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Metals
Special Issues
Recent Developments in Medium and High Manganese Steels
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Recent Developments in Medium and High Manganese Steels

A special issue of Metals (ISSN 2075-4701). This special issue belongs to the section "Metal Casting, Forming and Heat Treatment".

Deadline for manuscript submissions: closed (31 January 2022) | Viewed by 31611

Special Issue Editor

Dr. Colin Scott

E-Mail Website
Guest Editor
NRCAN-RNCAN, Natural Resources Canada, Hamilton, ON, Canada
Interests: AHSS development; thermomechanical processing; structure/properties modelling

Special Issue Information

Dear Colleagues,

Manganese steels have been continuously studied since the publication of an article entitled, “Hadfield’s Patent Manganese Steel” in the 8 February 1884 edition of “The Engineer”. Incredibly, 136 years on, few areas of physical metallurgy still generate as much excitement and activity as the study of medium and high manganese steels. The current wave of interest began in the mid-1990s/early 2000s, triggered by the development of cold-rolled 2G and later 3G sheets to meet more stringent automotive weight reduction and crash resistance requirements. Some of the new alloys have already found their way into the automotive marketplace; many others are close to commercial production. A multitude of other, non-automotive, applications is also being enthusiastically pursued. These include shape memory alloys, steels for ballistic protection, cryogenic containers, medical stents, tank cars, slurry pipes, alloys for additive manufacturing, etc.

This Special Issue is intended to provide a broad forum for the latest results in the physical metallurgy of these fascinating steels. This includes fundamental questions regarding phase transformations and strain hardening mechanisms such as the relative importance of mechanical twinning and DSA in TWIP steels, and the factors governing the nucleation, growth and stability of austenite islands in TRIP alloys. Strain partitioning, size effects, extended Lüders plateaus and the origin of negative strain rate sensitivities are all crucial aspects that require much better understanding. Fracture properties and especially the sensitivity to hydrogen embrittlement are equally important subjects for research. Contributions on these and other topics related to the processing, testing, characterization and applications of modern manganese steels are invited.

Dr. Colin Scott
Guest Editor

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Keywords

  • Manganese steels
  • Alloy design
  • Strain hardening mechanisms
  • Phase transformations
  • Mechanical properties

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Editorial

Jump to: Research, Review

3 pages, 162 KiB  
Editorial
Recent Developments in Medium and High Manganese Steels
by Colin P. Scott
Metals 2022, 12(5), 743; https://doi.org/10.3390/met12050743 - 27 Apr 2022
Cited by 2 | Viewed by 2396
Abstract
A huge amount of intellectual effort is currently being devoted to the study of medium and high manganese steels due to the diverse and impressive mechanical properties that can be achieved with these steels [...] Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)

Research

Jump to: Editorial, Review

20 pages, 9493 KiB  
Article
Effect of Intercritical Annealing Parameters and Starting Microstructure on the Microstructural Evolution and Mechanical Properties of a Medium-Mn Third Generation Advanced High Strength Steel
by Kazi M. H. Bhadhon, Xiang Wang, Elizabeth A. McNally and Joseph R. McDermid
Metals 2022, 12(2), 356; https://doi.org/10.3390/met12020356 - 18 Feb 2022
Cited by 13 | Viewed by 2912
Abstract
A prototype medium-Mn TRIP steel (0.2 C–6 Mn–1.7 Si–0.4 Al–0.5 Cr (wt %)) with a cold-rolled tempered martensite (CR) and martensitic (M) starting microstructures was subjected to continuous galvanizing line (CGL) compatible heat treatments. It was found that the M starting microstructures achieved [...] Read more.
A prototype medium-Mn TRIP steel (0.2 C–6 Mn–1.7 Si–0.4 Al–0.5 Cr (wt %)) with a cold-rolled tempered martensite (CR) and martensitic (M) starting microstructures was subjected to continuous galvanizing line (CGL) compatible heat treatments. It was found that the M starting microstructures achieved greater than 0.30 volume fraction of retained austenite and target 3G properties (UTS × TE ≥ 24,000 MPa%) using an intercritical annealing temperature (IAT) of 675 °C with an IA holding time of 60–360 s, whereas the CR microstructure required an IAT of 710 °C and annealing times of 360 s or greater to achieve comparable fractions of retained austenite and target 3G properties. This was attributed to the rapid austenite reversion kinetics for the M starting microstructures and rapid C partitioning from the C supersaturated martensite, providing chemical and mechanical stability to the retained austenite, thereby allowing for a gradual deformation-induced transformation of retained austenite to martensite—the TRIP effect—and the formation of nano-scale planar faults in the retained austenite (TWIP effect), such that a high work-hardening rate was maintained to elongation of greater than 0.20. Overall, it was concluded that the prototype steel with the M starting microstructure is a promising candidate for CGL processing for 3G AHSS properties. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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Figure 1

Figure 1
<p>Schematic diagram of the heat treatment profile.</p> Full article ">Figure 2
<p>(<b>a</b>) Load–unload–reload test; (<b>b</b>) enlarged portion of (<b>a</b>) showing process for determining <span class="html-italic">σ<sub>F</sub></span> and <span class="html-italic">σ<sub>R</sub></span>.</p> Full article ">Figure 3
<p>SEM micrographs of (<b>a</b>) as-received CR samples with tempered martensite and (<b>b</b>) heat-treated M samples with martensite; C = carbides, TM = tempered martensite, M = martensite, TT = through thickness, TD = transverse direction.</p> Full article ">Figure 4
<p>Montage of TEM micrographs showing carbide distribution in (<b>a</b>) CR and (<b>b</b>) M starting microstructures.</p> Full article ">Figure 5
<p>Retained austenite volume fraction as a function of intercritical annealing temperature and holding time for (<b>a</b>) CR and (<b>b</b>) M starting microstructures.</p> Full article ">Figure 6
<p>Estimated fresh martensite volume fraction as a function of holding time at 675 and 710 °C IATs for (<b>a</b>) CR and (<b>b</b>) M starting microstructures.</p> Full article ">Figure 7
<p>SEM micrographs of CR samples annealed at 710 °C for (<b>a</b>) 120 s, (<b>b</b>) 240 s, and (<b>c</b>) 600 s; C = carbide, F = ferrite, M = martensite, and A = retained austenite.</p> Full article ">Figure 8
<p>SEM micrographs of M samples annealed at 675 °C for (<b>a</b>) 120 s, (<b>b</b>) 360 s, (<b>c</b>) 600 s; C = carbide, F = ferrite, M = martensite, and A = retained austenite.</p> Full article ">Figure 8 Cont.
<p>SEM micrographs of M samples annealed at 675 °C for (<b>a</b>) 120 s, (<b>b</b>) 360 s, (<b>c</b>) 600 s; C = carbide, F = ferrite, M = martensite, and A = retained austenite.</p> Full article ">Figure 9
<p>TEM results for the M 675 °C + 120 s sample: (<b>a</b>) Montage of bright field (BF) TEM micrographs; (<b>b</b>) dark field (DF) TEM corresponding to <math display="inline"><semantics> <mrow> <mo>〈</mo> <mn>110</mn> <mo>〉</mo> </mrow> </semantics></math>γ; (<b>c</b>) SAD patterns corresponding to <math display="inline"><semantics> <mrow> <mfenced close="]" open="["> <mrow> <mn>100</mn> </mrow> </mfenced> <msup> <mi>α</mi> <mo>′</mo> </msup> <mn>1</mn> <mo>∥</mo> <mfenced close="]" open="["> <mrow> <mn>111</mn> </mrow> </mfenced> <msup> <mi>α</mi> <mo>′</mo> </msup> <mn>2</mn> <mo>∥</mo> <mfenced close="]" open="["> <mrow> <mn>110</mn> </mrow> </mfenced> <mi>γ</mi> </mrow> </semantics></math>.</p> Full article ">Figure 10
<p>Montage of BF TEM micrographs showing carbide distribution in (<b>a</b>) CR 710 °C + 120 s and (<b>b</b>) M 675 °C + 120 s samples.</p> Full article ">Figure 11
<p>Carbide area fraction as a function of microstructure/heat treatment.</p> Full article ">Figure 12
<p>(<b>a</b>) Engineering stress vs. engineering strain; (<b>b</b>) true stress vs. true strain curves; work hardening rate vs. true strain curves for selected (<b>c</b>) CR and (<b>d</b>) M starting microstructures.</p> Full article ">Figure 13
<p>Retained austenite transformation kinetics for selected CR and M samples.</p> Full article ">Figure 14
<p>(<b>a</b>) Bright field (BF) TEM; (<b>b</b>) dark field (DF) TEM corresponding to <math display="inline"><semantics> <mrow> <mo>〈</mo> <mn>110</mn> <mo>〉</mo> </mrow> </semantics></math>γ; (<b>c</b>) SAD patterns corresponding to <math display="inline"><semantics> <mrow> <mfenced close="]" open="["> <mrow> <mn>100</mn> </mrow> </mfenced> <msup> <mi>α</mi> <mo>′</mo> </msup> <mn>1</mn> <mo>∥</mo> <mfenced close="]" open="["> <mrow> <mn>111</mn> </mrow> </mfenced> <msup> <mi>α</mi> <mo>′</mo> </msup> <mn>2</mn> <mo>∥</mo> <mfenced close="]" open="["> <mrow> <mn>110</mn> </mrow> </mfenced> <mi>γ</mi> </mrow> </semantics></math> for M 675 °C + 120 s sample at <span class="html-italic">ε</span> = 0.10.</p> Full article ">Figure 15
<p>R (<span class="html-italic">ε</span>) versus true strain for the M 675 °C + 120 s sample.</p> Full article ">
15 pages, 3725 KiB  
Article
Study on High-Temperature Mechanical Properties of Fe–Mn–C–Al TWIP/TRIP Steel
by Guangkai Yang, Changling Zhuang, Changrong Li, Fangjie Lan and Hanjie Yao
Metals 2021, 11(5), 821; https://doi.org/10.3390/met11050821 - 18 May 2021
Cited by 10 | Viewed by 2247
Abstract
In this study, high-temperature tensile tests were carried out on a Gleeble-3500 thermal simulator under a strain rate of ε = 1 × 10−3 s−1 in the temperature range of 600–1310 °C. The hot deformation process of Fe–15.3Mn–0.58C–2.3Al TWIP/TRIP at different [...] Read more.
In this study, high-temperature tensile tests were carried out on a Gleeble-3500 thermal simulator under a strain rate of ε = 1 × 10−3 s−1 in the temperature range of 600–1310 °C. The hot deformation process of Fe–15.3Mn–0.58C–2.3Al TWIP/TRIP at different temperatures was studied. In the whole tested temperature range, the reduction of area ranged from 47.3 to 89.4% and reached the maximum value of 89.4% at 1275 °C. Assuming that 60% reduction of area is relative ductility trough, the high-temperature ductility trough was from 1275 °C to the melting point temperature, the medium-temperature ductility trough was 1000–1250 °C, and the low-temperature ductility trough was around 600 °C. The phase transformation process of the steel was analyzed by Thermo-Calc thermodynamics software. It was found that ferrite transformation occurred at 646 °C, and the austenite was softened by a small amount of ferrite, resulting in the reduction of thermoplastic and formation of the low-temperature ductility trough. However, the small difference in thermoplasticity in the low-temperature ductility trough was attributed to the small amount of ferrite and the low transformation temperature of ferrite. The tensile fracture at different temperatures was characterized by means of optical microscopy and scanning electron microscopy. It was found that there were Al2O3, AlN, MnO, and MnS(Se) impurities in the fracture. The abnormal points of thermoplasticity showed that the inclusions had a significant effect on the high-temperature mechanical properties. The results of EBSD local orientation difference analysis showed that the temperature range with good plasticity was around 1275 °C. Under large deformation extent, the phase difference in the internal position of the grain was larger than that in the grain boundary. The defect density in the grain was large, and the high dislocation density was the main deformation mechanism in the high-temperature tensile process. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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Figure 1

Figure 1
<p>Sampling diagram of high-temperature tensile sample.</p> Full article ">Figure 2
<p>Flowchart of high-temperature tensile specimen processing.</p> Full article ">Figure 3
<p>Fracture topography.</p> Full article ">Figure 4
<p>(<b>a</b>) True stress–strain curve and (<b>b</b>) curve of peak stress versus temperature.</p> Full article ">Figure 5
<p>Temperature dependence of reduction of area of Fe–15.3Mn–5.8C–2.3Al TWIP/TRIP steel and Fe–24.2Mn–3Al–2.6Si and 12Cr1MoVG steel.</p> Full article ">Figure 6
<p>Fracture morphology of samples at different test temperatures. (<b>a</b>) 600 °C; (<b>b</b>,<b>c</b>) 700 °C; (<b>d</b>) 800 °C; (<b>e</b>) 900 °C; (<b>f</b>) 1100 °C; (<b>g</b>) 1200 °C; (<b>h</b>,<b>i</b>) 1275 °C.</p> Full article ">Figure 7
<p>Microstructure morphology near the high-temperature tensile fracture. (<b>a</b>) 25 °C; (<b>b,c</b>) 700 °C; (<b>d</b>) 1200 °C; (<b>e</b>) 1250 °C; (<b>f</b>) 1275 °C.</p> Full article ">Figure 8
<p>Predicted phase change behavior of steel by Thermo-Calc thermodynamic calculation. (<b>a</b>) 600–1300 °C; (<b>b</b>) 1425–1460 °C.</p> Full article ">Figure 9
<p>EBSD analysis near tensile fracture at 1200 and 1275 °C. (<b>a</b>) 1200 °C; (<b>b</b>) 1275 °C.</p> Full article ">
17 pages, 3795 KiB  
Article
Temperature Effects on Tensile Deformation Behavior of a Medium Manganese TRIP Steel and a Quenched and Partitioned Steel
by Whitney A. Poling, Emmanuel De Moor, John G. Speer and Kip O. Findley
Metals 2021, 11(2), 375; https://doi.org/10.3390/met11020375 - 23 Feb 2021
Cited by 17 | Viewed by 2954
Abstract
Third-generation advanced high-strength steels (AHSS) containing metastable retained austenite are being developed for the structural components of vehicles to reduce vehicle weight and improve crash performance. The goal of this work was to compare the effect of temperature on austenite stability and tensile [...] Read more.
Third-generation advanced high-strength steels (AHSS) containing metastable retained austenite are being developed for the structural components of vehicles to reduce vehicle weight and improve crash performance. The goal of this work was to compare the effect of temperature on austenite stability and tensile mechanical properties of two steels, a quenched and partitioned (Q&P) steel with a martensite and retained austenite microstructure, and a medium manganese transformation-induced plasticity (TRIP) steel with a ferrite and retained austenite microstructure. Quasi-static tensile tests were performed at temperatures between −10 and 85 °C for the Q&P steel (0.28C-2.56Mn-1.56Si in wt.%), and between −10 and 115 °C for the medium manganese TRIP steel (0.14C-7.14Mn-0.23Si in wt.%). X-ray diffraction measurements as a function of strain were performed from interrupted tensile tests at all test temperatures. For the medium manganese TRIP steel, austenite stability increased significantly, serrated flow behavior changed, and tensile strength and elongation changed significantly with increasing temperature. For the Q&P steel, flow stress was mostly insensitive to temperature, uniform elongation decreased with increasing temperature, and austenite stability increased with increasing temperature. The Olson–Cohen model for the austenite-to-martensite transformation as a function of strain showed good agreement for the medium manganese TRIP steel data and fit most of the Q&P steel data above 1% strain. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>Schematic of tensile specimen geometry.</p> Full article ">Figure 2
<p>Combined EBSD phase map and image quality map showing the as-processed microstructures of (<b>a</b>) QP3Mn and (<b>b</b>) TRIP7Mn. Red indicates ferrite/martensite and green indicates austenite.</p> Full article ">Figure 3
<p>Representative engineering stress–strain curves from tests at 0.0005 s<sup>−1</sup> and various temperatures for (<b>a</b>) QP3Mn and (<b>c</b>) TRIP7Mn. Serrated flow behavior during necking is shown by the zoomed-in stress–strain plot for QP3Mn. Work hardening rate (d<span class="html-italic">σ</span>/d<span class="html-italic">ε</span>) as a function of strain from tests at 0.0005 s<sup>−1</sup> and various temperatures for (<b>b</b>) QP3Mn and (<b>d</b>) TRIP7Mn. Work hardening rate was calculated from a smoothed true stress–strain curve obtained by a polynomial fit to the plastic deformation region between yield point elongation (YPE)/yielding and necking.</p> Full article ">Figure 4
<p>Austenite volume fraction versus plastic strain from tests at 0.0005 s<sup>−1</sup> and various temperatures for (<b>a</b>) QP3Mn and (<b>b</b>) TRIP7Mn.</p> Full article ">Figure 5
<p>Fraction of transformed martensite versus plastic strain at various test temperatures (0.0005 s<sup>−1</sup> strain rate) for (<b>a</b>) QP3Mn and (<b>c</b>) TRIP7Mn. The points are experimental data, and the curves are the Olson–Cohen model fits. For QP3Mn, data at 2% and greater strain were used for model fits. Olson–Cohen model parameters <span class="html-italic">α</span> and <span class="html-italic">β</span> versus temperature with <span class="html-italic">n</span> equal to 2 for (<b>b</b>) QP3Mn and (<b>d</b>) TRIP7Mn.</p> Full article ">Figure 6
<p>Tensile mechanical properties versus temperature from replicate tests at 0.0005 s<sup>−1</sup>. (<b>a</b>) QP3Mn 0.2% YS and UTS; (<b>b</b>) QP3Mn UE and TE; (<b>c</b>) TRIP7Mn lower YS, upper YS, and UTS; and (<b>d</b>) TRIP7Mn YPE, UE, and TE.</p> Full article ">
18 pages, 15513 KiB  
Article
Effect of Processing Parameters on Mechanical Properties of Deformed and Partitioned (D&P) Medium Mn Steels
by Chengpeng Huang and Mingxin Huang
Metals 2021, 11(2), 356; https://doi.org/10.3390/met11020356 - 20 Feb 2021
Cited by 6 | Viewed by 2602
Abstract
Deformed and partitioned (D&P) medium Mn steels exhibiting high strength, large ductility, and excellent fracture toughness have been developed recently. The ultra-high dislocation density and transformation-induced plasticity (TRIP) effect are the main mechanisms for their exceptional mechanical properties. The simple processing route to [...] Read more.
Deformed and partitioned (D&P) medium Mn steels exhibiting high strength, large ductility, and excellent fracture toughness have been developed recently. The ultra-high dislocation density and transformation-induced plasticity (TRIP) effect are the main mechanisms for their exceptional mechanical properties. The simple processing route to manufacturing D&P steel makes it promising for large-scale industrial applications. However, the exact effect of each processing step on the final mechanical properties of D&P steel is not yet fully understood. In the present work, the effects of processing parameters on the mechanical properties of D&P steels are systematically investigated. The evolution of microstructure, tensile behavior and austenite fraction of warm rolled samples and D&P samples are revealed. Two D&P steels, with and without the intercritical annealing process, are both produced for comparison. It is revealed that the intercritical annealing process plays an insignificant role to the mechanical properties of D&P steel. The partitioning process is extremely important for obtaining large uniform elongation via slow but sustaining strain hardening by the TRIP effect in the partitioned austenite. The cold rolling process is also significant for acquiring high strength, and the cold rolling thickness reduction (CRTR) is extremely critical for the strength–ductility synergy of D&P steels. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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Figure 1

Figure 1
<p>Equilibrium phase fraction of studied steel at different temperatures calculated by Thermo-Calc.</p> Full article ">Figure 2
<p>Schematic illustration of the thermomechanical process for producing deformed and partitioned (D&amp;P) steel.</p> Full article ">Figure 3
<p>Optical micrographs showing the initial microstructures of the hot rolling (HR) + warm rolling (WR) sample (<b>a1</b>,<b>a2</b>) and the HR + WR + intercritical annealing (IA) + cold rolling with 25% thickness reduction (CR25%) + partitioning at 400 °C (PT400) sample (<b>b1</b>,<b>b2</b>). RD: rolling direction, ND: normal direction. M: martensite, RA: retained austenite.</p> Full article ">Figure 4
<p>Electron backscatter diffraction (EBSD) maps showing the initial microstructures of the HR + WR sample (<b>a1</b>–<b>a4</b>) and HR + WR + IA + CR25% + PT400 sample (<b>b1</b>–<b>b4</b>). Therein, (<b>a1</b>,<b>b1</b>) are the phase maps showing the phase distribution of the two samples. (<b>a2</b>,<b>b2</b>) are the inverse pole figure (IPF) maps showing the grains orientations of the two samples. (<b>a3</b>,<b>b3</b>) are the maps of geometrically necessary dislocation (GND) density, estimated based on the kernel average misorientation. (<b>a4</b>,<b>b4</b>) are the band contrast maps showing image quality of the mapping. RD: rolling direction, ND: normal direction. <span class="html-italic">γ</span>: austenite, <span class="html-italic">α</span>′: martensite.</p> Full article ">Figure 5
<p>(<b>a</b>) XRD profiles of the HR + WR sample before deformation and after fracture, normalized to (111) <span class="html-italic">γ</span>. (<b>b</b>) XRD profiles of the HR + WR + IA + CR25% + PT400 sample before deformation and after fracture, normalized to (220) <span class="html-italic">γ</span>. (<b>c</b>) Volume fraction of austenite of HR + WR sample and HR + WR + IA + CR25% + PT400 sample before deformation and after fracture.</p> Full article ">Figure 6
<p>(<b>a</b>) Engineering stress–strain curves of HR + WR sample and HR + WR + IA + CR25% + PT400 sample. (<b>b</b>) The corresponding true stress–strain curves (solid line) and the work hardening rate curves (dotted line).</p> Full article ">Figure 7
<p>(<b>a</b>) Engineering stress–strain curves of the HR + WR sample and HR + WR + IA sample. (<b>b</b>) The enlarged engineering stressstrain curve of the rectangle in (<b>a</b>).</p> Full article ">Figure 8
<p>EBSD maps showing the initial microstructures of the HR + WR + CR20% + PT350 sample. (<b>a</b>) Phase map showing the distribution of martensite phase (red color) and retained austenite phase (blue color). (<b>b</b>) IPF map showing the grains orientations of the sample. (<b>c</b>) The map of GND density estimated based on the kernel average misorientation. (<b>d</b>) Band contrast map showing image quality of the mapping. RD: rolling direction, ND: normal direction. <span class="html-italic">γ</span>: austenite, <span class="html-italic">α</span>′: martensite.</p> Full article ">Figure 9
<p>(<b>a</b>) Engineering stress–strain curves of the HR + WR + IA + CR25% + PT400 sample and the HR + WR + CR20% + PT350 sample. (<b>b</b>) The corresponding true stress–strain curves (solid line) and work hardening rate curves (dotted line).</p> Full article ">Figure 10
<p>(<b>a</b>) Engineering stress–strain curves of HR + WR + IA + CR25% + PT400 sample, HR + WR + CR20% + PT350 sample, and corresponding samples without the partitioning process. (<b>b</b>) Uniform elongation versus yield strength for samples in (<b>a</b>), the yield strength is determined as the upper yield point for D&amp;P samples, the square dots represent samples without partitioning process, the circular dots represent samples with partitioning process, the colors in (<b>b</b>) are in consistent with colors in (<b>a</b>).</p> Full article ">Figure 11
<p>(<b>a</b>) Engineering stress–strain curves of the HR + WR + IA + CR25% + PT samples with the partitioning temperatures (PT) set as 350 °C, 400 °C, 450 °C, 500 °C, respectively. (<b>b</b>) Uniform elongation versus yield strength for samples in (<b>a</b>), the yield strength is determined as the upper yield point for D&amp;P samples, the colors in (<b>b</b>) are in consistent with colors in (<b>a</b>).</p> Full article ">Figure 12
<p>(<b>a</b>) Engineering stress–strain curves of the HR + WR + CR samples with different cold rolling thickness reductions (0% to 35%), with and without the partitioning process. (<b>b</b>) Uniform elongation versus yield strength for samples in (<b>a</b>), the square dots represent samples without partitioning process, the circular dots represent samples with partitioning process, the colors in (<b>b</b>) are in consistent with colors in (<b>a</b>). (<b>c</b>) The evolution of yield strength with the cold rolling thickness reduction. (<b>d</b>) The evolution of the volume fraction of austenite with the cold rolling thickness reduction. (The yield strength is determined as the upper yield point for D&amp;P samples).</p> Full article ">
15 pages, 4750 KiB  
Article
Effect of Niobium on Inclusions in Fe-Mn-C-Al Twinning-Induced Plasticity Steel
by Fangjie Lan, Wenhui Du, Changling Zhuang and Changrong Li
Metals 2021, 11(1), 83; https://doi.org/10.3390/met11010083 - 3 Jan 2021
Cited by 9 | Viewed by 2307
Abstract
The effect of Nb addition on the composition, morphology, quantity, and size of inclusions in Fe-Mn-C-Al steel was studied by SEM, EDS, and thermodynamic analysis. The research shows that the number of inclusions in Fe-Mn-C-Al high manganese steel decreases obviously after adding 0.04% [...] Read more.
The effect of Nb addition on the composition, morphology, quantity, and size of inclusions in Fe-Mn-C-Al steel was studied by SEM, EDS, and thermodynamic analysis. The research shows that the number of inclusions in Fe-Mn-C-Al high manganese steel decreases obviously after adding 0.04% element Nb, and some inclusions in the steel evolve into complex niobium inclusions. When the niobium content increases to 0.08%, the influence of niobium on inclusions in steel becomes more obvious. The precipitation temperature of inclusions in Fe-Mn-C-Al steel was analyzed by thermodynamics. The results show that the nucleation core of the composite inclusions is AlN, and then NbC and MnS precipitate locally on its surface. With the increase of Nb, the amount and volume fraction of NbC inclusions precipitated in steel increase. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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Figure 1

Figure 1
<p>Sampling diagram.</p> Full article ">Figure 2
<p>Morphology of typical inclusions in TWIP steel (steel A) (the preset Al content is 0, and the actual Al content is 0.014%). (<b>a</b>) MnO inclusion; (<b>b</b>) Al<sub>2</sub>O<sub>3</sub>-MnS complex inclusion.</p> Full article ">Figure 3
<p>Morphology of typical inclusions in TWIP steel (steel B) (the mass fraction of Al is 1.49%). (<b>a</b>) AlN inclusion; (<b>b</b>) AlN-MnS complex inclusion; (<b>c</b>) The surface scan of AlN-MnS complex inclusion.</p> Full article ">Figure 4
<p>Morphology of typical inclusions in TWIP steel (steel C) (the mass fraction of Al and Nb is 1.5% and 0.04%, respectively). (<b>a</b>) MnS inclusion; (<b>b</b>) NbC inclusion; (<b>c</b>) AlN inclusion; (<b>d</b>) AlN-MnS-NbC complex inclusion; (<b>e</b>) the surface scan of niobium complex inclusion.</p> Full article ">Figure 5
<p>The proportion of inclusions in TWIP steel with different numbers of inclusions and different heats: (<b>a</b>) is the number of different inclusions in steel A, (<b>b</b>) is the number of different inclusions in steel B, (<b>c</b>) is the number of different inclusions in steel C, (<b>d</b>) is the number of different inclusions in steel D, (<b>e</b>) is the total number of inclusions in different heats, and (<b>f</b>) is the proportion of different inclusions in different heats.</p> Full article ">Figure 5 Cont.
<p>The proportion of inclusions in TWIP steel with different numbers of inclusions and different heats: (<b>a</b>) is the number of different inclusions in steel A, (<b>b</b>) is the number of different inclusions in steel B, (<b>c</b>) is the number of different inclusions in steel C, (<b>d</b>) is the number of different inclusions in steel D, (<b>e</b>) is the total number of inclusions in different heats, and (<b>f</b>) is the proportion of different inclusions in different heats.</p> Full article ">Figure 6
<p>Average size distribution of inclusions in different TWIP steels.</p> Full article ">Figure 7
<p>Temperature dependence of precipitation amount of second phase (NbC) in different niobium contents: (<b>a</b>) is the result of statistical analysis and (<b>b</b>) is the result of phase analysis.</p> Full article ">Figure 8
<p>Volume fraction of TWIP steel with different Nb content at any temperature: (<b>a</b>) is the result of statistical analysis and (<b>b</b>) is the result of phase analysis.</p> Full article ">
14 pages, 9430 KiB  
Article
Microstructural Influence on Mechanical Properties of a Lightweight Ultrahigh Strength Fe-18Mn-10Al-0.9C-5Ni (wt%) Steel
by Michael Piston, Laura Bartlett, Krista R. Limmer and Daniel M. Field
Metals 2020, 10(10), 1305; https://doi.org/10.3390/met10101305 - 29 Sep 2020
Cited by 14 | Viewed by 2730
Abstract
This study evaluates the role of thermomechanical processing and heat treatment on the microstructure and mechanical properties of a hot rolled, annealed, and aged Fe-18Mn-10Al-0.9C-5Ni (wt%) steel. The steel exhibited rapid age hardening kinetics when aged in the temperature range of 500–600 °C [...] Read more.
This study evaluates the role of thermomechanical processing and heat treatment on the microstructure and mechanical properties of a hot rolled, annealed, and aged Fe-18Mn-10Al-0.9C-5Ni (wt%) steel. The steel exhibited rapid age hardening kinetics when aged in the temperature range of 500–600 °C for up to 50 h, which has been shown in other work to be the result of B2 ordering in the ferrite and κ-carbide precipitation within the austenite matrix. The ultimate tensile strength increased from 1120 MPa in the annealed condition to 1230 MPa after 2 h of aging at 570 °C. Charpy V-notch toughness was evaluated at −40 °C in sub-sized specimens with a maximum in the annealed and quenched condition of 28.5 J in the L-T orientation. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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Figure 1
<p>Schematic diagram of thermomechanical processing and subsequent heat treatment for the Fe-18Mn-10Al-0.9C-5Ni steel.</p> Full article ">Figure 2
<p>Optical micrographs of Fe-18Mn-10Al-0.9C-5Ni steel in the (<b>a</b>) as-cast condition and (<b>b</b>) after hot rolling at 1050 °C.</p> Full article ">Figure 3
<p>(<b>a</b>) Optical micrograph and (<b>b</b>) secondary electron image of the hot rolled steel annealed at 900 °C for 2 h.</p> Full article ">Figure 4
<p>(<b>a</b>) Optical image of Fe-18Mn-10Al-0.9C-5Ni specimen after secondary annealing at 950 °C for 30 min. (<b>b</b>) Secondary electron image shows globular B2 on grain boundaries (white arrow) and platelets of matrix B2 that are coarsening in crystallographic directions (black arrow).</p> Full article ">Figure 5
<p>Optical micrographs and corresponding higher magnification secondary electron images of the Fe-18Mn-10Al-0.9C-5Ni alloy in the annealed condition and aged at 530 °C for (<b>a</b>) 1 h and (<b>b</b>) 50 h. Precipitation of ferrite and lamellar κ-carbide on austenite grain boundaries was observed at 50 h at 530 °C.</p> Full article ">Figure 6
<p>(<b>a</b>) X-ray diffraction (XRD) patterns from for the 1, 8, and 50 h aged at 530 °C specimen. (<b>b</b>) Truncated XRD spectra for the 1, 8, and 50 h aged at 530 °C specimen.</p> Full article ">Figure 7
<p>(<b>a</b>) Hardness as a function of aging time for the Fe-18Mn-10Al-0.9C-5Ni alloy. (<b>b</b>) Vickers microhardness (with a 50 g load) of ferrite/B2 stringers and austenite after aging at 500 °C. The first data point for each phase is the 950 °C annealed condition.</p> Full article ">Figure 8
<p>(<b>a</b>) Optical micrograph of the 950 °C annealed tensile specimen sectioned perpendicular to the fracture surface and parallel to the tensile loading direction. (<b>b</b>) Secondary electron image of tensile fracture surface in the annealed condition.</p> Full article ">Figure 9
<p>(<b>a</b>) Optical micrograph of the annealed and aged for 2 h at 570 °C tensile specimen sectioned perpendicular to the fracture surface and parallel to the tensile loading direction. (<b>b</b>) The secondary electron image of the tensile fracture surface in the aged condition.</p> Full article ">Figure 10
<p>Secondary electron images of the fracture surface of an annealed specimen broken at −40 °C in the (<b>a</b>) L-T and (<b>b</b>) T-L orientations.</p> Full article ">Figure 11
<p>(<b>a</b>,<b>b</b>) Secondary electron images of the CVN fracture surface of the specimen that was annealed and aged for 2 h at 570 °C and broken in the L-T orientation at −40 °C.</p> Full article ">Figure 12
<p>Ultimate tensile strength (UTS) as a function of total elongation (%) for various wrought hot rolled (HR) and cold rolled (CR) Fe-Mn-Al-C steels and annealing heat treatments.</p> Full article ">
13 pages, 3339 KiB  
Article
The Influence of Specimen Geometry and Strain Rate on the Portevin-Le Chatelier Effect and Fracture in an Austenitic FeMnC TWIP Steel
by Jidong Kang, Liting Shi, Jie Liang, Babak Shalchi-Amirkhiz and Colin Scott
Metals 2020, 10(9), 1201; https://doi.org/10.3390/met10091201 - 8 Sep 2020
Cited by 5 | Viewed by 2511
Abstract
We studied the Portevin-Le Chatelier effect and fracture behavior of a FeMnC TWIP steel using high speed digital image correlation by varying the specimen geometry (flat vs. round) and test strain rate (0.001 vs. 0.1 s−1). The results show that the [...] Read more.
We studied the Portevin-Le Chatelier effect and fracture behavior of a FeMnC TWIP steel using high speed digital image correlation by varying the specimen geometry (flat vs. round) and test strain rate (0.001 vs. 0.1 s−1). The results show that the mean flow stress, the mean strain hardening rate and the mean strain rate sensitivity parameters are all independent of the specimen geometry and are uncorrelated with the presence or not of Portevin-Le Chatelier (PLC) bands, the type of PLC bands observed or the critical strain for band formation. However, both the fracture strains and stresses and the PLC behavior are highly geometry and/or strain rate dependent. Dynamic strain aging (DSA) and in particular the presence of PLC instabilities appears to play an important but as yet unclear role in promoting premature necking and final fracture. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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Figure 1
<p>Optical micrograph showing Klemm color etching of the as-quenched Fe-0.9C-17Mn-0.5Si-0.3V-0.019N alloy.</p> Full article ">Figure 2
<p>TEM bright field micrographs of Fe-0.9C-17Mn-0.5Si-0.3V-0.019N in: (<b>a</b>) as-quenched condition; (<b>b</b>) after straining, from a region close to the fracture surface. Two secant deformation twinning systems are indicated by the dotted lines.</p> Full article ">Figure 3
<p>Tensile data from Fe-0.9C-17Mn-0.5Si-0.3V-0.019N flat and round specimens under low and high strain rates: (<b>a</b>) engineering stress-engineering strain curves; (<b>b</b>) true stress-true strain curves, fitted instantaneous strain hardening rates and strain rate sensitivities.</p> Full article ">Figure 4
<p>Digital image correction strain maps showing initiation of Type A Portevin-Le Chatelier (PLC) bands in round specimens: (<b>a</b>) strained at 0.001 s<sup>−1</sup>; (<b>b</b>) strained at 0.1 s<sup>−1</sup>.</p> Full article ">Figure 4 Cont.
<p>Digital image correction strain maps showing initiation of Type A Portevin-Le Chatelier (PLC) bands in round specimens: (<b>a</b>) strained at 0.001 s<sup>−1</sup>; (<b>b</b>) strained at 0.1 s<sup>−1</sup>.</p> Full article ">Figure 5
<p>Digital image correction strain maps showing the initiation of different types of PLC bands in flat specimens: (<b>a</b>) Type A bands, strain rate of 0.001 s<sup>−1</sup>; (<b>b</b>) Type C bands, strain rate of 0.1 s<sup>−1</sup>.</p> Full article ">Figure 6
<p>Digital image correlation results from line scans along the tensile direction showing final steps to fracture in round specimens at strain rates of (<b>a</b>) 0.001 s<sup>−1</sup>; (<b>b</b>) 0.1 s<sup>−1</sup>. The macroscopic true strain is given in parentheses.</p> Full article ">Figure 7
<p>Digital image correlation results from line scans along the tensile direction showing final steps to fracture in flat specimens at strain rates of: (<b>a</b>) 0.001 s<sup>−1</sup>; (<b>b</b>) 0.1 s<sup>−1</sup>. The macroscopic true strain is given in parentheses.</p> Full article ">
15 pages, 4823 KiB  
Article
On Mechanical Properties of Welded Joint in Novel High-Mn Cryogenic Steel in Terms of Microstructural Evolution and Solute Segregation
by Jia-Kuan Ren, Qi-Yuan Chen, Jun Chen and Zhen-Yu Liu
Metals 2020, 10(4), 478; https://doi.org/10.3390/met10040478 - 4 Apr 2020
Cited by 14 | Viewed by 3777
Abstract
There is a growing demand for high-manganese wide heavy steel plate with excellent welding performance for liquefied natural gas (LNG) tank building. However, studies on welding of high-Mn austenitic steel have mainly focused on the applications of automotive industry for a long time. [...] Read more.
There is a growing demand for high-manganese wide heavy steel plate with excellent welding performance for liquefied natural gas (LNG) tank building. However, studies on welding of high-Mn austenitic steel have mainly focused on the applications of automotive industry for a long time. In the present work, a high-Mn cryogenic steel was welded by multi-pass Shielded Metal Arc Welding (SMAW), and the microstructural evolution, solute segregation and its effect on the properties of welded joint (WJ) were studied. The yield strength, tensile strength and elongation of the WJ reached 804 MPa, 1027 MPa and 11.2% at −196 °C, respectively. The elongation of WJ was reduced with respect to the BM due to the poorer strain hardening capacity of weld metal (WM) at −196 °C. The WM and coarse-grained heat affected zone (CGHAZ) had the lowest cryogenic impact absorbed energy of ~55 J (at −196 °C). The inhibited twin formation caused by the higher critical resolved shear twinning stress ( τ T ) in the C-Mn-Si segregation band, the inhomogeneous microstructure caused by solute segregation, and the hardened austenite matrix deteriorated the plastic deformation capacity, finally resulting in the decreased cryogenic impact toughness of the CGHAZ. To summarize, the cryogenic toughness and tensile properties of the WJ meet the requirements for LNG tank building. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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<p>Schematic diagram of sampling location, welding groove, locations of partial Charpy V-notch (indicated by the yellow, red, and blue solid lines) and microstructure characterization locations (indicated by the red, blue and pink four-pointed star).</p> Full article ">Figure 2
<p>Optical microstructure of WJ at (<b>a</b>) FL, (<b>b</b>) FL + 2 mm location and (<b>c</b>) BM.</p> Full article ">Figure 3
<p>(<b>a</b>) Engineering stress-strain curves, (<b>b</b>) strain hardening rate curves of WJ and BM at room temperature and −196 °C.</p> Full article ">Figure 4
<p>Vickers macrohardness profile across the WJ. There are three measuring lines across the WJ which are illustrated on the top left.</p> Full article ">Figure 5
<p>(<b>a</b>) Cryogenic Charpy V-notch impact absorbed energy profile across the WJ, secondary cracks underneath fracture surface at the (<b>b</b>) FL location, (<b>c</b>) BM, and fracture morphology of impact samples at the (<b>d</b>) FL location, (<b>e</b>) BM. The mechanical twins are indicated by yellow arrows in (<b>c</b>) and the crack deflection at TB and HAGB are also indicated by orange arrows in (<b>c</b>). The facets in fractograph are also indicated by yellow arrows in (<b>d</b>).</p> Full article ">Figure 6
<p>TEM micrographs of the deformed microstructures close to fracture surface of Charpy V-notch impact samples at the FL location and BM, (<b>a</b>–<b>c</b>) FL, (<b>d</b>–<b>f</b>) BM: (<b>a</b>) bright-field image showing nano-twins formed in primary twinning system and (<b>b</b>,<b>c</b>) bright-field images showing nano-twins formed in two twinning systems, (<b>d</b>) bright-field image showing nano-twins formed in primary twinning system, (<b>e</b>) bright-field image and (<b>f</b>) corresponding dark-field image showing nano-twins formed in two twinning systems.</p> Full article ">Figure 7
<p>Distribution map of (<b>a</b>) mechanical twin thickness and (<b>b</b>) inter-twin spacing in deformed microstructures close to fracture surface of impact samples.</p> Full article ">Figure 8
<p>The correlation of cryogenic Charpy V-notch impact absorbed energy and average grain size at different locations of HAZ.</p> Full article ">Figure 9
<p>EMPA composition maps of (<b>a</b>) Fe, (<b>b</b>) Mn, (<b>c</b>) C, (<b>d</b>) Si elements and (<b>e</b>) EBSP in CGHAZ. The c1 and c2 arrows indicate the C-Mn-Si segregation band and C-Mn-Si depletion band, respectively.</p> Full article ">Figure 10
<p>KAM maps, showing local orientation gradients in grain interior of (<b>a</b>) HAZ and (<b>b</b>) BM.</p> Full article ">Figure 11
<p>EMPA composition maps of (<b>a</b>) Fe, (<b>b</b>) Mn, (<b>c</b>) C, (<b>d</b>) Si elements and (<b>e</b>) EBSP in BM. The b1 and b2 arrows indicate the Mn segregation band and Mn depletion band, respectively.</p> Full article ">

Review

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23 pages, 7539 KiB  
Review
Hydrogen Embrittlement of Medium Mn Steels
by Lawrence Cho, Yuran Kong, John G. Speer and Kip O. Findley
Metals 2021, 11(2), 358; https://doi.org/10.3390/met11020358 - 20 Feb 2021
Cited by 21 | Viewed by 5322
Abstract
Recent research efforts to develop advanced–/ultrahigh–strength medium-Mn steels have led to the development of a variety of alloying concepts, thermo-mechanical processing routes, and microstructural variants for these steel grades. However, certain grades of advanced–/ultrahigh–strength steels (A/UHSS) are known to be highly susceptible to [...] Read more.
Recent research efforts to develop advanced–/ultrahigh–strength medium-Mn steels have led to the development of a variety of alloying concepts, thermo-mechanical processing routes, and microstructural variants for these steel grades. However, certain grades of advanced–/ultrahigh–strength steels (A/UHSS) are known to be highly susceptible to hydrogen embrittlement, due to their high strength levels. Hydrogen embrittlement characteristics of medium–Mn steels are less understood compared to other classes of A/UHSS, such as high Mn twinning–induced plasticity steel, because of the relatively short history of the development of this steel class and the complex nature of multiphase, fine-grained microstructures that are present in medium–Mn steels. The motivation of this paper is to review the current understanding of the hydrogen embrittlement characteristics of medium or intermediate Mn (4 to 15 wt pct) multiphase steels and to address various alloying and processing strategies that are available to enhance the hydrogen-resistance of these steel grades. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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Figure 1
<p>Schematic of representative thermo-mechanical processing routes for medium-Mn steels and the resulting microstructures. IA, intercritical annealing; ART, austenite-reverted transformation; Q&amp;T, quenching and tempering; Q&amp;P, quenching and partitioning; UFG, ultra–fine–grained; α′, α′–martensite; γ, austenite; α, α–ferrite; and δ, δ–ferrite.</p> Full article ">Figure 2
<p>Pseudo-binary phase diagrams for (<b>a</b>) Fe–6.0Mn–xC, (<b>b</b>) Fe–6.0Mn–1.5Al–xC, (<b>c</b>) Fe–6.0Mn–3.0Al–xC, and (<b>d</b>) Fe–6.0Mn–1.5Al–2.0Si–xC systems (all in wt pct). γ, austenite; α, α–ferrite; δ, δ–ferrite; and θ, cementite. Thermodynamic calculations were performed, using the Thermo-Calc<sup>®</sup> software with database of TCFE9.</p> Full article ">Figure 3
<p>TEM micrographs of (<b>a</b>) lamellarized and (<b>b</b>) equiaxed morphologies, produced by IA of hot-rolled and cold-rolled medium-Mn steels (Fe–7Mn–0.1C–0.5Si, in wt pct), respectively; α<sub>L</sub> and γ<sub>L</sub> are lath–shaped (lamellarized) ferrite and retained austenite, respectively; α<sub>G</sub> and γ<sub>G</sub> are globular-shaped (equiaxed) ferrite and retained austenite, respectively. (<b>c</b>) Thermal desorption analysis (TDA) curves of the specimens pre-charged with H under the same charging conditions. (<b>d</b>) Engineering stress–strain curves obtained by SSRT. Reproduced from Reference [<a href="#B36-metals-11-00358" class="html-bibr">36</a>], with permission from Elsevier.</p> Full article ">Figure 4
<p>Effect of various H-charging conditions on SSRT properties of intercritically annealed medium-Mn steels (Fe–0.11C–7.2Mn–1.0Si, in wt pct) with (<b>a</b>) lamellarized (M900) and (<b>b</b>) equiaxed morphologies (M820). (<b>c</b>) H-induced elongation loss as a function of the diffusible H content. (<b>d</b>) Change of austenite fraction during tensile deformation. The specimen M900 was austenitized at 900 °C for 10 min, cooled to room temperature, and reheated to an IA temperature of 650 °C for 4 min. The specimen M820 was austenitized at 820 °C for 10 min, cooled to room temperature, and reheated to an IA temperature of 650 °C for 2 min. Reproduced from Reference [<a href="#B10-metals-11-00358" class="html-bibr">10</a>], with permission from Elsevier.</p> Full article ">Figure 5
<p>(<b>a</b>–<b>c</b>) Electron backscatter diffraction (EBSD) phase map for microstructures of a medium-Mn steel (Fe–0.01C–9Mn–3Ni–1.4Al, in wt pct). (<b>a</b>) As-quenched martensite (<span class="html-italic">M</span>). Austenite-reversion-treated microstructures consisting of martensite and austenite films with an average thickness of (<b>b</b>) 200 nm (<span class="html-italic">MA<sub>200nm</sub></span>) and (<b>c</b>) 500 nm (<span class="html-italic">MA<sub>500nm</sub></span>). (<b>d</b>) H–thermal–desorption analysis results for <span class="html-italic">M</span> and <span class="html-italic">MA<sub>500nm</sub></span> steel samples. (<b>e</b>) Engineering stress–strain curves showing SSRT properties. In (<b>d</b>,<b>e</b>), corresponding microstructures with H are referred to as (<span class="html-italic">M</span>)<span class="html-italic"><sub>H</sub></span>, (<span class="html-italic">MA<sub>200 nm</sub></span>)<span class="html-italic"><sub>H</sub></span>, and (<span class="html-italic">MA<sub>500nm</sub></span>)<span class="html-italic"><sub>H</sub></span>. Reproduced from Reference [<a href="#B42-metals-11-00358" class="html-bibr">42</a>], with permission from Springer Nature.</p> Full article ">Figure 6
<p>HE index, represented by elongation loss (EL<sub>loss</sub>) due to H, and H concentration as a function of the volume fraction of retained austenite in a medium-Mn steel (Fe–0.065C–0.2Si–5.45–Mn, in wt pct). Figure reproduced from Reference [<a href="#B43-metals-11-00358" class="html-bibr">43</a>], with permission from Elsevier.</p> Full article ">Figure 7
<p>(<b>a</b>) Variations in the volume fraction of retained austenite in a medium-Mn steel (Fe–0.2C–5.0Mn–0.6Si–3Al, in wt pct) as a function of engineering strain and hold time (10–360 min) at an IA temperature of 750 °C. (<b>b</b>) HE index (relative elongation loss due to H) evaluated by SSRT testing as a function of IA time. Reproduced from Reference [<a href="#B44-metals-11-00358" class="html-bibr">44</a>], with permission from Elsevier.</p> Full article ">Figure 8
<p>EBSD phase maps for BCC (white) and FCC (red) of (<b>a</b>) low–Al/medium-Mn steel (L–Al: Fe–0.12C–4.6Mn–0.55Si–1.1Al, in wt pct) and (<b>b</b>) high-Al/medium-Mn steel (H–Al: Fe–0.12C–5.8Mn–0.47Si–3.1Al, in wt pct). (<b>c</b>) Engineering stress–strain curves of low-Al and (<b>d</b>) high-Al alloys pre-charged with an H concentration of 0 to 9 ppm, tested at a slow strain rate of 10<sup>5</sup> s<sup>−1</sup>. In (<b>a</b>,<b>b</b>), green and black lines indicate low-angle (misorientation of 2–15°) and high-angle (misorientation &gt;15°) boundaries, respectively. Reproduced from Reference [<a href="#B48-metals-11-00358" class="html-bibr">48</a>], with permission from Elsevier.</p> Full article ">Figure 9
<p>SEM fractographs of H–charged SSRT specimens of (<b>a</b>) an intercritically annealed and (<b>b</b>) an intercritically warm-rolled medium–Mn steel (Fe–0.20C–4.9Mn–3.1Al–0.6Si, in wt pct). Reproduced from Reference [<a href="#B50-metals-11-00358" class="html-bibr">50</a>], with permission from Elsevier.</p> Full article ">Figure 10
<p>Three–dimensional atom probe tomography (3D–APT) maps for Ni, Al, Mn, and Cu in (<b>a</b>) the ferritic and (<b>b</b>) austenitic regions in a Cu-added steel (Fe–7Mn–2.5Ni–1.5Al–1.5Cu–0.01C, in wt pct). (<b>c</b>) Engineering stress–strain curves before and after H-charging in Cu-free (Fe–7Mn–2.5Ni–1.5Al–0.01C, in wt pct) and Cu-added steels. These two steels were two–stage annealed; the first IA was performed at 630 °C for 1 h, and the second step involved “tempering” at 500 °C. The hold times in the “tempering” step used for the Cu–free and Cu–added steels were 5 and 2 h, respectively. (<b>d</b>) Average hardness, measured by nano-indentation, of the ferrite (α) and austenite (γ) in the Cu-free and Cu-added steels. Reproduced from Reference [<a href="#B56-metals-11-00358" class="html-bibr">56</a>], with permission from Elsevier.</p> Full article ">Figure 11
<p>(<b>a</b>) SEM fractograph of H–charged medium–Mn steel specimen with an equiaxed ferrite–austenite microstructure. (<b>b</b>) Enlargement of the area indicated by a yellow box in (<b>a</b>). Reproduced from Reference [<a href="#B36-metals-11-00358" class="html-bibr">36</a>], with permission from Elsevier.</p> Full article ">Figure 12
<p>Comparison of H-induced cracks that formed in (<b>a</b>–<b>c</b>) an intercritically annealed medium–Mn steel and (<b>d</b>) a 2205 duplex stainless steel, both tensile–tested during electrochemical H charging. Reproduced from References [<a href="#B49-metals-11-00358" class="html-bibr">49</a>,<a href="#B61-metals-11-00358" class="html-bibr">61</a>], with permission from Elsevier.</p> Full article ">
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