Advertisement

Springer Nature is making SARS-CoV-2 and COVID-19 research free. View research | View latest news | Sign up for updates

Development of an Advanced Ultrahigh Strength TRIP Steel and Evaluation of Its Unique Strain Hardening Behavior

Abstract

A novel ferrite-53% austenite medium-Mn transformation induced plasticity (TRIP) steel with extraordinary combination of 1420 MPa tensile strength and 27% elongation was produced by a simple designed thermomechanical treatment. The deformation and strain hardening behavior of this material were studied in detail using field emission scanning electron microscopy and electron backscatter diffraction examinations of different strained specimens. Exclusive strain hardening behavior observed in this new TRIP steel was found to be related to the severe grain refinement to less than 200 nm and austenite grain disintegration within the first 1% of applied strain. Moreover, it was shown that the deformation of highly unstable retained austenite benefits the mechanical properties, since it could promote the severe grain refinement and subsequent uniform strain distribution within the microstructure of the TRIP steel. Finally, based on the strain hardening behavior analysis a new and simple measurable criterion was introduced as a reliable replacement of traditional austenite stability parameter or the K value.

Graphical Abstract

Introduction

Transformation induced plasticity (TRIP) steels show a significant combination of strength (1.0–1.6 GPa) and ductility (~ 20%) that is attributed to the value of the retained austenite, its stability, and its transformation to martensite during straining [1, 2]. In this context, modifying the composition [3, 4] and production treatment [5, 6] lead to a noticeable advancement in TRIP steels, from the viewpoint of deformation and mechanical properties. The introduction of medium Mn TRIP (MM-TRIP) steels [7] is considered a significant advancement in this context. This kind of steel is composed of 3%–10% Mn. The superior mechanical properties of this steel, along with its simple and low-cost production process, have made it a promising candidate for automobile structural components. An ultrafine grained microstructure and a high amount of retained austenite (10%–30%) are the main microstructural features of these steels [8]. Moreover, transformation of a high proportion of retained austenite to martensite was supposed to be the parameter influencing the ductility and elongation observed in these materials [2, 9]. However, the initial morphology of the incorporated ferrite and austenite or martensite phases within the microstructure affects the austenite-to-martensite transformation behavior [10]. This could also affect the strain hardening behavior of MM-TRIP steels. Rapid work hardening behavior is another characteristic of these materials; although some of them show static and/or dynamic strain aging, during deformation [11].

Numerous investigations have been conducted to improve the mechanical properties and strain hardening behavior of MM-TRIP steels. By modifying the chemical composition, Song et al. [4] developed ultra-high-strength, lightweight medium-Mn duplex steels. The improvement of strain hardening as well as yield strength observed in this steel was attributed to the effect of Cu addition, which raised the austenite volume fraction. Xu et al. [5] investigated the effect of one-step and two-step annealing processes on the deformation behavior and mechanical properties of MM-TRIP steels. They showed that the sample annealed in two steps is involved with the higher volume fraction of austenite-stimulated martensitic transformation (even at a low strain level of deformation), compared with the sample annealed in one step. This led to a rapid increase of the flow stress (high work hardening rate) at the early stages of deformation, and so to the significant increase of the tensile strength in this sample (i.e., comparing the two-step and one-step annealing processes). Yang et al. [6] studied the effects of intercritical annealing temperature and duration on the microstructures and mechanical properties of 7Mn–2Al–0.3C steel. They showed that the size and volume fraction of the granular grains of austenite and ferrite increased with increase of the intercritical annealing temperature and duration. They tried to interpret the work hardening rate of their medium-Mn steel; however, they did not determine the exact mechanisms involved with deformation in the different stages of the work hardening processes.

Given the literature review performed on the MM-TRIP steels, it can be seen that most previous investigations have been devoted to further improvement of the mechanical properties of these steels. To further enhance these achievements, it would be valuable to understand the mechanisms and parameters involved in their noticeable mechanical behavior. The current study aims to provide deep understanding of the deformation behavior of MM-TRIP steels. In this study, 4 different MM-TRIP steels were produced and their tensile mechanical properties were evaluated at room temperature. These examinations were then accompanied by microstructural characterization employing field emission scanning electron microscopy (FESEM) and electron backscatter diffraction (EBSD) examinations. The results revealed an unexpected strain hardening behavior in the steels studied which was analyzed and discussed based on the detailed microstructural characterizations.

Materials and Methods

The chemical composition of the steel produced was Fe–3.7Mn–0.2C. The alloy was made by arc melting and subsequent hot rolling at 1100 °C and air cooling to room temperature. This process was followed by 50% cold rolling, annealing at 700 °C and water quenching to room temperature. The heat treatments were done in muffle furnace. Tensile tests specimen were prepared according to Fig. 1 and the tests were performed following the ASTM A370 standard test [12] with a cross-head speed of 1 mm/min. The EBSD (Hikari EBSD Camera, EDAX) analysis was conducted to characterize microstructural and phase evolution during the tensile test. The step size for EBSD measurements was 40 nm. The EBSD measurements were carried out in the middle thickness section of the samples in a plane perpendicular to the normal direction of the steel sheet as presented in the previous paper of the author [13]. Specimen preparation was done by mechanical and colloidal silica polishing. To enhance the EBSD image quality, an ion milling process (Helios Nano Lab, FEI) was conducted on the surface of the samples. The von Misses equivalent strain (εeq) was calculated from current thickness (t) values and the initial thickness (t0) of the sheet using εeq= − 2 ln(t/t0) [14].

Fig. 1
figure1

Engineering drawing of tensile specimen (all dimensions are in mm)

In order to evaluate the effect of microstructure on the mechanical properties of the steel produced, different heat treatment cycles were designed and applied (Table 1). Heat treatment cycles were designed based on the phase diagram (Fig. 2) of the steel produced. Considering this diagram, 2 treatments of full austenitization (FA) and partial austenitization (PA) were supposed to be applied and then in each case, one part of specimens was treated by second short annealing process at 670 °C (2-step annealed) and one part of them was tensile tested without further treatments (1-step annealed). Partial austenitization treatment at the 670 °C, intercritical region, was done to promote austenite reverted transformation (ART) to achieve a microstructure composed of ferrite and austenite phases [15, 16]. Given these treatments, the specimens were designated as presented in Table 1.

Table 1 Thermomechanical treatments used for production of the MM-TRIP steels
Fig. 2
figure2

Phase diagram of the MM-TRIP steel produced, calculated by using Thermo-calc (2017—TCFE7 databases)

The temperatures A1 and A3 of the sample austenite transformation are measured on the calculated phase diagram to be 593 °C and 742 °C, respectively.

Results and Discussion

Tensile Behavior and Strain Hardening Analysis

The stress–strain curves and mechanical properties of the produced MM-TRIP steels are presented in Fig. 3 and Table 2. As seen the 1 step-FA MM-TRIP steel shows brittle fracture behavior that may be representative for the brittle fracture of martensitic microstructure. In contrast, the 2 step-FA MM-TRIP steel shows interesting elongation of ~ 40% along with the high tensile strength of ~ 1000 MPa. On the other hand, the 1 step-PA and 2 step-PA MM-TRIP steels show rather similar tensile and ductility behaviors.

Fig. 3
figure3

Stress–strain curves of the steels studied

Table 2 Tensile mechanical properties of the steels studied

Comparing the absorbed fracture energy (Table 2), the area under the stress–strain curve, it is obvious that the 2 step-FA steel is involved with the maximum absorbed fracture energy and the 1 step-PA steel is in the second place. Beyond the high value of the fracture energy of the 2 step-FA steel, its stress–strain curve is involved with the noticeable serrated behavior, attributed to the dynamic strain aging [17, 18]. This serration behavior is not appropriate for forming processes and it could be a negative point for the 2 step-FA steel. So, owing to the significant tensile strength, elongation and absorbed fracture energy of the 1 step-PA MM-TRIP steel, it was selected for further deformation and strain hardening analysis. This steel exhibits an extraordinary combination of ~ 1420 MPa tensile strength and 27% elongation. As seen, the stress–strain curve of this steel shows a step-like behavior. This noticeable step-like behavior cannot be seen in the ordinary TRIP steels [3, 19,20,21]; one example related to deformation behavior of ordinary TRIP steels is presented in Fig. 4b. To elucidate the origin of this step-like behavior, as well as the overall deformation behavior of the studied TRIP steel, the strain hardening rate (SHR) of this steel was calculated and is also presented in Fig. 4. As seen in this figure, the strain hardening curve generally involves four stages of deformation during tensile straining. These stages are discussed below.

Fig. 4
figure4

a Stress–strain curves and SHR diagrams of a of 1 step-PA MM-TRIP [3] and b O-TRIP steel

Stage I: This stage involves decreasing SHR. In other words, the softening rate is higher than the hardening rate. This was supposed to be associated with the plastic deformation of ferrite as well as the elastic–plastic deformation of martensite and phase transformation. In this stage, the dislocation pile-ups were annihilated by the initiation of martensite plastic deformation. Moreover, the strain-induced transformation of austenite to martensite would accelerate the strain relaxation within the microstructure [22], leading to subsequent decrease of the SHR. In the following sections, this subject will be further explained.

Stage II: This stage involves constant SHR behavior. In other words, equilibrium is achieved between the hardening and softening rates, or the strain relaxation is compensated by the hardening. It was said that in this stage, plastic deformation of ferrite and elastic–plastic deformation of martensite happens. The interaction of different slip systems and creation of dislocation tangles, as well as dislocation annihilation by relaxation mechanisms [22], was supposed to result in a fairly constant change in the SHR. Moreover, in TRIP steels, the strain-induced martensite phase would act as obstacles against the dislocation motion. This could cause an increase in the SHR in stage II.

Stage III: This stage involves an increase in SHR value. As seen in Fig. 4b, the effect of this stage is not very noticeable in the O-TRIP steel. However, in the 1 step-PA MM-TRIP steel, its variation is very sharp. It is expected that the hardening mechanisms activated in the previous stages should most affect the probable relaxation mechanisms. However, other hardening mechanisms might also have been activated.

Stage IV: This stage involves a slight decrease of the SHR value. This could be associated with the plastic deformation of the ferrite and martensite phases. The activation of different dislocation annihilation mechanisms, such as the plastic deformation of martensite [23], substructure formation [22], and void creation [24] are some of the causes of decreasing SHR behavior in this stage.

To clarify the mechanisms involved in each of the stages mentioned, the 1 step-PA MM-TRIP steel was microstructurally examined using EBSD analysis under different strain conditions. First, the SFE of the 1 step-PA MM-TRIP steel was calculated as follows. The SFE was calculated using a thermodynamic model [25,26,27] that is reliable for use with Fe–Mn–C steels, as follows [28]:

$$\begin{aligned} & \gamma = 2\rho \Delta G^{\gamma \to \varepsilon } + 2\sigma^{\gamma /\varepsilon } + 2\rho \Delta G_{ex} , \\ & \Delta G_{ex} = 170.06{ \exp }\left( { - \frac{{d_{\gamma } }}{18.55}} \right), \\ \end{aligned}$$
(1)

where ρ, \(\Delta G^{\gamma \to \varepsilon }\), \(\sigma^{\gamma /\varepsilon }\), and \(\Delta G_{ex}\) show a molar surface density on the {111} planes, free energy for \(\gamma /\varepsilon\) transformation in relation to the chemical composition, \(\gamma /\varepsilon\) interfacial energy, and excess SFE due to the effect of the grain size (d). Here, \(d_{\gamma }\) = austenite grain size (μm) = 1 μm → \(\Delta G_{ex} = 160.7 \to 2\rho \Delta G_{ex} = 2*2.53*10^{ - 9} *160.7 = 8.1 \,\,{\text{mJ/m}}^{2}\).

In this specimen, of which the ferrite fraction is about 47%, the austenite fraction contains 1.6 and 5.5 mol fraction of C and Mn, respectively, calculated by Thermo-calc. Thus, the SFE of the austenite of the specimen is higher than that of the average steel composition (16 mJ/m2). Based upon its low SFE level (~ 16 mJ/m2), the austenite shows planar slip characteristics during deformation; instead of deformation twinning such as frequently occurs in high-Mn steels [29,30,31,32]. Moreover, it was shown that with low SFEs, Taylor lattice formed by cross-slipping of planar slips occurring in different slip planes [29, 31]. The parameters used for the SFE calculation are listed in Table 3.

Table 3 Parameters for the calculation of SFE [33]

Microstructural Examination of Deformation Behavior at Different Strain Values

The microstructure of 1 step-PA MM-TRIP steel was examined when undeformed, under 0.008, 0.015, and 0.061 strain (equivalent strain), and under fracture strain (equivalent strain of 0.24). The strain value of 0.008 placed it in the stage I region. The strain value of 0.015 placed it in stage II and the strain values of 0.061 and 0.24 placed it in stage IV of the SHR diagram.

The IQ and Phase maps of these specimens are presented in Fig. 5. It is seen that the initial undeformed microstructure (Fig. 5a) is involved with about 53% retained austenite and ferrite. Moreover, it can be seen that in all of the strain conditions, the grains are approximately equiaxed and up to 0.061 strain does not show noticeable stretching behavior. To elucidate the grain size variation and distribution, grain size histograms were calculated and are presented in Fig. 6. As seen in the undeformed specimen, the diameter of the grains is mostly < 1 µm and about 20% of them are in the vicinity of 200 nm. As deformation is applied to the specimen, the grain diameter is seen to shift to about 200 nm, such that by applying an equivalent strain of 0.008, the fraction of these fine grains increases to 40% from the initial undeformed condition. By further increase in strain, the fine 200 nm grains increase by about 10% and after that; they did not change noticeably all the way up to the fracture strain.

Fig. 5
figure5

IQ and Phase maps related to different equivalent strain values of a 0, b 0.008, c 0.015, d 0.061, and e 0.24. The red regions indicate the retained austenite

Fig. 6
figure6

Grain size distribution related to different strains

From measurement of the retained austenite employing EBSD analysis, it was determined (Fig. 7) that the retained austenite (RA) decreased by about 50% in the first 0.008 (stage I) of applied strain and then decreased to about 90%, at the strain value of 0.015 (stage II). Subsequent straining did not change the retained austenite values significantly. Considering these results, it is seen that in the first 0.008 of strain, 50% of the RA is transformed to the martensite phase. At the same time, the grain fraction with a diameter of less than 200 nm increases to about 40%. Hence, it seems that the main mechanisms involved with the rapid decrease of SHR in stage I was RA transformation and substructure formation. It has been shown [34,35,36,37,38] that the center part of an austenite grain has the lowest stability and transforms to martensite at the early stages of deformation. Therefore, the initial grains of the retained austenite are disintegrated to form martensite and stable austenite. Therefore, it can be said that the observed severe grain refinement, with strain, is tightly related to the RA transformation. This behavior can be called retained austenite disintegration (RAD).

Fig. 7
figure7

Retained austenite distribution related to different strains

At the strain value of 0.015 (stage II of the SHR diagram), it was seen that the RA decreased by about 40% more than in the previous stage. However, the grain refinement was not significant, in this stage. It was found that in the previous stage, strain hardening was dominated by the softening mechanisms of austenite transformation and substructure formation. However, in that stage, the microstructure was refined significantly and reached a constant value of about 200 nm. This could be because of the low SFE value of 16 mJ/m2 calculated for this steel. The low SFE promotes planar slip deformation mechanisms along with the γ → ε phase transformation [39]. Synergistic effects of these two mechanisms may lead to the activation multiple planar slip and subsequent formation of Taylor Lattice at low strains. [40]. Taylor lattice originates from the planar slips occurring on two slip planes that crossed each other to form dislocation substructures. A Taylor lattice is a low-energy dislocation structure equivalent to a dislocation cell in a cell forming material [41].

Cuddy et al. [42] have also shown that the dislocation substructure size decreases with tensile strain up to the final cell size of 200–300 nm, in ferrite grains.

A noticeable increase of the length fraction of low angle grain boundaries (LAGB) emphasizes the high value of substructures created within stages I and II (Fig. 8). This ultra-fine substructure could act as a barrier against dislocation movement and would increase the SHR. Hence, it can be said that the ultra-fine substructure created, along with the effect of the newly created martensite grains, could be the cause of the sharp increase of the SHR in stage III. In other words, it is reasonable to suppose that in stage III, the high value of the strain-induced martensite phase created by the TRIP effect, along with the high value of the created substructures, act as significant barriers to dislocation motion. Therefore, rapid increase of SHR is expected to happen, as seen in Fig. 4a.

Fig. 8
figure8

Evolution of low angle grain boundaries (LAGB) and high angle grain boundaries (HAGB) with strain, in 1 step-PA MM-TRIP steel

Unlike in stages I and II, stage III cannot be seen in the O-TRIP steel. Considering the mentioned observations and inferences, it seems that two parameters were involved in this behavior:

  • A lower initial RA fraction and lower subsequent strain-induced martensite within the microstructure: The initial RA was found to be about 10 and 53%, in the O-TRIP and 1 step-PA MM-TRIP steels, respectively.

  • Considering the mentioned RAD phenomenon, the lower the initial RA fraction, the lower the grain refinement caused by strain: Lower grain refinement is expected in the O-TRIP than in the 1 step-PA MM-TRIP steel.

In stage IV, plastic deformation of the martensite phase could decrease the stored strain energy within the microstructure and decrease the SHR value. Moreover, it was seen that in this deformation stage, the LAGB did not change significantly (Fig. 8).

In order to obtain a better understanding of the deformation behavior of the material, IPF maps of the ferrite grains at different strained specimens were calculated and presented in Fig. 9. As seen the undeformed specimen is involved with tightly clustered points corresponding to specific orientations occupied by the individual grains. Increasing the strain led to increase of the orientation spread such that in the strain of 0.015 a total random orientation distribution can be observed; this is beneficial to improve the formability [41]. This kind of behavior is different to the common behavior of DP and TRIP steels [22] in which a preferred orientation created within the microstructure by strain.

Fig. 9
figure9

IPF images of the ferrite in the [001] direction (perpendicular to the rolling direction). These images are related to different equivalent strain values of a 0, b 0.008, c 0.015, d 0.061, and e 0.24

Total random orientation distribution created in 1 step-PA MM-TRIP steel could be due to the extreme grain refinement happened in the microstructure. Grain refinement could affect this behavior by two mechanisms:

  • Lack of shear band creation

Grain refinement impedes the creation and growth of shear bands; through the high density of grain boundaries serving as effective barriers against shear band evolution. So, strain localization and necking prohibited [43]. Lack of shear bands within the microstructure is clear within the microstructure of 1 step-PA MM-TRIP steel (Fig. 5). Therefore, in the absence of strain localization, strain distribution would be better within the material.

  • The high resistance to deformation

The lower the grain size, the higher the grain boundaries acting as barriers to the dislocation movement. So, stretching of the grains that could lead to creation of the preferred orientation would be less probable. The equiaxed grains observed at high strains in Fig. 5, support this statement.

Evaluation of Retained Austenite Stability

Hu et al. [44] showed that the austenite stability in medium-Mn steels is much higher than that in the conventional TRIP steels. In the present work, the parameter introducing the austenite stability (the K value in Eq. 2) was found to be 30.4, which is much higher than the value they found for Fe–5Mn–0.1C (K = 12.9). They mentioned that the low K value is beneficial in MM-TRIP steels because this would lead to homogenous strain distribution within the microstructure. However, in the current work, despite the high calculated value of K, homogenous strain distribution was observed within the material such that the necking phenomena did not occur in the specimen (Fig. 10). So, in the MM-TRIP steel studied, it can be said that deformation of highly unstable retained austenite benefits the mechanical properties, since it could promote the grain refinement significantly. The created ultrafine grained microstructure could improve the mechanical properties from the view point of homogenous strain distribution, elongation and strength.

Fig. 10
figure10

Cross section of the tensile strained and fractured specimen. No thickness necking can be observed in this specimen. Fracture surface are indicted by black arrow

$$K = \left[ {\ln \left( {\gamma_{f} } \right) - \ln \left( {\gamma_{0} } \right)} \right]/\varepsilon .$$
(2)

In this equation, γf stands for the retained austenite at a strain of ε, and γ0 stands for the retained austenite at ε = 0.

Given the results presented in Sect. 3.2 and Figs. 4a and 7, it can be said that the most of the retained austenite has been transformed before the strain value of εc (shown in Fig. 11), and after this point the transformed martensite and subsequent created high density of boundaries within the microstructure, led to the increase of strain hardening rate (stage III of SHR diagram). It is expected that the more the volume fraction of created martensite or the finer the created microstructure, the sharper the increase of the SHR in stage III. So, the strain εc could be assigned as a criterion for austenite stability. Considering this parameter, the hard work of calculating the traditional austenite stability parameter [44, 45], by measurement of the variation of the retained austenite with strain can be avoided.

Fig. 11
figure11

Variation of SHR with strain in 1 step-PA MM-TRIP steel and critical strain of εc in which the most of the retained austenite transformed to martensite

Conclusions

In this study, Fe–3.7Mn–0.2C TRIP steel was produced by arc melting and subsequent thermomechanical treatment. The deformation mechanism of this steel was studied in detail by analyzing the strain hardening rate and doing EBSD microstructural examinations. Overall, the following conclusions were drawn:

  • A novel ferrite–austenite MM-TRIP steel with attractive mechanical properties including about 1420 MPa tensile strength and 27% elongation was produced by thermomechanical treatment.

  • Different SHR behavior was observed in the MM-TRIP steel studied compared with the ordinary traditional TRIP steel in which the SHR seemed to decrease in all the stages of tensile straining.

  • Severe grain refinement (less than 200 nm) in the initial 1% of the applied strain was found to be responsible for the sharp increase of the SHR in the MM-TRIP steel.

  • Austenite grain disintegration appears to be the main phenomenon involved with the observed severe grain refinement.

  • Despite the high value of stability coefficient (K value) of the MM-TRIP steel, homogeneous strain distribution was observed in this material and it was attributed to the ultrafine grained microstructure created at low strains.

  • A new criterion (εc) was introduced as reliable replacement of traditional austenite stability parameter or the K value. Unlike the K value, the εc can simply be calculated by tensile testing.

References

  1. 1.

    H. Lee et al., Acta Mater. 147, 247–260 (2018)

  2. 2.

    D.W. Suh, S.J. Kim, Scr. Mater 126, 63–67 (2017)

  3. 3.

    N. Saeidi, M. Raeissi, M.M. Abdar, H. Vaghei, Mater. Sci. Eng. A 702, 1–456 (2017)

  4. 4.

    H. Song et al., Acta Mater. 135, 215–225 (2017)

  5. 5.

    Y. Bo Xu et al., Mater. Sci. Eng. A 688, 40–55 (2017)

  6. 6.

    F. Yang, H. Luo, C. Hu, E. Pu, H. Dong, Mater. Sci. Eng. A 685, 115–122 (2017)

  7. 7.

    Y.-K. Lee, J. Han, Mater. Sci. Technol. 31, 843–856 (2015)

  8. 8.

    Y. Ma, Mater. Sci. Technol. (United Kingdom) 33, 1713–1727 (2017)

  9. 9.

    J. Long Zhao, Y. Xi, W. Shi, L. Li, J. Iron Steel Res. Int. 19, 57–62 (2012)

  10. 10.

    J. Han, S.-J. Lee, J.-G. Jung, Y.-K. Lee, Acta Mater. 78, 369–377 (2014)

  11. 11.

    D.M. Field, D.C. Van Aken, Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 49, 1152–1166 (2018)

  12. 12.

    ASTM A370-12, Standard Test Methods and Definitions for Mechanical Testing of Steel Products (American Standards for Testing of Materials, 2001)

  13. 13.

    N. Saeidi, F. Ashrafizadeh, B. Niroumand, M.R. Forouzan, S.M. Mofidi, F. Barlat, Mater. Sci. Eng. A 644, 210–217 (2015)

  14. 14.

    T.W.J. de Geus, R.H.J. Peerlings, M.G.D. Geers, Int. J. Solids Struct. 67–68, 326–339 (2015)

  15. 15.

    W.Q. Cao, C. Wang, J. Shi, M.Q. Wang, W.J. Hui, H. Dong, Mater. Sci. Eng. A 528, 6661–6666 (2011)

  16. 16.

    H.F. Xu, J. Zhao, W.Q. Cao, J. Shi, C.Y. Wang, C. Wang, J. Li, H. Dong, Mater. Sci. Eng. A 532, 435–442 (2012)

  17. 17.

    Z.H. Cai, H. Ding, R.D.K. Misra, H. Kong, Scripta Mater. 71, 5–8 (2014)

  18. 18.

    D.M. Field, D.C. Van Aken, Metall. Mater. Trans. A 49, 1152–1166 (2018)

  19. 19.

    K. Hariharan, O. Majidi, C. Kim, M.G. Lee, F. Barlat, Mater. Des. 52, 284–288 (2013)

  20. 20.

    J. Mendiguren, F. Cortés, X. Gómez, L. Galdos, Mater. Sci. Eng. A 634, 147–152 (2015)

  21. 21.

    I. Hajiannia, M. Shamanian, M. Atapour, E. Ghassemali, N. Saeidi, Trans. Indian Inst. Met. 71, 1363–1370 (2018)

  22. 22.

    N. Saeidi, F. Ashrafizadeh, B. Niroumand, F. Barlat, Mater. Des. 87, 130–137 (2015)

  23. 23.

    N. Saeidi, F. Ashrafizadeh, B. Niroumand, M.R. Forouzan, S.M. Mofidi, F. Barlat, Mater. Chem. Phys. 172, 54–61 (2016)

  24. 24.

    N. Saeidi, F. Ashrafizadeh, B. Niroumand, F. Barlat, Steel Res. Int. 85, 1386–1392 (2014)

  25. 25.

    S. Allain, J.P. Chateau, O. Bouaziz, Mater. Sci. Eng. A 387–389, 143–147 (2004)

  26. 26.

    A. Dumay, J.P. Chateau, S. Allain, S. Migot, O. Bouaziz, Mater. Sci. Eng. A 483–484, 184–187 (2008)

  27. 27.

    G.B. Olson, M. Cohen, Metall. Trans. A 7, 1897–1904 (1976)

  28. 28.

    A. Saeed-Akbari, J. Imlau, U. Prahl, W. Bleck, Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 40, 3076–3090 (2009)

  29. 29.

    I. Gutierrez-Urrutia, D. Raabe, Acta Mater. 60, 5791–5802 (2012)

  30. 30.

    K.T. Park, K.G. Jin, S.H. Han, S.W. Hwang, K. Choi, C.S. Lee, Mater. Sci. Eng. A 527, 3651–3661 (2010)

  31. 31.

    H.R. Abedi, A. Zarei Hanzaki, K.L. Ou, C.H. Yu, Mater. Des. 116, 472–480 (2017)

  32. 32.

    F. Yang, R. Song, Y. Li, T. Sun, K. Wang, J. Mater. 76, 32–39 (2015)

  33. 33.

    Y.U. Heo, D.H. Kim, N.H. Heo, C.W. Hong, S.J. Kim, Metall. Mater. Trans. A 47A, 6004–6016 (2016)

  34. 34.

    M.N. Yoozbashi, S. Yazdani, Mater. Sci. Eng. A 527, 3200–3205 (2010)

  35. 35.

    M.N. Yoozbashi, S. Yazdani, T.S. Wang, Int. J. Iron Steel Soc. Iran 7, 6–11 (2010)

  36. 36.

    M.N. Yoozbashi, S. Yazdani, T.S. Wang, Mater. Des. 32, 3248–3253 (2011)

  37. 37.

    M.N. Yoozbashi, S. Yazdani, Solid State Phenom. 172, 214–220 (2011)

  38. 38.

    M.N. Yoozbashi, S. Yazdani, Mater. Chem. Phys. 160, 148–154 (2015)

  39. 39.

    S.S. Sohn, H. Song, B.C. Suh, J.H. Kwak, B.J. Lee, N.J. Kim, S. Lee, Acta Mater. 96, 301–310 (2015)

  40. 40.

    J. Zhang, D. Raabe, C.C. Tasan, Acta Mater. 141, 374–387 (2017)

  41. 41.

    S. Chen, R. Rana, A. Haldar, R.K. Ray, Prog. Mater Sci. 89, 345–391 (2017)

  42. 42.

    J.K. Cuddy, M.N. Bassim, Mater. Sci. Eng. A 125, 43–48 (1990)

  43. 43.

    M. Jafari, S. Ziaei-Rad, N. Saeidi, M. Jamshidian, Mater. Sci. Eng. A 670, 57–67 (2016)

  44. 44.

    J. Hu, W. Cao, C. Wang, H. Dong, J. Li, ISIJ Int. 54, 1952–1957 (2014)

  45. 45.

    M.Y. Sherif, C. Garcia Mateo, T. Sourmail, H.K.D.H. Bhadeshia, Mater. Sci. Technol. 20, 319–322 (2004)

Download references

Acknowledgements

The authors would like to thank INSF: Iran National Science Foundation (No. 95003392), IUT: Isfahan University of Technology, NRF: The National Research Foundation of Korea (No. 2014R1A2A1A10051322) for supporting of this project.

Author information

Correspondence to N. Saeidi.

Additional information

Publisher's Note

Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.

Rights and permissions

Reprints and Permissions

About this article

Verify currency and authenticity via CrossMark

Cite this article

Saeidi, N., Jafari, M., Kim, J.G. et al. Development of an Advanced Ultrahigh Strength TRIP Steel and Evaluation of Its Unique Strain Hardening Behavior. Met. Mater. Int. 26, 168–178 (2020). https://doi.org/10.1007/s12540-019-00322-2

Download citation

Keywords

  • Medium Mn TRIP steel
  • EBSD analysis
  • Strain hardening
  • Retained austenite stability
  • Nano scale grain size