Hydrogen Embrittlement of Medium Mn Steels
<p>Schematic of representative thermo-mechanical processing routes for medium-Mn steels and the resulting microstructures. IA, intercritical annealing; ART, austenite-reverted transformation; Q&T, quenching and tempering; Q&P, quenching and partitioning; UFG, ultra–fine–grained; α′, α′–martensite; γ, austenite; α, α–ferrite; and δ, δ–ferrite.</p> "> 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> "> 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> "> 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> "> 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> "> 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> "> 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> "> 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 >15°) boundaries, respectively. Reproduced from Reference [<a href="#B48-metals-11-00358" class="html-bibr">48</a>], with permission from Elsevier.</p> "> 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> "> 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> "> 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> "> 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> ">
Abstract
:1. Introduction
2. Thermomechanical Processing and Metallurgy of Medium–Mn Steels
3. Alloying and Microstructural Effects on HE Characteristics
3.1. Equiaxed Versus Lamellarized Morphology
3.2. Retained Austenite and Mechanically–Induced Martensitic Transformation
3.3. Al- and Si-Alloyed Medium-Mn Steels Containing Coarse δ-Ferrite Grains
3.4. Other Alloying Elements and Precipitates
3.5. H–Induced Crack Initiation and Propagation
4. Alloying and Microstructural Engineering Strategies to Improve H–Resistance
- The microstructural morphology (equiaxed versus lamellarized) that has better HE resistance is inconclusively defined in the literature, as there are contradictory reports (Table 1). For the lamellarized microstructure, care must be taken to avoid solute segregation to prior austenite grain boundaries, as this microstructure may be prone to H-induced intergranular fracture along the prior austenite grain boundaries [36]. The equiaxed microstructure does not preserve its prior austenite grain structure. It is therefore postulated that the level of impurities in the steel may be more influential for the lamellarized morphology.
- The effect of micro alloy precipitates is not clearly established for medium–Mn “duplex” steel, and future study is thus needed to evaluate their potential to mitigate HE. There are some reports on the beneficial effect of Cu–rich precipitates with respect to H-resistance.
- In general, the HE characteristics of medium–Mn steels are governed by the volume fraction and mechanical stability of retained austenite. Effective alloy and process design should target a sufficiently high fraction of retained austenite with a high mechanical stability and/or SFE, to avoid α′- or ε-martensitic transformation or even suppress planar slip. A finer grain size is often found to help mechanically stabilize the austenite.
- In theory, alloying elements that increase the SFE are expected to improve the HE resistance. C, Mn, Al, and Ni increase the SFE of austenite significantly. C is the most powerful in increasing SFE, but the amount is often limited below 0.6 wt pct, due to concerns related to the weldability or C segregation during casting [3]. Sufficiently high Al (≥3 wt pct) may be helpful to increase SFE of austenite, while stabilizing ferrite. Many investigations have focused on medium-Mn steels containing approximately 3 wt pct Al, as a greater amount of Al may cause difficulties during melting, secondary refining, and casting [3]. Cu appears to slightly increase SFE when in solid solution.
5. Conclusions
Author Contributions
Funding
Institutional Review Board Statement
Informed Consent Statement
Data Availability Statement
Acknowledgments
Conflicts of Interest
References
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Chemical Composition (wt.%) | Product Type or Starting Microstructure | Heat Treatment | Microstructure (Morphology and Austenite Fraction) | H content, wppm | H-induced Elongation Loss, Pct | Remarks | Ref. |
---|---|---|---|---|---|---|---|
Fe-0.01C -9Mn-3Ni- 1.4Al | Plate-type: as-quenched martensite | IA or ART at 600 C for 8 h | Lamellarized α′ + γ (33–36 vol pct) | 1.87 | ~92 | “Equiaxed” absorbed much greater H content for a given H-charging condition. | Cameron et al. [35] |
Cold-rolled, martensite | IA at 600 °C for 1 h | Equiaxed α + γ (~40 vol pct) | 15.6 | ~91 | |||
Fe-0.1C- 7Mn-0.5Si | Hot-rolled | IA at 640 °C for 30 min | Lamellarized α′ + γ (47 vol pct) | ~1.2 | ~87 | “Equiaxed” had a higher ultimate tensile strength and was more H-resistant than lamellarized. | Han et al. [36] |
Cold-rolled | IA at 640 °C for 30 min | Equiaxed α + γ (50 vol pct) | ~1.2 | ~74 | |||
Fe-0.06C- 11.7Mn-2.9Al-0.2Si | Cold-rolled | IA at 675 °C for 2 h | Larger, mixed lamellarized and equiaxed α + γ (55.2 vol pct) | 3.1 | ~58 | “Larger mixed” microstructure was more H-resistant than finer, lamellarized condition. | Shen et al. [37] |
10.0 | ~75 | ||||||
25.9 | ~83 | ||||||
Cold-rolled | Aus. at 800 °C for 20 min + IA at 650 °C for 15 min | Finer, lamellarized α + γ (53.1 vol pct) | 2.4 | ~86 | |||
7.6 | ~87 | ||||||
34.6 | ~87 | ||||||
Fe-0.11C- 7.2Mn-1.0Si | Cold-rolled | Aus. at 900 °C for 10 min + IA at 650 °C for 4 min | Lamellarized α + γ (32 vol pct) | 0.4 | 0 | “Lamellarized” was more H-resistant than equiaxed. Tested with samples having similar austenite fraction and mechanical stability. | Jeong et al. [10] |
0.9 | ~1 | ||||||
1.6 | ~3 | ||||||
2.6 | ~50 | ||||||
3.7 | ~85 | ||||||
4.2 | ~90 | ||||||
Cold-rolled | Aus. at 820 °C for 10 min + IA at 650 °C for 2 min | Equiaxed α + γ (32 vol pct) | 0.5 | ~38 | |||
1.5 | ~54 | ||||||
2.0 | ~75 | ||||||
3.4 | ~95 | ||||||
4.1 | ~98 | ||||||
4.4 | ~98 |
Product Type | Chemical Composition (wt.%) | IA Temperature /Hold Time | Microstructure (Morphology and Austenite Fraction) | H Content, Wppm | H-Induced Elongation Loss, Pct | Authors’ Interpretation | Ref. |
---|---|---|---|---|---|---|---|
Cold-rolled | Fe-0.12C- 4.6Mn- 0.55Si- 1.1Al | 720 °C/2 min | Equiaxed α + γ (26 vol%) | 0.1 | ~16 | HE is more pronounced for the low‑Al alloy containing less stable austenite. Martensitic decomposition of the austenite leaves the inherited H in a more mobile state. | Ryu et al. [48] |
0.6 | ~56 | ||||||
2.0 | ~77 | ||||||
3.3 | ~87 | ||||||
Fe-0.12C- 5.8Mn- 0.47Si- 3.1Al | 780 °C/2 min | Equiaxed α + γ (30 vol%) + coarse δ | 1.0 | ~31 | |||
1.2 | ~47 | ||||||
2.2 | ~65 | ||||||
3.1 | ~72 | ||||||
4.0 | ~78 | ||||||
6.1 | ~92 | ||||||
9.0 | ~96 | ||||||
Hot-rolled | Fe-0.22C- 6.1Mn- 3.1Al | 740 °C/3 min | Lamellarized α′ + γ (24.8 vol%) | 3.9 | 13.5 | The presence of δ can promote Mn enrichment in reverted γ. H-resistance increases with increasing stability and fraction of γ. H-induced cracking occurs along the boundaries of δ and UFG regions. | Wang et al. [49] |
5.2 | 25.8 | ||||||
7.9 | 39.8 | ||||||
740 °C/30 min | Lamellarized α′ + γ (37.4 vol%) | 3.2 | 79.2 | ||||
4.1 | 82.1 | ||||||
7.4 | 88.2 | ||||||
Fe-0.18C- 6.1Mn- 2.9Al-0.6Si | 740 °C/3 min | Lamellarized α′ + γ (15.2 vol%) + coarse δ | 2.2 | 46.7 | |||
2.8 | 68.3 | ||||||
5.8 | 70.3 | ||||||
740 °C/30 min | Lamellarized α′ + γ (31.4 vol%) + coarse δ | 5.0 | 76.5 | ||||
6.2 | 91.9 | ||||||
7.2 | 89.1 | ||||||
Warm-rolled at IA temperature | 0.20C- 5.0Mn- 3.0Al-0.6Si | 750 °C/10 min | Equiaxed α + γ (33.1 vol%) + coarse δ | 1.4 | ~16 | HE became increasingly significant with increasing γ grain size. H-resistance relates to the γ mechanical stability. | Shao et al. [44] |
750 °C/1 h | Equiaxed α + γ (34.2 vol%) + coarse δ | 1.3 | ~24 | ||||
750 °C/6 h | Equiaxed α + γ (35.7 vol%) + coarse δ | 1.1 | ~35 | ||||
Hot-rolled | 0.20C-4.9Mn-3.1Al-0.6Si | 750 °C/1 h | Lamellarized α′ + γ (~26 vol%) + coarse δ | 0.9 | ~78 | Warm rolling, i.e., fine lamellar structure, significantly enhances the H-resistance. | Zhang et al. [50] |
750 °C/1 h + 89%‑reduction warm-rolled | Fine, lamellarized α + γ (~15 vol%) + coarse δ | 1.6 | ~28 |
Chemical Composition (wt. %) | Product Type | Heat Treatment | Microstructure (Austenite Fraction) | Fracture Surface Appearances and Crack Initiation Sites | Refs. |
---|---|---|---|---|---|
Fe-0.1C- 7Mn-0.5Si | Hot-rolled | IA at 640 °C for 30 min | Lamellarized α′ + γ (47 vol%) | Cracking along prior γ grain boundaries. Rugged facets, likely associated with fracture of mechanically-induced α′. | Han et al. [36] |
Cold-rolled | IA at 640 °C for 30 min | Equiaxed α + γ (50 vol%) | Dimples with granular features.The granular features are likely associated with intergranular cracking along equiaxed γ grain boundaries. | ||
Fe-0.22C- 6.1Mn-3.1Al | Hot-rolled | IA at 740 °C for 3 min and 30 min | Lamellarized α′ + γ (24.8–37.4 vol%) | Cracking preferentially along γ/α phase boundaries. Cracking along prior γ grain boundaries or across the lamellar structure. | Wang et al. [49] |
Fe-0.18C- 6.1Mn-2.9Al-0.6Si | Hot-rolled | IA at 740 °C for 3 min and 30 min | Lamellarized α′ + γ (15.2–31.4 vol%) + coarse δ | Cracking at the phase boundaries, preferentially along (γ or α)/δ phase boundaries. | |
0.20C-5.0Mn-3.0Al-0.6Si | Warm-rolled at IA temperature | IA at 750 °C for 10 min, 1 h, and 6 h | Equiaxed α + γ (33.1–35.7 vol%) + coarse δ | Dimples with granular features. The granular features likely associated with cracking in the region of mechanically-induced α′. | Shao et al. [44] |
0.20C-4.9Mn-3.1Al-0.6Si | Hot-rolled | IA at 750 °C for 1 h | Lamellarized α′ + γ (~26 vol%) + coarse δ | Cracking across the lamellar structure or along γ/α′ phase boundaries. A few cracks along prior γ grain boundaries. | Zhang et al. [50] |
IA at 750 °C for 1 h + 89% reduction warm rolled | Fine, lamellarized α′ + γ (~15 vol%) + coarse δ | Micro-delamination cracking at γ/α′ interfaces along the rolling direction. Larger-scale crack deflections near δ-ferrite layers. |
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Cho, L.; Kong, Y.; Speer, J.G.; Findley, K.O. Hydrogen Embrittlement of Medium Mn Steels. Metals 2021, 11, 358. https://doi.org/10.3390/met11020358
Cho L, Kong Y, Speer JG, Findley KO. Hydrogen Embrittlement of Medium Mn Steels. Metals. 2021; 11(2):358. https://doi.org/10.3390/met11020358
Chicago/Turabian StyleCho, Lawrence, Yuran Kong, John G. Speer, and Kip O. Findley. 2021. "Hydrogen Embrittlement of Medium Mn Steels" Metals 11, no. 2: 358. https://doi.org/10.3390/met11020358