Microstructure, Phase Transformations, and Mechanical Properties of Shape Memory Fe-Mn-V-C Austenitic Steels
Institute of Metal Physics Ural Branch, RAS, S. Kovalevskoi St. 18, 620990 Ekaterinburg, Russia
*
Author to whom correspondence should be addressed.
Metals 2023, 13(2), 248; https://doi.org/10.3390/met13020248
Submission received: 28 November 2022
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Revised: 2 January 2023
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Accepted: 18 January 2023
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Published: 28 January 2023
(This article belongs to the Special Issue Microstructure and Mechanical Behaviour of Shape Memory Alloys)
Abstract
:The structure and mechanical properties of new dispersion-hardened Mn-(Cr)-Si-V-C steels with a shape memory effect (SME), which undergo strengthening due to the precipitation of VC carbides in a steel matrix, were analyzed. The stabilizing and the destabilizing carbide aging at different temperatures allows one to control and thus adjust the desired strength characteristics (σ0,2 = 250–880 MPa) and amount of reversible deformation (up to 2.7%). The recovery of the shape of the samples was carried out as a result of both the γ-ε transition in the course of heating after preliminary deformation in the initial austenitic state (i.e., due to the transformation of the γ to the ε phase) and the shear re-twinning of martensite in the martensitic ε phase (i.e., due to the occurrence of transformation of the ε to the twinned εtw phase).
1. Introduction
Metastable austenitic Fe-Mn-Si-based steels alloyed with a large amount of the manganese (up to 30 wt. %) and silicon (up to 6 wt. %) are increasingly used among alloys exhibiting the shape memory effect (SME). This effect is a result of the direct γ→ε and the reverse ε→γ martensitic transformation [1,2,3,4,5,6,7,8,9,10,11]. These elements (i.e., Mn and Si) reduce the energy of stacking faults (SF). Consequently, a large amount of ε-martensite develops at a small degree of plastic deformation which is accompanied by a significant magnitude of SME (usually with a reversible deformation of 2.0–3.5%) under the conditions of the ε→γ transformation during subsequent heating. One of the disadvantages of such alloys is their relatively low strength characteristics (usually no higher than 300–350 MPa in value). An increase in the strength characteristics of austenitic steels by heat treatment is achieved as a result of dispersion hardening with the precipitation of VC carbides of up to 10 nm in size [12]. The SME steels that we proposed to employ contained vanadium and 0.2–1.0 wt % carbon. These steels were first proposed in the year of 2005 [13,14]. Other methods of strengthening and-providing particles in SME steels are also seen in the literature [15,16,17,18,19], but their use in small quantities does not permit one to significantly increase hardening. The alloying of an SME steel with niobium and carbon (e.g., Fe-28Mn-6Si-5Cr-0.5NbC) [15,16] makes it possible to increase the yield strength through the precipitation of disperse NbC particles. The addition of titanium and carbon to steel also leads to a certain strengthening of alloys [17]. However, titanium often turns out to be in a bound state to form relatively large-size carbides, as in steel Kh18N10T (i.e., Cr18Ni10Ti), and only a small amount of titanium forms the nanosized strengthening particles TiC. Alloying of steels by a small amount of vanadium and nitrogen (up to 0.08 wt % N) leads to the precipitation of VN particles, which should also contribute to strengthening [18,19]. The greatest dispersion hardening of austenitic steels is usually provided by the precipitation of the disperse carbides VC [12]. One of the modern methods of synthesizing stainless steels with the special strength- and plastic properties consists in the obtaining of the samples using high-energy milling [20,21,22]. This work is devoted to the analysis of the SME Mn-Cr-Si-V-C steels developed by us [13,14,23,24,25,26,27,28] with different carbon and vanadium contents [13,14,23,24,25,26,27,28,29,30,31], in which the increase in strength is provided for by implementation of the carbide aging. Also discussed is the feasibility of obtaining a significant value of SME not only via realization of the phase transformation ε-γ, but also due to the shear re-twinning of ε-martensite.
2. Materials and Methods
Austenitic metastable SME steels, mainly containing (hereinafter, weight %) up to 18 Mn; (2–4)% Si; (1–14)% Cr; (0.2–1.0)% C; and (1–3)% V—were subjected to investigation. The steels were smelted in an induction furnace, homogenized and forged when heated to 1150–1200 °C. Then the workpieces were subjected to hot and the cold rolling in air. After the rolling and quenching of samples from 1150 °C, carbide aging was carried out at 600, 650, and 720 °C (the aging time was mainly 0.5, 3, and 12 h). In order to carry out the martensitic γ→ε transformation, heat-treated plate samples with a thickness of b = 1.38 mm, a width of 5.00 mm, and a length of 55.00 mm were bent at room temperature by 180 degrees around a cylinder with a diameter of D = 29 mm. The deformation on the outer stretched surface of the curved plate was determined by the well-known formula e = b/D. Under these conditions, the maximum deformation reached ~ 4.7%. Then, in a special dilatometer, at a heating rate of ~ 5 K/min, the temperature interval of the reverse martensitic transformation ε-γ and the value of the SME at the unbending of each sample, estimated by the divergence of the ends of the sample (in mm) during heating—were determined.
The structure of the alloys was analyzed using a JEM-200CX electron microscope at an accelerating voltage of 160 kV.
3. Experimental Results and Their Discussion
3.1. The Sparingly Alloyed Sme Steels Mn-Si-V-C with Strengthening by the Carbides VC
The austenitic Mn-Si-V-C steels [12,13,14,23,24,25,26,27,28], (namely, 0.4C-18Mn-2Si-2V, 0.2C-18Mn-2Si-1V, 0.7C-18Mn-2Si-3V, etc.) were proposed as sparingly alloyed dispersion-hardening SME steels with strengthening by VC carbides, which usually contained no more than 18–20 wt. % Mn. To obtain the required amount of the carbide phase VC, the concentration of vanadium in the steels was five times higher than the concentration of carbon. In order to reduce the energy of stacking faults and to provide for the formation of a larger amount of ε-martensite, steels were usually additionally alloyed with two or more wt % Si.
For the first time, the weak shape memory effect in steel with 18.5% manganese (18 Mn) was described in [5]. In the steel 18 Mn with low strength characteristics (σ0,2 = 242 MPa), pre-deformed by bending (to 2–3%), there is an insignificant effect of shape restoration (reversible deformation is ~ 0.2%). This result can be explained by the formation—during preliminary bending—of a small amount of oriented deformation-induced ε-martensite; this formation causes a realization of an insignificant value of the SME that takes place during the subsequent reverse transformation ε→γ in the heating process. In low-strength austenite, such deformation is mainly carried out by the sliding of perfect dislocations. In the alloys of more strength, with low energy of SF, for example, in aging-exhibiting SME steels [12,13,14,23,24,25,26,27,28], deformation is carried out due to the movement of partial dislocations with the formation of the stacking faults (SF) and crystals of ε-martensite, which contributes to an increase in the amount of deformation-induced ε-martensite and to the growth of the value of reversible deformation during the reverse ε→γ transformation.
It should be noted that VC particles, such as NbC carbides [32], when precipitating on stacking faults, promote the splitting (extending) of perfect dislocations into partial Shockley dislocations and entail the formation of a larger amount of deformation-induced ε-martensite with an HCP lattice. The alloying of the austenitic matrix of SME steels with vanadium and carbon complicates the formation of deformation-induced ε-martensite. So, in order to intensify the γ-ε transformation, it is advisable to remove V and C from the solid solution in (and via) the process of preliminary carbide aging. However, the disperse VC carbides formed during “low-temperature” aging (at 600, 650 °C) can hinder the nucleation and growth of ε-martensite crystals by their (VC) elastic stress fields [14,23]. For the first time, such stabilization of austenite was detected when the disperse γ′-phase Ni3Ti was precipitated in Fe-Ni-Ti alloys [33]. The precipitation of larger rarely located VC particles during ″high temperature″ aging (at 700–750 °C), on the contrary, does not restrain the formation of martensite [14,23]. In order to increase the values of reversible deformation, the destabilizing aging of austenite of SME steels was performed at elevated temperatures, for example, at 720 °C [23]. Figure 1a,b shows the structure of the 0.4C-18Mn-2Si-2V steel after destabilizing aging with the precipitation of enlarged VC carbides ((a) with an average size of 9 nm and a number density of 3 × 1016 cm-3) and after the martensitic γ→ε transformation with the formation of (b) 74% ε-martensite in the process of a small (4.7%) cold deformation.
The hardening by VC carbides (with a density ~ 1016 cm−3) was estimated according to the Orowan mechanism. We used the following reduced formula: σ0,2 = αDb/λ. Here, α = 0.9 is the coupling coefficient, D = 80,000 MPa is the shear modulus of austenite, and b = 0.25 nm is the Burgers vector of typical dislocation for fcc steels a/2<110>; λ = 50 nm is the distance between particles. Apparently, this distance is larger than the average particle size (9 nm). The calculated effect of dispersion hardening is 360 MPa, almost ~2.0 times less than the total strength, which is 728 MPa [12].
Figure 1c shows the microdiffraction pattern corresponding to the structure shown in (b), with the indexing of the reflections from the phases γ, ε, and VC. The subsequent heating of the steel with the structure of ε-martensite up to 400 °C, which ensures the realization of the reverse ε→γ martensitic transformation, leads to the manifestation of the SME with a reversible deformation of more than 2%.
Figure 2 shows the dilatometric heating curves of the steel 0.4C-18Mn-2Si-2V with the initial structure of the ε-martensite obtained during cold bending deformation (to 4.7%) after destabilizing aging at 720 °C for 0.5–12 h. An increase in the aging time at 720 °C contributes to the destabilization of austenite, which leads to the formation of a larger amount of deformation-induced ε-martensite and, accordingly, to an increase in the reversible deformation, when heated to 2% or more during the ε → γ transformation in the range of 100–350 °C. As follows from Figure 2, the shift of the ends of the unbending steel samples increases from 7 to 18 mm.
Figure 3 shows dilatometric curves obtained during the ε→γ transformation in the course of heating of the 0.7C-20Mn-2Si-3V (no. 1–3) and the 28Mn-6Si (no. 4) deformed (by ~4%) steel. It can be seen that the maximum reversible deformation (more than 2.2%) in steel with 0.7 wt % C is achieved after destabilizing aging at 720 °C (6 h), when rarely located enlarged VC carbides of about 9 nm in size are formed and the austenitic matrix is freed from V and C. The value of the SME in this case is comparable to the values of the reversible deformation of the well-known high-alloy SME steel 28Mn-6Si (compare Curves 1 and 4 in Figure 3). The formation of densely arranged more dispersed carbides (up to 5 nm in size [12]) during aging at lower temperatures of 600 and 650 °C prevents the growth of crystals of deformation-induced ε-martensite, which reduces the amount of reversible deformation during the ε→γ transformation to 0.3–0.8% (see Curves 2 and 3 in Figure 3).
Changing the regime of carbide aging allows not only to adjust the value of the SME of the studied steels, but also to significantly increase their strength characteristics. Table 1 presents the mechanical properties (σu is the conditional tensile strength limit (ultimate strength), σ0.2 is the yield strength, δ is the relative elongation, and ψ is the relative reduction) and values of reversible deformation of various Mn-Si-V-C SME steels. The aging of SME steels (with from 0.2 to 0.7% carbon) at 650 °C for 3–12 h allows the yield strength to be increasing to 585–883 MPa, while maintaining sufficiently high ductility characteristics, but the value of SME decreases to 0.8–1.4% (see Table 1). Reversible deformation above 2% can be maintained only after high-temperature destabilizing aging at 720 °C for 3 h (at the same time, the yield strength of SME steels with 0.4–0.7 wt % C remains at a high level (642–750 MPa).
Of considerable interest is the elucidation of the possibility of the realization of a noticeable SME in Mn-Si-V-C steels due to the deformation re-twinning of ε-martensite with its subsequent transformation into austenite upon heating [27]. After quenching from 1150 °C and destabilizing aging at 720 °C (6 h), more than 70% -vol. fraction martensite was obtained (on the stretched outer surface of the sample) as a result of cold bending deformation by ~−4%. Subsequent multiply -repeated deformation through the bending and unbending of the sample in the opposite directions caused an increase in the amount of deformation-induced ε-martensite. Five-fold bending—unbending treatment of the plate of 0.20C-20Mn-2Si-1V steel led to the formation of up to 95% of ε-martensite in the surface layer (while the total degree of deformation on the stretched surface of the sample was 24.5%). The heating of such a sample to 400 °C led to the development of εtw→γ transformation with a reversible deformation of 1.6%, which is less than after the initial deformation of austenite by 4% upon conventional treatment, where we have γ → ε → γ (2.6%).
3.2. Stainless SME Steels Strengthened by VC Carbides
Table 2 presents four groups of corrosion-resistant SME steels studied and shows the values of reversible deformation (e, %) after various treatments, including quenching and aging with the formation of VC carbides (see Figure 4a,b). The first group of steels includes corrosion-resistant alloys with 9–10 wt. % chromium, 4–5% silicon and different carbon content (from 0.06 to 0.45 wt. %). The second, third, and fourth groups of stainless steels contain mainly 13–14 wt. % chromium and differ in carbon content (0.20–0.25, 0.30–0.35, and 0.45 wt. %, respectively). The highest value of SME (2.02–2.70%) is achieved after destabilizing carbide aging at 720–750 °C only in the first group of steels with 9% chromium and 0.06–0.20% carbon (see Table 2). In silicon-containing steels with an increased amount of chromium (13–14%) and 0.2–0.45% carbon, the magnitude of reversible deformation in the destabilized aged state is significantly less (0.6–1.4%)—see Table 2. In steel without silicon (0.40C-14Cr-20Mn-2V), the SME value is the lowest—0.42%. In quenched steels that have not experienced destabilization during high-temperature aging, lesser high values of SME are observed. Aging at 650 °C slightly increases the value of SME, but noticeably to a lesser extent than that obtained in result of aging at 720–750 °C (see Table 2).
VC particles are nucleated on dislocations during high-temperature aging, becoming sources of splitting of perfect dislocations into partial ones with the formation of stacking faults. New small-sized VC particles are formed on the resulting stacking faults; these fine VC particles, with their elastic fields, contribute to further splitting of dislocations. In Figure 5a, one can see—in the dark-field image of the structure of the 0.20C-9Cr-15Mn-3Ni-4Si-1V steel aged at 720 °C for 6 h—the stacking faults on which VC particles were precipitated. Similarly, in [32], it was shown for the first time on the austenitic 0.1C-18Cr-10Ni-1Nb steel that its aging at 700 °C (5 h) reduces the energy of stacking faults and leads to the dissociation of perfect dislocations with the formation of stacking faults. An increase in the aging time at 700 °C also caused the formation of disperse carbides (NbC) on stacking faults [32].
Thus, an increase in the aging time of 0.20C-9Cr-15Mn-3Ni-4Si-1V steel to 6 h at 720 °C contributes to the change of the mechanism of carbides nucleation from homogeneous to heterogeneous. In this case, the partial destabilization of austenite occurs with respect to the martensitic γ-ε transformation (more ε-martensite is formed during the initial deformation) and, as a consequence, the magnitude of reversible deformation increases during the ε-γ transformation during heating (e = 2.7%).
As it follows from [34], a decrease in the quenching temperature from 1190 to 1100 °C leads to the partial preservation of insoluble chromium carbides Cr23C6 with a size of 100–400 nm in the stainless 0.30C-14Cr-15Mn-4Si-3Ni-1V and 0.30C-14Cr-15Mn-4Si-3Ni steels with shape memory effect. These carbides destabilize the austenite by reducing the chromium content in the austenite matrix and contribute to an increase in reversible deformation from 0.1–0.4% to 0.7–1.1%. An increase in the aging time at 720 °C to 12 h after quenching from 1100 °C (see Table 3) leads to the maximum magnitudes of SME (1.9%) in both steels strengthened with vanadium and chromium carbides. The article [11] shows the possibility of a sharp increase in the values of reversible deformation due to the enlargement of the initial austenitic grain. Thus, in the cast 19Mn-5.5Si-9Cr-4.5Ni SME steel, an increase in the size of the austenitic grain from 382 to 652 microns leads to a giant increase in reversible deformation from 5.4 to 7.7%. Although a similar situation occurs after the preliminary cold deformation of not 3.5%, but 10. It can be assumed that a significant enlargement of the austenite grain of polycrystalline SME steels at high temperatures is a powerful reserve for increasing the magnitude of SME.
4. Conclusions
1. The sparingly alloyed dispersion-hardening Mn-Si-V-C austenitic SME steels (e.g., 0.4 C-18Mn-2Si-2V, 0.2C-18Mn-2Si-1V, 0.7C-18Mn-2Si-3V, etc.) containing up to 20 wt. % Mn and 2–4% Si are proposed, in which the amount of reversible deformation (0.4–2.7%) during the ε → γ transformation is regulated as a result of preliminary stabilizing or destabilizing aging with the precipitation of VC carbides. As a result of carbide aging, there is also a significant increase in the yield strength in SME steels (up to 642–750 MPa) while maintaining their satisfactory ductility characteristics (a relative elongation of 14–39%).
2. In Mn-Si-V-C SME steels with the initial ε-martensitic structure, a reversible deformation of 1.6–1.7% can be obtained as a result of the shear re-twinning of ε-martensite during cold deformation (by 3–4%) and subsequent ε → γ transformation when heated to 400 °C.
3. In the Mn-Cr-Si-V-C corrosion-resistant SME steels with 13–14 wt. %Cr, significantly lower values of reversible deformation (mainly 1.0–1.4%) are observed than in SME steels with 9% chromium (up to 2.7%), even after low-temperature quenching and destabilizing aging at 720 °C (6 h) with the precipitation of VC and Cr23C6 carbides. The change of the mechanism of carbide precipitation from homogeneous to heterogeneous, with an increase in the temperature and aging time—contributes to an increase in the magnitude of reversible deformation in stainless SME steels.
Author Contributions
Conceptualization, V.S.; methodology, N.K. and S.A.; investigation, N.K., S.A. and Y.U.; writing—review and editing, V.S., N.K., S.A. and Y.U. All authors have read and agreed to the published version of the manuscript.
Funding
The research was carried out within the state assignment of Ministry of Science and Higher Education of the Russian Federation (theme “Structure” No. 122021000033-2).
Institutional Review Board Statement
Not applicable.
Informed Consent Statement
Not applicable.
Data Availability Statement
Not applicable.
Acknowledgments
Electron microscopic studies were performed in the Electron Microscopy Department of the Collective Usage Center at the M.N. Mikheev Institute of Metal Physics, Ural Division, Russian Academy of Sciences.
Conflicts of Interest
The authors declare no conflict of interest.
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Figure 1.
(a) Dark-field image of coarse carbides in the VC reflection and (b) crystals of ε—martensite in the 0.4C-Mn18-Si2-V2 type steel after aging (720 °C, 12 h) and subsequent cold deformation (ε~4.7%); (c)—SAED pattern corresponding to (b), from the phases γ, ε, and VC.
Figure 1.
(a) Dark-field image of coarse carbides in the VC reflection and (b) crystals of ε—martensite in the 0.4C-Mn18-Si2-V2 type steel after aging (720 °C, 12 h) and subsequent cold deformation (ε~4.7%); (c)—SAED pattern corresponding to (b), from the phases γ, ε, and VC.
Figure 2.
Variation of the recovery strain l (mm) of the 0.4C-18Mn-2Si-2V-steel sample in the process of the ε→γ transformation upon heating after preliminary quenching from 1150 °C and aging at 720 °C and cold deformation. Aging time: (1)—0; (2)—0.5; (3)—3; (4)—12; (5)—3 h (after the repeated quenching).
Figure 2.
Variation of the recovery strain l (mm) of the 0.4C-18Mn-2Si-2V-steel sample in the process of the ε→γ transformation upon heating after preliminary quenching from 1150 °C and aging at 720 °C and cold deformation. Aging time: (1)—0; (2)—0.5; (3)—3; (4)—12; (5)—3 h (after the repeated quenching).
Figure 3.
Change in reversible deformation (l, mm, or e, %) in the course of the ε → γ transformation at the heating of the steels 0.7C-20Mn-2Si-3V (curves 1–3) and 28Mn-6Si (curve 4). The preliminary treatment of the steel 0.7C-20Mn-2Si-V: quenching from 1150 °C; aging (for 6h) at (1) 720 °C; (2) 650 °C; and (3) 600 °C; cold deformation via bending e ~ 4%. The initial treatment of the 28Mn-6Si steel: quenching from 1150 °C (curve 4) and cold deformation e ~ 4%.
Figure 3.
Change in reversible deformation (l, mm, or e, %) in the course of the ε → γ transformation at the heating of the steels 0.7C-20Mn-2Si-3V (curves 1–3) and 28Mn-6Si (curve 4). The preliminary treatment of the steel 0.7C-20Mn-2Si-V: quenching from 1150 °C; aging (for 6h) at (1) 720 °C; (2) 650 °C; and (3) 600 °C; cold deformation via bending e ~ 4%. The initial treatment of the 28Mn-6Si steel: quenching from 1150 °C (curve 4) and cold deformation e ~ 4%.
Figure 4.
(a) SAED pattern and (b) dark-field image of the 0.30C-13Cr-15Mn-3Si-1V steel in the reflection of VC carbide. Treatment: quenching (Q) from 1100 °C; aging (A) at 720 °C, for 6 h; cold deformation (CD) ~4% (γ→ε transformation), and heating (H) to 500°C (ε→γ transformation). The zone axis [001]γ,VC.
Figure 4.
(a) SAED pattern and (b) dark-field image of the 0.30C-13Cr-15Mn-3Si-1V steel in the reflection of VC carbide. Treatment: quenching (Q) from 1100 °C; aging (A) at 720 °C, for 6 h; cold deformation (CD) ~4% (γ→ε transformation), and heating (H) to 500°C (ε→γ transformation). The zone axis [001]γ,VC.
Figure 5.
(a) Dark-field image of carbides in the reflection of VC carbide and (b) SAED pattern from the 0.20C-9Cr-15Mn-3Ni-4Si-1V steel. Treatment: quenching from 1100 °C, aging at 720 °C (6 h) and γ → ε → γ transformation. The zone axis [012]γ,VC.
Figure 5.
(a) Dark-field image of carbides in the reflection of VC carbide and (b) SAED pattern from the 0.20C-9Cr-15Mn-3Ni-4Si-1V steel. Treatment: quenching from 1100 °C, aging at 720 °C (6 h) and γ → ε → γ transformation. The zone axis [012]γ,VC.
Table 1.
Mechanical properties and reversible deformation eSME in Mn-Si-V-C steels.
| Type of Steel | Treatment | Ultimate Strength σu, MPa | Yield Strength σ0.2, MPa | Relative Elongation δ, % | Relative Reduction ψ, % | eSME, % |
|---|---|---|---|---|---|---|
| 0.2C-20Mn-2Si-2V | Quenching (Q) from 1150 °C (0.5h) in water | 984 | 257 | 32.2 | 32.0 | 1.13 |
| Q + 650 °C, 3 h | 1096 | 424 | 39.7 | 38.0 | 1.36 | |
| Q + 650 °C, 12 h | 1179 | 585 | 34.2 | 43.2 | - - | |
| Q + 720 °C, 0.5 h | 1090 | 349 | 38.3 | 37.0 | ||
| Q + 720 °C, 3 h | 1102 | 528 | 21.0 | 26.0 | 1.7 | |
| 3 + 720 °C, 6 h | 1097 | 497 | 19.1 | 21.1 | - | |
| 0.45C-20Mn-2Si-2V | Q + 650 °C, 12 h | 1281 | 805 | 17.3 | 20.4 | 1.4 |
| Q + 720 °C, 3 h | 1312 | 728 | 19.4 | 22.7 | 2.1 | |
| 0.7C-20Mn-2Si-3V | Q-1150 °C, 0.5 h, water | 989 | 547 | 11.7 | 18.4 | - |
| Q + 650 °C, 12 h | 1319 | 883 | 13.3 | 21.0 | 0.8 | |
| Q + 720 °C, 3 h | 1168 | 642 | 14.6 | 22.7 | 2.3 | |
| 1.0C-20Mn-2Si-4V | Q-1150 °C, 0.5 h, water | 1110 | 641 | 16.6 | 21.7 | - |
| Q + 650 °C, 3 h | 1217 | 819 | 14.3 | 18.6 | ||
| Q + 650 °C, 12 h | 1267 | 825 | 11.3 | 11.6 | - - | |
| Q + 720 °C, 3 h | 1120 | 665 | 10.0 | 16.0 | ||
| 0.45C-20Mn-2Si-2V-2W | Q-1150 °C, 0.5 h, water | 1046 | 517 | 31.4 | 34.2 | 1.0 |
| Q + 650 °C, 3 h | 1157 | 632 | 39.5 | 36.5 | 2.1 | |
| Q + 650 °C, 12 h | 1319 | 965 | 14.1 | 20.0 | - | |
| Q + 720 °C, 0.5 h | 1157 | 608 | 37.0 | 36.2 | - | |
| Q + 720 °C, 3 h | 1212 | 750 | 23.0 | 35.5 | 2.4 | |
| Q + 720 °C, 6 h | 1234 | 704 | 21.6 | 29.1 | - |
Table 2.
Magnitudes (in %) of the SME in various stainless steels after different heat treatments [28].
Table 2.
Magnitudes (in %) of the SME in various stainless steels after different heat treatments [28].
| no./no.pla | Type of Steel | e, % (Treatment: * Q + * A - 720, 750 °C) | e, % (Treatment: Q + A - 650 °C, 6 h) |
|---|---|---|---|
| Group 1 | |||
| 1/47 | 0.06C-9Cr-15Mn-5Ni-5Si-0.5Nb | 2.02 (720 °C, 6 h) | 1.85 |
| 2/52 | 0.06C-9Cr-15Mn-5Ni-5Si-1V | 2.50 (750 °C, 10 h) | |
| 3/56 | 0.20C-9Cr-15Mn-2Ni-5Si-1V | 2.45 (750 °C, 10 h) | |
| 4/59 | 0.20C-9Cr-15Mn-3Ni-4Si-1V | 2.70 (720 °C, 6 h) | |
| 5/48 | 0.25C-9Cr-15Mn-5Ni-3.5Si-1.2V | 0.78 (720 °C, 6 h) | 0.66 |
| 6/49 | 0.25C-9Cr-15Mn-3.5Si-1.2V | 1.25 (720 °C, 6 h) | 0.74 |
| 7/33 | 0.30C-10Cr-16Mn-2Si-1.5V | 1.20 (720 °C, 6 h) | |
| 44 | 0.45C-9Cr-12Mn-2Si-2V | 1.13 (720 °C, 6 h) | |
| Group 2 | |||
| 8/58 | 0.20C-14Cr-15Mn-1Ni-4Si-1V | 1.40 (720 °C, 6 h) | |
| 9/54 | 0.20C-13Cr-17Mn-5Si-1V | 0.74 (750 °C, 10 h) | |
| 10/51 | 0.25C-13Cr-14Mn-3Si-1.2V | 0.97 (720 °C, 6 h) | 0.44 |
| Group 3 | |||
| 11/60 | 0.30C-13Cr-15Mn-3Si-1V | 1.10 (720 °C, 6 h) | |
| 12/31 | 0.30C-14Cr-16Mn-2Si-1.5V | 0.60 (720 °C, 6 h) | |
| 13/34 | 0.30C-18Cr-16Mn-2Si-1.5V | fracture of sample | |
| 14/57 | 0.35C-14Cr-15Mn-3Ni-4Si-1V | 1.3 (720 °C, 3 h) | |
| Group 4 | |||
| 15/46 | 0.45C-12Cr-10Mn | 0.45 (720 °C, 6 h) | |
| 16/36 | 0.40C-14Cr-20Mn-2V | 0.42 (720 °C, 6 h) | |
| 17/35 | 0.45C-14Cr-16Mn-2Si-2V | 1.14 (720 °C, 6 h) | |
| 18/45 | 0.45C-12Cr-12Mn-2Si-2V | 1.20 (720 °C, 6 h) | |
| 19/50 | 0.45C-13Cr-14Mn-3.5Si-2V | 1.35 (720 °C, 6 h) | 0.26 |
* Q—quenching from 1150 °C, * A—aging.
Table 3.
Reversible deformation e in the stainless SME steels quenched from 1100 °C dependent on the aging time at 720 °C.
Table 3.
Reversible deformation e in the stainless SME steels quenched from 1100 °C dependent on the aging time at 720 °C.
| No. of Treat. | Type of Steel | Aging Time at 720 °C, Hours | Initial Deformation eo, % | Reversible Deformation e, % |
|---|---|---|---|---|
| 1 | 0.30C-14Cr-15Mn-4Si-3Ni1V | 0 | 3.5 | 0.7 |
| 2 | ---- // ---- | 1 | 3.2 | 0.9 |
| 3 | ---- // ---- | 3 | 3.3 | 1.2 |
| 4 | ---- // ---- | 6 | 3.4 | 1.3 |
| 5 | ---- // ---- | 12 | 3.3 | 1.9 |
| 6 | 0.30C-14Cr-15Mn-4Si-3Ni | 0 | 3.4 | 0.8 |
| 7 | ---- // ---- | 1 | 3.4 | 1.1 |
| 8 | ---- // ---- | 3 | 3.4 | 1.1 |
| 9 | ---- // ---- | 6 | 3.4 | 1.2 |
| 10 | ---- // ---- | 6 | 3.4 | 1.9 |
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Sagaradze, V.; Afanasiev, S.; Kataeva, N.; Ustyugov, Y. Microstructure, Phase Transformations, and Mechanical Properties of Shape Memory Fe-Mn-V-C Austenitic Steels. Metals 2023, 13, 248. https://doi.org/10.3390/met13020248
AMA Style
Sagaradze V, Afanasiev S, Kataeva N, Ustyugov Y. Microstructure, Phase Transformations, and Mechanical Properties of Shape Memory Fe-Mn-V-C Austenitic Steels. Metals. 2023; 13(2):248. https://doi.org/10.3390/met13020248
Chicago/Turabian StyleSagaradze, Victor, Sergey Afanasiev, Natalya Kataeva, and Yurii Ustyugov. 2023. "Microstructure, Phase Transformations, and Mechanical Properties of Shape Memory Fe-Mn-V-C Austenitic Steels" Metals 13, no. 2: 248. https://doi.org/10.3390/met13020248
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