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Volume 15
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Metals, Volume 15, Issue 3 (March 2025) – 77 articles

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17 pages, 6717 KiB  
Article
Seawater Corrosion of the Anodized A1050 Aluminum Plate for Heat Exchangers
by Hirofumi Arima
Metals 2025, 15(3), 300; https://doi.org/10.3390/met15030300 (registering DOI) - 9 Mar 2025
Abstract
To confirm the suitability of aluminum for the heat transfer surfaces as a heat exchanger material for ocean thermal energy conversion, the seawater corrosion resistance of aluminum plates in a plate heat exchanger was experimentally investigated. In this study, four different surface shapes [...] Read more.
To confirm the suitability of aluminum for the heat transfer surfaces as a heat exchanger material for ocean thermal energy conversion, the seawater corrosion resistance of aluminum plates in a plate heat exchanger was experimentally investigated. In this study, four different surface shapes with chevron angles of 45° and 60° and different treatment types of A1050 aluminum heat transfer surfaces were processed into herringbone patterns. Additionally, the surfaces of the test plates were either anodized or untreated. In continuously flowing deep ocean water, the surface conditions of the test plates were observed at 1, 3, 6, and 12 months using mass measurements, visual inspection, laser microscopy, and SEM. For the anodized A1050 plates, regardless of the surface shape, there was almost no change in the mass, laser microscopy, or SEM results even after 12 months. In contrast, the untreated plate mass decreased in the samples after 3 months or later, and the mass reduction rate was approximately 2–7%. In conclusion, untreated aluminum is not suitable for use in seawater and an anodizing treatment is necessary for its use in heat exchangers for ocean thermal energy conversion. Full article
13 pages, 4546 KiB  
Article
Efficient and Green Flotation Separation of Molybdenite from Chalcopyrite Using 1-Thioglycerol as Depressant
by Feng Jiang, Shuai He, Wei Sun, Yuanjia Luo and Honghu Tang
Metals 2025, 15(3), 299; https://doi.org/10.3390/met15030299 (registering DOI) - 9 Mar 2025
Viewed by 131
Abstract
The effective and environmental separation of chalcopyrite and molybdenite has long presented a challenge in mineral processing due to their similar floatability and close association at room temperature. This study explores the non-toxic 1-thioglycerol (1-TG) as a selective depressant for chalcopyrite–molybdenite flotation separation. [...] Read more.
The effective and environmental separation of chalcopyrite and molybdenite has long presented a challenge in mineral processing due to their similar floatability and close association at room temperature. This study explores the non-toxic 1-thioglycerol (1-TG) as a selective depressant for chalcopyrite–molybdenite flotation separation. An impressive separation effect was realized through single-mineral and mixed-mineral flotation experiments when using 1-TG as a depressant and kerosene as a collector. Contact angle measurements, zeta potential tests, and Fourier transform infrared spectroscopy (FT-IR) confirm the selective adsorption of 1-TG on the chalcopyrite surface, leading to enhanced surface hydrophilicity and the inhibition of collector adsorption. The depression mechanism is further elucidated through X-ray photoelectron spectroscopy (XPS), which demonstrates that it occurs via chemosorption between the thiol group in 1-TG and active iron sites on the chalcopyrite surface. These findings provide a potential efficient depressant for chalcopyrite–molybdenite flotation separation with low dosage, environmental friendliness, and human harmlessness. Full article
(This article belongs to the Special Issue Advances in Flotation Separation and Mineral Processing)
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<p>XRD patterns of (<b>a</b>) chalcopyrite and (<b>b</b>) molybdenite samples.</p> Full article ">Figure 2
<p>Flotation recovery rates of chalcopyrite and molybdenite vary with the (<b>a</b>) concentration of 1-TG and (<b>b</b>) pH value (2.8 × 10<sup>−4</sup> mol/L 1-TG, 2.5 × 10<sup>−4</sup> mol/L kerosene, and 4.5 × 10<sup>−4</sup> mol/L terpineol).</p> Full article ">Figure 3
<p>Recovery rate and grade of flotation concentrate for artificially mixed ore were measured at room temperature.</p> Full article ">Figure 4
<p>Contact angles of chalcopyrite and molybdenite before and after treatment with various reagents.</p> Full article ">Figure 5
<p>Zeta potentials of (<b>a</b>) chalcopyrite and (<b>b</b>) molybdenite versus pH value.</p> Full article ">Figure 6
<p>FT-IR spectra of (<b>a</b>) kerosine and 1-TG; (<b>b</b>) chalcopyrite and (<b>c</b>) molybdenite treated with different reagents.</p> Full article ">Figure 7
<p>(<b>a</b>) Cu 2p<sub>3/2</sub>, (<b>b</b>) Fe 2p<sub>3/2</sub>, and (<b>c</b>) S 2p high-resolution XPS spectra of chalcopyrite.</p> Full article ">Figure 8
<p>(<b>a</b>) Mo 3d and (<b>b</b>) S 2p high-resolution XPS spectra of molybdenite.</p> Full article ">Figure 9
<p>Illustration of the probable adsorption mechanism of 1-TG on chalcopyrite surface.</p> Full article ">
15 pages, 19069 KiB  
Article
Effect of Deep Cryogenic Treatment on Microstructure and Mechanical Properties of Friction Stir Welded TRIP590 Steel Joints
by Yashuai Hu, Weidong Liu, Liguo Wang, Yufeng Sun, Wenbo Cao and Shaokang Guan
Metals 2025, 15(3), 298; https://doi.org/10.3390/met15030298 (registering DOI) - 9 Mar 2025
Viewed by 107
Abstract
In this study, friction stir welding was first applied to the 1.4 mm thick TRIP590 steel sheets at a constant transverse speed of 100 mm/min and different rotation speeds from 200 to 500 rpm. Then, the obtained joints received deep cryogenic treatment in [...] Read more.
In this study, friction stir welding was first applied to the 1.4 mm thick TRIP590 steel sheets at a constant transverse speed of 100 mm/min and different rotation speeds from 200 to 500 rpm. Then, the obtained joints received deep cryogenic treatment in liquid nitrogen for 24 and 48 h, respectively. It was revealed that the content of retained austenite in the stir zone of the welded joints decreased from 3.3% to 0.2% when the rotation speed increased from 200 rpm to 500 rpm. The stability of retained austenite increased due to grain refinement and work hardening at low rotation speeds. After deep cryogenic treatment of the welded joints, the retained austenite in the stir zone partially transformed into martensite, which led to the precipitation of nano-sized carbide in the ferrite matrix and the release of local stress. As a result, both the strength and plasticity of the stir zone after 48 h of deep cryogenic treatment increased from 798 MPa, 15% to 927 MPa, 17% for the 200 rpm joint, and from 914 MPa, 14% to 1086 MPa, 16% for the 300 rpm joint during the tensile tests. Full article
(This article belongs to the Special Issue Editorial Board Members’ Collection Series: “Welding and Joining”)
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<p>Dimensions of tensile specimens (<b>a</b>) global; and (<b>b</b>) miniature tensile specimens.</p> Full article ">Figure 2
<p>The top surface of the samples welded at various rotation speeds (<b>a</b>) 200 rpm; (<b>b</b>) 300 rpm; (<b>c</b>) 400 rpm; (<b>d</b>) 500 rpm.</p> Full article ">Figure 3
<p>Overview of the transverse cross-section of the joints made at different rotation speeds (<b>a</b>) 200 rpm; (<b>b</b>) 300 rpm; (<b>c</b>) 400 rpm; (<b>d</b>) 500 rpm.</p> Full article ">Figure 4
<p>(<b>a</b>) SEM and (<b>b</b>) EBSD image-quality map of BM (F: Ferrite; B: Bainite).</p> Full article ">Figure 5
<p>Typical TEM images of RA in TRIP steels. (<b>a</b>) lengthy RA; (<b>b</b>,<b>c</b>) equiaxial RA at grain boundary.</p> Full article ">Figure 6
<p>SEM images showing the microstructure of joint SZs. (<b>a</b>) 200 rpm; (<b>b</b>) 300 rpm; (<b>c</b>) 400 rpm; and (<b>d</b>) 500 rpm (F: Ferrite; M: Martensite).</p> Full article ">Figure 7
<p>EBSD image-quality map showing the distribution of RA in the SZ. (<b>a</b>) 200 rpm; (<b>b</b>) 300 rpm; (<b>c</b>) 400 rpm; and (<b>d</b>) 500 rpm.</p> Full article ">Figure 8
<p>Bright field TEM images of SZ. (<b>a</b>) 200 rpm; (<b>1</b>,<b>2</b>) the SAED patterns from the white dashed circle in (<b>a</b>); (<b>b</b>) 300 rpm; (<b>3</b>,<b>4</b>) the SAED patterns from the white circle in (<b>b</b>) (PF: Polygonal Ferrite).</p> Full article ">Figure 9
<p>Microstructure after DCT. (<b>a</b>) BM, 24 h; (<b>b</b>) BM, 48 h; (<b>c</b>) 200 rpm, 24 h; (<b>d</b>) 200 rpm, 48 h; (<b>e</b>) 300 rpm, 24 h; (<b>f</b>) 300 rpm, 48 h (M: Martensite).</p> Full article ">Figure 10
<p>EBSD image-quality map showing RA phase after DCT. (<b>a</b>) BM, 24 h; (<b>b</b>) BM, 48 h; (<b>c</b>) 200 rpm, 24 h; (<b>d</b>) 200 rpm, 48 h; (<b>e</b>) 300 rpm, 24 h; (<b>f</b>) 300 rpm, 48 h.</p> Full article ">Figure 11
<p>Bright field TEM images of the 200 rpm joint SZ after DCT for (<b>a</b>) 24 h; (<b>b</b>) 48 h.</p> Full article ">Figure 12
<p>Bright field and dark field images of the 200 rpm joint SZ after DCT for (<b>a</b>,<b>b</b>) 24 h; (<b>c</b>,<b>d</b>) 48 h.</p> Full article ">Figure 13
<p>Kernel average misorientation maps (<b>a</b>) BM, 0 h; the 200 rpm joint SZ after DCT for (<b>b</b>) 0 h; (<b>c</b>) 24 h; and (<b>d</b>) 48 h.</p> Full article ">Figure 14
<p>Microhardness profile of the welds at different transverse speeds.</p> Full article ">Figure 15
<p>Microhardness variation on the center of the cross-section of joints treated after DCT. (<b>a</b>) 200 rpm; and (<b>b</b>) 300 rpm joints before and after DCT.</p> Full article ">Figure 16
<p>Tensile properties of BM and the FSW joints. (<b>a</b>) stress-strain curves, (<b>b</b>) fracture location of the joints.</p> Full article ">Figure 17
<p>Tensile stress-strain curves of SZs at different rotation speeds.</p> Full article ">Figure 18
<p>Tensile stress-strain curves of (<b>a</b>) BM; (<b>b</b>) 200 rpm; and (<b>c</b>) 300 rpm joint SZ after DCT with various time.</p> Full article ">
16 pages, 7742 KiB  
Article
Study on the Effects of Cryogenic Treatment on WC-Co Cemented Carbide at Different Scales Using an Indentation Technique
by Suparoj Premjarunan, Karuna Tuchinda and Kaweewat Worasaen
Metals 2025, 15(3), 297; https://doi.org/10.3390/met15030297 (registering DOI) - 8 Mar 2025
Viewed by 55
Abstract
Cemented carbide (WC-Co) combines high hardness, wear resistance, and toughness, making it ideal for tooling applications. This study investigated cryogenic treatment’s effects on the mechanical properties of samples from various suppliers prepared at different scales. Indentation tests were performed to assess the mechanical [...] Read more.
Cemented carbide (WC-Co) combines high hardness, wear resistance, and toughness, making it ideal for tooling applications. This study investigated cryogenic treatment’s effects on the mechanical properties of samples from various suppliers prepared at different scales. Indentation tests were performed to assess the mechanical properties at the microscale and nanoscale. Overall, the mean microhardness did not show a significant change after cryogenic treatment. Instead, nanoindentation testing was used to identify the improvement after cryogenic treatment. However, considering the mean nanohardness may not adequately capture improvements in the material’s resistance to deformation, the maximum nanoindentation depth and nanohardness were analyzed to elucidate the mechanisms underlying mechanical property improvements in the form of histograms of %frequency along with load–displacement curves. The results showed a decreased frequency of high maximum indentation depths from Co phase improvement. This agreed with an increased frequency of moderate and high nanohardness and a decreased frequency of low nanohardness representing different areas with different phase controls. These results indicate that an alternative interpretation of nanoindentation data, presenting nanohardness and nanoindentation depth in the form of histograms, can provide a more detailed representation of the data distribution. Full article
(This article belongs to the Special Issue Microstructure and Characterization of Metal Matrix Composites)
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<p>Drawing of samples: (<b>a</b>) bulk sample, B; (<b>b</b>) high-precision sample, P (dimensions in mm).</p> Full article ">Figure 2
<p>Arrangement for temperature sensor positions inside the cryogenic chamber (dimensions in mm, arrows show temperature sensor location for process control).</p> Full article ">Figure 3
<p>SEM image for B sample (<b>a</b>) before cryogenic treatment and (<b>b</b>) after cryogenic treatment.</p> Full article ">Figure 4
<p>SEM image for P sample (<b>a</b>) before cryogenic treatment and (<b>b</b>) after cryogenic treatment.</p> Full article ">Figure 5
<p>Grain size distribution of B samples before cryogenic treatment (NCT).</p> Full article ">Figure 6
<p>Grain size distribution of B samples after deep cryogenic treatment (DCT).</p> Full article ">Figure 7
<p>Grain size distribution of P samples before cryogenic treatment (NCT).</p> Full article ">Figure 8
<p>Grain size distribution of P samples after deep cryogenic treatment (DCT).</p> Full article ">Figure 9
<p>Load–displacement curves of B samples before (NCT) and after cryogenic treatment (DCT).</p> Full article ">Figure 10
<p>Examples of load–displacement curves of P samples before (NCT) and after deep cryogenic treatment (DCT).</p> Full article ">Figure 11
<p>Maximum nanoindentation depths of bulk samples (B) before (NCT) and after cryogenic treatment (DCT).</p> Full article ">Figure 12
<p>Maximum nanoindentation depths of high-precision samples (P) before and after cryogenic treatment.</p> Full article ">Figure 13
<p>Nanohardness of B samples.</p> Full article ">Figure 14
<p>Nanohardness of P samples.</p> Full article ">Figure 15
<p>AFM image of carbide phase indentation (<b>a</b>) before and (<b>b</b>) after indentation of B sample before deep cryogenic treatment.</p> Full article ">Figure 16
<p>AFM image of carbide phase indentation (<b>a</b>) before and (<b>b</b>) after indentation of B sample after deep cryogenic treatment.</p> Full article ">
17 pages, 5934 KiB  
Article
In Situ Observation by X-Ray Radioscopy of Liquid Decomposition During Directional Solidification of Al-Cu-Sn Alloys
by Sarah De Albuquerque, Guillaume Reinhart, Hadjer Soltani, Danielle Cristina Camilo Magalhães, José Eduardo Spinelli and Henri Nguyen-Thi
Metals 2025, 15(3), 296; https://doi.org/10.3390/met15030296 - 7 Mar 2025
Viewed by 287
Abstract
Immiscible Al–Sn–Cu alloys may offer attractive properties, attaining superior tribological and mechanical properties when Sn-rich soft particles are homogeneously distributed in the reinforced Al–Cu matrix. In this paper, the solidifications of both Al-10 wt.% Cu-10 wt.% Sn and Al-10 wt.% Cu-20 wt.% Sn [...] Read more.
Immiscible Al–Sn–Cu alloys may offer attractive properties, attaining superior tribological and mechanical properties when Sn-rich soft particles are homogeneously distributed in the reinforced Al–Cu matrix. In this paper, the solidifications of both Al-10 wt.% Cu-10 wt.% Sn and Al-10 wt.% Cu-20 wt.% Sn alloys were investigated to analyze the successive stages that occur during the controlled cooling of these alloys, from the initial formation of the α-Al dendritic array to the final eutectic reaction. In particular, we focus on the liquid-phase demixing occurring during the solidification path, which leads to the formation of Sn droplets in the melt through a nucleation-growth process. Horizontal directional solidifications were performed on thin samples in a Bridgman-type furnace, with in situ and real-time observations using X-ray radioscopy. Two different behaviors have been found concerning liquid separation: for the low-Sn-content alloy, liquid demixing occurs in one single step, whereas for the high-Sn-content alloy, it is a two-step process, with first the nucleation of a few small Sn droplets followed by a sudden formation of a large amount of wide Sn droplets. The possible causes of these different behaviors are discussed in relation to the literature, namely, either a switch from immiscible to miscible liquids or a transition from the binodal region to the spinodal region. Full article
(This article belongs to the Special Issue Development, Characterization and Properties of High-Performance Aluminum Alloys)
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<p>(<b>a</b>) Scheme of the furnace with its components; (<b>b</b>) scheme of the imaging system showing the distances of the source-sample (5 mm) and source-detector (25 mm) used to determine the geometric magnification factor.</p> Full article ">Figure 2
<p>Examples of images after the two image processing procedures used to improve the legibility of the images. Radiographs of Al–10Cu–20Sn alloy solidification (<span class="html-italic">R</span> = 0.15 °C/s and <span class="html-italic">G<sub>app</sub></span> = 5.55 °C/mm): (<b>a</b>) flat-field correction revealing the growth microstructure; (<b>b</b>) frame-differencing procedure showing the solidified α-Al during a time interval of 5 s.</p> Full article ">Figure 3
<p>Equilibrium solidification paths of the two investigated alloy compositions: (<b>a</b>) Al-10Cu-10Sn; (<b>b</b>) Al-10 Cu-20Sn. Calculated by Thermo-Calc software.</p> Full article ">Figure 4
<p>Series of isothermal sections of the ternary phase diagram for Al-Cu-Sn alloy, calculated at several temperatures by Thermo-Calc software: (<b>a</b>) T = 625 °C; (<b>b</b>) T = 570 °C; (<b>c</b>) T = 540 °C; (<b>d</b>) T = 529.5 °C; (<b>e</b>) T = 520 °C; (<b>f</b>) T = 229.4 °C. The red and blue dots indicate the positions of the Al-10Cu-10Sn and Al-10Cu-20Sn alloy compositions in the isothermal sections.</p> Full article ">Figure 5
<p>DSC cooling curves at a rate of 5 °C/min of the two Al-Cu-Sn alloys displaying the exothermic peaks of each transformation (α-Al dendrite growth, liquid-phase separation, monotectic reaction, and eutectic reaction). For Al-10Cu-10Sn (red line), the transformation temperatures are T(α-Al) = 600.5 °C; T (liquid-phase separation) = 525.5 °C; T (monotectic reaction) = 510.5 °C; and T (eutectic reaction) = 220.5 °C. For Al-10Cu-20Sn (black line), the corresponding temperatures are T(α-Al) = 572.8 °C; T (liquid-phase separation) = 517.8 °C; T (monotectic reaction) = 530.3 °C; and T (eutectic reaction) = 220.3 °C.</p> Full article ">Figure 6
<p>Sequence of radiographs processed by flat-field image processing technique showing the early stages of solidification experiments of Al-10Cu-10Sn (first row: (<b>a1</b>–<b>c1</b>)) and Al-10Cu-20Sn (second row: (<b>a2</b>–<b>c2</b>)) with <span class="html-italic">R</span> = −0.15 °C/s and <span class="html-italic">G<sub>app</sub></span> = 5.55 °C/mm. For each image, the average temperature <span class="html-italic">T<sub>avg</sub></span> at the center of the field of view is given. In these images, α-Al dendrites appear in white surrounded by a darker gray solute-rich liquid. The white dashed lines in (<b>b1</b>) and (<b>b2</b>) indicate the mean position of the dendrite tips.</p> Full article ">Figure 7
<p>Sequence of three radiographs (<b>a</b>–<b>c</b>) after the frame-differencing processing revealing the liquid-phase separation followed by the monotectic reaction during the directional solidification of Al-10Cu-10Sn (<span class="html-italic">R</span> = −0.15 °C/s, <span class="html-italic">G<sub>app</sub></span> = 5.55 °C/mm). <span class="html-italic">T<sub>avg</sub></span> is the temperature at the center of the field of view. The white dashed line in Figure (a) indicates the altitude where the first droplets are detected. (<b>d</b>) Vertical average gray level profile along the field of view in arbitrary unit (a.u.).</p> Full article ">Figure 8
<p>Sequence of three radiographs (<b>a</b>–<b>c</b>) after the frame-differencing processing revealing the liquid-phase separation followed by the monotectic reaction during the directional solidification of Al-10Cu-20Sn (<span class="html-italic">R</span> = −0.15 °C/s, <span class="html-italic">G<sub>app</sub></span> = 5.55 °C/mm). <span class="html-italic">T<sub>avg</sub></span> is the temperature at the center of the field of view. The white dashed line in Figure (a) indicates the altitude where the first droplets are detected. (<b>d</b>) Vertical average gray level profile along the field of view in arbitrary unit (a.u.). The initial stage of liquid demixing, which gives rise to the formation of small L″ droplets, is observable in <a href="#metals-15-00296-f006" class="html-fig">Figure 6</a>a. In the second stage, the abrupt generation of a copious quantity of the Sn-rich liquid L″ is discernible from t = 65 s.</p> Full article ">Figure 9
<p>(Reproduced from Schaffer et al. [<a href="#B21-metals-15-00296" class="html-bibr">21</a>]) Al–6 wt.% Bi–8 wt.% Zn alloy solidified in a thermal gradient of 60 K/mm and at a velocity of 17.5 µm/s. The transition from immiscible to miscible liquids when the Zn-rich boundary layer approaches can be seen from t = 0 to 16.8 s. As the diffuse Bi domains move closer to the monotectic front, Zn concentration reduces, and immiscibility re-establishes leading to secondary nucleation of Zn droplets (t = 25.2 to 42 s). Image size corresponds to 1.3 × 1.3 mm<sup>2</sup>. © Deutsche Physikalische Gesellschaft. Reprinted with permission from ref. [<a href="#B21-metals-15-00296" class="html-bibr">21</a>]. 2008 IOP Publishing. CC BY-NC-SA.</p> Full article ">Figure 10
<p>Typical temperature–composition graph for a binary alloy, showing the miscibility gap and spinodal line in a regular solution system. Uniform liquid within the spinodal curve is unstable and can decompose without overcoming an energy activation barrier. Uniform liquid between the binodal and spinodal curves are metastable and decomposition must proceed by a process of nucleation and growth.</p> Full article ">
14 pages, 5541 KiB  
Article
Dendrite Structure Refinement and Mechanical Property Improvement of a Single-Crystal Superalloy
by Hongyuan Sun, Dexin Ma, Yunxing Zhao, Jianhui Wei, Xiaoyi Gong and Zhongyuan Sun
Metals 2025, 15(3), 295; https://doi.org/10.3390/met15030295 - 7 Mar 2025
Viewed by 98
Abstract
In the present work, the effect of different casting processes on the microstructure and creep properties of the second-generation single-crystal superalloy DD419 was investigated. Under conventional production conditions and a contour-suited thermal insulation method, single-crystal rods of types A and B were fabricated, [...] Read more.
In the present work, the effect of different casting processes on the microstructure and creep properties of the second-generation single-crystal superalloy DD419 was investigated. Under conventional production conditions and a contour-suited thermal insulation method, single-crystal rods of types A and B were fabricated, respectively. In comparison to rod type A, the solidification process of rod type B featured a 1.6-fold increase in the temperature gradient and a 32% reduction in primary dendrite spacing. The γ/γ′ eutectic in the as-cast microstructure, the residual eutectic phase, and porosity after heat treatment were also significantly reduced, resulting in the improved homogeneity of the single crystal castings. Under the testing conditions of 850 °C/650 MPa and 1050 °C/190 MPa, the stress rupture life of sample B was enhanced by 25% and 5.2%, respectively, compared to sample A. Therefore, due to dendrite structure refinement, the stress rupture life of the superalloy was evidently improved, especially at medium temperatures. Full article
(This article belongs to the Special Issue Research Progress of Crystal in Metallic Materials)
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<p>(<b>a</b>) Schematic of the conventional furnace condition for experiment A; (<b>b</b>) the internal insulation and contour-suited baffle used for experiment B.</p> Full article ">Figure 2
<p>The cross-sections of SC rod A (<b>a1</b>–<b>a3</b>) and rod B (<b>b1</b>–<b>b3</b>), showing the as-cast microstructure at the bottom, middle, and top portions, respectively. Diagram ruler: 100 μm.</p> Full article ">Figure 3
<p>The measured dendrite spacing (λ) at the bottom, middle, and top cross-sections of rods A and B.</p> Full article ">Figure 4
<p>The fraction of the worst γ/γ′ eutectic microstructures (f<sub>E</sub>) measured at the bottom, middle, and top of as-cast rods A and B.</p> Full article ">Figure 5
<p>The cross-sections of SC rod A (<b>a1</b>–<b>a3</b>) and rod B (<b>b1</b>–<b>b3</b>) after heat treatment, showing the worst residual eutectic microstructures at the bottom, middle, and top portions, respectively. Diagram ruler: 100 μm.</p> Full article ">Figure 6
<p>The average residual eutectic fraction f<sub>H</sub> measured on each cross-section of heat-treated rods A and B.</p> Full article ">Figure 7
<p>The residual eutectic fraction f<sub>W</sub> in the worst view fields on each cross-section of the heat-treated rods A and B.</p> Full article ">Figure 8
<p>Size distribution of residual eutectic in heat-treated samples A (<b>a</b>) and B (<b>b</b>).</p> Full article ">Figure 9
<p>The worst micro-porosity images on the bottom, middle, and top cross-sections of heat-treated samples A (<b>a1</b>–<b>a3</b>) and B (<b>b1</b>–<b>b3</b>), respectively.</p> Full article ">Figure 10
<p>The measured average micro-porosity ratio ρ on each cross-section of the heat-treated samples A and B.</p> Full article ">Figure 11
<p>The local porosity ratio in the worst view field of each cross-section of the heat-treated samples A and B.</p> Full article ">Figure 12
<p>The area distribution statistics of porosity holes in heat-treated samples A (<b>a</b>) and B (<b>b</b>).</p> Full article ">
14 pages, 4760 KiB  
Article
Machine Learning-Assisted Hardness Prediction of Dispersion-Strengthened Tungsten Alloy
by Shaowu Dai, Chong Chen, Cong Zhang, Shizhong Wei, Beibei Han, Changji Wang, Kunming Pan, Liujie Xu, Feng Mao and Hua Yu
Metals 2025, 15(3), 294; https://doi.org/10.3390/met15030294 - 7 Mar 2025
Viewed by 217
Abstract
Hardness, as a typical mechanical property of dispersion-strengthened tungsten alloy, is influenced by various coupled factors. This paper aims to identify the key factors affecting the hardness of the dispersion-strengthened tungsten alloys with different carbides and oxides as the reinforcement phase in order [...] Read more.
Hardness, as a typical mechanical property of dispersion-strengthened tungsten alloy, is influenced by various coupled factors. This paper aims to identify the key factors affecting the hardness of the dispersion-strengthened tungsten alloys with different carbides and oxides as the reinforcement phase in order to enable the high-throughput prediction of hardness. A dataset was established with alloy hardness as the target variable, and the features included the content of reinforcement phase, the Vickers hardness of reinforcement phase, the melting point of the reinforcement phase, the valence electron number of the reinforcement phase, the sintering temperature, the sintering time, pressure, relative density, and grain size. Seven regression models were trained, and we selected random forest, support vector regression, and XGBoost regression machine learning models with better performance to construct a hardness prediction model of the dispersion-strengthened tungsten alloy. SHAP analysis, based on random forests, shows that the content of reinforcement phase, grain size, and relative density have the most significant impact on the hardness. A random forest model is the most suitable machine learning method for predicting the hardness of dispersion-strengthened tungsten alloys in this work. The R2 values of the training and test sets are 0.93 and 0.80, and the MAE values of the training and test sets are 22.72 and 38.37. The influence of the most important features on the hardness was also discussed based on the random forest model. This study provides a data-driven approach for the accurate and efficient prediction of the hardness of dispersion-strengthened tungsten alloys, offering an important reference for the design and development of high-performance tungsten alloy materials. Full article
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<p>An illustration of the machine learning process for dispersion-strengthened tungsten alloys.</p> Full article ">Figure 2
<p>Diagram illustrating the 10-fold cross-validation used in this work. The blue rectangle represents the training set, and the green rectangle represents the test set.</p> Full article ">Figure 3
<p>A Pearson correlation coefficient diagram of the features affecting the hardness of the dispersion-strengthened tungsten alloy.</p> Full article ">Figure 4
<p>A comparison of R<sup>2</sup> and MAE for the seven regression models.</p> Full article ">Figure 5
<p>SHAP feature importance rank for target variable hardness.</p> Full article ">Figure 6
<p>Predicted hardness vs. actual hardness for RF algorithm model: (<b>a</b>) training set; (<b>b</b>) test set.</p> Full article ">Figure 7
<p>Predicted hardness vs. actual hardness for SVR algorithm model: (<b>a</b>) training set; (<b>b</b>) test set.</p> Full article ">Figure 8
<p>Predicted hardness vs. actual hardness for XGB model: (<b>a</b>) training set; (<b>b</b>) test set.</p> Full article ">Figure 9
<p>Comparison of model metrics of the training sets and test sets for the RF, SVR and XGBoost models: (<b>a</b>) R<sup>2</sup>; (<b>b</b>) MAE.</p> Full article ">Figure 10
<p>The hardness of tungsten alloy varies with the grain size and reinforcement phase content of different carbon oxides: (<b>a</b>) actual values; (<b>b</b>) predicted values.</p> Full article ">Figure 11
<p>A comparison of predicted hardness based on the current RF model with only RC and GZ features used and actual hardness, along with the percentage error.</p> Full article ">
11 pages, 4192 KiB  
Article
Arc Erosion Properties of the Ag-Cr2AlC Contact Material
by Xiaochen Huang, Jinlong Ge, Zijue Zhou, Bin Chen, Feng Zhuo and Hongdi Zhang
Metals 2025, 15(3), 293; https://doi.org/10.3390/met15030293 - 7 Mar 2025
Viewed by 118
Abstract
This study investigates the arc performance of Ag-Cr2AlC composite materials. Spark plasma sintering method was employed to prepare the Ag-Cr2AlC composite material. A self-made arc erosion device was utilized to erode the material with different times of arc. The [...] Read more.
This study investigates the arc performance of Ag-Cr2AlC composite materials. Spark plasma sintering method was employed to prepare the Ag-Cr2AlC composite material. A self-made arc erosion device was utilized to erode the material with different times of arc. The surface of the material was categorized into three distinct areas: the eroded center area, the eroded edge area, and the heat-affected area. After one time of arc erosion, the material exhibits a relatively flat surface with a small erosion area. However, after one hundred arc erosions, the eroded area has significantly increased, accompanied by numerous splashes, protrusions, and pores. The action of the arc leads to the decomposition and oxidation of the Ag-Cr2AlC composite material, resulting in the formation of Ag2O, Al2O3, and Cr2O3 on the surface. During the process of 100 arc erosions, the breakdown current value remains relatively stable, ranging from 20 to 35 A. From the first to the 70th arc erosion, the breakdown strength consistently varies between 3 × 106 V/m and 6 × 106 V/m. Subsequently, there is an observed enhancement in breakdown strength, leading to the appearance of ageing. These findings establish a theoretical foundation for the application of silver-based electrical contact materials. Full article
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<p>(<b>a</b>) X-ray diffraction (XRD), (<b>b</b>) scanning electron micrograph (SEM), (<b>c</b>) point-scanning results of Ag-Cr<sub>2</sub>AlC composite.</p> Full article ">Figure 2
<p>Micro-indentation (inside red boxes) images of Ag-Cr<sub>2</sub>AlC composites (<b>a</b>) test point a; (<b>b</b>) test point b; (<b>c</b>) test point c.</p> Full article ">Figure 3
<p>The surface of Ag-Cr<sub>2</sub>AlC composite eroded by (<b>a</b>) 1 time and (<b>b</b>) 100 times of arc erosion. (<b>c</b>,<b>d</b>) corresponding to the enlarged image inside circles 1 of (<b>a</b>,<b>b</b>), respectively.</p> Full article ">Figure 4
<p>SEM images of Ag-Cr<sub>2</sub>AlC composite surface after eroded for (<b>a</b>,<b>b</b>) 1 time; (<b>c</b>,<b>d</b>) 100 times.</p> Full article ">Figure 5
<p>(<b>a</b>–<b>d</b>) corresponding to point-scanning results of rectangles 1–4 in <a href="#metals-15-00293-f004" class="html-fig">Figure 4</a>.</p> Full article ">Figure 6
<p>Three-dimensional morphology of Ag-Cr<sub>2</sub>AlC composite surface after eroded for (<b>a</b>) 1 time and (<b>b</b>) 100 times.</p> Full article ">Figure 7
<p>Raman spectrum of Ag-Cr<sub>2</sub>AlC eroded with (<b>a</b>) 1 and (<b>b</b>) 100 times of arc erosion.</p> Full article ">Figure 8
<p>Curve of current-time after one discharge.</p> Full article ">Figure 9
<p>Curves of (<b>a</b>) breakdown current and (<b>b</b>) breakdown strength with the number of arc erosion.</p> Full article ">
17 pages, 11964 KiB  
Article
Effects of Heat Treatment on Microstructures and Corrosion Properties of AlxCrFeNi Medium-Entropy Alloy
by Pushan Guo, Yuan Pang, Qingke Zhang, Lijing Yang, Zhenlun Song and Yi Zhang
Metals 2025, 15(3), 292; https://doi.org/10.3390/met15030292 - 7 Mar 2025
Viewed by 69
Abstract
This study designed AlxCrFeNi (x = 0.8, 1.0, 1.2) medium-entropy alloys featuring a BCC + B2 dual-phase structure to systematically investigate the effects of Al content variation and heat treatment on microstructure evolution and corrosion behavior. Microstructural characterization revealed that [...] Read more.
This study designed AlxCrFeNi (x = 0.8, 1.0, 1.2) medium-entropy alloys featuring a BCC + B2 dual-phase structure to systematically investigate the effects of Al content variation and heat treatment on microstructure evolution and corrosion behavior. Microstructural characterization revealed that all investigated alloys maintained the BCC + B2 dual-phase labyrinth structure. Electrochemical tests showed that as the Al content increased, the corrosion current density and corrosion rate in a 3.5 wt% NaCl solution increased. Synergistic analysis of post-corrosion morphology (through electrochemical testing and in-situ immersion) combined with XPS analysis of the passive films revealed that the initial stage of corrosion was primarily pitting. Subsequently, due to the loose and porous Al2O3 passive layer formed by the NiAl-rich phase, which was easily attacked by Cl− ions, the corrosion progressed into selective corrosion of the NiAl phase. Notably, heat treatment at 1000 °C induced microstructural refinement with enhanced coupling between chunky and labyrinth structures, resulting in improved corrosion resistance despite a 4–6% reduction in Vickers hardness due to elemental homogenization. Among the investigated alloys, the heat-treated Al0.8CrFeNi exhibited the most promising corrosion resistance. Full article
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>The XRD pattern and partially enlarged view of (<b>a</b>) as-cast and (<b>b</b>) heat-treated Al<span class="html-italic"><sub>x</sub></span>CrFeNi Medium entropy alloys.</p> Full article ">Figure 2
<p>The microstructure of (<b>a</b>) Al<sub>0.8</sub>CrFeNi, (<b>b</b>) Al<sub>1.0</sub>CrFeNi, and (<b>c</b>) Al<sub>1.2</sub>CrFeNi medium entropy alloy.</p> Full article ">Figure 3
<p>Backscattered SEM images and phase elements proportion of (<b>a</b>) Al<sub>0.8</sub>CrFeNi, (<b>b</b>) Al<sub>1.0</sub>CrFeNi, and (<b>c</b>) Al<sub>1.2</sub>CrFeNi medium-entropy alloy.</p> Full article ">Figure 4
<p>Potentiodynamic polarization curves of as-cast and heat-treated Al<span class="html-italic"><sub>x</sub></span>CrFeNi alloys in 3.5 wt% NaCl corrosive solution.</p> Full article ">Figure 5
<p>Backscattered SEM surface morphologies of (<b>a</b>) Al<sub>0.8</sub>CrFeNi, (<b>b</b>) Al<sub>1.0</sub>CrFeNi, and (<b>c</b>) Al<sub>1.2</sub>CrFeNi medium entropy alloy after potentiodynamic polarization in 3.5 wt% NaCl corrosive solution at room temperature.</p> Full article ">Figure 6
<p>In-situ immersion corrosion morphologies of the as-cast (<b>a</b>) Al<sub>0.8</sub>CrFeNi, (<b>b</b>) Al<sub>1.0</sub>CrFeNi, and (<b>c</b>) Al<sub>1.2</sub>CrFeNi medium-entropy alloy immersed in 3.5 wt% NaCl corrosive solution at room temperature for 0 d, 3 d, 7 d, and 15 d.</p> Full article ">Figure 7
<p>In-situ immersion corrosion morphologies of the heat-treated (<b>a</b>) Al<sub>0.8</sub>CrFeNi, (<b>b</b>) Al<sub>1.0</sub>CrFeNi, and (<b>c</b>) Al<sub>1.2</sub>CrFeNi medium-entropy alloy immersed in 3.5 wt% NaCl corrosive solution at room temperature for 0 d, 3 d, 7 d, and 15 d.</p> Full article ">Figure 8
<p>High-resolution XPS spectra of Al 2p, Cr 2p, Fe 2p, Ni 2p, C 1s, and O 1s of passive film on the surfaces of the as-cast (<b>a</b>) Al<sub>0.8</sub>CrFeNi, (<b>b</b>) Al<sub>1.0</sub>CrFeNi, and (<b>c</b>) Al<sub>1.2</sub>CrFeNi medium entropy alloy immersed in 3.5 wt% NaCl corrosive solution for 15 d.</p> Full article ">Figure 9
<p>High-resolution XPS spectra of Al 2p, Cr 2p, Fe 2p, Ni 2p, C 1s, and O 1s of passive film on the surfaces of the heat-treated (<b>a</b>) Al<sub>0.8</sub>CrFeNi, (<b>b</b>) Al<sub>1.0</sub>CrFeNi, and (<b>c</b>) Al<sub>1.2</sub>CrFeNi medium-entropy alloy immersed in 3.5 wt% NaCl solution for 15 d.</p> Full article ">Figure 10
<p>Schematic of the corrosion mechanism of Al<span class="html-italic"><sub>x</sub></span>CrFeNi medium-entropy alloys in 3.5 wt% NaCl corrosive solution.</p> Full article ">
15 pages, 2665 KiB  
Article
Fluid Dynamics Analysis of Coherent Jet with a Mixed Shrouding H2-CO2/N2 for EAF Steelmaking
by Songtao Yan, Fuhai Liu, Rong Zhu, Guangsheng Wei and Kai Dong
Metals 2025, 15(3), 291; https://doi.org/10.3390/met15030291 - 7 Mar 2025
Viewed by 173
Abstract
In order to suppress the rapid combustion effect and consumption rate of pure hydrogen gas, N2 or CO2 at flow rates of 0, 80, and 240 Nm3/h was pre-mixed with shrouding H2 at flow rates of 800, 720, [...] Read more.
In order to suppress the rapid combustion effect and consumption rate of pure hydrogen gas, N2 or CO2 at flow rates of 0, 80, and 240 Nm3/h was pre-mixed with shrouding H2 at flow rates of 800, 720, and 560 Nm3/h at room temperature, and the behaviors of the main oxygen jet and shrouding flame were analyzed by both numerical simulation and combustion experiments. The results showed that, because of the participation of CO2 in the H2 combustion reaction, the length of the axial velocity potential core was reduced using the CO2 shrouding mixed injection method, compared to the same mixed rate of N2. This trend would be further enhanced as N2 and CO2 mixing ratio increased. Meanwhile, when the shrouding mixed rate is 30%, the maximum axial and radial expansion rate generated by N2-H2 shrouding method is 1.28 and 1.04 times longer than that by the CO2-H2 shrouding method. The Fo-a, theoretical impaction depth and area generated by the 10% N2 shrouding mixed rate was 84.0, 95.5 and 86.4% of those generated by the traditional coherent jet, respectively, which indicated that the 10% N2 shrouding mixed rate method might lead to comparable production indexes in the EAF steelmaking process. Full article
(This article belongs to the Special Issue Advanced Metal Smelting Technology and Prospects)
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<p>Cross view (<b>a</b>) and front view (<b>b</b>) of the coherent lance.</p> Full article ">Figure 2
<p>The (<b>a</b>) physical diagram, (<b>b</b>) front view and (<b>c</b>) cross view of the high temperature combustion furnace. (Unit: mm).</p> Full article ">Figure 3
<p>(<b>a</b>) Geometric construction of the numerical model. (<b>b</b>) Mesh profile of the numerical model with boundary conditions.</p> Full article ">Figure 4
<p>The axial velocity profiles of the main oxygen jet using different fuel mixed injection methods at the centerline of the Laval nozzle: (<b>a</b>) 10% shrouding fuel mixed method. (<b>b</b>) 30% shrouding fuel mixed method.</p> Full article ">Figure 5
<p>The total temperature of main oxygen profiles using different fuel mixed injection methods at centerline of the Laval nozzle: (<b>a</b>) 10% shrouding fuel mixed method. (<b>b</b>) 30% shrouding fuel mixed method.</p> Full article ">Figure 6
<p>The total temperature of coherent jet profiles using different fuel mixed injection methods.</p> Full article ">Figure 7
<p>The theoretical impaction depth and area generated by different fuel mixed injection methods.</p> Full article ">Figure 8
<p>The effective oxygen flow rate through the theoretical impaction area using different fuel mixed injection methods.</p> Full article ">
15 pages, 8614 KiB  
Article
Microstructural Modification by Large Pre-Deformation and Post-Aging to Improve Properties in Al-Mg-Li Alloy
by Zeyu Zheng, Peipei Ma, Longhui Chen and Chunhui Liu
Metals 2025, 15(3), 290; https://doi.org/10.3390/met15030290 - 6 Mar 2025
Viewed by 210
Abstract
Al-Mg-Li alloy is an ideal lightweight structural material for aerospace applications due to its low density, high specific strength, and excellent low-temperature performance. This study examines the mechanical properties and microstructural evolution of Al-Mg-Li alloy subjected to cryogenic and room temperature cold rolling, [...] Read more.
Al-Mg-Li alloy is an ideal lightweight structural material for aerospace applications due to its low density, high specific strength, and excellent low-temperature performance. This study examines the mechanical properties and microstructural evolution of Al-Mg-Li alloy subjected to cryogenic and room temperature cold rolling, which induces large plastic deformation. Compared with room temperature rolling, cryogenic rolling significantly reduces surface cavity formation, thereby enhancing the alloy’s rolling surface quality. After cryogenic rolling by 80% and subsequent natural aging, the yield strength of artificially aged Al-Mg-Li alloy reaches 560 MPa, delivering a 60% increase compared to the traditional T6 state with a slight reduction in elongation from 6.5% to 4.6%. The specific strength achieves 2.23 × 105 N·m/kg, outperforming conventional Al-Cu-Li and 7xxx-series Al alloys. The depth of intergranular corrosion decreases from 100 µm to 10 µm, demonstrating excellent corrosion resistance enabled by the new method. Transmission electron microscopy reveals that finely distributed δ′ (Al3Li) is the primary strengthening phase, with high-density dislocations further enhancing strength. However, coarsening of δ′ (from ~2.9 nm to >6 nm) induced by ensuing artificial aging results in coplanar slip and reduced elongation. Lowering the post-aging temperature inhibits δ′ coarsening, thereby improving both strength and elongation. Our results provide valuable insights into optimizing the properties of Al-Mg-Li alloys for advanced lightweight applications. Full article
(This article belongs to the Special Issue Structural Quality and Its Effects on the Performance of Light Alloy Castings)
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<p>Thermo-mechanical schedules for preparation of Al-Mg-Li alloy.</p> Full article ">Figure 2
<p>SEM images of the surface morphology of the largely pre-deformed Al-Mg-Li alloy. (<b>a</b>,<b>b</b>) Macroscopic appearance of the Al-Mg-Li alloy prepared by cryogenic rolling and RT rolling; (<b>c</b>–<b>e</b>) surface of the alloy after cryogenic rolling; and (<b>f</b>–<b>h</b>) surface of the alloy after RT rolling.</p> Full article ">Figure 3
<p>SEM images of different positions of largely pre-deformed Al-Mg-Li alloy. (<b>a</b>) 1 mm from crack; (<b>b</b>) 2 mm from crack; (<b>c</b>) 3 mm from crack; and (<b>d</b>) 4 mm from crack.</p> Full article ">Figure 4
<p>Subsequent age hardening curves of largely pre-deformed Al-Mg-Li alloy. (<b>a</b>) Natural aging (NA) process and (<b>b</b>–<b>d</b>) artificial aging (AA) process at 125 °C, 145 °C, and 165 °C.</p> Full article ">Figure 5
<p>Tensile curves of largely pre-deformed and peak-aged Al-Mg-Li alloy. (<b>a</b>) Room temperature-rolled (RTR) alloy and (<b>b</b>) cryogenic temperature-rolled (CR) alloy. (<b>c</b>) Dimension of tensile specimen.</p> Full article ">Figure 6
<p>EBSD-IPF images of Al-Mg-Li alloy in different states. (<b>a</b>) T4; (<b>b</b>) as CR80%.</p> Full article ">Figure 7
<p>EBSD-PF images of Al-Mg-Li alloy in different states. (<b>a</b>) T4; (<b>b</b>) cold-rolled by 80%.</p> Full article ">Figure 8
<p>Dislocation morphology of T6 state and largely pre-deformed Al-Mg-Li alloy after different aging treatments. (<b>a</b>) CR; (<b>b</b>) CR-125 °C; (<b>c</b>) CR-165 °C; and (<b>d</b>) T6.</p> Full article ">Figure 9
<p>TEM and high-resolution HADDF images of Al<sub>3</sub>Li precipitated phase in different states of largely pre-deformed Al-Mg-Li alloy. (<b>a</b>) T6; (<b>b</b>) CR-125 °C; (<b>c</b>) CR-165 °C; (<b>d</b>) CR-NA; (<b>e</b>) magnified section of (<b>d</b>); and (<b>f</b>) schematic diagram of Al<sub>3</sub>Li atomic structure.</p> Full article ">Figure 10
<p>Statistical diagram of precipitation phases of Al-Mg-Li alloy in different states. (<b>a</b>) T6; (<b>b</b>) CR-NA; (<b>c</b>) CR-125 °C; and (<b>d</b>) CR-165 °C.</p> Full article ">Figure 11
<p>SEM images of intergranular corrosion depth of Al-Mg-Li alloy in different states. (<b>a</b>) T6 and (<b>b</b>) CR-NA.</p> Full article ">Figure 12
<p>SEM images of edge cracks in room temperature-rolled alloy. (<b>a</b>) Crack propagation path and (<b>b</b>) magnified image of initial cavity.</p> Full article ">Figure 13
<p>Synchrotron radiation XRD patterns. (<b>a</b>) T6 and (<b>b</b>) CR-NA.</p> Full article ">Figure 14
<p>Schematic diagram of microstructure evolution of Al-Mg-Li alloy prepared by large pre-deformation combined with aging process.</p> Full article ">
30 pages, 65330 KiB  
Article
Experiments and Simulations on the Low-Temperature Reduction of Iron Ore Oxide Pellets with Hydrogen
by Róbert Findorák, Zuzana Miškovičová, Jaroslav Legemza, Róbert Dzurňák, Branislav Buľko, Peter Demeter, Andrea Egryová and Róbert Maliňák
Metals 2025, 15(3), 289; https://doi.org/10.3390/met15030289 - 6 Mar 2025
Viewed by 188
Abstract
This article examines the low-temperature reducibility of four types of iron ore pellets in a pure hydrogen atmosphere, with the aim of understanding the thermodynamic aspects of the process. The research focuses on optimizing conditions for pellet reduction in order to reduce CO [...] Read more.
This article examines the low-temperature reducibility of four types of iron ore pellets in a pure hydrogen atmosphere, with the aim of understanding the thermodynamic aspects of the process. The research focuses on optimizing conditions for pellet reduction in order to reduce CO2 emissions and improve iron production efficiency. Experimental tests were conducted at temperatures of 600 °C and 800 °C, supplemented by thermodynamic simulations predicting the equilibrium composition and energy requirements. Chemical and microstructural analyses revealed that porosity, mineralogical composition, and phase distribution homogeneity significantly affect reduction efficiency. High-quality pellets with low SiO2 content demonstrated the best reduction ability, while fluxed pellets with the presence of calcium silicate ferrites and pellets with a higher content of SiO2 showed lower reduction potential due to the presence of hard-to-reduce phases such as calcium silicate ferrites and iron silicates. The results highlight the importance of controlling process conditions and optimizing pellet properties to enhance the reduction process and minimize environmental impacts. This study provides valuable insights for the application of hydrogen reduction in industrial conditions, contributing to the decarbonization of the metallurgical industry. Full article
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<p>Schematic of the hydrogen pellet reduction experiments.</p> Full article ">Figure 2
<p>High-temperature observation of iron pellet samples. DT = deformation temperature (°C); ST = shrinkage temperature (°C); HT = hemisphere temperature or melting point (°C); NA—not analyzed at higher temperatures due to the temperature limit being reached, with a maximum of about 1515 °C.</p> Full article ">Figure 3
<p>Microstructures of the Fe pellet samples: (<b>a</b>) A; (<b>b</b>) B; (<b>c</b>) C; (<b>d</b>) D.</p> Full article ">Figure 4
<p>Image conversion of the microstructural images for porosity assessment (black areas on the image to the right). (<b>a</b>) Pellet A; (<b>b</b>) Pellet B; (<b>c</b>) Pellet C; (<b>d</b>) Pellet D.</p> Full article ">Figure 5
<p>Correlation of pellet porosities. MAO—porosity based on microstructure; PA—calculation based on pycnometric and volumetric true density.</p> Full article ">Figure 6
<p>Cross-sections of the analyzed pellets.</p> Full article ">Figure 7
<p>Mechanical properties of Fe pellets under compression.</p> Full article ">Figure 8
<p>Baur–Glessner diagram with the Boudouard reaction (BR) and the CO and H<sub>2</sub> reactions.</p> Full article ">Figure 9
<p>Equilibrium system of hydrogen reduction for Fe pellets: (<b>a</b>) A; (<b>b</b>) B; (<b>c</b>) C; (<b>d</b>) D. Explanation: *2FeO*SiO<sub>2</sub>—standardized notation of a complex compound fayalite (2FeO·SiO<sub>2</sub>) in the HSC Chemistry thermodynamic program.</p> Full article ">Figure 9 Cont.
<p>Equilibrium system of hydrogen reduction for Fe pellets: (<b>a</b>) A; (<b>b</b>) B; (<b>c</b>) C; (<b>d</b>) D. Explanation: *2FeO*SiO<sub>2</sub>—standardized notation of a complex compound fayalite (2FeO·SiO<sub>2</sub>) in the HSC Chemistry thermodynamic program.</p> Full article ">Figure 10
<p>Predicted yield during the reduction of pellets in an H<sub>2</sub> atmosphere.</p> Full article ">Figure 11
<p>Equilibrium composition of calcium silicates with varying CaO content. Explanation: *2CaO*Fe<sub>2</sub>O<sub>3</sub>—standardized notation of a complex dicalcium ferrite (2CaO·Fe<sub>2</sub>O<sub>3</sub>) in the HSC Chemistry thermodynamic program.</p> Full article ">Figure 12
<p>Comparison of Fe yield for the pure hematite and tested pellets in an H<sub>2</sub> atmosphere with respect to the reducing agent ratio and reduction temperature: (<b>a</b>) pure hematite at 600 °C; (<b>b</b>) pure hematite at 800 °C; (<b>c</b>) Pellet A; (<b>d</b>) Pellet B.</p> Full article ">Figure 12 Cont.
<p>Comparison of Fe yield for the pure hematite and tested pellets in an H<sub>2</sub> atmosphere with respect to the reducing agent ratio and reduction temperature: (<b>a</b>) pure hematite at 600 °C; (<b>b</b>) pure hematite at 800 °C; (<b>c</b>) Pellet A; (<b>d</b>) Pellet B.</p> Full article ">Figure 13
<p>Hydrogen reduction indexes of pellets and the effect of temperature on the reduction degree: (<b>a</b>) 600 °C; (<b>b</b>) 800 °C.</p> Full article ">Figure 14
<p>Comparison of Fe yield for the tested pellets in an H<sub>2</sub> atmosphere—thermodynamic simulation.</p> Full article ">Figure 15
<p>Temperature profile in the ceramic tube during hydrogen reduction (H<sub>2</sub> input from the left).</p> Full article ">Figure 16
<p>Graphical representation of the temperature distribution on the analyzed sample.</p> Full article ">Figure 17
<p>Flow simulation in the tube.</p> Full article ">Figure 18
<p>Comparison of the change in % FeTOT in the pellets after reduction in the H<sub>2</sub> experiments.</p> Full article ">Figure 19
<p>The structure of Pellet A after reduction at 600 °C in an H<sub>2</sub> atmosphere.</p> Full article ">Figure 20
<p>The structure of Pellet B after reduction at (<b>a</b>) 600 °C and (<b>b</b>) 800 °C in an H<sub>2</sub> atmosphere.</p> Full article ">Figure 21
<p>The structure of Pellet D after reduction at 600 °C in an H<sub>2</sub> atmosphere.</p> Full article ">Figure 22
<p>The structure of Pellet C after reduction at 600 °C in an H<sub>2</sub> atmosphere.</p> Full article ">Figure 23
<p>Schematic comparison of the reduction of iron ore pellets.</p> Full article ">
15 pages, 3466 KiB  
Article
Prediction of Creep Rupture Life of 5Cr-0.5Mo Steel Using Machine Learning Models
by Muhammad Ishtiaq, Hafiz Muhammad Rehan Tariq, Devarapalli Yuva Charan Reddy, Sung-Gyu Kang and Nagireddy Gari Subba Reddy
Metals 2025, 15(3), 288; https://doi.org/10.3390/met15030288 - 6 Mar 2025
Viewed by 165
Abstract
The creep rupture life of 5Cr-0.5Mo steels used in high-temperature applications is significantly influenced by factors such as minor alloying elements, hardness, austenite grain size, non-metallic inclusions, service temperature, and applied stress. The relationship of these variables with the creep rupture life is [...] Read more.
The creep rupture life of 5Cr-0.5Mo steels used in high-temperature applications is significantly influenced by factors such as minor alloying elements, hardness, austenite grain size, non-metallic inclusions, service temperature, and applied stress. The relationship of these variables with the creep rupture life is quite complex. In this study, the creep rupture life of 5Cr-0.5Mo steel was predicted using various machine learning (ML) models. To achieve higher accuracy, various ML techniques, including random forest (RF), gradient boosting (GB), linear regression (LR), artificial neural network (ANN), AdaBoost (AB), and extreme gradient boosting (XGB), were applied with careful optimization of hidden parameters. Among these, the ANN-based model demonstrated superior performance, yielding high accuracy with minimal prediction errors for the test dataset (RMSE = 0.069, MAE = 0.053, MAPE = 0.014, and R2 = 1). Additionally, we developed a user-friendly graphical user interface (GUI) for the ANN model, enabling users to predict and optimize creep rupture life. This tool helps materials scientists and industrialists prevent failures in high-temperature applications and design steel compositions with enhanced creep resistance. Full article
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<p>Heatmap illustrating the Pearson’s correlation coefficient for the variables analyzed in this study. Each square contains the corresponding coefficient value, while the color intensity represents the strength of the relationship: darker shades indicate stronger correlations, and lighter shades signify weaker relationships.</p> Full article ">Figure 2
<p>Schematic representation showing the steps involved in the present work.</p> Full article ">Figure 3
<p>Graphical representation of the performance evaluation of various machine learning methods, (<b>a</b>) for training data (<b>b</b>) for testing data. (<b>c</b>) R<sup>2</sup> for testing and training data. The smaller values of RMSE, MAE, and MAPE represent higher accuracy. The higher values~1 for R<sup>2</sup> represent better predictability.</p> Full article ">Figure 4
<p>Graphical representation of the effect of applied stress on the creep rupture life of 5Cr-0.5Mo steel: (<b>a</b>) at 500 °C, (<b>b</b>) at 550 °C, (<b>c</b>) at 600 °C, (<b>d</b>) at 650 °C. The red color is for experimental, and the black color is for prediction from ANN model.</p> Full article ">Figure 5
<p>Graphical representation of the effect of temperature on the creep rupture life under a stress of 98 MPa. The red color is for experimental, and the black color is for prediction from ANN model.</p> Full article ">Figure 6
<p>Graphical representation of the effect of: (<b>a</b>) carbon content, (<b>b</b>) silicon content on the creep rupture life at 550 °C and under a stress of 98 MPa. The red color is for experimental, and the black color is for prediction from ANN model.</p> Full article ">Figure 7
<p>Contour plots showing the combined effect of (<b>a</b>) temperature and stress, (<b>b</b>) carbon and silicon content on the creep rupture life of 5Cr-0.5Mo steel.</p> Full article ">Figure 8
<p>Screenshot of the graphical user interface (GUI) of the developed model.</p> Full article ">Figure 9
<p>Screenshot of the optimized inputs for the maximum creep rupture life.</p> Full article ">
15 pages, 10067 KiB  
Article
Effects of a Welding Wire Containing Er or Sc on the Microstructure, Mechanical Properties, and Corrosion Resistance of the 5xxx Aluminum Alloy MIG Joint
by Cunwei Zou, Ruizhi Wu, Xinhe Yang, Zhikun Ma and Legan Hou
Metals 2025, 15(3), 287; https://doi.org/10.3390/met15030287 - 6 Mar 2025
Viewed by 133
Abstract
The development of MIG (metal inert gas) welding for five-series aluminum alloys primarily involves the improvement and optimization of welding processes. Building upon research findings regarding the enhancement of aluminum alloy properties through the use of scandium (Sc) and erbium (Er), our study [...] Read more.
The development of MIG (metal inert gas) welding for five-series aluminum alloys primarily involves the improvement and optimization of welding processes. Building upon research findings regarding the enhancement of aluminum alloy properties through the use of scandium (Sc) and erbium (Er), our study incorporates Sc and Er into the welding wire to examine their impact on welding quality. The results show that the introduction of Er and Sc results in grain refinement from 47 µm to 29 µm and 31 µm, respectively. Grain refinement is mainly attributed to the heterogeneous nucleation of submicron-sized, coherent Al3Er and Al3Sc phases with L12 structure. The ultimate tensile strength (UTS), fracture elongation EI [%], and microhardness of joints welded with Er-containing and Sc-containing filler wires exhibit significant enhancements due to the refinement strengthening and dispersion strengthening. Joints welded with the filler wires containing Er and Sc display reduced corrosion current density and higher corrosion potential. The enhanced corrosion resistance comes from the formation of a denser oxide film and the equilibrium in the potential difference between the precipitated phases (Al3Er and Al3Sc) and the matrix. Filler wires containing Er and Sc have almost similar effects on improvements of the MIG welding joints. Full article
(This article belongs to the Special Issue Manufacturing Processes of Metallic Materials)
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<p>Welding wire performance test diagram.</p> Full article ">Figure 2
<p>Surface morphology of welded joints: (<b>a</b>) Wire-1; (<b>b</b>) Wire-2; and (<b>c</b>) Wire-3.</p> Full article ">Figure 3
<p>Optical microstructure of cross-section of welded joint: (<b>a</b>) appearance of welded joint of Wire-2; (<b>b</b>) WZ of Wire-1; (<b>c</b>) WZ of Wire-2; and (<b>d</b>) WZ of Wire-3.</p> Full article ">Figure 4
<p>Secondary electron morphology in WZ of different welded joints: (<b>a</b>) Wire-1; (<b>b</b>) Wire-2; and (<b>c</b>) Wire-3.</p> Full article ">Figure 5
<p>Back-scattered electron images for the WZ of welded joints: (<b>a</b>) Wire-1; (<b>b</b>) Wire-2; and (<b>c</b>) Wire-3.</p> Full article ">Figure 6
<p>Energy spectrum of the SEM image in <a href="#metals-15-00287-f005" class="html-fig">Figure 5</a>c.</p> Full article ">Figure 7
<p>HADDF-STEM image of the Wire-2 weld zone and its element mapping and diffraction pattern of this region with Al<sub>3</sub>Er stripes.</p> Full article ">Figure 8
<p>HADDF-STEM image of the Wire-3 weld zone and its element mapping and diffraction pattern of this region with Al<sub>3</sub>Sc stripes. The yellow circles highlight the Al<sub>3</sub>Sc phases.</p> Full article ">Figure 9
<p>Distribution of microhardness in three kinds of welding wire joints.</p> Full article ">Figure 10
<p>Tensile stress–strain curves of three kinds of wire-welded joints.</p> Full article ">Figure 11
<p>SEM micrographs of fracture surfaces of tensile samples: (<b>a</b>) Wire-1-welded joint; (<b>b</b>) Wire-2-welded joint; and (<b>c</b>) Wire-3-welded joint.</p> Full article ">Figure 12
<p>Corrosion morphology of different welded joints: (<b>a</b>) Wire-1; (<b>b</b>) Wire-2; (<b>c</b>) Wire-3; and (<b>d</b>) polarization curves of different welded joints.</p> Full article ">
17 pages, 6445 KiB  
Article
Influence of B2O3 on the Viscosity and Melt Structure of CaO-SiO2-M2O (M = Li, Na)-Based Slags
by Jinhui Wang, Jie Qi, Yuanxin Shi, Yingying Dou and Chengjun Liu
Metals 2025, 15(3), 286; https://doi.org/10.3390/met15030286 - 6 Mar 2025
Viewed by 211
Abstract
In the process of continuous casting, especially high-speed continuous casting, the inflow state of the mold flux is particularly important. The fluxing agent is one of the most important factors affecting the flow state. The influence of the typical fluxing agent B2 [...] Read more.
In the process of continuous casting, especially high-speed continuous casting, the inflow state of the mold flux is particularly important. The fluxing agent is one of the most important factors affecting the flow state. The influence of the typical fluxing agent B2O3 on the viscous characteristics and melt structure of the fluorine-free CaO-SiO2-M2O (M = Li, Na) system was analyzed. The following conclusions were drawn. In the CaO-SiO2-Na2O slags, with the increasing addition of B2O3, the viscosity, breaking temperature, and polymerization degree of the slag show a gradually decreasing trend. When the mass fraction of B2O3 increased from 0 to 10%, the increase in two-dimensional [BO3] structural units played a dominant role. When the mass fraction of B2O3 reached 15%, the network was affected by the increase in [BO3] and the low-polymerized [SiO4] tetrahedrons. The CaO-SiO2-Li2O slag system had a lower breaking temperature due to the formation of phases such as Li2O·2B2O3, of a low melting temperature. The initial degree of depolymerization of the network was high. Upon increasing the addition of B2O3, the relative proportion of the network modifier structural units significantly increased, resulting in the enhanced instability of the network structure. As a result, the effect of [SiO4]-polymerization was stronger than that of [BO3]-depolymerization in maintaining the stability of the network structure. Full article
(This article belongs to the Special Issue Advanced Tundish Metallurgy and Clean Steel Technology—Second Edition)
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<p>Isotherms of CaO-SiO<sub>2</sub>-Na<sub>2</sub>O slag.</p> Full article ">Figure 2
<p>Isotherms of CaO-SiO<sub>2</sub>-Li<sub>2</sub>O slag.</p> Full article ">Figure 3
<p>Diagram of crucible structure.</p> Full article ">Figure 4
<p>Viscosity curves of CaO-SiO<sub>2</sub>-Na<sub>2</sub>O slags with different B<sub>2</sub>O<sub>3</sub> additions.</p> Full article ">Figure 5
<p>Viscosity at 1300 °C and breaking temperature of CaO-SiO<sub>2</sub>-Na<sub>2</sub>O slags with different B<sub>2</sub>O<sub>3</sub> additions.</p> Full article ">Figure 6
<p>Viscosity curves of CaO-SiO<sub>2</sub>-Li<sub>2</sub>O slags with different B<sub>2</sub>O<sub>3</sub> additions.</p> Full article ">Figure 7
<p>Viscosity at 1300 °C and breaking temperature of CaO-SiO<sub>2</sub>-Li<sub>2</sub>O slags with different B<sub>2</sub>O<sub>3</sub> additions.</p> Full article ">Figure 8
<p>Raman spectra of CaO-SiO<sub>2</sub>-Na<sub>2</sub>O slags with different B<sub>2</sub>O<sub>3</sub> additions after smoothing to baseline.</p> Full article ">Figure 9
<p>Gaussian spectral solution results of CaO-SiO<sub>2</sub>-Na<sub>2</sub>O-B<sub>2</sub>O<sub>3</sub> slag system.</p> Full article ">Figure 10
<p>Variation in mole fraction of <math display="inline"><semantics> <msubsup> <mi>Q</mi> <mrow> <mi>S</mi> <mi>i</mi> </mrow> <mi>i</mi> </msubsup> </semantics></math> in CaO-SiO<sub>2</sub>-Na<sub>2</sub>O-B<sub>2</sub>O<sub>3</sub> slag system.</p> Full article ">Figure 11
<p>Raman spectra of CaO-SiO<sub>2</sub>-Li<sub>2</sub>O slags with different B<sub>2</sub>O<sub>3</sub> additions after smoothing to baseline.</p> Full article ">Figure 12
<p>Gaussian spectral solution results of CaO-SiO<sub>2</sub>-Li<sub>2</sub>O-B<sub>2</sub>O<sub>3</sub> slag system.</p> Full article ">Figure 13
<p>Variation in mole fraction of <math display="inline"><semantics> <msubsup> <mi>Q</mi> <mrow> <mi>S</mi> <mi>i</mi> </mrow> <mi>i</mi> </msubsup> </semantics></math> in CaO-SiO<sub>2</sub>-Li<sub>2</sub>O-B<sub>2</sub>O<sub>3</sub> slag system.</p> Full article ">
23 pages, 10069 KiB  
Article
Microstructural Evolution, Strengthening Mechanisms, and Fracture Behavior of Aluminum Composites Reinforced with Graphene Nanoplatelets and In Situ–Formed Nano-Carbides
by Rumyana Lazarova, Lubomir Anestiev, Yana Mourdjeva, Kateryna Valuiska and Veselin Petkov
Metals 2025, 15(3), 285; https://doi.org/10.3390/met15030285 - 5 Mar 2025
Viewed by 127
Abstract
The microstructure and mechanical properties of GNP-reinforced aluminum composites obtained by powder metallurgy and hot extrusion (at 400 °C, 500 °C, and annealing at 3 h at 610 °C), were investigated. It was found that: (i) depending on the processing applied, the composites [...] Read more.
The microstructure and mechanical properties of GNP-reinforced aluminum composites obtained by powder metallurgy and hot extrusion (at 400 °C, 500 °C, and annealing at 3 h at 610 °C), were investigated. It was found that: (i) depending on the processing applied, the composites showed an increase in yield strength (YS) and ultimate strength (US) of up to 283%, and 78%, respectively; (ii) depending on the size of the ex situ GNP and in situ Al4C3 reinforcements, two fracture mechanisms are observed: ductile and brittle–ductile; (iii) annealing for 3 h at 610 °C did not improve the mechanical properties; (iv) the plot of YS vs. the volume fraction of the GNP introduced showed a peculiar pattern not been reported so far. Theoretical analysis of the results showed: (1) the major contributor to the YS increase is the Hall–Petch mechanism; (2) the reinforcements contribution to YS, complements that of Hall–Petch; (3) the main contributor to the composite strength is GNP; (4) a critical size of the reinforcement exists, 1.43 nm, at which the YS is maximal, 260 MPa; (5) the increase in the processing temperature and time leads to Ostwald ripening and increase of Al4C3 size and deterioration of mechanical properties. Full article
(This article belongs to the Special Issue Powder Metallurgy of Metals and Alloys)
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<p>Microstructure of a sample containing 1.1 wt. % GNPs after extrusion at 500 °C. Carbides formed on the surface of GNP are shown with red arrows.</p> Full article ">Figure 2
<p>SEM image of a specimen containing 0.5 wt. % GNP, showing GNPs and nano-sized Al<sub>4</sub>C<sub>3</sub> embedded in the aluminum matrix.</p> Full article ">Figure 3
<p>SEM image of a GNP particle and nano-sized Al<sub>4</sub>C<sub>3</sub> formed on its surface taken from a specimen containing 0.3 wt. % GNPs.</p> Full article ">Figure 4
<p>(<b>a</b>) TEM SAED pattern of a layer, identified as Al<sub>2</sub>C<sub>2</sub>, with an octahedral structure of Al<sub>4</sub>C<sub>3</sub> carbide; (<b>b</b>) TEM SAED pattern representing the orientation relationship between the carbide Al<sub>4</sub>C<sub>3</sub> [110] (red) and the Al [110] (white) of the Al matrix.</p> Full article ">Figure 5
<p>Microstructure of a sample taken from a longitudinal section of a specimen reinforced with 0.7 wt. % GNPs and extruded at 500 °C, showing the differences in the specimen’s microstructure before (<b>a</b>) and after (<b>b</b>) tensile loading.</p> Full article ">Figure 6
<p>Microstructure of the same specimen (0.7 wt. % GNP, extruded at 500 °C), but the image is taken from a location near the edge of the neck, showing the formation of dislocation pile-ups near the grain walls.</p> Full article ">Figure 7
<p>Dependence of yield strength (<b>a</b>) and ultimate strength (<b>b</b>) on the content of GNPs in the aluminum composite.</p> Full article ">Figure 8
<p>Dependence of relative elongation on the content of GNPs in aluminum composite.</p> Full article ">Figure 9
<p>Microstructure of a sample taken from a tensile-loaded specimen containing 0 wt. % GNPs and extruded at 500 °C. GBs denote the grain boundaries.</p> Full article ">Figure 10
<p>Fracture surface after tensile testing of specimens containing 0 wt. % GNPs and extruded at 500 °C—(<b>a</b>,<b>c</b>,<b>e</b>), and containing 0.7 wt. % GNPs and extruded at 500 °C—(<b>b</b>,<b>d</b>,<b>f</b>). Dimples with nano-sized carbides at the bottom are marked with red circles.</p> Full article ">Figure 11
<p>TEM BF image of a sample taken from a longitudinal section of a specimen containing 0.7 wt. % GNPs after tensile testing.</p> Full article ">Figure 12
<p>Net contribution of the ex situ and in situ reinforcements to the composite’s yield strength plotted vs. VF of the GNP.</p> Full article ">Figure 13
<p>Plots of <math display="inline"><semantics> <mrow> <mrow> <mi>τ</mi> <mo>/</mo> <mrow> <msub> <mi>G</mi> <mrow> <mi>A</mi> <mi>l</mi> </mrow> </msub> </mrow> </mrow> <mo> </mo> <mi>vs.</mi> <mo> </mo> <mrow> <mrow> <mfenced close="&#x232A;" open="&#x2329;"> <mrow> <msub> <mi>r</mi> <mi>p</mi> </msub> </mrow> </mfenced> </mrow> <mo>/</mo> <mrow> <msub> <mi>b</mi> <mrow> <mi>A</mi> <mi>l</mi> </mrow> </msub> </mrow> </mrow> </mrow> </semantics></math> calculated with (2a) and (2b) (<math display="inline"><semantics> <mrow> <msubsup> <mi>f</mi> <mn>0</mn> <mrow> <mi>G</mi> <mi>N</mi> <mi>P</mi> </mrow> </msubsup> </mrow> </semantics></math> = 0.005), showing (i) the contributions of the OSM and OBM mechanisms to the strength and (ii) the critical particulate radii of the two reinforcements at which the transition between these two strengthening mechanisms occurs. The arrows indicate the highest possible attainable yield strength (hpas) for the ex-situ and the in situ reinforcements. The noticeable strengthening effect of the GNP near the critical size of the GNP compared to that of Al<sub>4</sub>C<sub>3</sub> is readily distinguishable.</p> Full article ">Figure 14
<p>Plots of <math display="inline"><semantics> <mrow> <msubsup> <mi>f</mi> <mn>0</mn> <mrow> <mi>G</mi> <mi>N</mi> <mi>P</mi> </mrow> </msubsup> </mrow> </semantics></math>, <math display="inline"><semantics> <mrow> <msup> <mi>f</mi> <mrow> <mi>A</mi> <msub> <mi>l</mi> <mn>4</mn> </msub> <msub> <mi>C</mi> <mn>3</mn> </msub> </mrow> </msup> </mrow> </semantics></math>, and <math display="inline"><semantics> <mrow> <msubsup> <mi>f</mi> <mn>0</mn> <mrow> <mi>G</mi> <mi>N</mi> <mi>P</mi> </mrow> </msubsup> </mrow> </semantics></math>, calculated with (5). Notice the almost linear dependence of <math display="inline"><semantics> <mrow> <msup> <mi>f</mi> <mrow> <mi>A</mi> <msub> <mi>l</mi> <mn>4</mn> </msub> <msub> <mi>C</mi> <mn>3</mn> </msub> </mrow> </msup> </mrow> </semantics></math> on <math display="inline"><semantics> <mrow> <msubsup> <mi>f</mi> <mn>0</mn> <mrow> <mi>G</mi> <mi>N</mi> <mi>P</mi> </mrow> </msubsup> </mrow> </semantics></math>, which is confirmed by the correlation coefficient, R<sup>2</sup> = 0.9911.</p> Full article ">
17 pages, 7697 KiB  
Article
Dynamic Compression and Blast Failure Behavior of a Biomimetic Novel Lattice with Vertex Modifications Made of 316L Stainless Steel
by Fei Zhou, Zhihua Xue and Xiaofei Cao
Metals 2025, 15(3), 284; https://doi.org/10.3390/met15030284 - 5 Mar 2025
Viewed by 102
Abstract
A novel 316L stainless steel Vertex Modified BCC (VM-BCC) lattice unit cell with attractive performance characteristics is developed. Lattice structure, as well as the sandwich panel, are constructed. Numerical simulation is utilized to simulate the quasi-static compression, dynamic compression and blast behavior considering [...] Read more.
A novel 316L stainless steel Vertex Modified BCC (VM-BCC) lattice unit cell with attractive performance characteristics is developed. Lattice structure, as well as the sandwich panel, are constructed. Numerical simulation is utilized to simulate the quasi-static compression, dynamic compression and blast behavior considering the rate-dependent properties, elastoplastic response and nonlinear contact. Finite element results are validated by comparing with the experimental results. Parametric studies are conducted to gain insight into the effects of loading velocity, equivalent TNT load and explosion distance on the dynamic behavior of the lattice pattern and sandwich panel. Testing results indicate that the proposed 316L stainless steel VM-BCC structure exhibits more superior plateau stress and specific energy absorption (SEA) than those of the BCC or Octet one. The proposed novel lattice will provide reference for improving the protective efficiency in key equipment fields and enhancing overall safety. Full article
(This article belongs to the Special Issue Fracture Mechanics of Materials—the State of the Art)
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<p>(<b>a</b>) Geometric configurations of the BCC, CM-BCC, V-BCC, Octet and VM-BCC lattice patterns; (<b>b</b>) VM-BCC design inspired by deep-sea glass sponge.</p> Full article ">Figure 2
<p>(<b>a</b>) Front-view schematic of the compression computational model; (<b>b</b>) Uniaxial tensile true stress–strain curve for static compression simulation. Herein, the true stress–strain curve is that of the average of the three dog-bone specimens.</p> Full article ">Figure 3
<p>Air blast symmetry model on the lattice sandwich panel.</p> Full article ">Figure 4
<p>(<b>a</b>) Effect of element size on the stress–strain curves of the V-BCC lattice pattern; (<b>b</b>) Plateau stress values of the V-BCC lattice pattern at different element sizes.</p> Full article ">Figure 5
<p>Comparison of the (<b>a</b>) compression stress–strain curves and the (<b>b</b>) deformation state between simulation and experiment.</p> Full article ">Figure 6
<p>Quasi-static compression (<b>a</b>) stress–strain curves; (<b>b</b>) plateau stress and SEA; (<b>c</b>) deformation processes of the BCC, CM-BCC, V-BCC, Octet and VM-BCC lattice patterns.</p> Full article ">Figure 7
<p>Low-speed impact compression (<b>a</b>) stress–strain curves; (<b>b</b>) plateau stress and SEA; (<b>c</b>) deformation processes of the BCC, CM-BCC, V-BCC, Octet and VM-BCC lattice patterns.</p> Full article ">Figure 8
<p>Medium-speed impact compression (<b>a</b>) stress–strain curves; (<b>b</b>) plateau stress and SEA; (<b>c</b>) deformation processes of the BCC, CM-BCC, V-BCC, Octet and VM-BCC lattice patterns.</p> Full article ">Figure 9
<p>High-speed impact compression (<b>a</b>) stress–strain curves; (<b>b</b>) plateau stress and SEA; (<b>c</b>) deformation processes of the BCC, CM-BCC, V-BCC, Octet and VM-BCC lattice patterns.</p> Full article ">Figure 10
<p>(<b>a</b>) Plateau stress and (<b>b</b>) SEA of the BCC, CM-BCC, V-BCC, Octet and VM-BCC lattice patterns under different loading velocities.</p> Full article ">Figure 11
<p>Back face deflection curves of the BCC, CM-BCC, V-BCC, Octet and VM-BCC lattice patterns at (<b>a</b>) 10 g; (<b>b</b>) 12 g; (<b>c</b>) 14 g equivalent TNT loads. (<b>d</b>) Maximum back face deflection values at different equivalent TNT loads.</p> Full article ">Figure 12
<p>Back face deflection curves of the BCC, CM-BCC, V-BCC, Octet and VM-BCC lattice patterns at (<b>a</b>) 14 mm; (<b>b</b>) 18 mm; (<b>c</b>) 22 mm explosion distances. (<b>d</b>) Maximum back face deflection values at different explosion distances.</p> Full article ">Figure 13
<p>Stress distribution characteristics of the structural patterns, intermediate core layer and back face under 10 g equivalent TNT load and 14 mm explosion distance.</p> Full article ">
22 pages, 20741 KiB  
Article
Microstructure and Properties of Resistance Element Welded Joints of DP780 Steel and 6061 Aluminum Alloy
by Qinglong Wu, Yue Yang, Yingzhe Li, Qing Guo, Shuyue Luo and Zhen Luo
Metals 2025, 15(3), 283; https://doi.org/10.3390/met15030283 - 5 Mar 2025
Viewed by 223
Abstract
This study developed a metallurgical and mechanical hybrid resistance element welding (REW) method to fabricate lightweight Al/steel joints between 2.0 mm 6061 aluminum alloy and 1.2 mm DP780 steel, addressing critical challenges of interfacial intermetallic compounds (IMC layer thickness: 4.6–8.3 μm) in dissimilar [...] Read more.
This study developed a metallurgical and mechanical hybrid resistance element welding (REW) method to fabricate lightweight Al/steel joints between 2.0 mm 6061 aluminum alloy and 1.2 mm DP780 steel, addressing critical challenges of interfacial intermetallic compounds (IMC layer thickness: 4.6–8.3 μm) in dissimilar metal welding. In addition, the scanning electron microscope (SEM), electron backscatter diffraction (EBSD), and electron probe microanalysis (EPMA) were used to observe the microstructure characteristics and element distribution. The lath martensite and solidification microstructure were observed in the steel-nugget zone and Al-nugget zone, respectively. Furthermore, the microhardness distribution, volume fraction of the α phase, tensile–shear load, and failure mode of REWed joint were studied. Process optimization demonstrated welding current’s pivotal role in joint performance, achieving a maximum tensile–shear load of 6914.1 N under 10 kA conditions with a button pull-out failure (BPF) mechanism. Full article
(This article belongs to the Special Issue Modeling and Mechanism Analysis of Welding Process for Metals)
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<p>A schematic diagram of tensile–shear specimens for Al/steel REWed joints.</p> Full article ">Figure 2
<p>Microstructure and XRD results of base metals: (<b>a</b>) microstructure of Q235 steel, (<b>b</b>) XRD results of Q235 steel, (<b>c</b>) microstructure of DP780 steel, (<b>d</b>) XRD results of DP780 steel, (<b>e</b>) microstructure of 6061-T6 Al alloy, and (<b>f</b>) XRD results of 6061-T6 Al alloy.</p> Full article ">Figure 3
<p>Influence of welding currents on Al/steel REWed joints formed with flat head-shaped element weld formation: (<b>a</b>–<b>o</b>) 6 kA~20 kA.</p> Full article ">Figure 4
<p>Cross-section morphology of the Al/steel REWed joints using the flat element with cap at various welding currents: (<b>a</b>–<b>o</b>) 6 kA~20 kA.</p> Full article ">Figure 5
<p>Effect of welding current on the depth of cap penetration into the Al sheet.</p> Full article ">Figure 6
<p>Nugget diameter and the bearing area of Al/steel REWed joints using the flat element with cap as a function of welding current.</p> Full article ">Figure 7
<p>Cross-sectional morphologies of REW joints with the distribution of HAZ.</p> Full article ">Figure 8
<p>SEM images of the REW joint: (<b>a</b>) NZ, (<b>b</b>) FZ of 6061-T6 Al alloy, and (<b>c</b>) FGHAZ near element side.</p> Full article ">Figure 9
<p>EBSD analysis results of the NZ of the REWed welding joint by 8 kA: (<b>a</b>) IPF map, (<b>b</b>) KAM map, and (<b>c</b>) GB map.</p> Full article ">Figure 10
<p>Results of flat-head unit fusion energy spectrum (EDS) line scanning.</p> Full article ">Figure 11
<p>Element distribution in the steel-NZ with flat head-shaped element with cap: (<b>a</b>) SEM image showing steel-NZ of joint; (<b>b</b>–<b>f</b>) elemental distributions of Fe, Mn, Al, C and Si, respectively.</p> Full article ">Figure 12
<p>SEM images of the interfaces in different regions of the REW joint: (<b>a</b>) macroscopic morphology of the unit cross-section, (<b>b</b>) region b in (<b>a</b>), (<b>c</b>) region c in (<b>a</b>), (<b>d</b>) region d in (<b>a</b>), (<b>e</b>) region e in (<b>a</b>), (<b>f</b>) region f in (<b>a</b>), and (<b>g</b>) region g in (<b>a</b>).</p> Full article ">Figure 13
<p>Interfacial compound of 6061 Al alloy/Q235 steel with a flat-head element shape: (<b>a</b>) SEM image of the interface between the element and aluminum alloy, (<b>b</b>) Iron distribution at (<b>a</b>), (<b>c</b>) Manganese distribution at (<b>a</b>), (<b>d</b>) Aluminum distribution at (<b>a</b>), (<b>e</b>) Magnesium distribution at (<b>a</b>), and (<b>f</b>) Silicon distribution at (<b>a</b>).</p> Full article ">Figure 14
<p>Interfacial compound of 6061 Al alloy/DP780 steel with a flat-head element shape: (<b>a</b>) SEM image of the interface between the element and DP780 steel, (<b>b</b>) Iron distribution at (<b>a</b>), (<b>c</b>) Manganese distribution at (<b>a</b>), (<b>d</b>) Aluminum distribution at (<b>a</b>), (<b>e</b>) Magnesium distribution at (<b>a</b>), and (<b>f</b>) Silicon distribution at (<b>a</b>).</p> Full article ">Figure 15
<p>Microhardness distribution of the joint: (<b>a</b>) nugget and steel, (<b>b</b>) aluminum.</p> Full article ">Figure 16
<p>Variation of α-phase fraction: (<b>a</b>) vertical direction path, (<b>b</b>) horizontal direction path.</p> Full article ">Figure 17
<p><b>The</b> relationship between the nugget diameter, peak load, and energy absorption with welding current.</p> Full article ">Figure 18
<p>Failure modes of the REW joint: (<b>a</b>) IF, (<b>b</b>) POF, and (<b>c</b>) BPF.</p> Full article ">Figure 19
<p>Fracture morphology in IF mode: (<b>a</b>) macroscopic morphology, (<b>b</b>) 3D image of the steel side surface, and (<b>c</b>) 3D image of the aluminum side surface.</p> Full article ">Figure 20
<p>Fracture morphology in POF mode: (<b>a</b>) macroscopic morphology, (<b>b</b>) 3D surface image.</p> Full article ">Figure 21
<p>Fracture morphology in BPF mode: (<b>a</b>) macroscopic morphology, (<b>b</b>) 3D surface image.</p> Full article ">Figure 22
<p>Cross-sectional images of different areas of the Al/steel REWed joint using the flat element with a cap by 7 kA welding current: (<b>a</b>) OM image, (<b>b</b>) B region, (<b>c</b>) C region, (<b>d</b>) D region, (<b>e</b>) E region, (<b>f</b>) F region, and (<b>g</b>) G region.</p> Full article ">Figure 23
<p>Cross-sectional images of different areas of the Al/steel REWed joint using the flat element with a cap by 9 kA welding current: (<b>a</b>) OM image, (<b>b</b>) B region, (<b>c</b>) C region, (<b>d</b>) D region, (<b>e</b>) E region, (<b>f</b>) F region, and (<b>g</b>) G region.</p> Full article ">Figure 24
<p>Cross-sectional images of different areas of the Al/steel element welded joint using the flat element with a cap by 16 kA welding current: (<b>a</b>) optical microscope image, (<b>b</b>) B region, (<b>c</b>) C region, (<b>d</b>) D region, and (<b>e</b>) E region.</p> Full article ">Figure 25
<p>Evolution behavior of Al/steel REWed joint using the flat element with a cap at various welding current.</p> Full article ">
21 pages, 12165 KiB  
Article
Microscopic Modeling of Interfaces in Cu-Mo Nanocomposites: The Case Study of Nanometric Metallic Multilayers
by Abdelhafid Akarou, Florence Baras and Olivier Politano
Metals 2025, 15(3), 282; https://doi.org/10.3390/met15030282 - 5 Mar 2025
Viewed by 220
Abstract
Nanocomposites composed of Cu and Mo were investigated by means of molecular dynamics (MD) simulations to study the incoherent interface between Cu and Mo. In order to select an appropriate potential capable of accurately describing the Cu-Mo system, five many-body potentials were compared: [...] Read more.
Nanocomposites composed of Cu and Mo were investigated by means of molecular dynamics (MD) simulations to study the incoherent interface between Cu and Mo. In order to select an appropriate potential capable of accurately describing the Cu-Mo system, five many-body potentials were compared: three Embedded Atom Method (EAM) potentials, a Tight Binding Second Moment Approximation (TB-SMA) potential, and a Modified Embedded Atom Method (MEAM) potential. Among these, the EAM potential proposed by Zhou in 2001 was determined to provide the best compromise for the current study. The simulated system was constructed with two layers of Cu and Mo forming an incoherent fcc-Cu(111)/bcc-Mo(110) interface, based on the Nishiyama–Wassermann (NW) and Kurdjumov–Sachs (KS) orientation relationships (OR). The interfacial energies were calculated for each orientation relationship. The NW configuration emerged as the most stable, with an interfacial energy of 1.83 J/m², compared to 1.97 J/m² for the KS orientation. Subsequent simulations were dedicated to modeling Cu atomic deposition onto a Mo(110) substrate at 300 K. These simulations resulted in the formation of a dense layer with only a few defects in the two Cu planes closest to the interface. The interfacial structures were characterized by computing selected area electron diffraction (SAED) patterns. A direct comparison of theoretical and numerical SAED patterns confirmed the presence of the NW orientation relationship in the nanocomposites formed during deposition, corroborating the results obtained with the model fcc-Cu(111)/bcc-Mo(110) interfaces. Full article
(This article belongs to the Special Issue Design and Development of Metal Matrix Composites)
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<p>Comparison between the MD results obtained for Cu and Mo with experimental results from Touloukian [<a href="#B38-metals-15-00282" class="html-bibr">38</a>].</p> Full article ">Figure 2
<p>The evolution of the coefficient of thermal expansion as a function of temperature of copper and molybdenum.</p> Full article ">Figure 3
<p>Evolution of the ratio of the thermal expansion coefficient of Cu and Mo as a function of temperature.</p> Full article ">Figure 4
<p>The variation of the young modulus <span class="html-italic">E</span>, Poisson ratio <math display="inline"><semantics> <mi>υ</mi> </semantics></math>, shear modulus <span class="html-italic">G</span>, and bulk modulus <span class="html-italic">K</span> of Cu as a function of temperature.</p> Full article ">Figure 5
<p>The variation of the young modulus <span class="html-italic">E</span>, Poisson ratio <math display="inline"><semantics> <mi>υ</mi> </semantics></math>, shear modulus <span class="html-italic">G</span>, and bulk modulus <span class="html-italic">K</span> of Mo as a function of temperature.</p> Full article ">Figure 6
<p>KS and NW boxes. Mo and Cu atoms are represented by blue and orange spheres, respectively.</p> Full article ">Figure 7
<p>Two dimensional partial radial distribution function computed in the first atomic plane of Cu and of Mo at the interface. The theoretical positions of the first, second, and third neighbor peaks in a 2D plane from a bulk system are indicated by the lines with a star (∗), a diamond (⧫), and a cross (x), respectively.</p> Full article ">Figure 8
<p>Energy, atomic volume and local atomic environment of KS box. For direct comparison, cohesive energy and volume per atom for bulk systems are also reported.</p> Full article ">Figure 9
<p>Energy, atomic volume and local atomic envirronement of NW box. For direct comparison, cohesive energy and volume per atom for bulk systems are also reported.</p> Full article ">Figure 10
<p>Excess of potential energy at the interface at 0 K.</p> Full article ">Figure 11
<p>MD model of the deposition of sputtered Cu atoms onto Mo surface.</p> Full article ">Figure 12
<p>Two-dimensional radial distribution functions computed for the first three planes of Cu and the Mo plane located at the interface. The positions of the peaks computed in Cu(111) and Mo(110) planes in bulk systems are presented for comparison by lines with a cross (x) and a star (∗), respectively. Side views of the final configurations from simulations at 300 K. Thin slices of Cu atoms at different heights are shown on (<b>a</b>–<b>d</b>). In snapshots (<b>a</b>–<b>c</b>), the atoms are colored according to their local environment (i.e., <span class="html-italic">fcc</span> (green), <span class="html-italic">bcc</span> (blue), <span class="html-italic">hcp</span> (red), and <span class="html-italic">unk</span> (gray)), whereas in snapshot (<b>d</b>), the atoms are colored according to their relative height.</p> Full article ">Figure 13
<p>(<b>a</b>,<b>b</b>) Calculated diffraction pattern of <span class="html-italic">bcc</span>-Mo substrate with zone axis aligned along the [011] and [110] directions. (<b>c</b>) Calculated diffraction pattern of <span class="html-italic">fcc</span>-Cu with zone axis aligned along the [111]-direction.</p> Full article ">Figure 14
<p>SAED for Nishiyama–Wassermann (NW) relationship. (<b>a</b>) Theoretical diffraction patterns. (<b>b</b>) Virtual diffraction patterns. Mo and Cu peaks are represented in blue and orange, respectively.</p> Full article ">Figure 15
<p>SAED for Kurdjumov–Sachs (KS) relationship. (<b>a</b>) Theoretical diffraction patterns. (<b>b</b>) Virtual diffraction patterns. Mo and Cu peaks are represented in blue and orange, respectively.</p> Full article ">Figure 16
<p>(<b>a</b>) SAED patterns computed for the full system at the end of deposition simulation. The black square defined by the Mo peaks, and the black hexagon, defined by the Cu peaks, are drawn as guides to identify the structure. (<b>b</b>) superposition of theoretical and virtual SAED patterns (deposition simulation). Theoretical Mo and Cu peaks are represented in blue and orange, respectively.</p> Full article ">Figure 17
<p>(<b>a</b>) SAED patterns computed for the full system at the end of deposition simulation. The black square defined by the Mo peaks, and the black hexagon, defined by the Cu peaks, are drawn as guides to identify the structure. (<b>b</b>) superposition of theoretical and virtual SAED patterns (deposition simulation). Theoretical Mo and Cu peaks are represented in blue and orange, respectively.</p> Full article ">
17 pages, 1643 KiB  
Article
An Innovative Algorithm for Damage Mapping in Multiaxial Fatigue Using the Stress Scale Factor (SSF) Concept
by Francisco Bumba, Vitor Anes and Luis Reis
Metals 2025, 15(3), 281; https://doi.org/10.3390/met15030281 - 5 Mar 2025
Viewed by 120
Abstract
Predicting damage to materials under multiaxial fatigue is a complex challenge, especially when normal and shear stresses interact in dynamic and non-linear ways. Traditional methods often oversimplify these interactions, leading to less reliable fatigue predictions and limiting their usefulness in real-world applications. To [...] Read more.
Predicting damage to materials under multiaxial fatigue is a complex challenge, especially when normal and shear stresses interact in dynamic and non-linear ways. Traditional methods often oversimplify these interactions, leading to less reliable fatigue predictions and limiting their usefulness in real-world applications. To address this, we present a novel algorithm based on the principles of the Stress Scale Factor (SSF), designed to dynamically evaluate the relative contributions of normal and shear stresses to fatigue damage. By providing a more accurate mapping of multiaxial fatigue damage, this approach enables improved predictions of fatigue life. The methodology combines experimental insights with mathematical modeling to create a flexible and adaptive framework. By making it possible to map multiaxial fatigue damage with greater precision, this SSF-based approach not only enhances the understanding of fatigue behavior but also enables better design decisions. The result is safer, more reliable, and efficient structures across a range of applications. This study bridges the gap between theoretical methods and practical needs, offering engineers and researchers a powerful tool to improve fatigue analysis and optimize structural performance. Full article
(This article belongs to the Special Issue Novel Insights into Fatigue and Fracture Behavior of Metallic Materials)
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<p>Correlation between normal stress amplitude and loading path: (<b>a</b>) pure tensile loading, (<b>b</b>) proportional loading with SAR = 30°, (<b>c</b>) proportional loading with SAR = 45°, and (<b>d</b>) proportional loading with SAR = 60°.</p> Full article ">Figure 2
<p>Regression analysis of the polynomial constants from <a href="#metals-15-00281-t005" class="html-table">Table 5</a>: (<b>a</b>) third-degree polynomial regression for parameter a, (<b>b</b>) third-degree polynomial regression for parameter b.</p> Full article ">Figure 3
<p>Regression analysis of the polynomial constant a: (<b>a</b>) two-degree polynomial regression, (<b>b</b>) third-degree polynomial regression.</p> Full article ">Figure 4
<p>Regression analysis of the polynomial constant b: (<b>a</b>) two-degree polynomial regression, (<b>b</b>) third-degree polynomial regression.</p> Full article ">
23 pages, 6226 KiB  
Article
Optimizing FSP Parameters for AA5083/SiC Composites: A Comparative Analysis of Taguchi and Regression
by Oritonda Muribwathoho, Velaphi Msomi and Sipokazi Mabuwa
Metals 2025, 15(3), 280; https://doi.org/10.3390/met15030280 - 5 Mar 2025
Viewed by 152
Abstract
The fabrication of AA5083/SiC composites by the friction stir processing (FSP) method is the main objective of this study. The study looks at how the mechanical properties of the composites are affected by three important process parameters: traversal speed, rotational speed, and tilt [...] Read more.
The fabrication of AA5083/SiC composites by the friction stir processing (FSP) method is the main objective of this study. The study looks at how the mechanical properties of the composites are affected by three important process parameters: traversal speed, rotational speed, and tilt angle. The Taguchi L9 design matrix was used to effectively investigate parameter effects, decreasing experimental trials and cutting expenses. Tensile testing measured tensile strength, whereas microhardness tests evaluated hardness. The findings showed that a maximum tensile strength of 243 MPa and a maximum microhardness of 94.80 HV were attained. The findings also showed that the optimal ultimate tensile strength (UTS) and percentage elongation (PE) were achieved at a tilt angle of 2°, a traverse speed of 30 mm per minute, and a rotating speed of 900 rev/min. On the other hand, a slightly greater traverse speed of 45 mm per minute was required to reach maximal microhardness (MH) with the same rotational speed and tilt angle. Analysis of variance (ANOVA) showed that rotational speed has a substantial impact on all mechanical properties, highlighting how important it is for particle dispersion and grain refining. This work is unique in that it systematically optimizes FSP parameters by using regression analysis and the Taguchi technique in addition to ANOVA. This allows for a better understanding of how these factors affect the mechanical properties of SiC-reinforced composites. The findings contribute to advancing the cost-effective fabrication of high-performance metal matrix composites for industrial applications requiring enhanced strength and durability. Full article
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<p>(<b>a</b>) FSW procedure; (<b>b</b>) Drilling of holes and filling them with SiC particles; (<b>c</b>) Using a pinless tool to close the hole; (<b>d</b>) FSP single-pass procedure; (<b>e</b>) Tool with pin tool; (<b>f</b>) Pinless.</p> Full article ">Figure 2
<p>Dimensions and arrangement of the hardness and tensile specimens.</p> Full article ">Figure 3
<p>Results from experiments and regressions for (<b>a</b>) Microhardness (MH), (<b>b</b>) Percentage elongation (PE), and (<b>c</b>) Ultimate tensile strength (UTS).</p> Full article ">Figure 3 Cont.
<p>Results from experiments and regressions for (<b>a</b>) Microhardness (MH), (<b>b</b>) Percentage elongation (PE), and (<b>c</b>) Ultimate tensile strength (UTS).</p> Full article ">Figure 4
<p>AA5083/SiC composite optical microstructures photographed at 20 × 100 µm magnification with a 100 µm scale bar. (<b>a<sub>1</sub></b>–<b>a<sub>3</sub></b>) Microstructures at 600 rev/min with traverse rates of 30 mm per min, 45 mm per min, and 60 mm per min, respectively, and tilt angles of 1°, 1.75°, and 2°. (<b>b<sub>1</sub></b>–<b>b<sub>3</sub></b>) Microstructures at 900 rev/min with traverse rates of 30 mm per min, 45 mm per min, and 60 mm per min, respectively, and tilt angles of 1.75°, 2°, and 1°. (<b>c<sub>1</sub></b>–<b>c<sub>3</sub></b>) Microstructures having traverse rates of 30 mm per min, 45 mm per min, and 60 mm per min at 1200 rev/min with tilt angles of 2°, 1°, and 1.75°, respectively.</p> Full article ">Figure 5
<p>AA5083/SiC composite S/N ratio and mean plot: (<b>a</b>) Microhardness; (<b>b</b>) Percentage elongation; (<b>c</b>) Ultimate tensile strength.</p> Full article ">Figure 6
<p>Percentage contribution for process parameters.</p> Full article ">Figure 7
<p>Probability Plots; (<b>a</b>) Microhardness, (<b>b</b>) Percentage elongation, (<b>c</b>) Ultimate tensile strength.</p> Full article ">Figure 7 Cont.
<p>Probability Plots; (<b>a</b>) Microhardness, (<b>b</b>) Percentage elongation, (<b>c</b>) Ultimate tensile strength.</p> Full article ">
18 pages, 2081 KiB  
Article
Characterization of EAF and LF Slags Through an Upgraded Stationary Flowsheet Model of the Electric Steelmaking Route
by Ismael Matino, Alice Petrucciani, Antonella Zaccara, Valentina Colla, Maria Ferrer Prieto and Raquel Arias Pérez
Metals 2025, 15(3), 279; https://doi.org/10.3390/met15030279 - 4 Mar 2025
Viewed by 239
Abstract
The current, continuous increase in attention toward preservation of the environment and natural resources is forcing resource-intensive industries like steelworks to investigate new solutions to improve resource efficiency and promote the growth of a circular economy. In this context, electric steelworks, which inherently [...] Read more.
The current, continuous increase in attention toward preservation of the environment and natural resources is forcing resource-intensive industries like steelworks to investigate new solutions to improve resource efficiency and promote the growth of a circular economy. In this context, electric steelworks, which inherently implement circularity principles, are spending efforts to enhance valorization of their main by-product, namely slags. A reliable characterization of the slag’s composition is crucial for the identification of the best valorization pathway, but, currently, slag monitoring is often discontinuous. Furthermore, in the current period of transformation of steel production, preliminary knowledge of the effect of modifications of operating practices on slags composition is crucial to assessing the viability of these modifications. In this paper, a stationary flowsheet model of the electric steelmaking route is presented; this model enables joint monitoring of key variables related to process, steel and slags. For the estimation of the content of most compounds in slags, the average relative percentage error is below 20% for most of the considered steel families. Thus, the tool can be considered suitable for scenario analyses supporting slag valorization. Higher performance is achievable by exploiting more reliable data for model tuning. These data can be obtained via novel devices that gather more numerous and representative data on the amount and composition of slags. Full article
(This article belongs to the Special Issue Recent Advances in the Recycling and Reuse of Metallurgical Wastes and By-Products)
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<p>Main sections, inputs and outputs of upgraded model.</p> Full article ">Figure 2
<p>Pareto diagrams of RPEs of tested heats for the content of main EAF slag compounds.</p> Full article ">Figure 3
<p>Pareto diagrams of RPEs of tested heats for the content of main LF slag compounds.</p> Full article ">Figure 4
<p>RPEs for main compounds of EAF (<b>top</b>) and LF (<b>bottom</b>) slags belonging to a single simulated heat.</p> Full article ">
19 pages, 3694 KiB  
Review
Review of the Properties and Degradation Mechanisms of Refractories in Aluminum Reduction Cells
by Mohamed Hassen Ben Salem, Gervais Soucy, Daniel Marceau, Antoine Godefroy and Sébastien Charest
Metals 2025, 15(3), 278; https://doi.org/10.3390/met15030278 - 4 Mar 2025
Viewed by 222
Abstract
This review examines the degradation of refractory materials in aluminum reduction cells, focusing specifically on contamination caused by the cryolite-based bath. Aluminosilicate refractories, particularly Ordinary Refractory Bricks, play a vital role in maintaining the structural integrity and thermal balance of these cells under [...] Read more.
This review examines the degradation of refractory materials in aluminum reduction cells, focusing specifically on contamination caused by the cryolite-based bath. Aluminosilicate refractories, particularly Ordinary Refractory Bricks, play a vital role in maintaining the structural integrity and thermal balance of these cells under demanding operational conditions. The interaction between the molten bath and refractory linings leads to chemical reactions and mineralogical changes that modify the mechanical and thermal properties of the material over time. The study integrates findings from industrial autopsies, laboratory experiments, and a comprehensive review of the existing literature to identify and analyze the mechanisms of degradation. By analyzing the findings obtained from these methodologies, this review explores how cryolitic infiltration triggers transformations that compromise performance and reduce the lifespan of refractory linings. Covering a broad temperature range (665–960 °C), the study addresses key challenges in understanding bath-induced contamination and provides insights into how to improve the durability and efficiency of refractory materials in aluminum production. Full article
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<p>Diagram of a reduction cell [<a href="#B18-metals-15-00278" class="html-bibr">18</a>]. Reproduced with permission from Springer Nature.</p> Full article ">Figure 2
<p>Schematic representation of (<b>a</b>) open and closed porosity in refractory materials and (<b>b</b>) their influence on thermal conductivity and resistance to infiltration [<a href="#B1-metals-15-00278" class="html-bibr">1</a>]. Reproduced with permission from Springer Nature.</p> Full article ">Figure 3
<p>Thermal conductivity of lightweight fireclay brick (1), fireclay brick (2), silica brick (3), mullite brick (4), high alumina brick (85% Al<sub>2</sub>O<sub>3</sub>) (5), magnesia (6), zircon (7), chromite (8), and alumina (approx. &gt; 90%)(9) according to [<a href="#B1-metals-15-00278" class="html-bibr">1</a>]. Reproduced with permission from Springer Nature.</p> Full article ">Figure 4
<p>Diagram of interplay between thermal, chemical, and mechanical phenomena in refractory materials.</p> Full article ">Figure 5
<p>Na<sub>2</sub>O-SiO<sub>2</sub>-Al<sub>2</sub>O<sub>3</sub> phase diagram.</p> Full article ">Figure 6
<p>Phase compositions due to chemical reactions between alumina–silicate materials and sodium fluoride [<a href="#B25-metals-15-00278" class="html-bibr">25</a>]. Reproduced with permission from Minerals, Metals and Materials Society.</p> Full article ">Figure 7
<p>A cross-sectional image of degraded ORBs showing distinct layers, as described in industrial autopsies [<a href="#B10-metals-15-00278" class="html-bibr">10</a>]. Reproduced with permission from Springer Nature.</p> Full article ">
19 pages, 6589 KiB  
Article
Atmospheric Corrosion Behavior of Typical Aluminum Alloys in Low-Temperature Environment
by Tengfei Cui, Jianguo Wu, Jian Song, Di Meng, Xiaoli Jin, Huiyun Tian and Zhongyu Cui
Metals 2025, 15(3), 277; https://doi.org/10.3390/met15030277 - 4 Mar 2025
Viewed by 119
Abstract
The atmospheric corrosion behavior of type 2024, 5083, 6061, and 7075 aluminum alloys in the Antarctic environment was investigated by outdoor exposure tests and indoor characterization. After one year of exposure to the Antarctic atmosphere, significant differences in surface corrosion states were observed [...] Read more.
The atmospheric corrosion behavior of type 2024, 5083, 6061, and 7075 aluminum alloys in the Antarctic environment was investigated by outdoor exposure tests and indoor characterization. After one year of exposure to the Antarctic atmosphere, significant differences in surface corrosion states were observed among the specimens. The results revealed that the corrosion rate of the 2024 aluminum alloy was the highest, reaching 14.5 g/(m2·year), while the 5083 aluminum alloy exhibited the lowest corrosion rate of 1.36 g/(m2·year). The corrosion products formed on the aluminum alloys exposed to the Antarctic environment were primarily composed of AlOOH and Al2O3. In the Antarctic atmosphere environment, the pits were dominated by a freezing–thawing cycle and salt deposition. The freezing–thawing cycle promotes the wedge effect of corrosion products at the grain boundary, resulting in exfoliation corrosion of high-strength aluminum alloys. Full article
(This article belongs to the Special Issue Corrosion of Metals: Behaviors and Mechanisms)
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<p>Microstructure characteristics of 2024 (<b>a</b>,<b>a1</b>), 5083 (<b>b</b>,<b>b1</b>), 6061 (<b>c</b>,<b>c1</b>), and 7075 (<b>d</b>,<b>d1</b>) aluminum alloys, as well as the fraction of the granular intermetallic compounds of a typical aluminum alloy field area (<b>e</b>). The red boxes are the EDS test area.</p> Full article ">Figure 2
<p>Corrosion rates of 2024, 5083, 6061, and 7075 aluminum alloys exposed to Antarctic atmospheric environment.</p> Full article ">Figure 3
<p>XRD spectra of 2024, 5083, 6061, and 7075 aluminum alloys exposed to Antarctic atmospheric environment.</p> Full article ">Figure 4
<p>Macromorphology of the skyward and groundward surfaces of 2024 aluminum alloy (<b>a</b>,<b>b</b>), 5083 aluminum alloy (<b>c</b>,<b>d</b>), 6061 aluminum alloy (<b>e</b>,<b>f</b>), and 7075 aluminum alloy (<b>g</b>,<b>h</b>) under Antarctic atmospheric environment.</p> Full article ">Figure 5
<p>The magnified surface morphologies at 100× and 500×, along with the results of EDS (energy-dispersive spectroscopy, point analysis), of 2024 (<b>a<sub>1</sub></b>,<b>b<sub>1</sub></b>), 5083 (<b>a<sub>2</sub></b>,<b>b<sub>2</sub></b>), 6061 (<b>a<sub>3</sub></b>,<b>b<sub>3</sub></b>), and 7075 (<b>a<sub>4</sub></b>,<b>b<sub>4</sub></b>) aluminum alloys after one year of exposure to the Antarctic atmospheric environment.</p> Full article ">Figure 6
<p>The cross-sectional morphologies and EDS mapping results of 2024 (<b>a</b>), 5083 (<b>b</b>), 6061 (<b>c</b>), and 7075 (<b>d</b>) aluminum alloys after one year of exposure to the Antarctic atmospheric environment. The light-gray sections represent the substrate, while the areas marked with blue dashed lines indicate the corrosion products.</p> Full article ">Figure 7
<p>XPS spectra of 2024 (<b>a</b>), 5083 (<b>b</b>), 6061 (<b>c</b>), and 7075 (<b>d</b>) aluminum alloys exposed to the Antarctic atmospheric environment.</p> Full article ">Figure 8
<p>Surface morphologies of 2024 (<b>a</b>,<b>b</b>), 5083 (<b>c</b>,<b>d</b>), 6061 (<b>e</b>,<b>f</b>), and 7075 (<b>g</b>,<b>h</b>) aluminum alloys after removing corrosion products.</p> Full article ">Figure 9
<p>3D corrosion morphologies of 2024 (<b>a</b>,<b>b</b>), 5083 (<b>c</b>,<b>d</b>), 6061 (<b>e</b>,<b>f</b>), and 7075 (<b>g</b>,<b>h</b>) aluminum alloys after removing corrosion products.</p> Full article ">Figure 10
<p>Pit cumulative probability (<b>a</b>), pit depth statistics (<b>b</b>), average depth of corrosion defects (<b>c</b>), and maximum depth of corrosion defects (<b>d</b>) for 2024, 5083, 6061, and 7075 aluminum alloys.</p> Full article ">Figure 11
<p>Schematic illustration of corrosion process of aluminum alloy in Antarctic environment (<b>a</b>), pit initiation process (<b>b</b>), intergranular corrosion process (<b>c</b>), exfoliation corrosion process (<b>d</b>).</p> Full article ">
12 pages, 468 KiB  
Review
Recent Hydrometallurgical Investigations to Recover Antimony from Wastes
by Francisco Jose Alguacil
Metals 2025, 15(3), 276; https://doi.org/10.3390/met15030276 - 3 Mar 2025
Viewed by 121
Abstract
Antimony is a chemical element with diverse uses that falls into the range of a critical raw material. Although it appears in nature as stibnite, the mining of this mineralogical species is rare or uncommon, and it is the element that is basically [...] Read more.
Antimony is a chemical element with diverse uses that falls into the range of a critical raw material. Although it appears in nature as stibnite, the mining of this mineralogical species is rare or uncommon, and it is the element that is basically recovered as a secondary material in the processing of various elements (such as gold and copper). Another source for the recovery of this element is the recycling of Sb-bearing wastes such as batteries and alloys. Once dissolved and in order to recover it from the different leachates, adsorption processes are the ones that seem to have, at least for the scientific community, the highest acceptance. This work reviews the most recent advances (in 2024) in the recovery of antimony from different sources using not only adsorption processes but also other technologies of practical interest. Full article
(This article belongs to the Special Issue Hydrometallurgical Processes for the Recovery of Critical Metals)
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<p>Main hydrometallurgical operational units.</p> Full article ">
11 pages, 4938 KiB  
Article
Influence of Heat Treatment Temperature on the Electrochemical Properties of Cold-Rolled 0.2%C–3%Al–6/8.5%Mn–Fe Medium-Manganese Steel
by Jihui Luo and Huixin Zuo
Metals 2025, 15(3), 275; https://doi.org/10.3390/met15030275 - 3 Mar 2025
Viewed by 160
Abstract
The microstructure evolution, polarization curve and impedance of cold-rolled 0.2%C–3%Al–6/8.5%Mn–Fe steel under heat treatment temperatures of 600–800 °C holding 10 min were tested. The results show that the cold-rolled texture of the steel does not completely disappear at 600 °C and 650 °C, [...] Read more.
The microstructure evolution, polarization curve and impedance of cold-rolled 0.2%C–3%Al–6/8.5%Mn–Fe steel under heat treatment temperatures of 600–800 °C holding 10 min were tested. The results show that the cold-rolled texture of the steel does not completely disappear at 600 °C and 650 °C, exhibiting high charge transfer resistance Rc and corresponding corrosion potential Ecorr. When the heat treatment temperature rises to 700 °C, the texture begins to be eliminated and the Rc begins to decrease, indicating a decrease in corrosion resistance. When the heat treatment temperature rises to 750 °C and 800 °C, it was found that the proportion of austenite begins to increase and the number of grain boundaries decreases, resulting in an increase in Rc and an improvement in the corrosion resistance of the steel. Compared to 6.5 Mn steel, the higher Mn content in 8.5 Mn steel results in better corrosion resistance after high-temperature heat treatment. Full article
(This article belongs to the Section Corrosion and Protection)
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<p>SEM of 0.2%C–3%Al–6%Mn–Fe steel with different heat treatment temperatures. (<b>a</b>) 600 °C; (<b>b</b>) 650 °C; (<b>c</b>) 700 °C; (<b>d</b>) 750 °C; (<b>e</b>) 800 °C.</p> Full article ">Figure 2
<p>SEM of 0.2%C–3%Al–8.5%Mn–Fe steel with different heat treatment temperatures. (<b>a</b>) 600 °C; (<b>b</b>) 650 °C; (<b>c</b>) 700 °C; (<b>d</b>) 750 °C; (<b>e</b>) 800 °C.</p> Full article ">Figure 3
<p>Polarization curves of 0.2%C–3%Al–6/8.5%Mn–Fe steel with different heat treatment temperatures. (<b>a</b>) 6 Mn; (<b>b</b>) 8.5 Mn.</p> Full article ">Figure 4
<p>Nyquist diagrams of 0.2%C–3%Al–6/8.5%Mn–Fe steel with different heat treatment temperatures. (<b>a</b>) 6 Mn; (<b>b</b>) 8.5 Mn.</p> Full article ">Figure 5
<p>Bode curves of 0.2%C–3%Al–6/8.5%Mn–Fe steel with different heat treatment temperatures. (<b>a</b>,<b>b</b>) 6 Mn; (<b>c</b>,<b>d</b>) 8.5 Mn.</p> Full article ">Figure 6
<p>Equivalent circuit for the tested steel in 3.5% NaCl solution.</p> Full article ">
17 pages, 4214 KiB  
Article
Metallic Metamaterials for Reducing the Magnetic Signatures of Ships
by Fabio Distefano, Roberto Zivieri, Gabriella Epasto, Antonio Pantano and Vincenzo Crupi
Metals 2025, 15(3), 274; https://doi.org/10.3390/met15030274 - 3 Mar 2025
Viewed by 222
Abstract
In this study, the magnetic signatures of ship structures were investigated. The magnetic signature impacts both navigation safety and the health of the marine ecosystem. Reducing this signature is essential for minimising risks associated with navigation and protecting marine biodiversity. A finite element [...] Read more.
In this study, the magnetic signatures of ship structures were investigated. The magnetic signature impacts both navigation safety and the health of the marine ecosystem. Reducing this signature is essential for minimising risks associated with navigation and protecting marine biodiversity. A finite element model was developed to assess the magnetic signature of honeycomb sandwich panels for ship structures. A theoretical approach was proposed, and the predicted results were compared with the values obtained by the finite element analyses. Different types of structures were compared to evaluate the combined effect of materials and geometry on the magnetic signature. The finite element results and the theoretical predictions indicate that the use of metamaterial structures, consisting of honeycomb sandwich panels with a steel core and aluminium skins, produces a significant reduction of the ship magnetic signature compared to the one arising from a steel panel with the same bending stiffness. Full article
(This article belongs to the Special Issue Metallic Magnetic Materials: Manufacture, Properties and Applications)
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<p>Honeycomb geometrical parameters: a and <span class="html-italic">b</span> represent respectively the specimen’s length and width, <span class="html-italic">t</span> is the skin thickness, <span class="html-italic">c</span> is the core thickness, <span class="html-italic">t<sub>c</sub></span> is the foil thickness and <span class="html-italic">d<sub>c</sub></span> represents the cell dimension. (<b>a</b>) honeycomb sandwich panel; (<b>b</b>) unit cell of honeycomb core.</p> Full article ">Figure 2
<p>Results of the mesh sensitivity study.</p> Full article ">Figure 3
<p>(<b>a</b>) Domain simulating the earth’s magnetic field. (<b>b</b>) Honeycomb panel coordinate system.</p> Full article ">Figure 4
<p>Trend of the induced magnetic signature on a path along the <span class="html-italic">z</span>-axis passing on the honeycomb’s centreline for BS 60 mm × 60 mm × 8.5 mm (black line), SHS 60 mm × 60 mm × 9 mm (red line), and ASHS 60 mm × 60 mm × 9 mm (blue line) calculated according to FEM at a distance of 10 m from the panel’s surface.</p> Full article ">Figure 5
<p>Induced magnetic signature vs. <span class="html-italic">z</span> coordinate till <span class="html-italic">z</span> = 10 m for BS 60 mm × 60 mm × 8.5 mm (black line), SHS 60 mm × 60 mm × 9 mm (red line), and ASHS 60 mm × 60 mm × 9 mm (blue line). Inset: induced magnetic signature vs. <span class="html-italic">z</span> coordinate till <span class="html-italic">z</span> = 10 m for SHS (red line) and ASHS (blue line).</p> Full article ">Figure 6
<p>Trend of the induced magnetic signature on a path along the <span class="html-italic">z</span>-axis passing on the honeycomb’s centreline for BS 60 mm × 60 mm × 8.5 mm (black line), SHS 60 mm × 60 mm × 9 mm (red line), and ASHS 60 mm × 60 mm × 9 mm (blue line) calculated according to FEM at a distance of 1 m from the panel’s surface.</p> Full article ">Figure 7
<p>Induced magnetic signature vs. <span class="html-italic">z</span> coordinate till <span class="html-italic">z</span> = 1 m for BS 60 mm × 60 mm × 8.5 mm (black line), SHS 60 mm × 60 mm × 9 mm (red line), and ASHS 60 mm × 60 mm × 9 mm (blue line). Inset: induced magnetic signature vs. <span class="html-italic">z</span> coordinate till <span class="html-italic">z</span> = 1 m for SHS (red line) and ASHS (blue line).</p> Full article ">Figure 8
<p>Induced magnetic signature vs. <span class="html-italic">z</span> coordinate till <span class="html-italic">z</span> = 10 m for BS 600 mm × 600 mm × 8.5 mm (black line), SHS 600 mm × 600 mm × 9 mm (red line), and 600 mm × 600 mm × 9 mm (blue line). Inset: induced magnetic signature vs. <span class="html-italic">z</span> coordinate till <span class="html-italic">z</span> = 10 m for the honeycomb with steel skins (red line) and the honeycomb with aluminium skins (blue line).</p> Full article ">Figure 9
<p>Induced magnetic signature vs. <span class="html-italic">z</span> coordinate till <span class="html-italic">z</span> = 1 m for BS 600 mm × 600 mm × 8.5 mm (black line), SHS 600 mm × 600 mm × 9 mm (red line), and ASHS 600 mm × 600 mm × 9 mm (blue line). Inset: induced magnetic signature vs. <span class="html-italic">z</span> coordinate till <span class="html-italic">z</span> = 1 m for SHS (red line) and ASHS (blue line).</p> Full article ">Figure 10
<p>Comparison between the induced magnetic signature vs. <span class="html-italic">z</span> coordinate of the FEM (red lines) and of the TM (black lines) till <span class="html-italic">z</span> = 0.2 m for (<b>a</b>) BS 60 mm × 60 mm × 8.5 mm, (<b>b</b>) SHS 60 mm × 60 mm × 9 mm, and (<b>c</b>) ASHS 60 mm × 60 mm × 9 mm.</p> Full article ">
17 pages, 3797 KiB  
Article
Influence of Sulfide Concentration on the Properties of Cr3C2-25(Ni20Cr) Cermet Coating on Al7075 Substrate
by Mieczyslaw Scendo
Metals 2025, 15(3), 273; https://doi.org/10.3390/met15030273 - 2 Mar 2025
Viewed by 341
Abstract
The influence of sulfide (S2−) concentration on the corrosion resistance of Cr3C2-25(Ni20Cr) cermet coating on Al7075 (EN, AW-7075) substrate (Cr3C2-25(Ni20Cr)/Al7075) was investigated. The coating was produced by the cold-sprayed (CS) method. The Cr [...] Read more.
The influence of sulfide (S2−) concentration on the corrosion resistance of Cr3C2-25(Ni20Cr) cermet coating on Al7075 (EN, AW-7075) substrate (Cr3C2-25(Ni20Cr)/Al7075) was investigated. The coating was produced by the cold-sprayed (CS) method. The Cr3C2-25(Ni20Cr)/Al7075 coatings were modified chemically in solutions containing thioacetic acid amide (TAA). The surface and microstructure of the specimens were both observed by a scanning electron microscope (SEM). The mechanical properties of the Cr3C2-25(Ni20Cr) coatings were characterized using microhardness (HV) measurements. The corrosion tests of the materials were carried out using the electrochemical method in a acidic chloride solution. The adsorbed (MemSn)ads layer effectively separates the Cr3C2-25(Ni20Cr)/Al7075 coating surface from contact with the aggressive corrosive environment. More than a twice lower value of corrosion rate (CW) was obtained for the Cr3C2-25(Ni20Cr)/Al7075 coating after exposure to the environment with 0.15 M TAA. Full article
(This article belongs to the Special Issue Corrosion Behavior of Alloys in Water Environments)
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Figure 1
<p>SEM micrograph: (<b>a</b>) Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) powder morphology, (<b>b</b>) X-ray diffraction pattern of powder.</p> Full article ">Figure 2
<p>Surface morphology of cold-sprayed cermet coating on the Al7075 substrate: (<b>a</b>) Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr)/Al7075, (<b>b</b>) X-ray diffraction pattern of as-sprayed coating.</p> Full article ">Figure 3
<p>SEM surfaces: (<b>a</b>) Al7075 substrate, (<b>b</b>) Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) cermet coating, and (<b>c</b>) cross-section of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) cermet coating on the Al7075 substrate.</p> Full article ">Figure 4
<p>SEM surface: (<b>a</b>) Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) cermet coating on the Al7075 substrate, (<b>b</b>) cross-section of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr)/Al7075 after exposure to an environment of 1.2 M Cl<sup>−</sup> (pH 1.5).</p> Full article ">Figure 5
<p>SEM of surface morphology of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) cermet coatings on the Al7075 substrate after exposure in thioacetic acid amide: (<b>a</b>) 0.05 M, (<b>b</b>) 0.10 M, (<b>c</b>) 0.15 M, and (<b>d</b>) cross-section of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr)/Al7075 for the 0.15 M TAA.</p> Full article ">Figure 6
<p>Evolution of open circuit potential of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) cermet coating on the Al7075 substrate after exposure in (a) 0 M, (b) 0.05 M, (c) 0.10 M, and (d) 0.15 M TAA. The test solution contained 1.2 M Cl<sup>−</sup> (pH 1.5).</p> Full article ">Figure 7
<p>Potentiodynamic polarization curves of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) cermet coatings on the Al7075 substrate after exposure in (<b>a</b>) 0 M, (<b>b</b>) 0.05 M, (<b>c</b>) 0.10 M, and (<b>d</b>) 0.15 M TAA. Test solution contained 1.2 M Cl<sup>−</sup> (pH 1.5). d<span class="html-italic">E</span>/d<span class="html-italic">t</span> 1 mV/s.</p> Full article ">Figure 8
<p>Chronoamperometric curves of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) cermet coatings on the Al7075 substrate after exposure in 0.15 M TAA. The test solution contained 1.2 M Cl<sup>−</sup> (pH 1.5). Potential values were as follows: (<b>a</b>) −900 mV, (<b>b</b>) −260 mV, (<b>c</b>) +100 mV (average current density values are marked with a dashed line).</p> Full article ">Figure 9
<p>Potentiodynamic polarization curves on a semi-logarithmic (Tafel) scale of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) cermet coatings on the Al7075 substrate after exposure in (<b>a</b>) 0 M, (<b>b</b>) 0.05 M, (<b>c</b>) 0.10 M, and (<b>d</b>) 0.15 M TAA. The test solution contained 1.2 M Cl<sup>−</sup> (pH 1.5).</p> Full article ">Figure 10
<p>SEM surface of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr) cermet coatings on the Al7075 substrate after exposure in (<b>a</b>) 0.05 M, (<b>b</b>) 0.10 M, and (<b>c</b>) 0.15 M TAA. (<b>d</b>) Cross-section of Cr<sub>3</sub>C<sub>2</sub>-25(Ni20Cr)/Al7075 for the 0.15 M TAA test solution containing 1.2 M Cl<sup>−</sup> (pH 1.5).</p> Full article ">
14 pages, 8263 KiB  
Article
Microstructural, Electrochemical, Mechanical, and Biocompatibility Characterization of ReN Thin Films Synthesized by DC Sputtering on Ti6Al4V Substrates
by Willian Aperador, Giovany Orozco-Hernández, Jonnathan Aperador and Jorge Bautista-Ruiz
Metals 2025, 15(3), 272; https://doi.org/10.3390/met15030272 - 1 Mar 2025
Viewed by 289
Abstract
Thin films of ReN were synthesized by DC sputtering at different nitrogen pressures (120, 140, 160, and 180 mTorr) on silicon and Ti6Al4V substrates. The coatings were evaluated for their microstructural and mechanical properties. Additionally, the biocompatibility and electrochemical properties of the films [...] Read more.
Thin films of ReN were synthesized by DC sputtering at different nitrogen pressures (120, 140, 160, and 180 mTorr) on silicon and Ti6Al4V substrates. The coatings were evaluated for their microstructural and mechanical properties. Additionally, the biocompatibility and electrochemical properties of the films were studied using Hanks’ lactate solution at 37 °C. X-ray diffraction (XRD) confirmed the formation of cubic ReN with higher nitrogen content. The optimized nitrogen pressure (180 mTorr) allowed the complete formation of the cubic phase of ReN. Regarding electrochemical behavior, ReN coatings significantly improve corrosion resistance, reducing the corrosion rate as nitrogen content increases, reaching 0.0145 µm/year at 180 mTorr. Regarding mechanical properties, the deposited ReN films presented an optimal combination of hardness and elastic modulus for the highest nitrogen contents. Cell viability was assessed by comparing uncoated and coated samples using a live/dead staining assay, demonstrating the biocompatibility of the coatings. To complement this study, scanning electron microscopy (SEM) was used to analyze the protein–coating interaction and cell morphology on the surface of the samples. Full article
(This article belongs to the Special Issue Corrosion Behavior and Surface Engineering of Metallic Materials)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>XRD diffraction patterns for DC sputter-deposited ReN thin films: (<b>a</b>) 120 mTorr; (<b>b</b>) 140 mTorr; (<b>c</b>) 160 mTorr; and (<b>d</b>) 180 mTorr.</p> Full article ">Figure 2
<p>SEM images of rhenium coatings deposited at various nitrogen flows: (<b>a</b>) 180 mTorr, (<b>b</b>) 160 mTorr, (<b>c</b>) 140 mTorr, and (<b>d</b>) 120 mTorr.</p> Full article ">Figure 3
<p>AFM images used for the quantitative evaluation of the surface morphology of coatings deposited at different nitrogen pressures.</p> Full article ">Figure 3 Cont.
<p>AFM images used for the quantitative evaluation of the surface morphology of coatings deposited at different nitrogen pressures.</p> Full article ">Figure 4
<p>Potentiodynamic polarization curves for the uncoated and ReN-coated Ti6Al4V substrate at different nitrogen flows.</p> Full article ">Figure 5
<p>Nanoindentation test results for ReN films as a function of nitrogen variation.</p> Full article ">Figure 6
<p>Graph of the elastic modulus and hardness of ReN coatings as a function of the nitrogen flow used in the synthesis.</p> Full article ">Figure 7
<p>Cell viability results on Ti6Al4V substrates and ReN coatings as a function of nitrogen flow in the synthesis process.</p> Full article ">Figure 8
<p>SEM micrographs showing cell growth. (<b>a</b>) Ti6Al4V substrate, (<b>b</b>) ReN coating at 120 mTorr, (<b>c</b>) ReN coating at 140 mTorr, (<b>d</b>) ReN coating at 160 mTorr, and (<b>e</b>) ReN coating at 180 mTorr with cells.</p> Full article ">
18 pages, 25726 KiB  
Article
Effect of Grain Size on Mechanical Properties and Deformation Mechanism of Nano-Polycrystalline Pure Ti Simulated by Molecular Dynamics
by Xiao Zhang, Adam Ibrahem Abdalrsoul Alduma, Faqi Zhan, Wei Zhang, Junqiang Ren and Xuefeng Lu
Metals 2025, 15(3), 271; https://doi.org/10.3390/met15030271 - 1 Mar 2025
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Abstract
Nano- and microscale titanium and its alloys have potential applications in semiconductor-based micro-electromechanical systems due to their excellent mechanical properties. The uniaxial tensile mechanical properties and deformation mechanism of polycrystalline pure Ti with five different grain sizes measuring 6.74–19.69 nm were studied via [...] Read more.
Nano- and microscale titanium and its alloys have potential applications in semiconductor-based micro-electromechanical systems due to their excellent mechanical properties. The uniaxial tensile mechanical properties and deformation mechanism of polycrystalline pure Ti with five different grain sizes measuring 6.74–19.69 nm were studied via molecular dynamics simulation using the embedded-atom potential function method. The Hall–Petch relationships and the critical grain size of the polycrystalline pure Ti are given. The dislocation migration of grain boundaries is the main deformation mechanism when the grain size exceeds 16.61 nm, which causes a direct Hall–Petch effect. When grain sizes are smaller than 16.61 nm, grain boundary sliding is the preferred deformation mechanism, which causes an inverse Hall–Petch effect. The polycrystalline pure Ti shows the highest tensile strength and average flow stress of 2.70 GPa and 2.15 GPa, respectively, at the 16.61 nm grain size, which is the critical grain size in the Hall–Petch relationships. The polycrystalline Ti is at its highest strength when its grain size ranges from 16 to 17 nm. The current research provides a theoretical basis for the use of pure titanium in emerging technologies at the nanoscale. Full article
(This article belongs to the Special Issue Microstructure, Crystallography, and Mechanical Properties of Metallic Materials)
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<p>The polycrystalline pure Ti models. (<b>a</b>) The atomic model. (<b>b</b>) The structural model, where the green atoms refer to pure Ti hexagonal close-packed (HCP) phase inside grains and the blue-black to GBs atoms. (<b>c</b>) The structural dyeing model. (<b>d</b>–<b>h</b>) The structural models of the different grain sizes, where the red atoms refer to pure Ti HCP phase inside grains and the whitish-gray to GBs atoms.</p> Full article ">Figure 2
<p>The characteristic curves of polycrystalline pure titanium model of different grain sizes under the deformation process. (<b>a</b>) The stress–strain curve. (<b>b</b>) The relationship curve between the grain size and tensile strength, when <math display="inline"><semantics> <mrow> <msup> <mrow> <mi>d</mi> </mrow> <mrow> <mo>−</mo> <mstyle scriptlevel="0" displaystyle="true"> <mfrac> <mrow> <mn>1</mn> </mrow> <mrow> <mn>2</mn> </mrow> </mfrac> </mstyle> </mrow> </msup> </mrow> </semantics></math> is less than 0.25 and the grain size and yield strength exhibit a Hall–Petch relationship; otherwise, they exhibit an inverse Hall–Petch relationship. (<b>c</b>) The relationship curve between the grain size and average flow stress. (<b>d</b>) The relationship curve between the pair separation distance and radial distribution function. The colors of the symbols and curves in (<b>b</b>–<b>d</b>) are the same as those in <a href="#metals-15-00271-f002" class="html-fig">Figure 2</a>a.</p> Full article ">Figure 3
<p>The structural changes in GB deformation in polycrystalline pure Ti with average grain sizes of 19.69 nm (<b>a</b>–<b>e</b>), 16.61 nm (<b>a1</b>–<b>e1</b>), and 6.74 nm (<b>a2</b>–<b>e2</b>) after exerting uniaxial tensile stress during the strain increase process. The red color denotes the atoms in the perfect structure, the green color refers to stacking faults (SFs) in atoms, the scattered blue atoms denote body-centered cubic (BCC) atoms, and the white-gray to GBs atoms; Figure (<b>f</b>) illustrates the rotation and stretching of the grain within the blue box in Figure (<b>a2</b>). The yellow dotted line is the part of the GB where migration occurs.</p> Full article ">Figure 4
<p>The structural changes in GB deformation and its corresponding potential energy changes in polycrystalline pure Ti after exerting uniaxial tensile under strains of 0.0, 0.15, and 0.30: (<b>a</b>–<b>c</b>) and (<b>d</b>–<b>f</b>). The structural changes in GB deformation of average grain sizes 19.69 nm and 6.74 nm, respectively; the red is the structure of the perfect atoms, whitish-gray are GBs, and blue lines are SFs: (<b>a1</b>–<b>c1</b>) and (<b>d1</b>–<b>f1</b>). The corresponding potential energy changes in average grain sizes 19.69 nm and 6.74 nm, respectively. (The color grading denotes the energy levels of the perfect structure and GBs atoms: blue for low-energy and red for high-energy atoms.)</p> Full article ">Figure 5
<p>Curve of the change in the numbers of atoms in various phase structures of polycrystalline pure Ti of different grain sizes under uniaxial tensile deformation. Here, (<b>a</b>) is BCC phase, (<b>b</b>) is FCC phase, (<b>c</b>) is HCP phase, and (<b>d</b>) is the number of Other-phase atoms.</p> Full article ">Figure 6
<p>The clusters of FCC atom phase structures in polycrystalline pure Ti of different grain sizes under strains of 0.0, 0.1, 0.2, and 0.3. Here, (<b>a</b>–<b>d</b>) is 19.69 nm, (<b>e</b>–<b>h</b>) is 16.61 nm, and (<b>i</b>–<b>l</b>) is 6.74 nm, respectively. (The green atoms are the SFs formation, where the other structures contents were ignored).</p> Full article ">Figure 7
<p>The variation in the dislocation lengths in polycrystalline pure Ti with grain sizes of (<b>a</b>) 19.69 nm, (<b>b</b>) 16.61 nm, (<b>c</b>) 13.18 nm, (<b>d</b>) 9.71 nm, and (<b>e</b>) 6.74 nm.</p> Full article ">Figure 8
<p>The evolution of the dislocation network morphology of polycrystalline pure Ti with different grain sizes under strains of 5%, 10%, and 15%, respectively. (<b>a</b>–<b>c</b>) 19.69 nm, (<b>d</b>–<b>f</b>) 16.61 nm, and (<b>g</b>–<b>i</b>) 6.74 nm. The green atoms are SFs atoms and the red lines are dislocations.</p> Full article ">Figure 9
<p>The stress concentration distribution diagram of polycrystalline pure Ti with different grain sizes under the strain of 5%, 10%, and 15%, respectively. (<b>a</b>–<b>c</b>) 19.69 nm, (<b>d</b>–<b>f</b>) 16.61 nm, and (<b>g</b>–<b>i</b>) 6.74 nm. (The blue color represents the stress concentration of atoms in the perfect structure, and the scaled red represents the stress concentration at GBs).</p> Full article ">Figure 10
<p>The structure scheme of the stacking faults evolution process and interactions with dislocation movement. Where (<b>a</b>–<b>d</b>) is the stacking faults growth and transformation process, (<b>a1</b>–<b>d1</b>) is a side view of the process (the HCP phase structure atoms were deleted), and (<b>a2</b>–<b>d2</b>) is the dislocation of the slip movements and interactions (the red line represents Other, the green line represents 1/3&lt;1-210&gt; dislocation, and the blue line represents &lt;0001&gt; dislocation.).</p> Full article ">
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