Temperature effect on deformation mechanisms and ...

Materials Science & Engineering A 642 (2015) 71–83

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Temperature effect on deformation mechanisms and mechanical properties of a high manganese C þN alloyed austenitic stainless steel L. Mosecker a,n, D.T. Pierce b,c,nn, A. Schwedt d, M. Beighmohamadi d, J. Mayer d, W. Bleck a, J.E. Wittig c a

Department of Ferrous Metallurgy, RWTH Aachen University, 52072 Aachen, Germany Advanced Steel Processing and Products Research Center, Colorado School of Mines, CO 80401, USA c Vanderbilt University, PMB 351683, 2301 Vanderbilt Place, Nashville, TN 37232, USA d Central Facility for Electron Microscopy (GFE), RWTH Aachen University, 52074 Aachen, Germany b

art ic l e i nf o

a b s t r a c t

Article history: Received 16 January 2015 Received in revised form 15 June 2015 Accepted 16 June 2015 Available online 20 June 2015

Recently developed high-manganese stainless Fe–Cr–Mn–CN steels exhibit an exceptional combination of strength and ductility and show great promise for structural applications. Understanding the relationships between temperature, stacking fault energy (SFE) and strain-hardening behavior is critical for alloying, design, and further optimization of these steels. The present study investigates the influence of temperature and SFE on the microstructural evolution to explain the deformation behavior and mechanical properties of an austenitic Fe–14Cr–16Mn–0.3C–0.3N alloy. The flow behavior is homogenous and no serrations in the flow stress occur during tensile deformation in the temperature range from 150 to 250 °C. Mechanical twinning and the formation of (planar) dislocation substructures strongly influence the mechanical properties and work-hardening behavior in the intermediate temperature range from 40 to 45 °C (SFE range from 17 to 24 mJ m 2). In the high temperature interval from 100 to 250 °C the SFE ranges from 29 to 44 mJ m 2 and the initiation of mechanical twinning is delayed leading to reduced work-hardening in the intermediate and final stages of strain-hardening. In the low temperature regime from 150 to 100 °C (SFE approximately 15 mJ m 2), εh.c.p.-martensite is the dominant secondary deformation mechanism, contributing to the enhanced work-hardening in the early and intermediate stages of deformation and slightly lower total elongations. The yield strength of the studied alloy is significantly larger and exhibits greater sensitivity to temperature within the thermal and athermal ranges for dislocation motion compared to conventional Fe–Mn–(Al)–C TWIP or austenitic stainless steels, which may be attributed to phenomena such as short range ordering. & 2015 Elsevier B.V. All rights reserved.

Keywords: Austenitic stainless steels Stacking fault energy Twinning Strain hardening EBSD Electron microscopy

1. Introduction Stainless Fe–Cr–Mn–CN twinning induced plasticity (TWIP) steels alloyed with C þN exhibit increased strength and ductility compared to conventional stainless steels [1], along with high impact toughness [2,3], longer fatigue life [4] and improved wet corrosion resistance [5]. The synergistic effect of C þN alloying has been shown to improve the fatigue life of austenitic stainless steels [4]. With respect to conventional stainless steels, the substitution of nickel by manganese increases the interstitial solubility of C and

n

Corresponding author. Corresponding author at: Advanced Steel Processing and Products Research Center, Colorado School of Mines, CO 80401, USA. E-mail addresses: [email protected] (L. Mosecker), [email protected] (D.T. Pierce). nn

http://dx.doi.org/10.1016/j.msea.2015.06.047 0921-5093/& 2015 Elsevier B.V. All rights reserved.

N [6,7]. Several authors have investigated the mechanisms by which additions of C þ N improve strength in austenitic stainless. According to Gavriljuk et al. [3,7], stabilizing the austenitic phase with C þN enhances the concentration of free electrons – more effectively than alloying only with N – promoting the metallic character of interatomic bonding and short-range ordering (SRO) rather than atomic-clustering. The increased concentration of conduction electrons strengthens the binding between immobile interstitial atoms and dislocations, enhancing the strength, ductility and impact toughness [2,7]. In comparison to conventional high-Mn TWIP steels [8–10] the homogenous flow and work hardening characteristics of Fe–Cr– Mn–CN steels indicate differences in the strain-induced hardening mechanisms. The occurrence of Cr–N SRO phenomena [11–13] and the resultant interactions with dislocations and stacking faults are believed to play a major role in the deformation behavior of these

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L. Mosecker et al. / Materials Science & Engineering A 642 (2015) 71–83

alloys. The strong affinity between Cr and N atoms generates Cr–N SRO clusters. Oda et al. [11] suggested that N in the austenitic Fe– 1.5Mn–15Cr–15Ni alloy gathers around Cr atoms to form interstitial-substitutional complexes based on X-ray absorption fine structure analysis. More recently, Li et al. [14] reported Cr–N SRO in the Co–29Cr–6Mo–(0.1–0.16)N alloy based on results of atom probe tomography. Moreover, Cr–N SRO is assumed to influence the energy barrier for the γ- 4 ε phase transition [14], which along with the stacking fault energy (SFE), is relevant to the deformation mechanisms and mechanical behavior of Fe–Cr–Mn–CN steels. The composition and temperature dependent SFE strongly influences the plasticity mechanisms and work-hardening behavior of austenitic high-Mn steels [8–10,15]. Increasing the SFE causes the active deformation mechanisms to change from strain-induced martensite formation and dislocation glide, to mechanical twinning and dislocation glide, and to pure dislocation glide [16]. The effect of N additions on the SFE in Fe–Cr–Mn steels is reported to be non-monotonic, exhibiting a minimum SFE at 0.4 wt% N in Fe–15Cr–17Mn [17,18] and Fe–20Cr–17Mn [19] alloys, from quantitative SFE measurements on extended three-fold dislocation nodes using transmission electron microscopy (TEM). Including these results in a thermodynamics based model, along with a composition dependent interfacial energy term, allows for calculation of the SFE [20]. The decrease in SFE at low N contents was explained due to the segregation of interstitial N atoms to stacking faults [21]. While at higher N contents the SFE increases as the bulk effect becomes more pronounced [18]. Limited available data on the influence of C on SFE in the Fe–Cr–Mn system exists since C contents are generally kept low ( o0.1 wt%) in austenitic stainless steels to avoid intergranular corrosion during welding. Remy [22] measured a room temperature (RT) SFE of 11 mJ m 2 for a Fe–4Cr– 20Mn–0.48C alloy. C acts as a strong austenite stabilizer that lowers the εh.c.p.-martensite-start and the γf.c.c.-εh.c.p. transition temperatures [23,24] in agreement with thermodynamic models that report increases in the SFE with C in the Fe–Mn–C system [8,25,26]. Experimental SFE values reported by Petrov [17,27] on Fe–22Mn–(0.1–0.7)C exhibit a linear increase in SFE with C content, while at dilute concentrations the segregation effects of C leads to a decrease in SFE. There are conflicting reports about the influence of C þN alloying on SFE in Fe–Mn–Cr steels. According to Lee et al. [28,29] the SFE increases linearly with increasing Cþ N content for a Fe–18Cr–10Mn alloy. However, experimental investigations by Roncery et al. [1,30] in the Fe–14Cr–(21–30)Mn– (0.57–0.77)CþN system imply a drop in SFE with increase in N content, which is explained by a decrease in the density of states at the Fermi level which is inversely proportional to the SFE [18]. Furthermore, segregation effects of N to dislocations and stacking faults [17,21], Cr–N SRO [31] and the distribution of alloying elements within the f.c.c. lattice are also reported to affect the SFE [13]. In addition, the experimental method of SFE measurement may also affect the SFE values. For instance, the method of measuring SFE by analyzing the geometry of extended three-fold nodes using TEM, which is utilized in several of the aforementioned studies, overestimated the SFE in Fe–Cr–Ni alloys [32]. The goal of this research is to provide a detailed description of the influence of temperature and SFE on the deformation mechanisms and strain-hardening behavior of a newly developed Fe– 14Cr–16Mn–0.3C–0.3N alloy. Analysis of partial dislocation spacings using weak-beam dark-field (WBDF) TEM, comparisons of the active deformation mechanisms with testing temperature, and the results of previous SFE investigations on similar materials yielded the temperature dependence of the SFE. The microstructure was characterized by electron back-scatter diffraction and TEM at different levels of strain. The strain-hardening behavior was assessed from uniaxial tensile tests conducted at temperatures ranging from 150 to 250 °C. The intrinsic properties

are correlated to the temperature dependent deformation mechanisms, with respect to the twin and dislocation substructure evolution, to explain the flow and work-hardening behavior.

2. Experimental procedure 2.1. Material The investigated steel with composition of 14.6 wt% Cr, 15.9 wt% Mn, 0.31 wt% C and 0.29 wt% N was strip cast (thickness of 2.2 mm), cold rolled to a final thickness of 1.1 mm and recrystallized at 1150 °C for 4 min. The samples were fully austenitic in the as-received state and after deformation to fracture at RT as determined with X-ray diffraction, with an average grain size of 20 mm approximating a log-normal distribution. 2.2. Ultrasonic pulse velocity measurements The ultrasonic pulse velocities (longitudinal and shear) of the Fe–14Cr–16Mn–0.3C–0.3N wt% steel and type 304 stainless steel for comparison were measured. A Tektronix TDS 2001C oscilloscope, Olympus 5072PR pulser/receiver, 20 MHz normal incident longitudinal transducer and a 10 MHz normal incident shear transducer were used to measure the transit time of longitudinal and shear pulses through the specimens. Soundsafe longitudinal and shear wave ultrasound couplants were used to couple specimens to the longitudinal and shear wave transducers, respectively. The transducers were used in a pitch/catch method. Transit time measurements were made by measuring the time between the crest of leading cycles of consecutive echoes [33]. Sheets with area between 1–2 cm2 were gently grinded, using a tool to ensure each surface is nearly parallel (standard deviation in thickness across the samples did not exceed 0.009 mm), with successively finer SiC paper up to 1200 grit. A transit time error associated with electronics and transducer specimen bond was determined using different thicknesses of stainless steel type 304 [33,34]. The longitudinal and transverse pulse velocities in the stainless steel type 304 were 5739 715 μm μs 1 and 3116 78 μm μs 1, respectively, agreeing well with that of 5759 715 μm μs 1 and 31347 18 μm μs 1 reported by Ledbetter et al. [33]. 2.3. Stacking fault energy measurements Three millimeters diameter disks were cut from the gage length of samples deformed to 0.015 plastic tensile strain using electro discharge machining. The disks were ground to 100 mm thickness and electro-polished to electron transparency with a TenuPol-5 using a solution of 95% ethanol and 5% perchloric acid at 20 °C. Partial dislocations were analyzed with a Philips CM20T TEM operating at 200 kV. Measurements of Shockley partial-dislocation separation distances were made with a beam direction near [111] on defects in the (111) habit plane using o 220 4 type g-vectors. Weak-beam dark-field (WBDF) imaging was employed using g(3 g) diffracting conditions. Multiple measurements of the separation distances were made on each dislocation pair to achieve an average separation distance for each dislocation. A correction was applied to account for the slightly smaller separation distance of the dislocation cores in relation to the intensity peaks, which arises due to asymmetries in the strain fields outside and between partial dislocations [35]. An average dactual and standard deviation of the measurements were obtained for each partial-dislocation pair. The perfect dislocation character angle, β, was determined from Burgers vector analysis on the partial dislocations in WBDF imaging mode. For Shockley partial dislocations in the [111]/(111) zone/habit plane configuration, |g∙bp| (where bp


73

is the partial dislocation Burgers vector) values are 1 or 0 and |g∙bp | ¼1 for both partials at only one g-vector. Stereographic analysis, from images taken at different zones, was employed to confirm the habit plane of the dislocations [36]. 2.4. Mechanical testing Tensile specimens of the sheet material were machined according to DIN 50124 with a gage length of 50 mm (A50) and 30 mm (A30). Uniaxial tensile tests were performed after DIN EN 10002-1 with applied engineering strain rate of 10 3 s 1 using a Zwick 1484 tensile machine (Zwick Reoll GmbH). The stress–strain rates were determined over the crosshead speed as change of length per time interval. Tensile tests at room temperature were interrupted at logarithmic strains of 0.04, 0.14, 0.21 and 0.31 using A50 samples. Tensile tests conducted at temperatures ranging from 150 °C to 250 °C were performed in a climate chamber. For the temperature measurements a thermocouple was applied directly on the specimen surface. A30 samples have been used to better control the homogenous temperature distribution in the samples. The accuracy of the temperature control was in a region of 72 °C. 2.5. Electron back-scatter diffraction The Electron back-scatter diffraction (EBSD) measurements were performed using a JEOL JSM 7000F FEG-SEM with an EDAXTSL Hikari camera and OIM DataCollection 6.2/OIM Analysis 6.2 software. The step size was adjusted to 100 nm. The criterion for the definition of twin boundaries was 60° misorientation about the o111 4 axis, with an angular tolerance of 5° within the γf.c.c.-matrix according to Saeed-Akbari et al. [16]. The vertical axis in the datasets/images was assigned as tensile axis (TA). The EBSD data are represented as image quality (IQ) maps and as inverse pole figure (IPF) maps with TA as reference direction. 2.6. Transmission electron microscopy Tensile samples deformed at 300 K with a strain rate of 10 3 s 1 to strain of 0.015 and 0.21 were prepared for TEM by standard twin-jet electro-polishing in 70% methanol 30% perchloric acid at 20 V and 30 °C. The TEM analysis was performed using a Philips CM20 T operating at 200 kV and a FEI Tecnai F20 operating at 200 kV. The twin width was measured from TEM bright-field images and high resolution TEM images from at least three different samples strained to 0.21. High resolution TEM characterization was performed using FEI Titan 80–300 field emission transmission electron microscope equipped with an imaging spherical aberration corrector element; accelerating voltage of 300 kV.

3. Results 3.1. Determination of ultrasound velocities and elastic constants Ultrasonic pulse echo patterns yield the longitudinal and transverse sound velocities. Fig. 1 shows the longitudinal ultrasound pulse echo pattern from the Fe–14Cr–16Mn–0.3C–0.3N alloy with the transit time measured from the apex of the leading cycle of consecutive echoes. Some dispersion is observed in the pulse echo patterns after the leading cycle. The Young's modulus (E) and shear modulus (G) are determined by Eq. (1) and (2) respectively [33]

Fig. 1. Longitudinal ultrasound pulse echoes in the Fe–14Cr–16Mn–0.3C–0.3N alloy showing the transit time as measured from the apex of the leading edge of each echo.

E=

(

4

)

3ρv 2t v 2l - 3 v 2t

(

v 2l -v 2t

)

G=ρv 2t

(1) (2)

Where vl and vt are the longitudinal and transverse pulse velocities, respectively, and ρ is the density of the material, calculated from the alloy composition and lattice parameter as measured using X-ray diffraction. The longitudinal and transverse sound velocities, density, Young's modulus, shear modulus and Poisson's ratio at room temperature are presented in Table 1, along with same data from other steel compositions. The calculated and or experimentally determined Néel temperatures for the f.c.c. phase of the alloys are provided for comparison.

3.2. Experimental measurement of SFE The SFE was determined from the separation distance of Shockley partial dislocation pairs. The classical relationship between SFE and partial dislocation separation provided by Hirth [41] is modified to include a correction factor (α) when using isotropic elasticity for SFE measurements, as shown in Eq. (3)

dactual =

Gbp2 α 2-ν ⎛ 2υcos2β ⎞ ⎜1⎟ 8πγ 1-ν ⎝ 2-ν ⎠

(3)

Where G and ν are the shear modulus (80 GPa) and Poisson's ratio (0.27), respectively, and taken from the above described ultrasound velocity measurements. dactual is the actual partial dislocation separation, β is the total dislocation character angle, bp is the partial dislocation Burger's vector (0.146 nm) and γ is the SFE. The use of isotropic elastic constants for SFE measurements slightly overestimates the SFE in f.c.c. materials [36,42]. A material dependent correction factor (α), which is equal to approximately 0.94 for high-Mn steels [36], is applied to correct the overestimation of the SFE. Fig. 2a shows a WBDF image of a partial dislocation pair exhibiting separation distances in the range of 8– 10 nm in a sample after light deformation (1.5%). The values of dactual for several dislocations are plotted against β in Fig. 2b. The fitted SFE curves in Fig. 2b indicate a SFE of 21 76 mJ m 2 for the Fe–14Cr–16Mn–0.3C–0.3N steel.

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Table 1 Néel temperatures for the f.c.c. phase, longitudinal and transverse ultrasound velocities, density and elastic constants for the Fe–14Cr–16Mn–0.3C–0.3N alloy and other TWIP steels at 293 K. Material Fe–14Cr–16Mn–0.3C–0.3N Fe–15Cr–20Mn–1.5N–0.1C [37] Fe–19Cr–18Mn–2.2N–0.6C [37] Fe–4Cr–20Mn–0.5C [22] Fe–5Cr–26Mn–0.2C [38] Fe–26Mn–0.2C [22] Type 304 stainless steel [33] 1 2

fcc T Neel ́

1

(K)

283 335 374

fcc T Neel ́

254 259 166 298 371 380

2

(K)

vl (μm μs 1)

vt (μm μs 1)

ρ (g cm 3)

E (GPa)

G (GPa)

ν ()

5759 723 5607 5655

3209 7 15 3188 3172

5759 715

31347 18

7.76 7.71 7.62 7.85 7.82 7.87 7.90

204 72 199 197 1997 2 183 176 72 2007 2

807 1 78.9 77.6 767 1 71.5 677 1 777 1

0.27 0.26 0.27 0.31 0.28 0.31 0.29

Experimental. Calculated TNéel after [39,40].

3.3. Mechanical properties and flow behavior

3.4. Deformation microstructure characterization

Stress–strain curves and the strain hardening rates (SHR, ds/ dε) as function of deformation temperature, obtained by uniaxial tensile testing, are plotted in Fig. 3. Within the tested temperature range the shape of the flow curves changes from linear to parabolic with increases in temperature up to 250 °C. The SHRs show a general drop down behavior, with pronounced multistage character at temperatures of 150 °C and above 45 °C. Between 40 °C and 45 °C the SHR remains nearly identically. At temperatures of 100 °C and above, a marked inflection in the SHR occurs at 0.13 logarithmic strain that is not observed in the lower temperature SHRs. Furthermore, a second inflection occurs between 0.2 and 0.3 logarithmic strain, and the strain at which such inflection occurs increases with temperature. The mechanical properties are plotted as a function of temperature in Fig. 4. The yield strength (YS) and ultimate tensile strength (UTS) decrease with increasing temperature more severely in the low temperature regime, than at higher temperatures (4 50–100 °C), which may indicate a change in the active deformation mechanism [43,44]. The yield to tensile ratio (YS/UTS) decreases with increasing temperature. The maximum total elongation (TE) of 65% occurs at 0 °C. In the temperature range from 100 to 150 °C the uniform elongation is relatively constant, ranging from 53% to 58%. A maximum post-uniform elongation (UE) of approximately 10% occurs between 0 and 25 °C and decreases to about 3% with increasing temperature to 250 °C. At 150 °C a significant loss of ductility is observed and no postuniform elongation occurs.

3.4.1. RT microstructure characterization Fig. 5 shows the IPF and IQ maps from the as-received and fracture strained tensile samples at RT. In the as-received microstructure numerous annealing twins are observed (Fig. 5a,c) which is possibly due to a low stacking fault energy [47] related to the enhanced Mn and N alloying [45,46]. In the fractured sample Σ3 twin boundaries (Fig. 5d blue colored) were determined by EBSD, revealing intense mechanical twinning. The activation of two non-coplanar twin systems is observed in a significant amount of grains in the fractured samples (e.g., an example is indicated by white and yellow arrows in Fig. 5d). Some grains exhibit twinning in one preferred twin system, which may be primary twinning initiated in the lower stress regime [28,48]. Until fracture the grains progressively rotate towards the o111 4//TA direction, and to a lesser extent, the o001 4//TA direction [49,50]. In accordance with Schmid's law, grains with a o111 4//TA orientation exhibit significantly more mechanical twins, as shown in Fig. 5b and d. With increased twin density, subboundaries developed within the grains that delimit different twin bundles, according to Barbier et al. [9]. As indicated by white arrows in Fig. 5b, these sub-grains show orientation in o001 4//TA direction. 3.4.2. Microstructure characterization as function of temperature The deformation twin activity was semi-quantified by EBSD as the relative fraction of detected Σ3 twin boundaries (ratio of Σ3 twin boundary length to the total length of boundary segments with more than 1° misorientation [51]). The applied method was

Fig. 2. (a) WBDF image of partial dislocations and (b) actual partial dislocation separation vs character angle in the Fe–14Cr–16Mn–0.3C–0.3N alloy. Measurements were made on samples in the recrystallized condition (open symbol) and after 0.015 strain (filled symbols). The large symbol indicates data taken from the partial dislocation pair in (a).


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Fig. 3. True stress (a) and strain-hardening rate (b) vs logarithmic strain for temperatures from 150 °C to 250 °C for the Fe–14Cr–16Mn–0.3C–0.3N steel.

Fig. 4. Variation of tensile strength, elongation and yield to tensile ratio (YS/UTS) over temperature.

Fig. 5. (a,b) IPF maps of the γf.c.c. phase and (c,d) IQ maps (blue-colored marks indicate Σ3 twin boundaries) of the as-received (a,c) and fractured (b,d) sample at RT. (For interpretation of the references to color in this figure legend,the reader is referred to the web version of this article.)

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4. Discussion 4.1. Elastic stiffness

Fig. 6. Relative fraction of Σ3 twin boundary length and martensite volume fraction as function of temperature; fractured samples.

The Fe–14Cr–16Mn–0.3C–0.3N steel has a lower density than that of traditional stainless and Fe–Mn–C TWIP steels and exhibits excellent values of Young's modulus and shear modulus, 204 GPa and 80 GPa, respectively, that are slightly larger than type 304 stainless steel [33,53] and vastly superior to Fe–Mn–C–(Al) TWIP steel [54–56]. A well-known phenomenon of Fe–Mn based steels is the substantial increase in the stiffness that occurs upon the antiferromagnetic to paramagnetic transition of the γ-phase, due to the anomaly in temperature dependence of elastic constants (around TNéel) [37,47,57]. The empirical relationship of TNéel proposed by Jin et al. [58] modified to include a 33.8 K reduction in TNéel per wt% addition of N [39] predicts a Néel temperature of 254 K and indicates that the alloy is paramagnetic at RT. In comparison, the Young's moduli of the steels which are

Fig. 7. IQ maps showing Σ3 twin boundaries (blue), εh.c.p.-martensite (yellow) and α′b.c.c.-martensite (red) in the fractured samples deformed at (a) 150 °C, (b) 100 °C, (c) 40 °C and (d) 250 °C. (For interpretation of the references to color in this figure legend,the reader is referred to the web version of this article.)

adopted after Randle [51] for a comparative use of length fraction statistics of Σ3 grain boundary analysis from orientation maps generated by EBSD; used in the current work to provide reliable information about the quantitative deformation twin activity as function of temperature (Fig. 6) and strain (Fig. 10). The highest twin activity is observed near RT (Fig. 6). The Σ3 twin boundary fraction is lower in the sample deformed to failure at 40 °C than at RT. Decreases in the test temperature below 40 °C results in further reductions in the twin boundary fraction and simultaneous increases in εh.c.p.-martensite (yellow) and α′b.c.c. -martensite (red) (Fig. 7a,b). The α′b.c.c.-martensite occurs preferentially at the intersections of ε-martensite bands. The volume fraction of ε-martensite increases with decreasing test temperature from 6 vol% at 100 °C to 17.6 vol% at 150 °C. In the intermediate temperature range, ε h.c.p.-martensite and Σ3 twin boundaries are both indexed as shown in Fig. 7b (white arrow), similar to other works [52]. Increasing the deformation temperature above RT results in less mechanical twinning and an increase in the spacing between individual twin bundles (Fig. 7d).

antiferromagnetic at RT are lower, ranging from 176 to 183 GPa (Table 1). The stiffness of the Fe–14Cr–16Mn–0.3C–0.3N alloy is bolstered by additions of Cr [59,60] and N. N additions to steel containing larger Cr concentrations ( 26 wt%) increases both E and G, the result being attributed to the strong affinity between Cr and N atoms [31]. The Cr–N SRO zones have been reported to locally increase the elastic shear modulus [37]. Depending on the probability and distribution of SRO within the austenitic matrix, the SRO is believed to also affect the bulk shear modulus, which may explain the larger shear modulus of the investigated alloy relative to type 304 stainless steel. 4.2. Evaluation of SFE 4.2.1. SFE as function of chemical composition The experimental SFE measurements for the Fe–14Cr–16Mn– 0.3C–0.3N alloy and those for Fe–Cr–Mn–CN steels reported in the literature are presented in Table 2 along with the observed microstructure before and after deformation. The results indicate that εh.c.p. martensite is suppressed when the SFE is increased above approximately 20 mJ m 2 in Fe–Cr–Mn–CN steels. An


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Table 2 Stacking fault energy for the Fe–14Cr–16Mn–0.3C–0.3N alloy and other austenitic stainless Fe–Cr–Mn–CN steels at RT. Reference

Present work Lee et al. [28] Lee et al. [29]

Lee et al. [64] Mujica et al. [30]

Kang et al. [62] Bracke et al. [67] Riedner et al. [63] a b

Chemical composition (wt%)

Microstructure

SFE [mJ m 2]

Cr

Mn

N

C

C þN (wt%)

C/N ( )

Before

After deformationa

Experimental

Empirical [29]

14.6 18.3 18.0 18.1 18.1 18.1 2.5 11.9 12.0 11.8 6.8 8.8 5.6 9.7 18.2

16.0 9.7 10.2 9.5 9.6 9.7 14.9 21.0 25.4 30.0 18.2 18.2 16.2 19.0 18.9

0.29 0.61 0.42 0.35 0.32 0.38 0.21 0.33 0.45 0.44 0.29 0.17 0.11 0.20 0.61

0.31 0.02 0.15 0.24 0.30 0.38 0.56 0.24 0.32 0.33 0.59 0.78 0.25 0.24 0.34

0.60 0.63 0.57 0.59 0.62 0.76 0.77 0.57 0.77 0.77 0.89 0.95 0.36 0.44 0.95

1.00 0.03 0.36 0.69 0.94 1.00 2.67 0.73 0.71 0.75 2.03 4.59 2.27 1.20 0.65

γ γ γ γ γ γ γ γ γ γ γ γ γ γ γ

γT γT þ εD γþ α′D γT γT γT γT γT γT γT γT – γþ α′D þ εD γT γT

21.0 ( 7 6) 19.3 ( 7 0.2) 20.1 ( 7 1.0) 20.8 ( 7 0.1) 22.5 ( 7 0.1) 27.6 ( 7 0.4) 33.1 ( 70.3) 36 31 43.3 33.2b 40.9b 5 20 37b

22.1 19.3 18.2 20.2 22.4 28.2 34.9 19.6 27.5 27.6 36.9 49.5 17.1 16.2 34.1

Fracture strained samples. Approximated SFE from calculations.

empirical relation of SFE depending on C þ N content and C/N ratio, proposed by Lee et al. [29] for Fe–18Cr–10Mn–CN alloys, reveals that the relative increase in SFE due to addition of C is larger than compared to N, while referring to a study by Reed and Austin [61] in Fe–20Cr–10Ni–CN steels. The proposed relationship is in good agreement with experimental SFE values for the present Fe–14Cr– 16Mn–0.3C–0.3N steel and other Fe–Cr–Mn–CN alloys with higher Mn concentrations up to 19 wt% [62,63] and lower Cr contents to 42.5 wt% [64] (compare Table 2). The good agreement suggests that Cr and Mn may have only a small influence on the SFE in the ranges of 2.5–18 wt% Cr and 10–19 wt% Mn content. The influence of Cr on the SFE is reported to be typically within þ/ 1 mJ m 2 per wt% depending on alloy composition, according to experimental [32], thermodynamic [20] and empirical SFE investigations of austenitic steels [65]. On the other hand, the influence of Mn on the SFE is still in discussion, with many authors reporting a parabolic evolution of the SFE with increasing Mn content based on theoretical calculations and experimental measurements [26,36,65,66]. The experimental SFE values reported by Roncery

Fig. 8. Experimental and calculated SFEs as function of temperature after model calculation for Fe–22Mn–0.6C [8,25], Fe–5Cr–20Mn–0.5C [8,71] and type 304 stainless steel [72] with respect to the deformation microstructures of the Fe–14Cr– 16Mn–0.3C–0.3N determined with EBSD. (γT – mechanical twinning, εD and αD – deformation induced h.c.p. and b.c.c/b.c.t. martensite).

et al. [1,30] in the Fe–14Cr–(21–30)Mn system with constant C/N 0.75 ranges from 31 to 43.3 mJ m 2, which is substantially larger than the range of 19.6–27.6 mJ m 2 predicted empirically, suggesting that at greater Mn contents the SFE becomes more sensitive to Mn additions, which is consistent with previous observations [20,36]. 4.2.2. SFE as function of temperature The SFE is strongly influenced by temperature. The SFE increases with temperature causing a change in the active deformation mechanisms from: (1) dislocation glide and εh.c.p. martensite to (2) dislocation glide and mechanical twinning to (3) pure dislocation glide. Several reports indicate εh.c.p. martensite is suppressed above a SFE of 18 mJ m 2 in Fe–22Mn–0.6C [8,68] and Fe–Cr–Mn–N [20] systems, and 15 mJ m 2 in the Fe–5Cr– 20Mn–0.5C [22]. At SFE values 440 mJ m 2, mechanical twinning gets progressively suppressed and dislocation glide becomes the dominant deformation mechanism [36,69,70]. However a strict division between the SFE values in the transition regions is not yet clear and often depends on the investigated alloying system. Fig. 8 represents the change of SFE with temperature based on thermodynamic calculations [8,25] and experimental observations [8,71,72] with respect to the measured SFE of the investigated Fe– 14Cr–16Mn–0.3C–0.3N alloy at RT and the observed deformation microstructures with EBSD at various temperatures. At T 4RT the linear temperature coefficient of 0.1 mJ m 2 K 1 [72] is found convenient to reproduce the SFE variation with temperature, as the SFE would be overestimated by the thermodynamic calculations. In the applied thermodynamic subregular models [8,25] the interstitial elements are considered to be in a substitutional solution without taking vacancies into account. Consequently, the effect of interstitial elements on the SFE is insufficiently described [20]. The SFE at 250 °C is calculated to 44.5 mJ m 2, correlating with the decreasing volume fraction of twin boundaries compared to RT (see Fig. 6). However, at T oRT the SFE over temperature follows the thermodynamic based model calculations by SaeedAkbari et al. [25] that consider the contribution of magnetic ordering. At 40 °C mechanical twinning is the only type of secondary deformation mechanism observed and the SFE is assumed to be 15.5 mJ m 2. In the low temperature range to 150 °C, where ε h.c.p-/α′b.c.c.-martensite and mechanical twins are detected, the SFE is approximately 15 mJ m 2. With a decrease in temperature below TNéel the austenite phase transitions from the paramagnetic

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Fig. 9. Normalized yield strength (YS) as function of temperature for the Fe–14Cr– 16Mn–0.3C–0.3N alloy in comparison to the overall trend of Fe–(13–36)Mn–(0–1.2) C steels [74], Fe–17Cr–7Mn–4Ni–0.2N alloy [43] and type 316 stainless steel [75].

to antiferromagnetic state, providing a positive magnetic conγ ε tribution to ΔG - 4 that partially counteracts the negative chemical contribution [8,36,68]. According to the present results, the dominant secondary deformation mechanisms changes from mechanical twinning to deformation induced εh.c.p- and α′b.c.c.-martensite formation below a SFE of approximately 15 mJ m 2 for the Fe–14Cr–16Mn–0.3C–0.3N steel and the tranγ ε γ ε sition temperature T0 - is between 100 °C oT0 - o 40 °C. The evolution of deformation mechanisms in austenitic steel is also influenced by factors other than the SFE. The crystallographic orientation [10], dislocation density [70] and dislocation-solute interactions [31] are further relevant microstructure parameter that play a major role in the critical stress for mechanical twinning and the control of dislocation dynamics. 4.3. Yield strength The slope of YS as a function of temperature is more pronounced at lower temperatures, as result of the solid solution strengthening of interstitial N, which is approximately 40% more effective than C [73], influencing the thermally activated dislocation glide dynamics [31]. However, alloying C þN is found to increase the YS similar to N. In this regard, the YS is the most representative parameter to define the thermal activated nature of dislocation motion with respect to the effect of interstitial C and N [55]. The normalized YS (normalized by the room temperature YS), which excludes microstructural or solid solution hardening effects, is plotted as function of temperature in Fig. 9. As presented by Allain et al. [55,74] from various reports, the Fe–Mn–C alloys reveal very similar activation transition temperatures, defining the transition from thermally to athermal activated dislocation motion with increase in temperature (Fig. 9 Gy area). Within the thermal activated range the temperature sensitivity is strongly influence by the carbon content while the manganese content and the magnetic state had no significant effect [74]. In the athermal range the influence of temperature and carbon content is low. The minor slope above RT of 0.25 is attributed to the linear decrease of the elastic bulk modulus with increasing temperature for the paramagnetic Fe–Mn–C alloys in the considered temperature range [74]. The investigated Fe–14Cr–16Mn– 0.3C–0.3N alloy exhibits high temperature sensitivity in the variation of yield strength, both within the thermal and athermal temperature (slope 0.84) range, while the transition temperature is comparable to the alloys in the Fe–Mn–C system. The

significant high athermal flow stress in austenitic stainless Fe–Ni– Cr–Mo steels containing 0.04–0.36 wt% N is attributed to Cr–N SRO, increasing with N content [31]. Byrnes et al. [31] proposed that the SRO zones locally increase the elastic shear modulus and induce anisotropic elastic strain in the nearest surrounding of N atoms. The thermally activated component of the flow stress dependents mainly on the N content [31]. Due to the increased conduction electrons around the N atoms compared to C [7], nitrogen atoms carry an effective negative electric charge, whereas the core of screw dislocations reveal shortage of electrons. The large electrostatic attraction between the N atoms and dislocations leads to enhanced binding and increased sensitivity of the thermal activated dislocation motion in comparison to Fe–Mn–C alloys. Similarly, using a statistical thermodynamics based model of SRO, Zhou and Grujicic [76] reported that N induced short range order significantly enhances the strength of Fe–Cr–Ni alloys. The effect of nitrogen on temperature dependence of thermal stress was identified for the strengthening at low temperatures rather than its role in promoting the SRO [76]. The synergistic enhanced solid solution strengthening of C þN alloying is believed to be even more effective, following the investigations by Gavriljuk et al. [3,7]. Although the electron density is decreased in the vicinity of C-atoms, in the case of C þN alloying, the spatial distribution and density of conduction electrons is more homogeneous and increases in the vicinity of C-atoms, intensifying the effect of N [3,7]. The results by Hamada et al. [43] on the carbon free Fe–17Cr– 7Mn–4Ni–0.2N alloy showing a similar trend of the YS(T)–YS (293 K) over temperature as the Fe–14Cr–16Mn–0.3C–0.3N alloy (Fig. 9), which supports the discussion on Cr–N SRO complexes on the dislocation dynamics in these materials. 4.4. Flow behavior and strain-hardening f(ε–s,T) 4.4.1. Flow behavior The stress–strain curves exhibit a homogenous flow behavior over the tested temperature range at strain rate of 10 3 s 1 and no periodic serrations have been detected in the flow curves. The phenomena of serrated yielding or jerky flow is described in the literature as macroscopic evidence of dynamic strain aging (DSA), known to occur in C alloyed austenitic steels [55,77,78]. Proposed by Lee et al. [79], the DSA effect arises from interactions between stacking faults and C–Mn SRO complexes, by breaking away of stacking faults and dislocations from the ordered zones, which requires higher stresses. In the presence of C–Mn complexes, the dislocation movement not only requires additional force to pass the local ordered zone, but also causes a disordering effect of the C–Mn complexes, which may occur by a single diffusive jump of C atoms [80]. In low-SFE material with wider stacking faults the reorientation of C–Mn complexes within the stacking fault region is favored and occurs before the stacking fault is removed by the trailing partial, leading to higher stresses for reinitiating of the dislocation glide, enhancing the DSA mechanism [80]. The DSA effect however is reduced in high-Mn TWIP steels containing Al (Fig. 13), which decreases C activity [55], increases the activation energy for reorientation of the C–Mn point defect [81], and results in smaller stacking faults widths due to a higher SFE. The addition of N was reported to increase the critical strain for the onset of serrations or even suppress the serrated yielding in Fe–Cr–Ni–Mo [31,82–84], Fe–Mn–C [77,79,85,86] and Fe–Cr–Mn–C [81] systems. Additions of 0.09 wt%N to the Fe–18Mn–0.6C alloy increased the SFE and reduced the stacking fault width, reducing the probability that the C–Mn complexes can reorient before the stacking fault is removed by the trailing partial, which delays the onset of DSA [79]. The critical strain for DSA appears to be related to both the stacking fault width and the type of SRO complexes [54,79,80]. In the Fe–14Cr–16Mn–0.3C–0.3N steel, with SFE of approximately


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interstitial C, reducing the formation of Mn–C dipoles and limiting them as a possible source of DSA. In Fe–Cr–Mn–C–N steels, Cr–N SRO is assumed to be the preferred ordered structure [88] rather than Mn–C SRO [89,90]. Furthermore, Cr additions, which significantly reduce the diffusivity of C in austenite [91], may also increase the activation energy of reorientation of the point defect complexes, similar to the effect of Al additions [81]. Therefore, the type of SRO and the activation energy for reorientation of Cr–C and Cr–N point defects, rather than the SFE may be the reason for the homogenous flow behavior in the present steel.

Fig. 10. True stress, strain-hardening rate and evolution of relative length of Σ3 twin boundaries over logarithmic strain at 300 K and microstructure IPF maps of the interrupted samples.

21 mJ m 2 and equilibrium spacing of stacking faults ranging from approximately 5.5 to 9 nm, no DSA is observed during deformation at quasi static strain rates. In comparison, DSA is observed in a Fe– 18Mn–0.6C–1.5Al alloy with SFE of 30 mJ m 2 and equilibrium partial dislocations spacing ranging from 3.5 to 6 nm [54]. The finding suggests that the absence of DSA in the present Fe–14Cr– 16Mn–0.3C–0.3N alloy is more strongly influenced by the type of SRO complexes rather than SFE or stacking fault width. In Cr–N steels, the strong affinity between Cr and N atoms result in the formation of Cr–N SRO phenomena as extensively discussed in literature [11–13]. The interatomic attraction (I) of the interstitial C and N to Mn or Cr–substitutes have been calculated by Xie et al. [87] to be ICr–N 4IMn–N 4ICr–C 4IMn–C, with the attraction of Cr–N and Mn–N significantly greater than the Cr–C and Mn–C. These results support the idea proposed by Bracke [86] that in Fe–Mn–C alloys, N atoms interfere with point-defect complexes involving

4.4.2. Strain-hardening f(ε–s) The strain-hardening rate (SHR) as function of strain at 300 K is plotted in Fig. 10, showing three SH-stages. The strain-hardening (SH) behavior of the Fe–14Cr–16Mn–0.3C–0.3N alloy is characterized by a single intermediate hardening stage associated with high SH, similar to Al-containing TWIP [68,91,92] and austenitic stainless steels [43,75,93]. The initial stage (I) hardening is characterized by a sharp drop of the SHR. Within stage (II) the SHR steadily decreases. In stage (III) the SHR rapidly drops until fracture. Stage (I) is characterized by dynamic recovery of dislocations and formation of stacking faults [9,48] along with a significant drop in the work-hardening rate. No mechanical twins were detected by EBSD (Fig. 10) and TEM observations of the microstructure at ε ¼0.015 indicate minor dislocation cross-slip, numerous single dislocations on at least two active slip planes and intrinsic stacking faults; as observed by typical diffraction contrast methods [35] (Fig. 11a and b). The observed planar glide is facilitated by a low SFE and the N content, which both promote planar glide of dislocations by reducing the frequency of cross slip. The activation of cross slip in high-N steels occurs only over short-range distance due to the existence of Cr–N SRO [77,94,95]. The transition between stage (I) and (II) is reported to correspond to the onset of deformation twinning [9] and or evolution of the dislocation substructure [10,69]. In stage (II) hardening, the fraction of Σ3 twin boundaries increases slightly in a linear manner up to at least 0.31 logarithmic strain (Fig. 10) likely facilitated by increasing stress and dislocation density [70]. The observed mechanical twinning is mainly confined to one preferred twin system. As can be seen from the IPF maps in Fig. 10, with increasing strain, more grains show orientation between the o0014 //TA and o111 4//TA direction, characteristic for grains with a well-developed twin substructure of one active twinning system [10]. Jin and Lee [40] suggest the existence of a single intermediate hardening stage is due to less

Fig. 11. Planar dislocation structure (a) and two-beam dark-field image corresponding to g ¼[200] indicating intrinsic stacking faults (b) at strain of 0.015%.

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Fig. 12. Bright-field image of (a) extended stacking faults blocked at twin lamella and (b) o 111 4 twin plane reflection showing dislocations pile-ups at planar defects at ε3 ¼ 0.21.

active secondary mechanical twinning in Fe–Mn–C–Al TWIP steels. Other authors explain planar slip and the resulting anisotropic and homogeneously spread dislocation substructures as a sufficient condition for the linear hardening behavior in f.c.c. materials [55]. The primary mechanical twins contribute to the strainhardening by acting as obstacles to dislocations gliding on different slip systems, as illustrated in Fig. 12a which shows mechanical twins after 0.21 logarithmic strain. Dislocation accumulation at the twin boundaries is evident in Fig. 12b which is a higher magnification image of the twinned area in Fig. 12a. According to Müllner et al. [96] N changes the glide distribution leading to an increased local dislocation density. By that, mechanical twinning not only occurs earlier but also at more positions causing a finer deformation substructure. The twin width was measured to be 5.9 nm with standard deviation of 3.9 nm. A reduced twin thickness with increasing N content is also reported by Jung et al. [48] in austenitic stainless steels. Thinner twins

provide stronger barriers for dislocation glide with more dislocation pile-ups required to generate the required stress to propagate dislocations across the twin leading to increased SH level, as observed for the Fe–14Cr–16Mn–0.3C–0.3N alloy. In stage (II), a gradual decrease in the strain-hardening rate occurs. This workhardening behavior differs significantly from that of Fe–Mn–C TWIP as illustrated in Fig. 13a and b. The decreasing effective stress-hardening rate over stage (II) can be discussed after Feaugas [97] as a result of the enhanced dynamic competition between storage and annihilation of dislocations, pronounced due to the N assisted development of a fine deformation substructure. The high dislocation activity is further indicated to enhance glide plane softening diminishing the hardening effect of deformation twinning [98,99]. The flow stress of the Fe–14Cr–16Mn–0.3C–0.3N is superior to the Fe–24Mn–0.6C steel in the initial and intermediate stages of deformation, which can be attributed to the larger solid solution

Fig. 13. Stress–strain curves (a) and strain hardening rates (b) for different high-Mn systems FeMnC [16], FeMnAlC [92] and FeCrMnN [93,100] compared to alloy Fe–14Cr– 16Mn–0.3C–0.3N and the type 316 stainless steel [75].


strengthening of C þN relative to N [83], SRO [31], a higher shear modulus and to mechanical twinning (Fig. 13). The Cr–N SRO results in increased planar glide [77] and twinning stress [95] as the SRO zone needs to be destroyed in the successive twin planes during twin growth. Conversely, the SHR rate of the Fe–14Cr–16Mn–0.3C–0.3N steel is substantially lower than for the Fe–24Mn–0.6C steel in the intermediate and final stages of deformation. One main reason, besides the differences in the evolution and characteristics of the deformation twin substructures, is the occurrence of serrated yielding or DSA in Fe–Mn–C systems as described earlier (Fig. 13a). The DSA results in inhomogeneous plastic flow and unstable flow behavior that significantly enhances hardening at lower strain rates [16,17,99]. In stage (III) the twin fraction and thickness of twin bundles increases significantly and twinning occurring in more than one active twin system is more predominant (Figs. 5b and d and 10), which contributes to the large post-uniform elongation. The development of intense dislocation accumulations within the refined microstructure results in pronounced deformation localization, which finally leads to the samples fracture [9]. Mechanical twinning has been identified as the main secondary deformation mechanism of the investigated Fe–14Cr–16Mn– 0.3C–0.3N alloy at RT. Due to the C þN alloying the deformation substructure is characterized by planar dislocation glide and a local increased dislocation density causing a fine deformation twin structure, that leads to increased strengthening. The associated enhanced dynamic recovery diminishes the hardening effect of the deformation twinning, which can explain the hardening behavior of the investigated alloy and other Cr–Mn austenitic stainless steel grades [43]. However, further anaylsis is requird to investigate the evolution of the twin/twin inter-bundle spacing as important parameter for the dynamic Hall–Petch effect on dislocations affecting the strain hardening behavior of this material. 4.4.3. Influence of temperature on strain-hardening The flow stress and SHRs for samples tested from 150 to 100 °C are significantly larger than the higher test temperatures, stage (III) of the SHR is absent, and minimal post-uniform elongation occurs. This temperature range corresponds to the region of thermally activated dislocation motion where larger stresses are required for dislocation glide, enhancing the flow stress and SH level compared to higher temperatures. In this temperature region the estimated SFE values are near 15 mJ m 2 (Fig. 8) promoting the formation of strain-induced εh.c.p.-martensite as the dominant secondary deformation mechanism besides mechanical twinning and αb.c.c-martensite which are also present but in smaller quantities (Fig. 6). Previous investigations on a Fe–22Mn–3Al–3Si alloy indicates dislocation cross slip is strongly impeded at SFE of 15 mJ m 2, as evidenced by TEM observations of planar dislocation structures [69]. The reduced ability for dislocation cross slip and the formation of εh.c.p.-martensite laths acting as barriers to dislocations gliding on other slip system [69], further hinder dislocation mobility and are likely enhancing the flow stress and SHR of the Fe–14Cr–16Mn–0.3C–0.3N alloy at test temperatures from 150 to 100 °C. The hardening associated with the formation ε h.c.p.-and αb.c.c-martensite typically occurs in the early and intermediate ranges of strain [36] which may reduce the capacity for the steel to further harden after uniform elongation is reached, resulting in the low post-uniform elongation values relative to those observed at temperatures where mechanical twinning is the dominant secondary deformation mechanism. The SHRs of the samples tested from 40 to 45 °C exhibit remarkable similarity, characterized by a single intermediate hardening stage and enhanced elongation. The flow stresses and SHR are significantly lower than those observed at test temperatures

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from 150 to 100 °C. The SFE ranges about 17–24 mJ m 2 (Fig. 8) and mechanical twinning is the dominant secondary deformation mechanism and no strain-induced martensite is detected. In the low-SFE range deformation twinning occurs gradually over the entire deformation process, while the majority of mechanical twin formation and mechanical twinning in multiple slip systems occurred after 0.31 logarithmic strain (Fig. 9), leading to a homogeneous strain-hardening, prolongation of local necking, and further enhancement of the post-uniform elongation [44]. The Fe–14Cr–16Mn–0.3C–0.3N appears to show optimized ductility around 0 °C. Towards higher and lower temperatures the elongation values decreases, revealing a typical behavior for low SFE high-Mn TWIP steels [44,68]. The flow stresses and SHR rates corresponding to samples tested from 100 to 250 °C exhibit similar behavior in stage (I) but different characteristics in the intermediate and final SH stages compared to samples tested at temperatures from 40 to 45 °C. At approximately 0.13 logarithmic strain the SHR of the samples tested from 100 to 250 °C experience an inflection followed by a rapid decrease in the strain-hardening rate. A second inflection in the SHR occurs at logarithmic strains of approximately 0.2, 0.25 and 0.3 for test temperatures of 100, 150 and 250 °C, respectively. The second inflection is thought to correspond to the onset of mechanical twinning and its associated contribution to hardening [69,70]. The increase in temperature delays the onset of mechanical twinning by reducing the flow stress, increasing the rate of dislocation annihilation and increasing the SFE, which raises the critical stress for activating mechanical twinning [70]. Consequently, as temperature increases lower twin boundary fractions are measured (Fig. 6), observed (Fig. 7), and mechanical twinning becomes progressively suppressed in favor of pure dislocation glide. The reduction in mechanical twinning with increasing temperature causes a slight reduction in TE of 7%, which is primarily attributed to decreased post-uniform elongation. In comparison, the hardening contribution due to mechanical twinning is almost completely suppressed in lower C TWIP alloys at 200 °C, such as Fe–17Cr–7Mn–3Ni–0.2N [43] and Fe–25Mn–3Al– 3Si [44,69], which results in a remarkable decrease in TE. Additional hardening in the athermal range is associated to DSA [55,101] generated by interactions of dislocations with SRO and planar defects [102]. According to Almeida et al. [103] DSA occurs even without macroscopic consequence of serrated flow, supporting the presence of DSA mechanism in the Fe–14Cr–16Mn– 0.3C–0.3N alloy despite of the homogenous flow behavior.

5. Summary and conclusion The deformation mechanisms and mechanical properties of the investigated Fe–14Cr–16Mn–0.3C–0.3N alloy depend on temperature, SFE and ordered microstructural phenomena like SRO. Deformation substructures like mechanical twins or ε h.c.p.-martensite laths form barriers to dislocation glide, enhancing the work-hardening behavior. The development of planar dislocation arrays and dislocation barriers are further governed by the interactions of dislocations and planar defects with SRO complexes. From this work, the following conclusions were drawn: (1) A RT SFE of 21 7 6 mJ m 2, measured from the separations of Shockley partial dislocation pairs, results in mechanical twinning as the dominant secondary deformation mechanism. Reducing the deformation temperature to 150 °C lowers the SFE to approximately 15 mJ m 2 and changes the dominant secondary deformation mechanism from mechanical twinning to strain-induced εh.c.p.-martensite formation.

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(2) Distinct temperature sensitivity of the yield strength within the thermal and athermal temperature range was observed and is attributed to N additions and Cr–N SRO. The yield strength of the studied alloy is significantly larger than conventional Fe–Mn-(Al)–C TWIP or austenitic stainless steels. (3) The flow behavior is homogenous and no serrations in the flow stress occur during tensile deformation in the temperature range from 150 to 250 °C. The apparent absence of dynamic strain aging is attributed to the type of SRO rather than the SFE. (4) The largest ductility is observed in the intermediate temperature interval from 45 to 40 °C corresponding to a SFE range from 17 to 24 mJ m 2. The SH behavior was found to be controlled by mainly dislocation glide at low strain, competition between dislocation dynamics and mechanical twinning in the intermediate strain regime and mainly mechanical twinning and dislocation accumulation at fracture strain. (5) In the high temperature interval from 100 to 250 °C the SFE ranges from 29 to 44 mJ m 2 and the initiation of mechanical twinning is delayed, leading to reduced work-hardening in the intermediate and final stages of strain-hardening. (6) In the low temperature regime from 150 to 100 °C (SFE approximately 15 mJ m 2), εh.c.p.-martensite is the dominant secondary deformation, resulting in enhanced work-hardening in the early and intermediate stages of deformation and slightly lower total elongations. The enhanced thermally activated glide component of the flow stress contributes to the strain hardening at low temperatures.

Acknowledgments The authors gratefully acknowledge the financial support of the German Research Foundation (DFG) within the Collaborative Research Center (SFB) 761 “Steel–ab initio” and Dr. JeeHyun Kang (Department of Ferrous Metallurgy, RWTH Aachen University) for the scientific discussion. The support of Prof. Dierk Raabe and the Max-Planck-Institut für Eisenforschung, Düsseldorf is also gratefully acknowledged.

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