipitates was enhanced in the electron microscopic images, but the diffraction patterns still contained reflections of the ␥ solid solution only. The strength of the steels was considerably improved already at this stage of aging. The selected area diffraction patterns changed after aging at 650 ◦ C for 5 h or longer. Additional reflections appeared near matrix reflections. The former were indexed as reflections of dispersed particles of the vanadium carbide with the matrix-isomorphous FCC lattice. The “cube–cube” orientation relations (0 0 1)VC || (0 0 1)␥ and [0 0 1]VC || [0 0 1]␥ of the lattices of VC and the ␥ solid solution were noted. The qualitative changes in the fine structure of the steels under study were insignificant after the time of aging at 600–700 ◦ C increased to 10–20 h and the carbon concentration increased from 0.2 to 0.7 mass%. A homogeneous matrix precipitation of MeC carbides was observed in all the steels at hand at the aging temperatures of 600–800 ◦ C. The maximum strengthening of the steels after aging at 650 ◦ C for 10 h was due to a high precipitation density of carbide particles and their dispersion (Table 1). A characteristic deformation contrast (in the form of a pair of dark segments separated by the zero-contrast line) was seen around the VC particles. This contrast, which usually is observed near coherent particles, suggested the spherical symmetry of the matrix distortion fields. The deformation contrast around particles, which indicated to the coherence of the VC particles in the austenite matrix, was preserved after aging at 700 ◦ C for 3 h (Fig. 1a). The aging treatment at 750 ◦ C for 12 h was followed by the attenuation of the deformation contrast around the particles, the loss of the coherence between the lattices of the vanadium carbide and the matrix, and the growth of the particle size (Fig. 1b). The precipitation of molybdenum carbides in the steels investigated was confirmed by the X-ray diffraction analysis of the carbide precipitates. However, fine Mo2 C carbides could not be detected during the electron microscopy examination because of a large density of VC precipitates. Separate coarse Me6 C carbides were observed on grain boundaries in the high-carbon Mn–V–Mo steel. The strengthening effect of molybdenum in the austenitic manganese steel was due not so much to the formation of intrinsic carbides as to the influence of molybdenum on the quantity, the size, the stability and the composition of the main strengthening phase, namely the vanadium carbide.
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Measurements of the lattice spacing in the vanadium carbide during the X-ray spectrum analysis of the carbide precipitate showed that molybdenum passed from the ␥ solid solution to the vanadium carbide during aging. The lattice spacing of the vanadium carbide in the molybdenum-free steel was a = 0.4153 ± 0.0001 nm, which corresponded to VC0.82 [5]. The lattice spacing of the vanadium carbide in the Fe–Mn–Mo–V steel increased to a = 0.4175 ± 0.0001 nm. The lattice with this spacing did not fit the homogeneity region of the unalloyed vanadium carbide. The anomalously large lattice spacing was explained by alloying of the vanadium carbide with molybdenum, because, in accordance with [5], up to 25% molybdenum could dissolve in VC. As a result, the lattice spacing of the carbide increased up to a = 0.4220 nm. Alloying of the vanadium carbide with molybdenum led to improvement of the strength, the hardness and the heat resistance of the austenitic manganese steels, which were alloyed simultaneously with vanadium and molybdenum. The strength of the molybdenum-free 0.29Mn20V2Si steel was less after the steel was treated under analogous thermal conditions (Table 1). The quantitative analysis demonstrated that the precipitation density of vanadium carbide particles depended on the carbon concentration of the steels, the presence of molybdenum, and the temperature-and-time parameters of aging. For example, the number of carbides in the aged Mn20V2Mo2 steel with 0.20 mass% C was smaller (9.0 × 1015 cm−3 ) than in the steel with 0.26 mass% C (3.6 × 1016 cm−3 ), and all the smaller than in the steel with 0.73 mass% C (9 × 1016 cm−3 ). Strengthening of the steels changed correspondingly (Table 1). If the carbon concentration was the same, the effect of molybdenum was seen. The number of MeC carbides increased in the molybdenumcontaining steels. For example, the number of particles per unit volume (after aging at 650 ◦ C for 10 h) was 1.8 × 1016 cm−3 in the 0.29Mn20V2Si steel and 3.6 × 1016 cm−3 in the 0.26Mn20V2Mo2Si steel or 3.8 × 1016 cm−3 in the 0.75Mn20V2Si steel and 9.2 × 1016 cm−3 cm−3 in the 0.78Mn20V2Mo2Si steel. Alloying of the vanadium carbide with molybdenum especially retarded the growth of particles at high aging temperatures. For example, in the case of overaging, the size of the (V, Mo)C carbide increased insignificantly in the Mn–V–Mo steels, while the size of some VC particles increased four to five times in the 0.75Mn20V2Si steel. It is seen from Table 1 that the maximum strengthening of the austenitic manganese steels with 0.2–0.4 mass% C was achieved after aging at 650 ◦ C for 10 h and was equal to 1040–1290 MPa, which was much larger than strengthening of analogous molybdenum-free steels. 3.2. The change of the SME magnitude in aged Mn–V–Mo steels Martensite phases were not detected in all the heat treatment conditions. However, tion of a few percent led to the formation of Fig. 1c and d presents the steel structure
the steels under plastic deformathe -martensite. after the ␥ →
750
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Fig. 1. Microstructure (TEM) of the 0.18Mn20V2Mo2Si (a and b) and 0.78Mn20V2Mo2Si (c and d) steels after aging in different conditions: (a) aging at 700 ◦ C, 3 h; (b) aging at 750 ◦ C, 12 h; (c) aging at 750 ◦ C, 12 h, deformation, dark-field image in the reflection of (V, Mo)C; (d) aging at 750 ◦ C, 12 h, deformation, dark-field image in the reflection of the -phase.
transformation, which was caused by 2–4% plastic deformation. The SME shows up in austenitic Mn–V–Mo steels as a partial restoration of the shape of a bend-strained sample under heating. The restoration is caused by the forward martensitic ␥ → transformation during cold deformation and the reverse shear → ␥ transformation during subsequent heating to a temperature above Af . According to the resistometric analysis, Af was 300 ◦ C in the alloys studied. It was found (Table 1) that the SME was enhanced with the steel aging temperature and its magnitude depended on the quantity of the -martensite. The absolute values of the reversible deformation under heating were as large as 1.7% and 0.9% in the steels, which were aged at 750 ◦ C for 12 h and 700 ◦ C for 3 h, respectively. The concentration of the -martensite in the deformed steel after destabilization aging at 750 ◦ C for 12 h was much larger than its concentration after stabilization aging at 700 ◦ C for 3 h. This
fact was explained by different stability (which was determined by changes in the coherence and the morphology of vanadium carbide precipitates) of the aged austenite with respect to the ␥ ↔ transformation and the saturation of the ␥-matrix with austenite-forming elements. The elastic field of the fine coherent (V, Mo)C carbides, which precipitated during aging, hampered the formation and the growth of the -martensite and stabilized the austenite phase. The loss of the coherence, the growth of (V, Mo)C particles and the increase in their volume fraction during aging at 750 ◦ C for 12 h decreased the stabilization effect of the carbides. Moreover, the martensite point was elevated due to the depletion of the matrix in carbon, vanadium and molybdenum after precipitation of carbides and the austenite was destabilized. As a result, the SME was enhanced after aging at 750 ◦ C for 12 h. It should be noted that 2% Mo added to the austenitic Mn–V steel stabilized it with respect to the formation of the strain-induced -martensite.
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In this case, the SME diminished in general. The thermal or thermomechanical treatment conditions need be optimized further so as to increase the SME in high-strength Mn–V–Mo steels. 4. Conclusion High-strength austenitic manganese steels showing SME and reaching the yield stress of 1100–1500 MPa were proposed. The proposed steels can be strengthened and their SME may be controlled by changing the temperature-and-time parameters of aging. In this case, the austenite is stabilized or destabilized with respect to the formation of the -martensite.
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Acknowledgment This study was supported by the RFBR (Project No. 06-03 32715). References [1] K. Enami, F. Nagasavo, S. Nenno, Scripta Metall. 9 (9) (1975) 941– 948. [2] K. Murakami, H. Otsuka, H.G. Suzuki, S. Matsuda, ICOMAT-96. Nara Imao Tamura Eds. (1987) 985. [3] V.V. Sagaradze, Ye.V. Beloserov, N.L. Pecherkina, et al., ICOMAT-2005, Abstract Book, Shanghai, 2005, p. 153. [4] V.V. Sagaradze, Ye.V. Beloserov, M.L. Mukhin, et al., Phys. Met. Metallogr. 101 (5) (2006) 506–512. [5] I.I. Kositsyna, V.V. Sagaradze, Metals 6 (2001) 65–74.