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Metals
Volume 14
Issue 4
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Metals, Volume 14, Issue 4 (April 2024) – 117 articles

Cover Story (view full-size image): Porous tungsten gradient materials with ordered gradient variations in pore size are valuable in the field of vacuum electronic devices. This work combines tape casting and dealloying methods to achieve the integrated preparation of porous tungsten gradient materials with a wide range of controllable porosity. The study focuses on the phase composition and microstructure evolution during the preparation of porous tungsten gradient materials. This work provides a design concept for the integrated preparation of porous metal gradient materials. View this paper
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19 pages, 8512 KiB  
Article
Semi-Analytical Solution Model for Bending Deformation of T-Shaped Aviation Aluminium Alloy Components under Residual Stress
by Ning Li, Shouhua Yi, Wanyi Tian and Qun Wang
Metals 2024, 14(4), 486; https://doi.org/10.3390/met14040486 - 22 Apr 2024
Viewed by 804
Abstract
Structures composed of aviation aluminium alloys, characterized by their limited rigidity and thin-walled configurations, frequently exhibit deformation after processing. This paper presents an investigation into T-shaped components fabricated from pre-stretched 7075-T7451 aviation aluminium alloy sheets, examining the effects of residual stress and the [...] Read more.
Structures composed of aviation aluminium alloys, characterized by their limited rigidity and thin-walled configurations, frequently exhibit deformation after processing. This paper presents an investigation into T-shaped components fabricated from pre-stretched 7075-T7451 aviation aluminium alloy sheets, examining the effects of residual stress and the geometrical parameters of T-shaped components on their deformational behavior. A semi-analytical model, developed to elucidate the bending deformation of T-shaped components subjected to residual stress, was validated through finite element analysis and empirical cutting experiments. The experimental results revealed that the bending deformation deflection of the T-shaped specimen was 0.920 mm, deviating by a mere 0.011 mm from the prediction provided by the semi-analytical model, resulting in an inconsequential error margin of 1.2%. This concordance underscores the precision and accuracy of the semi-analytical model specifically designed for T-shaped components. Moreover, the model’s simplicity and ease of application make it an effective tool for predicting the bending deformation of thin-walled T-shaped components under a range of residual stresses and dimensional variations, thereby demonstrating its significant utility in engineering applications. Full article
(This article belongs to the Special Issue Modeling, Characterization and Controlling of Residual Stress in Metal Components)
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Figure 1

Figure 1
<p>Thin plate model.</p> Full article ">Figure 2
<p>Rectangular thin plate loaded with bending moments on two opposite edges.</p> Full article ">Figure 3
<p>Stress model of t-shaped structural component.</p> Full article ">Figure 4
<p>Model of rectangular cantilever thin plate.</p> Full article ">Figure 5
<p>Finite element model of cantilever plate.</p> Full article ">Figure 6
<p>Deflection of rectangular cantilever thin plates under different <math display="inline"><semantics> <mrow> <msub> <mrow> <mi>M</mi> </mrow> <mrow> <mi>y</mi> </mrow> </msub> </mrow> </semantics></math> values.</p> Full article ">Figure 7
<p>Deformation of cantilever plate under <math display="inline"><semantics> <mrow> <msub> <mrow> <mi>M</mi> </mrow> <mrow> <mi>y</mi> </mrow> </msub> <mo>=</mo> <mi>μ</mi> <msub> <mrow> <mi>M</mi> </mrow> <mrow> <mi>x</mi> </mrow> </msub> </mrow> </semantics></math> conditions.</p> Full article ">Figure 8
<p>Deformation Deflection Values of Cantilever Plates with Varying Width Dimensions.</p> Full article ">Figure 9
<p>Relationship between plate dimension ratio (<math display="inline"><semantics> <mrow> <mi>W</mi> <mo>/</mo> <mi>L</mi> </mrow> </semantics></math>) and deflection.</p> Full article ">Figure 10
<p>Simulation results for model verification.</p> Full article ">Figure 11
<p>T-shaped component.</p> Full article ">Figure 12
<p>Initial Residual Stress Detection Diagram.</p> Full article ">Figure 13
<p>Initial Residual Stress of the Specimen.</p> Full article ">Figure 14
<p>Bending Deformation of T-shaped Component.</p> Full article ">Figure 15
<p>Measurement of Machining Deformation of T-shaped Component.</p> Full article ">
11 pages, 49133 KiB  
Article
Investigation on the Duration of Action of Mg3N2 as a Grain Refiner for AZ80 Alloy
by Thomas Hösele, Ernst Neunteufl and Jiehua Li
Metals 2024, 14(4), 485; https://doi.org/10.3390/met14040485 - 22 Apr 2024
Viewed by 816
Abstract
In magnesium alloys with aluminum as an alloying component, zirconium loses its grain refinement effect as a grain refiner. Instead of zirconium, Mg3N2 can be used, and promising results have already been obtained. However, the duration of action of Mg [...] Read more.
In magnesium alloys with aluminum as an alloying component, zirconium loses its grain refinement effect as a grain refiner. Instead of zirconium, Mg3N2 can be used, and promising results have already been obtained. However, the duration of action of Mg3N2 has not been elucidated yet. The aim of this work is therefore to determine the grain size of the AZ80 alloy as a function of the duration of action of Mg3N2 and thus the economically reasonable duration of use. It was found that the Mg3N2 reaches its full effect from 30 min after a complete remelting and does not lose this grain refinement effect even after 90 min. It thus proves to be a stable and reliable grain refiner. A grain size of 146.3 ± 10.3 µm was achieved. Furthermore, a minimum tensile strength of 205 MPa with a break elongation of 5.99% was achieved. Full article
(This article belongs to the Section Metal Casting, Forming and Heat Treatment)
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Figure 1

Figure 1
<p>Thermal analysis of the reference specimen (at the beginning of the test series) (<b>a</b>,<b>b</b>) and the specimen at 90 min after the addition of Mg<sub>3</sub>N<sub>2</sub> (at the end of the test series) (<b>c</b>,<b>d</b>), respectively. <a href="#metals-14-00485-f001" class="html-fig">Figure 1</a>b,d is enlarged from <a href="#metals-14-00485-f001" class="html-fig">Figure 1</a>a,c. The red line is taken from the temperature, while the green line is taken from the first derivative of the temperature.</p> Full article ">Figure 2
<p>Thermal analysis of the reference specimen (at the beginning of the test series) (<b>a</b>,<b>b</b>) and the specimen at 90 min after the addition of Mg<sub>3</sub>N<sub>2</sub> (at the end of the test series) (<b>c</b>,<b>d</b>), respectively. <a href="#metals-14-00485-f002" class="html-fig">Figure 2</a>b,d is enlarged from <a href="#metals-14-00485-f002" class="html-fig">Figure 2</a>a,c. The red line is taken from the temperature, while the green line is taken from the first derivative of the temperature.</p> Full article ">Figure 3
<p>Microstructure evolution using the TP1 test after the addition of Mg<sub>3</sub>N<sub>2</sub> as a function of the duration of action: (<b>a</b>) 0 min, (<b>b</b>) 10 min, (<b>c</b>) 20 min, (<b>d</b>) 30 min, (<b>e</b>) 40 min, (<b>f</b>) 50 min, (<b>g</b>) 60 min, and (<b>h</b>) 90 min. All samples were etched and taken in the bright field image mode. The intermetallic phase along grain boundaries is indexed as Mg<sub>17</sub>Al<sub>12</sub>. Mg<sub>3</sub>N<sub>2</sub> particles were not observed both within the Mg matrix and along grain boundaries due to the limited resolution of optical microscopy.</p> Full article ">Figure 3 Cont.
<p>Microstructure evolution using the TP1 test after the addition of Mg<sub>3</sub>N<sub>2</sub> as a function of the duration of action: (<b>a</b>) 0 min, (<b>b</b>) 10 min, (<b>c</b>) 20 min, (<b>d</b>) 30 min, (<b>e</b>) 40 min, (<b>f</b>) 50 min, (<b>g</b>) 60 min, and (<b>h</b>) 90 min. All samples were etched and taken in the bright field image mode. The intermetallic phase along grain boundaries is indexed as Mg<sub>17</sub>Al<sub>12</sub>. Mg<sub>3</sub>N<sub>2</sub> particles were not observed both within the Mg matrix and along grain boundaries due to the limited resolution of optical microscopy.</p> Full article ">Figure 4
<p>Grain size evolution as a function of duration of action.</p> Full article ">Figure 5
<p>SEM secondary electron (SE) images (<b>a</b>,<b>c</b>) and backscattering (BS) images (<b>b</b>,<b>d</b>) of the sample at 30 min after the addition of the Mg<sub>3</sub>N<sub>2</sub> nanoparticles. (<b>a</b>,<b>b</b>) Low magnification images. Mg<sub>17</sub>Al<sub>12</sub> and Al<sub>6</sub>Mn phases are indexed. (<b>c</b>,<b>d</b>) High magnification images. Mg<sub>3</sub>N<sub>2</sub> was observed together with the eutectic phase (Mg<sub>17</sub>Al<sub>12</sub>).</p> Full article ">Figure 6
<p>SEM EBSD results. (<b>a</b>) Inverse pole figure (IPF), (<b>b</b>) phase map, (<b>c</b>) size distribution of α-Mg grains, (<b>d</b>) size distribution of Mg<sub>3</sub>N<sub>2</sub> particles. Mg<sub>3</sub>N<sub>2</sub> was indexed by EBSD (blue in (<b>b</b>)).</p> Full article ">Figure 7
<p>Mechanical properties of four samples (0 min (<b>a</b>), 48 min (<b>b</b>), 76 min (<b>c</b>), 94 min (<b>d</b>)). Tensile strength, yield strength, and elongation at break are marked. A tensile strength of 205 MPa and a maximum elongation at a break of 5.99% could be achieved after 76 min.</p> Full article ">Figure 8
<p>SEM SE images of the fracture surfaces of four samples (0 min (<b>a</b>,<b>b</b>), 48 min (<b>c</b>,<b>d</b>), 76 min (<b>e</b>,<b>f</b>), 94 min (<b>g</b>,<b>h</b>)). (<b>b</b>,<b>d</b>,<b>f</b>,<b>h</b>) are enlarged from (<b>a</b>,<b>c</b>,<b>e</b>,<b>g</b>) to show more details. All samples show a ductile fracture surface.</p> Full article ">Figure 8 Cont.
<p>SEM SE images of the fracture surfaces of four samples (0 min (<b>a</b>,<b>b</b>), 48 min (<b>c</b>,<b>d</b>), 76 min (<b>e</b>,<b>f</b>), 94 min (<b>g</b>,<b>h</b>)). (<b>b</b>,<b>d</b>,<b>f</b>,<b>h</b>) are enlarged from (<b>a</b>,<b>c</b>,<b>e</b>,<b>g</b>) to show more details. All samples show a ductile fracture surface.</p> Full article ">
21 pages, 18973 KiB  
Article
Evaluation of Melting Efficiency in Cold Wire Gas Metal Arc Welding Using 1020 Steel as Substrate
by R. A. Ribeiro, P. D. C. Assunção and A. P. Gerlich
Metals 2024, 14(4), 484; https://doi.org/10.3390/met14040484 - 21 Apr 2024
Viewed by 1072
Abstract
A key welding parameter to quantify in the welding process is the ratio of the heat required to melt the weld metal versus the total energy delivered to the weld, and this is referred to as the melting efficiency. It is generally expected [...] Read more.
A key welding parameter to quantify in the welding process is the ratio of the heat required to melt the weld metal versus the total energy delivered to the weld, and this is referred to as the melting efficiency. It is generally expected that the productivity of the welding process is linked to this melting efficiency, with more productive processes typically having higher melting efficiency. A comparison is made between the melting efficiency in standard gas metal arc welding (GMAW) and cold wire gas metal arc welding (CW-GMAW) for the three primary transfer modes: short-circuit, globular, and spray regime. CW-GMAW specimens presented higher melting efficiency than GMAW for all transfer modes. Moreover, an increase in plate thickness in the spray transfer regime caused the melting efficiency to increase, contrary to what is expected. Full article
(This article belongs to the Special Issue Editorial Board Members’ Collection Series: Additive Manufacturing Technology)
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Figure 1
<p>Calorimeter used to perform the experiments: (<b>a</b>) schematic drawing; (<b>b</b>) photograph of the experimentation set-up. Numbers indicate (1) welding power source, (2) water inlet, (3) welding torch and cold wire feeder, (4) welding substrate (steel plate), and (5) negative cable (calorimeter and substrate were negative during welding).</p> Full article ">Figure 2
<p>A schematic of a cross-section of a weld bead. This figure is only a schematic representation of a general welding profile, and it was not designed to mirror actual profiles as shown in <a href="#metals-14-00484-f003" class="html-fig">Figure 3</a>.</p> Full article ">Figure 3
<p>Cross-sections of the welds for different metal transfer regimes: (<b>a</b>) short-circuit, (<b>b</b>) globular, and (<b>c</b>) spray. Thin plate (6 mm) and thick plate (9.5 mm), and for both 5 L/min was the calorimeter flow rate employed.</p> Full article ">Figure 4
<p>Melting efficiency in short-circuit transfer mode: (<b>a</b>) standard GMAW; (<b>b</b>) CW-GMAW-60%; and (<b>c</b>) CW-GMAW-10%. In this figure, the conditions were thin plate (6 mm) and thick plate (9.5 mm), and for both, 5 L/min was the calorimeter flow rate employed. The error bars represent 95% confidence intervals.</p> Full article ">Figure 5
<p>Melting efficiency in globular transfer mode: (<b>a</b>) standard GMAW; (<b>b</b>) CW-GMAW-60%; and (<b>c</b>) CW-GMAW-120%. In this figure, the conditions were thin plate (6 mm) and thick plate (9.5 mm), and for both, 5 L/min was the calorimeter flow rate employed. The error bars represent 95% confidence intervals.</p> Full article ">Figure 6
<p>Melting efficiency in spray transfer mode: (<b>a</b>) standard GMAW; (<b>b</b>) CW-GMAW-60%; and (<b>c</b>) CW-GMAW-120%. In this figure, the conditions were thin plate (6 mm) and thick plate (9.5 mm), and for both, 5 L/min was the calorimeter flow rate employed. The error bars represent 95% confidence intervals.</p> Full article ">Figure 7
<p>(<b>a</b>) Rykalin and (<b>b</b>) Christensen numbers for different cold wire feed rates for the short-circuit metal transfer regime. The error bars are 95% confidence intervals.</p> Full article ">Figure 8
<p>(<b>a</b>) Rykalin and (<b>b</b>) Christensen numbers for different cold wire feed rates for the globular metal transfer regime. The error bars are 95% confidence intervals.</p> Full article ">Figure 9
<p>(<b>a</b>) Rykalin and (<b>b</b>) Christensen numbers for different cold wire feed rates for the spray metal transfer regime. The error bars are 95% confidence intervals.</p> Full article ">Figure 10
<p>A comparison between the melting efficiencies calculated by the two approaches discussed: (<b>a</b>) short-circuit, (<b>b</b>) globular, and (<b>c</b>) spray.</p> Full article ">Figure 11
<p>Christensen versus Rykalin number for thin and thick plates: (<b>a</b>) no cold wire feed rates (GMAW), (<b>b</b>) intermediate cold wire feed rates (60%), and (<b>c</b>) high cold wire feed rates (100% and 120%). The <math display="inline"><semantics> <msup> <mi>R</mi> <mn>2</mn> </msup> </semantics></math> values that assess the quality of the regression are shown in <a href="#metals-14-00484-t005" class="html-table">Table 5</a>.</p> Full article ">
20 pages, 10851 KiB  
Article
A Novel Design of a Molten Salt Bath Structure and Its Quenching Effect on Wire Transformation from Austenite to Sorbite
by Jun Li, Bo Wang and Jieyu Zhang
Metals 2024, 14(4), 483; https://doi.org/10.3390/met14040483 - 20 Apr 2024
Viewed by 1060
Abstract
The technology for obtaining sorbite by isothermal quenching of high-temperature molten salt has been used by more and more factories to produce wires with high tensile strength. In this paper, the controlling cap and bottom pipeline of the original salt bath are redesigned. [...] Read more.
The technology for obtaining sorbite by isothermal quenching of high-temperature molten salt has been used by more and more factories to produce wires with high tensile strength. In this paper, the controlling cap and bottom pipeline of the original salt bath are redesigned. The mathematical model previously proposed is used to simulate the redesigned salt bath model, and the flow field is analyzed in detail. The redesigned and original controlling cap are compared in detail by applying third-generation vortex identification technology. Then, by using the inverse heat transfer method, the heat transfer coefficient (HTC) during the boiling heat transfer stage of the wire rod in molten salt is calculated by taking advantage of quenching experimental data, on the basis of which the original model is corrected. Finally, a new salt bath design is proposed, which divides the salt bath into two parts. The first salt bath at 515 °C is used to cool the austenitized wire and complete the initial phase transformation. The second salt bath at 560 °C is used to prevent the transformation from retained austenite to bainite, and to induce its transformation from retained austenite to sorbite. Full article
(This article belongs to the Special Issue Advanced Rolling, Heat Treatment and Electromagnetic Processing Technology of High Performance Metals (Second Edition))
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Figure 1

Figure 1
<p>Salt bath heat treatment process.</p> Full article ">Figure 2
<p>Actual photo of the salt bath.</p> Full article ">Figure 3
<p>Actual production photo of the salt bath.</p> Full article ">Figure 4
<p>The salt bath model for simulation.</p> Full article ">Figure 5
<p>The salt bath mesh.</p> Full article ">Figure 6
<p>Comparison between the original and redesigned salt bath.</p> Full article ">Figure 7
<p>The flow field of the original cap: (<b>a</b>) velocity diagram; (<b>b</b>) vector diagram.</p> Full article ">Figure 8
<p>The streamline of the original salt bath.</p> Full article ">Figure 9
<p>The flow field of the redesigned cap.</p> Full article ">Figure 10
<p>The streamline of the redesigned salt bath.</p> Full article ">Figure 11
<p>The cross-sectional flow field of the redesigned salt bath.</p> Full article ">Figure 12
<p>Comparison of the SEM images between the original and redesigned salt bath: (<b>a</b>) the original salt bath; (<b>b</b>) the redesigned salt bath.</p> Full article ">Figure 13
<p>Comparison of the vortex between the original and redesigned salt bath: (<b>a</b>) the whole salt bath; (<b>b</b>) partial enlargement of (<b>a</b>).</p> Full article ">Figure 14
<p>The position of the vortex in the redesigned salt bath.</p> Full article ">Figure 15
<p>The sample treated in the salt bath.</p> Full article ">Figure 16
<p>The experiments’ temperature measurement results.</p> Full article ">Figure 17
<p>The HTC for the salt temperatures of 515 and 560 °C.</p> Full article ">Figure 18
<p>The cooling curves obtained from the redesigned mathematic model.</p> Full article ">Figure 19
<p>The salt bath is divided into the first and second salt baths.</p> Full article ">Figure 20
<p>Detailed explanation of the first salt bath.</p> Full article ">Figure 21
<p>Detailed explanation of the second salt bath.</p> Full article ">Figure 22
<p>Comparison of cooling curves between the redesigned salt bath, 515 °C salt bath, and 560 °C salt bath.</p> Full article ">
23 pages, 18628 KiB  
Article
TiAl Alloy Fabricated Using Election Beam Selective Melting: Process, Microstructure, and Tensile Performance
by Yu Zhang, Yan Li, Meihui Song, Yanchun Li, Shulin Gong and Bin Zhang
Metals 2024, 14(4), 482; https://doi.org/10.3390/met14040482 - 20 Apr 2024
Cited by 3 | Viewed by 844
Abstract
TiAl alloy is one of the most attractive candidates for a new generation of high-temperature structural materials and has broad application prospects in the aerospace field. As a typical intermetallic material, TiAl is inevitably difficult to process using conventional methods. Election beam selective [...] Read more.
TiAl alloy is one of the most attractive candidates for a new generation of high-temperature structural materials and has broad application prospects in the aerospace field. As a typical intermetallic material, TiAl is inevitably difficult to process using conventional methods. Election beam selective melting (EBSM) is an effective method of addictive manufacturing to prepare TiAl alloy with a complex structure. However, the microstructure of TiAl alloy formed using EBSM often contains defects such as pores, which seriously reduces the mechanical properties of the material. In this work, the effects of EBSM and post-processing procedures on the microstructure and mechanical properties of Ti-48Al-2Cr-2Nb alloy were studied. The results show that the microstructure of Ti-48Al-2Cr-2Nb alloy formed using the EBSM process was dense and composed of equiaxed γ-phase and double-phase regions. A large number of dislocations that formed due to thermal stress were clearly observed inside the Ti-48Al-2Cr-2Nb alloy. When the EBSM process parameters were 13.5 mA, 4.0 m/s, and 40.50 J/mm3, as the current intensity increased, the Al content decreased, the content of α2 phase increased, and the microstructure of the material was coarse. The results of the tensile test fracture morphology indicate that the Ti-48Al-2Cr-2Nb alloy exhibited brittle fracture during tensile deformation, lacking the typical yield deformation of metal materials. As the energy density of the EBSM process increased, the mechanical properties of the Ti-48Al-2Cr-2Nb alloy first increased and then decreased. The samples prepared with an energy density of 34.50~40.50 J/mm3 had excellent mechanical properties, of which the maximum tensile strength and maximum elongation reached 643 MPa and 2.09%, respectively. The phase composition of the Ti-48Al-2Cr-2Nb alloy after hot isostatic pressing (HIP) treatment remained unchanged from the EBSM samples, but there was a slight difference in content. There was an increase in the amount of γ phase and a decrease in B2 phase, accompanied by the generation of a massive γ phase after HIP treatment. Moreover, the number of dislocations inside the material increased. The Ti-48Al-2Cr-2Nb alloy after HIP treatment exhibited obvious plastic deformation characteristics, with a tensile strength of 679 MPa and elongation of 2.5%. A heat treatment of 900 °C/5 h was performed on the Ti-48Al-2Cr-2Nb alloy after HIP. The dislocation density of the Ti-48Al-2Cr-2Nb alloy decreased, and the B2 phase transformed from massive to lamellar. Full article
(This article belongs to the Special Issue Advances in 3D Printing Technologies of Metals—2nd Edition)
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Figure 1
<p>Images of Ti-48Al-2Cr-2Nb powder: (<b>a</b>) particle size distribution; (<b>b</b>) SEM photograph.</p> Full article ">Figure 2
<p>Schematic illustration of the scanning method.</p> Full article ">Figure 3
<p>Tensile test samples: (<b>a</b>) sampling; (<b>b</b>) dimensions (unit: mm); (<b>c</b>) image.</p> Full article ">Figure 4
<p>The effect of the EBSM process parameters on the XRD diffraction pattern of Ti-48Al-2Cr-2Nb alloy.</p> Full article ">Figure 5
<p>The effect of the EBSM process on the pores in the Ti-48Al-2Cr-2Nb alloy: (<b>a</b>) 12.5 mA, 4.5 m/s, and 33.33 J/mm<sup>3</sup>; (<b>b</b>) 11.5 mA, 4.0 m/s, and 34.5 J/mm<sup>3</sup>; (<b>c</b>) 14.5 mA, 4.5 m/s, and 38.67 J/mm<sup>3</sup>; (<b>d</b>) 13.5 mA, 4.0 m/s, and 40.50 J/mm<sup>3</sup>.</p> Full article ">Figure 6
<p>The effect of the EBSM process on the metallographic structure of the Ti-48Al-2Cr-2Nb alloy: (<b>a</b>) 12.5 mA, 4.5 m/s, and 33.33 J/mm<sup>3</sup>; (<b>b</b>) 11.5 mA, 4.0 m/s, and 34.5 J/mm<sup>3</sup>; (<b>c</b>) 14.5 mA, 4.5 m/s, and 38.67 J/mm<sup>3</sup>; (<b>d</b>) 13.5 mA, 4.0 m/s, and 40.50 J/mm<sup>3</sup>.</p> Full article ">Figure 7
<p>TEM microstructure of Ti-48Al-2Cr-2Nb alloy with process parameters of 13.5 mA, 4.0 m/s, and 40.50 J/mm<sup>3</sup>: (<b>a</b>) lamellar structure; (<b>b</b>) distribution of α<sub>2</sub> phases; (<b>c</b>) dislocations.</p> Full article ">Figure 8
<p>Equiaxed γ phases in Ti−48Al−2Cr−2Nb alloy with process parameters of 13.5 mA, 4.0 m/s, and 40.50 J/mm<sup>3</sup>: (<b>a</b>) dark field images; (<b>b</b>) EDS analysis of the lineation zone.</p> Full article ">Figure 9
<p>Effect of EBSM process on element distribution of Ti-48Al-2Cr-2Nb alloy: (<b>a</b>) 12.5 mA, 4.5 m/s, and 33.33 J/mm<sup>3</sup>; (<b>b</b>) 11.5 mA, 4.0 m/s, and 34.5 J/mm<sup>3</sup>; (<b>c</b>) 14.5 mA, 4.5 m/s, and 38.67 J/mm<sup>3</sup>; (<b>d</b>) 13.5 mA, 4.0 m/s, and 40.50 J/mm<sup>3</sup>.</p> Full article ">Figure 9 Cont.
<p>Effect of EBSM process on element distribution of Ti-48Al-2Cr-2Nb alloy: (<b>a</b>) 12.5 mA, 4.5 m/s, and 33.33 J/mm<sup>3</sup>; (<b>b</b>) 11.5 mA, 4.0 m/s, and 34.5 J/mm<sup>3</sup>; (<b>c</b>) 14.5 mA, 4.5 m/s, and 38.67 J/mm<sup>3</sup>; (<b>d</b>) 13.5 mA, 4.0 m/s, and 40.50 J/mm<sup>3</sup>.</p> Full article ">Figure 10
<p>Comparison of room-temperature tensile properties of Ti-48Al-2Cr-2Nb alloy formed using EBSM process with different input energy densities: (1) 12.5 mA, 4.5 m/s, and 33.33 J/mm<sup>3</sup>; (2) 11.5 mA, 4.0 m/s, and 34.5 J/mm<sup>3</sup>; (3) 14.5 mA, 4.5 m/s, and 38.67 J/mm<sup>3</sup>; (4) 13.5 mA, 4.0 m/s, and 40.50 J/mm<sup>3</sup>.</p> Full article ">Figure 11
<p>Tensile fracture morphology of Ti-48Al-2Cr-2Nb alloy: (<b>a</b>) 12.5 mA, 4.5 m/s, and 33.33 J/mm<sup>3</sup>; (<b>b</b>) 11.5 mA, 4.0 m/s, and 34.5 J/mm<sup>3</sup>; (<b>c</b>) 14.5 mA, 4.5 m/s, and 38.67 J/mm<sup>3</sup>; (<b>d</b>) 13.5 mA, 4.0 m/s, and 40.50 J/mm<sup>3</sup>.</p> Full article ">Figure 11 Cont.
<p>Tensile fracture morphology of Ti-48Al-2Cr-2Nb alloy: (<b>a</b>) 12.5 mA, 4.5 m/s, and 33.33 J/mm<sup>3</sup>; (<b>b</b>) 11.5 mA, 4.0 m/s, and 34.5 J/mm<sup>3</sup>; (<b>c</b>) 14.5 mA, 4.5 m/s, and 38.67 J/mm<sup>3</sup>; (<b>d</b>) 13.5 mA, 4.0 m/s, and 40.50 J/mm<sup>3</sup>.</p> Full article ">Figure 12
<p>Metallographic structure of Ti-48Al-2Cr-2Nb alloy formed using EBSM in different directions: (<b>a</b>) transverse, low magnification; (<b>b</b>) transverse, high magnification; (<b>c</b>) longitudinal, low magnification; (<b>d</b>) longitudinal, high magnification.</p> Full article ">Figure 13
<p>Analysis of grain sizes in different directions of Ti-48Al-2Cr-2Nb alloy: (<b>a</b>) transverse direction; (<b>b</b>) longitudinal direction.</p> Full article ">Figure 14
<p>Tensile properties of Ti-48Al-2Cr-2Nb alloy formed using EBSM processes in different directions.</p> Full article ">Figure 15
<p>Microstructure of Ti-48Al-2Cr-2Nb alloy before and after hot isostatic pressing treatment: (<b>a</b>) transverse; (<b>b</b>) longitudinal.</p> Full article ">Figure 16
<p>The effect of post-processing on the microstructure of Ti-48Al-2Cr-2Nb alloy: (<b>a</b>) EBSM; (<b>b</b>) HT; (<b>c</b>) HIP; (<b>d</b>) HIP + HT.</p> Full article ">Figure 17
<p>Distribution of elements inside Ti-48Al-2Cr-2Nb alloy after different post-processing methods: (<b>a</b>) HT; (<b>b</b>) HIP + HT.</p> Full article ">Figure 18
<p>EBSD analysis of Ti-48Al-2Cr-2Nb alloy before and after post-processing (yellow is γ, red is α<sub>2</sub>): (<b>a</b>) EBSM; (<b>b</b>) HT; (<b>c</b>) HIP; (<b>d</b>) HIP + HT.</p> Full article ">Figure 19
<p>XRD analysis of phase structure in Ti-48Al-2Cr-2Nb alloy before and after HIP process.</p> Full article ">Figure 20
<p>TEM micrographs of the Ti-48Al-2Cr-2Nb alloy after heat treatment process: (<b>a</b>) double-phase microstructure; (<b>b</b>) orientation relationship of B2 phases and γ phases; (<b>c</b>) γ phases and the diffraction spots; (<b>d</b>) dislocations.</p> Full article ">Figure 20 Cont.
<p>TEM micrographs of the Ti-48Al-2Cr-2Nb alloy after heat treatment process: (<b>a</b>) double-phase microstructure; (<b>b</b>) orientation relationship of B2 phases and γ phases; (<b>c</b>) γ phases and the diffraction spots; (<b>d</b>) dislocations.</p> Full article ">Figure 21
<p>TEM micrographs of Ti-48Al-2Cr-2Nb alloy after hot isostatic pressing treatment process: (<b>a</b>) dual-phase microstructure; (<b>b</b>) B2 phase and γ phase; (<b>c</b>) γ phase twins and their diffraction spots; (<b>d</b>) dislocations.</p> Full article ">Figure 21 Cont.
<p>TEM micrographs of Ti-48Al-2Cr-2Nb alloy after hot isostatic pressing treatment process: (<b>a</b>) dual-phase microstructure; (<b>b</b>) B2 phase and γ phase; (<b>c</b>) γ phase twins and their diffraction spots; (<b>d</b>) dislocations.</p> Full article ">Figure 22
<p>TEM micrographs of Ti-48Al-2Cr-2Nb alloy after HIP + HT process: (<b>a</b>) dual-phase microstructure; (<b>b</b>) diffraction spots of γ phase; (<b>c</b>) diffraction spots of B2 phase.</p> Full article ">Figure 23
<p>Tensile curves of Ti-48Al-2Cr-2Nb alloy before and after post-processing.</p> Full article ">
12 pages, 7617 KiB  
Article
Study on the Optimization of Investment Casting Process of Exhaust Elbow for High-Power Engine
by Shiyu Xie, Zhaozhao Lv and Shengquan Dong
Metals 2024, 14(4), 481; https://doi.org/10.3390/met14040481 - 20 Apr 2024
Viewed by 908
Abstract
The high-power engine exhaust elbow has a complex construction, which makes it susceptible to casting flaws that could negatively impact its functionality. Therefore, the investment casting scheme was established and optimized in this study in order to cast structurally complete exhaust elbows for [...] Read more.
The high-power engine exhaust elbow has a complex construction, which makes it susceptible to casting flaws that could negatively impact its functionality. Therefore, the investment casting scheme was established and optimized in this study in order to cast structurally complete exhaust elbows for high-horsepower engines. ProCAST software was used to simulate and optimize the casting and solidification processes. The optimal process parameters were determined as follows: pouring temperature of 1650 °C, pouring speed of 1.5 kg/s, and shell preheating temperature of 1050 °C. The optimization of the primary parameters of the casting process, along with the results of dimensional accuracy analysis, shape and positional deviation, and defect detection, were validated through testing. The results indicated that the optimized castings had no casting defects and complied with the design specifications. Full article
(This article belongs to the Special Issue Casting Alloy Design and Characterization)
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<p>Three-dimensional diagram of the exhaust elbow and its cross-section: (<b>a</b>) three-dimensional drawing of exhaust elbow; (<b>b</b>) side view; (<b>c</b>) front view.</p> Full article ">Figure 2
<p>Schematic diagram of the pouring system model.</p> Full article ">Figure 3
<p>Face and body meshing of the 3D group tree model of the pouring system: (<b>a</b>) Meshing of castings; (<b>b</b>) Shell face meshing; (<b>c</b>) Mesh division of castings and molded shells.</p> Full article ">Figure 4
<p>Simulation of the filling process of the exhaust elbow: (<b>a</b>) 25% of the charge; (<b>b</b>) 50% of the charge; (<b>c</b>) 75% of the charge; (<b>d</b>) 98% of the charge.</p> Full article ">Figure 5
<p>Simulation of the solidification process of the exhaust elbow: (<b>a</b>) 25% of the charge; (<b>b</b>) 50% of the charge; (<b>c</b>) 75% of the charge; (<b>d</b>) 100% of the charge.</p> Full article ">Figure 6
<p>Exhaust elbow shrinkage distribution results: (<b>a</b>) Shrinkage distribution with casting system; (<b>b</b>) Distribution of shrinkage in castings without pouring system.</p> Full article ">Figure 7
<p>Thermal cracking tendency diagram and effective stress of exhaust pipe casting: (<b>a</b>) Exhaust pipe bend maximum position HTI value: 0.000163; (<b>b</b>) Maximum position of exhaust elbow HTI: 0.000760; (<b>c</b>) Effective stress at area A: 439.5 MPa; (<b>d</b>) Effective stress at area B: 258 MPa.</p> Full article ">Figure 8
<p>Shrinkage distribution inside the exhaust pipe casting: (<b>a</b>) Shrinkage distribution with casting system; (<b>b</b>) Shrinkage distribution without casting system.</p> Full article ">Figure 9
<p>Dynamic changes in mechanical behavior at specific regions A and B on the exhaust bend with optimal process parameters: (<b>a</b>) Temperature–Time; (<b>b</b>) Fraction Solid Rate–Time; (<b>c</b>) Effective stress–Time; (<b>d</b>) Effective Plastic Stress–time.</p> Full article ">Figure 10
<p>Exhaust Bend Module Shell Making: (<b>a</b>) Surface sanding; (<b>b</b>) Sanding of the transition layer; (<b>c</b>) Backing slurry.</p> Full article ">Figure 11
<p>Comparison of different process parameters: (<b>a</b>) Traditional machining–welding process; (<b>b</b>) Initial casting process parameters castings; (<b>c</b>) Optimized process parameters castings.</p> Full article ">Figure 12
<p>Schematic diagram of industrial CT inspection.</p> Full article ">Figure 13
<p>CT inspection results of exhaust elbow castings under initial pouring process parameters: (<b>a</b>) Shrinkage of bent pipe sections; (<b>b</b>) Dimensions of blowhole defects in straight sections; (<b>c</b>) (Analogue) Bend section corresponding to shrinkage position.</p> Full article ">Figure 14
<p>CT section of exhaust elbow for casting scheme with optimized process parameters: (<b>a</b>) Transverse section of a bend section of pipe; (<b>b</b>) Transverse section of a straight section of pipe; (<b>c</b>) DR result; (<b>d</b>) Transverse section of flange.</p> Full article ">
15 pages, 7097 KiB  
Review
Recent Status of Production, Administration Policies, and Low-Carbon Technology Development of China’s Steel Industry
by Yufeng Qiao and Guang Wang
Metals 2024, 14(4), 480; https://doi.org/10.3390/met14040480 - 20 Apr 2024
Cited by 4 | Viewed by 1813
Abstract
In 2023, China’s crude steel production amount reached 1.019 billion tons, and the energy consumption of China’s steel industry amount reached 561 million tons of coal. China’s steel industry, with its dominant reliance on coal for energy and the primary use of blast [...] Read more.
In 2023, China’s crude steel production amount reached 1.019 billion tons, and the energy consumption of China’s steel industry amount reached 561 million tons of coal. China’s steel industry, with its dominant reliance on coal for energy and the primary use of blast furnaces and converters in production processes, as well as its massive output, has become the main field for achieving China’s “carbon peaking” and “carbon neutrality” goals. Firstly, this article summarizes the current production status of the steel industry and the situation of carbon emissions in the steel industry. Secondly, it discusses the dual-carbon policies based on the national and steel industry levels and outlines the future directions for China’s steel industry. Subsequently, it analyzes the current state of research and application of mature and emerging low-carbon technology in China’s steel industry and details the low-carbon plans of China’s steel companies using the low-carbon technology roadmaps of two representative steel companies as examples. Finally, the article gives policy suggestions for the further carbon reduction of China’s steel industry. The purpose of this paper is to show the efforts and contributions of China’s steel industry to the early realization of its “carbon peaking” and “carbon neutrality” goals. Full article
(This article belongs to the Special Issue Feature Papers in Extractive Metallurgy)
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<p>China’s crude steel yield and the proportion of the world in recent years.</p> Full article ">Figure 2
<p>Carbon emission factors of main raw material.</p> Full article ">Figure 3
<p>Changes in carbon emissions of China’s steel industry in different years.</p> Full article ">Figure 4
<p>Forecast of China’s crude steel production.</p> Full article ">Figure 5
<p>Baowu group’s low-carbon technology roadmap.</p> Full article ">Figure 6
<p>Hebei iron and steel group’s low-carbon technology roadmap.</p> Full article ">
12 pages, 27439 KiB  
Article
Characteristics of Oxide Films on Zr702 and Their Corrosion Performance in Boiling Fluorinated Nitric Acid
by Hangbiao Su, Yaning Li, Yongqing Zhao, Weidong Zeng and Jianping Xu
Metals 2024, 14(4), 479; https://doi.org/10.3390/met14040479 - 19 Apr 2024
Viewed by 700
Abstract
Fluoride ions, which interfere with the oxide formation on zirconium have been over-looked until recently. The effect of fluoride ions on oxide formation and dissolution behaviors in zirconium was investigated in this study. A detailed quantitative characterization of the oxide films formed on [...] Read more.
Fluoride ions, which interfere with the oxide formation on zirconium have been over-looked until recently. The effect of fluoride ions on oxide formation and dissolution behaviors in zirconium was investigated in this study. A detailed quantitative characterization of the oxide films formed on Zr702 immersed in a fluorinated nitric acid solution was performed using X-ray photoelectron spectroscopy, high-resolution transmission electron microscopy, and representative high-angle annular dark-field scanning Transmission Electron Microscope, (TEM). The corrosion performance in a fluorinated nitric acid solution was discussed. The results reveal that the thickness of the oxide films immersed in the fluorinated nitric acid solution was between 42–48 nm, which is much thinner than that of the oxide layer (~98.85 nm thickness) in the F− free HNO3 solution. The oxide film was identified to be a nanocrystalline cluster, comprised of outermost HfO2 and HfF4 layers, sub-outer ZrO2 and ZrF4 layers, and an innermost Zr (F, O)3.6 layer. This fluoride species penetration through the oxide films indicated that the fluoride ions are responsible for the dissolution of the oxide film of Zr702. Full article
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<p>Typical SEM images of Zr702 specimens immersed in 6 M HNO<sub>3</sub> containing (<b>a</b>) 0 ppm; (<b>b</b>) 10 ppm; (<b>c</b>) 50 ppm; (<b>d</b>) 100 ppm; and (<b>e</b>) 200 ppm NaF.</p> Full article ">Figure 2
<p>XRD pattern of the Zr702 sample impregnated in 6 M HNO<sub>3</sub> containing different fluoride ion concentrations.</p> Full article ">Figure 3
<p>(<b>A</b>) HAADF-STEM image of a corroded specimen immersed for 240 h in a boiling 6 M HNO<sub>3</sub> solution with an F-concentration of 10 ppm, and EDX mapping of (<b>B</b>) O, (<b>C</b>) Zr and (<b>D</b>) Hf. (<b>E</b>) the EDX line scan profiles of O, Zr, and Hf.</p> Full article ">Figure 4
<p>TEM characterization of the oxide films formed on the specimens’ surfaces after being subjected to 240 h of corrosion in boiling 6 M HNO<sub>3</sub> with F<sup>−</sup> concentrations of (<b>a</b>) 0 ppm, (<b>b</b>) 10 ppm, (<b>c</b>) 50 ppm, (<b>d</b>) 100 ppm and (<b>e</b>) 200 ppm.</p> Full article ">Figure 5
<p>High resolution XPS spectra of the Zr702 specimens immersed in 6 M HNO<sub>3</sub> (<b>a</b>) Zr 3d; (<b>b</b>) Hf 4f; (<b>c</b>) O1s; and (<b>d</b>) F1s. after they were subjected to corrosion tests.</p> Full article ">Figure 6
<p>HRTEM images of three different zones in oxide films of the specimens immersed in HNO<sub>3</sub> solutions with 200 ppm F<sup>−</sup> for 240 h (<b>A</b>) outer layer 1# zone, (<b>B</b>) middle layer 2# zone, and (<b>C</b>) inner layer 3# zone (<b>D</b>).</p> Full article ">Figure 7
<p>Enlarged view of the high-resolution phases shown in <a href="#metals-14-00479-f006" class="html-fig">Figure 6</a>. (<b>1-A</b>–<b>1-C</b>) are different areas in 1#, (<b>2-A</b>–<b>2-D</b>) are different areas in 2#, (<b>3-A</b>–<b>3-E</b>) are different areas in 3#.</p> Full article ">Figure 8
<p>Corrosion rates (corrosion weight loss) versus exposure time for the Zr702 specimens during each of the five periods.</p> Full article ">
15 pages, 7292 KiB  
Article
A 3D Non-Linear FE Model and Optimization of Cavity Die Sheet Hydroforming Process
by Arun Achuthankutty, Ajith Ramesh and Ratna Kishore Velamati
Metals 2024, 14(4), 478; https://doi.org/10.3390/met14040478 - 19 Apr 2024
Viewed by 791
Abstract
Cryo-rolled aluminum alloys have a much higher strength-to-weight ratio than cold-rolled alloys, which makes them invaluable in the aerospace and automotive industries. However, this strength gain is frequently accompanied by a formability loss. When uniformly applied to the blank surface, hydroforming provides a [...] Read more.
Cryo-rolled aluminum alloys have a much higher strength-to-weight ratio than cold-rolled alloys, which makes them invaluable in the aerospace and automotive industries. However, this strength gain is frequently accompanied by a formability loss. When uniformly applied to the blank surface, hydroforming provides a solution by generating geometries with constant thickness, making it possible to produce complex structures with “near-net dimensions”, which are difficult to achieve with conventional approaches. This study delves into the cavity die sheet hydroforming (CDSHF) process for high-strength cryo-rolled AA5083 aluminum alloy, focusing on two primary research questions. Firstly, we explored the utilization of a nonlinear 3D finite-element (FE) model to understand its impact on the dimensional accuracy of hydroformed components within the CDSHF process. Specifically, we investigated how decreasing fluid pressure and increasing the holding time of peak fluid pressure can be quantitatively assessed. Secondly, we delved into the optimization of process parameters—fluid pressure (FP), blank holding force (BHF), coefficient of friction (CoF), and flange radius (FR)—to achieve dimensional accuracy in hydroformed square cups through the CDSHF process. Our findings reveal that our efforts, such as reducing peak fluid pressure to 22 MPa, implementing a 30 s holding period, and utilizing an unloading path, enhanced component quality. We demonstrated this with a 35 mm deep square cup exhibiting a 16.1 mm corner radius and reduced material thinning to 5.5%. Leveraging a sophisticated nonlinear 3D FE model coupled with response surface methodology (RSM) and multi-objective optimization techniques, we systematically identified the optimal process configurations, accounting for parameter interactions. Our results underscore the quantitative efficacy of these optimization strategies, as the optimized RSM model closely aligns with finite-element (FE) simulation results, predicting a thinning percentage of 5.27 and a corner radius of 18.64 mm. Overall, our study provides valuable insights into enhancing dimensional accuracy and process optimization in CDSHF, with far-reaching implications for advancing metal-forming technologies. Full article
(This article belongs to the Section Metal Casting, Forming and Heat Treatment)
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<p>Various stages of the CDSHF process of a square cup using FE modeling.</p> Full article ">Figure 2
<p>Blank Geometry and Dimensions.</p> Full article ">Figure 3
<p>Mesh convergence study.</p> Full article ">Figure 4
<p>FP path and BHF vs. time.</p> Full article ">Figure 5
<p>Loading and unloading FP path and BHF vs. time.</p> Full article ">Figure 6
<p>Optimization Process Flow Chart.</p> Full article ">Figure 7
<p>Fully formed cup depicting X-X and X-Y direction.</p> Full article ">Figure 8
<p>(<b>a</b>) Comparison of % thinning in X-X direction. (<b>b</b>) Comparison of % thinning in X-Y direction.</p> Full article ">Figure 9
<p>(<b>a</b>) Von Mises stress distribution in fully formed cup at 22 MPa peak FP. (<b>b</b>) Equivalent plastic strain in fully formed cup at 22 MPa peak FP. (<b>c</b>) Von Mises stress distribution in partially formed cup at 18 MPa peak FP.</p> Full article ">Figure 9 Cont.
<p>(<b>a</b>) Von Mises stress distribution in fully formed cup at 22 MPa peak FP. (<b>b</b>) Equivalent plastic strain in fully formed cup at 22 MPa peak FP. (<b>c</b>) Von Mises stress distribution in partially formed cup at 18 MPa peak FP.</p> Full article ">Figure 10
<p>(<b>a</b>) Main effect plot for a thinning percentage. (<b>b</b>) Main effect plot for corner radius.</p> Full article ">Figure 10 Cont.
<p>(<b>a</b>) Main effect plot for a thinning percentage. (<b>b</b>) Main effect plot for corner radius.</p> Full article ">Figure 11
<p>Contour Plots—effect of two parameters on Thinning Percentage. (<b>a</b>) BHF × CoF, (<b>b</b>) FP × CoF, (<b>c</b>) FR × CoF, (<b>d</b>) FP × BHF, (<b>e</b>) FR × BHF, (<b>f</b>) FR × FP.</p> Full article ">Figure 12
<p>Contour Plots—effect of two parameters on Corner Radius. (<b>a</b>) BHF × CoF, (<b>b</b>) FP × CoF, (<b>c</b>) FR × CoF, (<b>d</b>) FP × BHF, (<b>e</b>) FR × BHF, (<b>f</b>) FR × FP.</p> Full article ">
13 pages, 5140 KiB  
Article
Influence of Microstructure on the Mechanical Properties and Polishing Performance of Large Prehardened Plastic Mold Steel Blocks
by Hongxiao Chi, Jihao Liu, Jian Zhou, Dangshen Ma and Jinbo Gu
Metals 2024, 14(4), 477; https://doi.org/10.3390/met14040477 - 19 Apr 2024
Cited by 1 | Viewed by 711
Abstract
The microstructures throughout a 696 × 1360 mm cross-section of an ISO 1.2738 prehardened steel block for a plastic mold were characterized via optical and electron microscopy and electron backscatter diffraction. The hardness, strength, and polishing performance of the steel block were also [...] Read more.
The microstructures throughout a 696 × 1360 mm cross-section of an ISO 1.2738 prehardened steel block for a plastic mold were characterized via optical and electron microscopy and electron backscatter diffraction. The hardness, strength, and polishing performance of the steel block were also tested. The results showed that the microstructure of the steel bloom from the edge to the core consisted of tempered sorbite, tempered bainite, and pearlite microstructures. Abnormal upper bainite and coarse carbides were also found. The bloom sections with hardness values of 37.4 to 39.3 HRC comprised tempered sorbite and bainite. The hardness of the core was approximately 36.5 HRC due to the presence of pearlite. The tensile and yield strengths were the same in the edge and middle areas mainly owing to tempered sorbite. The polishing performance was affected by the microstructure. Tempered sorbite produced the best polishing performance due to its fine and uniform microstructure, whereas that of tempered bainite and pearlite, which contained large carbide particles and mixed phases, was worse. Full article
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<p>Dimensions of the bloom and examined sample.</p> Full article ">Figure 2
<p>OM images of microstructure of at different positions in steel block (50×): (<b>a</b>) edge; (<b>b</b>) 1/4 of diagonal; (<b>c</b>) core.</p> Full article ">Figure 3
<p>OM images of microstructure at different positions in steel block: (<b>a</b>) edge; (<b>b</b>,<b>c</b>) 1/4 of diagonal; (<b>d</b>) core.</p> Full article ">Figure 4
<p>SEM image of microstructure at different positions in steel block: (<b>a</b>) edge; (<b>b</b>,<b>c</b>) 1/4 of diagonal; (<b>d</b>) core.</p> Full article ">Figure 5
<p>TEM images of microstructure at different positions in the steel block: (<b>a</b>) edge; (<b>b</b>,<b>c</b>) 1/4 of diagonal; (<b>d</b>) core.</p> Full article ">Figure 6
<p>Microstructure displayed as unique color maps of a 1.2738 steel bloom (IPF map): (<b>a</b>) tempered sorbite; (<b>b</b>) tempered bainite; (<b>c</b>) tempered bainite + pearlite; (<b>d</b>) inverse pole figure for color code.</p> Full article ">Figure 7
<p>Hardness distribution across the transverse 1/4 section of a 1.2738 steel bloom.</p> Full article ">Figure 8
<p>Tensile properties at different positions of a 1.2738 steel bloom: (<b>a</b>) tensile strength (UTS) and yield strength (YS); (<b>b</b>) reduction (A) and elongation (Z) of the section.</p> Full article ">Figure 9
<p>Polishing performance characterized using a three-dimensional AFM map of different microstructures (50 × 50 μm): (<b>a</b>) tempered sorbite (edge); (<b>b</b>) tempered bainite (diagonal 1/4 position); (<b>c</b>) pearlite + tempered bainite (core); (<b>d</b>) pearlite (core); (<b>e</b>) sorbite + granular carbide (segregation band position).</p> Full article ">Figure 9 Cont.
<p>Polishing performance characterized using a three-dimensional AFM map of different microstructures (50 × 50 μm): (<b>a</b>) tempered sorbite (edge); (<b>b</b>) tempered bainite (diagonal 1/4 position); (<b>c</b>) pearlite + tempered bainite (core); (<b>d</b>) pearlite (core); (<b>e</b>) sorbite + granular carbide (segregation band position).</p> Full article ">Figure 10
<p>Partial five-fold-enlarged view of <a href="#metals-14-00477-f009" class="html-fig">Figure 9</a>: (<b>a</b>) tempered sorbite (edge); (<b>b</b>) tempered bainite (diagonal 1/4 position); (<b>c</b>) pearlite + tempered bainite (core); (<b>d</b>) pearlite (core); (<b>e</b>) sorbite + granular carbide (segregation band position).</p> Full article ">Figure 11
<p>Calculated CCT (<b>a</b>) and TTT (<b>b</b>) curves for the studied steel.</p> Full article ">
16 pages, 5217 KiB  
Article
Numerical Simulation of Temperature Evolution, Solid Phase Transformation, and Residual Stress Distribution during Multi-Pass Welding Process of EH36 Marine Steel
by Pengyu Wen, Jiaji Wang, Zhenbo Jiao, Kuijun Fu, Lili Li and Jing Guo
Metals 2024, 14(4), 476; https://doi.org/10.3390/met14040476 - 19 Apr 2024
Viewed by 910
Abstract
An investigation into the evolution of temperature and stress fields, as well as the phase transformation in marine steel EH36 during multi-pass welding, and their subsequent effects on Charpy impact toughness, remains in great lack. In this study, submerged arc welding (SAW) was [...] Read more.
An investigation into the evolution of temperature and stress fields, as well as the phase transformation in marine steel EH36 during multi-pass welding, and their subsequent effects on Charpy impact toughness, remains in great lack. In this study, submerged arc welding (SAW) was employed to carry out multi-pass welding on EH36 steel plates, followed by the low-temperature toughness test of weldments. Comsol software version 6.2 and finite element analysis are utilized to simulate the evolution of the microstructure, temperature, and residual stress fields throughout the multi-pass welding process. As welding progressed, the heat absorption along the vertical direction was enhanced; in contrast, a decrease is observed in the horizontal direction away from the heat source. This complicated temperature history favors the bainite transformation in the vicinity to the heat source, whereas areas more remote from the weld zone exhibit a higher prevalence of acicular ferrite due to the reduced cooling rate. The concentration of residual stress is predicted to occur at the boundary of the melt pool and at the interface between the weld and the heat-affected zone, with the greatest deformation observed near the fusion line at the top surface of the model. Furthermore, multi-pass welding may alleviate the residual stress, especially when coupled with the formation of acicular ferrite upon cooling, leading to improved low-temperature impact toughness in regions remote from the heat source. These findings offer valuable insights for the design and optimization of multi-pass welding in future applications. Full article
(This article belongs to the Special Issue Offshore Engineering Steel: Welding Performance and Microstructure Analysis)
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<p>The illustration of geometrical model.</p> Full article ">Figure 2
<p>Microstructure morphology at several typical cooling rates (B: Bainite, F: Ferrite; P: Pearlite; M: Martensite).</p> Full article ">Figure 3
<p>Continuous heating and cooling curves in the heat-affected zone of welding (SH-CCT).</p> Full article ">Figure 4
<p>Microstructure evolution as a function of time at different welding energy densities: (<b>a</b>) 100 kJ/cm, and (<b>b</b>) 250 kJ/cm.</p> Full article ">Figure 5
<p>Temperature field distribution during multi-pass welding process: (<b>a</b>) 250 s and (<b>b</b>) 350 s spatial distribution of temperature, and (<b>c</b>) temperature at different locations over time.</p> Full article ">Figure 6
<p>Microstructure evolution at different locations: (<b>a</b>) P1, (<b>b</b>) P2, (<b>c</b>) P3, (<b>d</b>) P4, and (<b>e</b>) P5.</p> Full article ">Figure 7
<p>Fraction of phases at different positions at the end of welding.</p> Full article ">Figure 8
<p>Regulation of residual stress by multi-pass welding: the three-dimension residual stress evolution at multi-pass welding time periods of (<b>a</b>) 100 s, (<b>b</b>) 200 s, (<b>c</b>) 300 s, and (<b>d</b>) 400 s; (<b>e</b>) residual stress distribution in the vertical section of the heat-affected zone; (<b>f</b>) residual stress distribution on special line segments.</p> Full article ">Figure 9
<p>Effect of multi-pass welding on material distortion: material at 400 s: (<b>a</b>) total distortion, (<b>b</b>) X-direction, (<b>c</b>) Y-direction, (<b>d</b>) Z-direction, and (<b>e</b>) distortion at positions P1 to P5 versus time.</p> Full article ">Figure 10
<p>Effect of cooling rate on the impact toughness of multi-pass weldments under different temperatures: (<b>a</b>) impact energy at 253 K (−20 °C); (<b>b</b>) impact energy at 233 K (−40 °C).</p> Full article ">
14 pages, 3998 KiB  
Article
Long-Term Oxidation Studies on Porous Stainless Steel 430L Substrate Relevant to Its Application in Metal-Supported SOFC
by Kai Xu and Liangzhu Zhu
Metals 2024, 14(4), 475; https://doi.org/10.3390/met14040475 - 18 Apr 2024
Cited by 1 | Viewed by 919
Abstract
Metal-supported solid oxide fuel cells (MS-SOFCs) can be used in portable mobile power generators due to their excellent thermal cycling performance, low cost, and strong mechanical strength. The selection and lifetime of the support material are crucial factors that affect the cell’s performance [...] Read more.
Metal-supported solid oxide fuel cells (MS-SOFCs) can be used in portable mobile power generators due to their excellent thermal cycling performance, low cost, and strong mechanical strength. The selection and lifetime of the support material are crucial factors that affect the cell’s performance and long-term stability. The oxidizability of porous 430L stainless steel in a dry air atmosphere at 800 °C was systematically studied and reported for up to 1500 h. The aim was to investigate the lifetime of porous stainless steel as a support skeleton in a symmetric MS-SOFC. The substrates were characterized and analyzed using scanning electron microscopy, energy spectroscopy, and X-ray diffractometry after different periods of oxidation. The analysis indicated that the porous substrate’s surface oxides, under dry air conditions, consisted primarily of Fe2O3 and Cr2O3, with small amounts of Fe3O4 and MnCr2O4 spinel. The long-term oxidation process can be divided into two stages with distinct characteristics. However, the oxide flaking phenomenon occurred after 1500 h of exposure. The estimated service life of the stainless steel was consistent with the experimental results, which were around 1500 h. This estimation was based on the measured weight gain and thickness data. Full article
(This article belongs to the Section Metallic Functional Materials)
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<p>SEM images showing surface (<b>a</b>,<b>b</b>) and cross-sectional structure (<b>c</b>,<b>d</b>) at different magnifications.</p> Full article ">Figure 2
<p>(<b>a</b>) Weight gain as a function of time, (<b>b</b>) square of weight gain as a function of time.</p> Full article ">Figure 3
<p>XRD patterns of 430L substrate oxidized at 800 °C under dry air for different times.</p> Full article ">Figure 4
<p>The surface SEM images of the porous substrate oxidized at 800 °C under dry air for different times. EDS elemental analysis corresponds to the yellow box in the SEM image.</p> Full article ">Figure 5
<p>SEM image of the oxide cracking and flaking for 1500 h oxidation exposure under dry air. EDS elemental analysis corresponds to the position of yellow solid dot in the SEM image.</p> Full article ">Figure 6
<p>Cross-section SEM images, EDS mapping, and line scanning of the oxide scale following oxidation exposure under dry air for various durations. Note the starting position of the line scan is marked by green solid dot, and the line scan direction is not always from the steel matrix.</p> Full article ">Figure 7
<p>The cross-section SEM images of the porous substrate oxidized at 800 °C under dry air for (<b>a</b>) 0 h and (<b>b</b>–<b>f</b>) 1500 h under different magnifications.</p> Full article ">Figure 8
<p>SEM (FIB-lamellae cut) images and EDS mapping of porous 430L alloy oxidized at 800 °C under dry air for 1500 h.</p> Full article ">Figure 9
<p>(<b>a</b>) The average thickness of oxide scale as a function of time, and (<b>b</b>) log–log plot (the dash line is the auxiliary line with a slope of 0.5).</p> Full article ">
14 pages, 7005 KiB  
Article
Comprehensive Unveiling of the Oxidation Resistance and Corrosion Protection of an Oxide Layer Formed on the Gd-Alloyed AZ80 Alloy Surface
by Chunlong Cheng, Gaolin Zhou, Bo Qu, Liang Wang, Abdul Malik and Zheng Chen
Metals 2024, 14(4), 474; https://doi.org/10.3390/met14040474 - 18 Apr 2024
Viewed by 754
Abstract
In our previous work, the effect of Gd alloying on the oxidation resistance of AZ80 alloy was revealed briefly. However, a comprehensive understanding of the oxidation and corrosion resistance of the oxide layer formed on the Gd alloying AZ80 alloy surface needs to [...] Read more.
In our previous work, the effect of Gd alloying on the oxidation resistance of AZ80 alloy was revealed briefly. However, a comprehensive understanding of the oxidation and corrosion resistance of the oxide layer formed on the Gd alloying AZ80 alloy surface needs to be developed. Thus, in this research, the high-temperature oxidation behaviors, oxidation products, and oxide layer characteristics of AZ80, AZ80-0.47Gd, and AZ80-0.75Gd (wt%) alloys were investigated at 420 °C. The corrosion protection of the oxide layer formed on the alloy surface was evaluated. The results showed that Gd alloying eliminated the content of the low melting point phase of β-Mg17Al12 and promoted the generation of a high melting point phase of Al2Gd. Gd2O3 appeared in the oxide layer and facilitated the propagation of homogeneous oxidation as well as densification of the oxide layer. In addition, the firm oxide layer showed characteristics of a blurred boundary with the magnesium matrix. After immersion of the oxide layer containing gadolinium oxide, the products of corrosion were massively nodulated, leading to the passivation of corrosion. This research provides new ideas for magnesium alloy protective layer preparation via a high-temperature oxidation technique. Full article
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<p>SEM microstructures of the AZ80 (<b>a</b>), AZ80-0.47Gd (<b>b</b>), and AZ80-0.75Gd (<b>c</b>) alloys with β-Mg<sub>17</sub>Al<sub>12</sub> and Al<sub>2</sub>Gd phases energy dispersive X-ray spectroscopy analysis as well as normalized X-ray diffraction patterns (<b>d</b>).</p> Full article ">Figure 2
<p>The DSC curves of the AZ80, AZ80-0.47Gd, and AZ80-0.75Gd alloys.</p> Full article ">Figure 3
<p>The weight gain behaviors of alloy oxidation for 1 h (<b>a</b>), 3 h (<b>b</b>), and 4.5 h (<b>c</b>) at 420 °C in air and the n values (<b>d</b>).</p> Full article ">Figure 4
<p>The surface three-dimensional micrographs of the AZ80 (<b>a</b>), AZ80-0.47Gd (<b>b</b>), and AZ80-0.75Gd (<b>c</b>) alloys after 3 h exposure at 420 °C.</p> Full article ">Figure 5
<p>The micrographs of oxide films formed on the AZ80 (<b>a</b>,<b>b</b>), AZ80-0.47Gd (<b>c</b>,<b>d</b>), and AZ80-0.75Gd (<b>e</b>,<b>f</b>) alloys after 4.5 h of exposure at 420 °C.</p> Full article ">Figure 6
<p>The cross-sectional microstructures of the AZ80 (<b>a</b>,<b>b</b>), AZ80-0.47Gd (<b>c</b>), and AZ80-0.75Gd (<b>d</b>,<b>e</b>) alloys after 4.5 h of exposure at 420 °C.</p> Full article ">Figure 7
<p>The normalized XRD patterns (<b>a</b>) and partially enlarged image (<b>b</b>) of the oxide products of the AZ80 and AZ80-0.75Gd alloys.</p> Full article ">Figure 8
<p>The crystal structure of Gd<sub>2</sub>O<sub>3</sub>.</p> Full article ">Figure 9
<p>The potentiodynamic polarization curves of the AZ80-0.75Gd alloy with and without an oxide layer while soaking in 3.5 wt% NaCl solution for 1 h.</p> Full article ">Figure 10
<p>The EIS diagrams of the AZ80-0.75Gd samples after 1 h of soaking in 3.5 wt% NaCl solution: (<b>a</b>) Nyquist plots and (<b>b</b>) Bode plots of impedance modulus and phase angle.</p> Full article ">Figure 11
<p>Equivalent electrical circuits used for the simulation of EIS data in <a href="#metals-14-00474-f010" class="html-fig">Figure 10</a>: (<b>a</b>) sample without oxide film and (<b>b</b>) sample with oxide film.</p> Full article ">Figure 12
<p>Hydrogen evolution curves of AZ80-0.75Gd samples with and without an oxide layer immersed in 3.5 wt% NaCl solution.</p> Full article ">Figure 13
<p>The microscopic corrosion morphologies of the AZ80-0.75Gd samples without (<b>a</b>–<b>c</b>) and with (<b>d</b>–<b>f</b>) an oxide layer after soaking for 12 h in 3.5 wt% NaCl solution.</p> Full article ">Figure 14
<p>Schematic illustration of the mechanism of Gd<sub>2</sub>O<sub>3</sub> involvement in oxide film growth and nodulizing corrosion products.</p> Full article ">
10 pages, 213 KiB  
Editorial
Microstructure and Properties in Metals and Alloys (Volume 2)
by Andrea Di Schino and Claudio Testani
Metals 2024, 14(4), 473; https://doi.org/10.3390/met14040473 - 18 Apr 2024
Viewed by 833
Abstract
Microstructure design is key in targeting the desired material’s properties [...] Full article
(This article belongs to the Topic Microstructure and Properties in Metals and Alloys, 2nd Volume)
12 pages, 2127 KiB  
Article
The Structure and Magnetic Properties of Sm2Fe17Cx Compounds Prepared from Ball-Milled Mixtures of Sm2Fe17 and Carbon Nanotubes or Graphite
by Vladislav A. Mikheev, Igor G. Bordyuzhin, Mikhail V. Gorshenkov, Elena S. Savchenko, Irina V. Dorofievich and Igor V. Shchetinin
Metals 2024, 14(4), 472; https://doi.org/10.3390/met14040472 - 18 Apr 2024
Viewed by 875
Abstract
The processing route of Sm2Fe17 carbides is shorter than that of nitrides, which can potentially be used for cost-effective mid-performance magnets’ production. The magnetic properties of Sm2Fe17Cx compounds can be controlled at the annealing step, [...] Read more.
The processing route of Sm2Fe17 carbides is shorter than that of nitrides, which can potentially be used for cost-effective mid-performance magnets’ production. The magnetic properties of Sm2Fe17Cx compounds can be controlled at the annealing step, which allows them to be used for a variety of applications. In this work, X-ray diffraction (XRD) analysis, Mössbauer spectroscopy, scanning and transmission electron microscopy (SEM, TEM) and vibrating sample magnetometry (VSM) were used for characterization of the structure and magnetic properties of Sm2Fe17Cx compounds. The powder samples were prepared by high-energy ball milling of Sm2Fe17 mixtures with carbon nanotubes (CNT) or graphite with subsequent annealing. The formation of Sm2Fe17Cx compounds after annealing was followed by the formation of α-Fe and amorphous Sm2O3. The hyperfine field values of Fe atoms of all the Sm2Fe17 lattice sites increased by 12% on average after annealing that was caused by carbon diffusion. The coercivity of the samples peaked after annealing at 375 °C. The samples with CNT demonstrated an increase of up to 14% in coercivity and 5% in specific remanence in the range of 250–375 °C annealing temperatures. Full article
(This article belongs to the Special Issue Advances in Metallic Glass and Metallic Glass Composite: Preparation, Structures, Properties and Applications)
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<p>XRD patterns of mixtures of Sm<sub>2</sub>Fe<sub>17</sub> with: (<b>a</b>) carbon nanotubes (CNT) after milling, (<b>b</b>) CNT after milling and annealing (375 °C 1 h), (<b>c</b>) graphite after milling and (<b>d</b>) graphite after milling and annealing (375 °C 1 h).</p> Full article ">Figure 2
<p>SEM images of mixtures (<b>a</b>) with CNT after milling (<b>b</b>) with CNT after milling and annealing (400 °C 1 h) (<b>c</b>) with graphite after milling and (<b>d</b>) with graphite after milling and annealing (375 °C 1 h).</p> Full article ">Figure 3
<p>Mössbauer spectra of mixtures with CNT after: (<b>a</b>) milling, and (<b>b</b>) milling and annealing (375 °C 1 h).</p> Full article ">Figure 4
<p>Phase composition of mixtures (<b>a</b>) with CNT and (<b>b</b>) with graphite.</p> Full article ">Figure 5
<p>Carbon content and coercivity values of mixtures (<b>a</b>) with CNT and (<b>b</b>) with graphite.</p> Full article ">Figure 6
<p>TEM dark-field images of annealed mixtures: (<b>a</b>) with CNT, (<b>b</b>) with graphite.</p> Full article ">
16 pages, 11153 KiB  
Article
First Principles Study of the Effects of Si, P, and S on the ∑5 (210)[001] Grain Boundary of γ-Fe
by Ying Xu, Weigang Cao, Mengzhe Huang and Fucheng Zhang
Metals 2024, 14(4), 471; https://doi.org/10.3390/met14040471 - 17 Apr 2024
Cited by 1 | Viewed by 611
Abstract
Solutes segregating at the grain boundary (GB) have a significant influence on the mechanical and chemical properties of steel. In this study, the segregation effects of Si, P, and S on γ-Fe ∑5 (210)[001] GB were systematically analyzed with solution energy, segregation energy, [...] Read more.
Solutes segregating at the grain boundary (GB) have a significant influence on the mechanical and chemical properties of steel. In this study, the segregation effects of Si, P, and S on γ-Fe ∑5 (210)[001] GB were systematically analyzed with solution energy, segregation energy, and tensile tests by using a first principles calculation. Si, P, and S are preferred to segregate at substitutional sites in the first layer near the GB. The variation in atomic configuration and electron distribution were investigated by the analysis of bond lengths, charge density, charge density difference, and density of states (DOS), which is caused by the atomic size and electronegativity of solute atoms. Through tensile tests, it was found that Si has a strengthening effect on GB, while P and S exhibit embrittlement effects at low concentration. As the concentration of solutes increase, the segregation sites of P are different from the others owing to the tendency to form Fe3P. The exhibited embrittlement effect is mitigated at first and then aggravated. However, in both cases Si and S show aggravating embrittlement effects on GB cohesion, while the effect of Si changes from strengthening to embrittlement. This work provides comprehensive insights into the effects of Si, P, and S, which will be a useful guidance in steel design. Full article
(This article belongs to the Section Computation and Simulation on Metals)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>Schematic diagram of the γ-Fe Σ5(210)[001] GB model with the possible interstitial sites and substitutional sites of impurity atoms (<b>a</b>), front view (<b>b</b>) and top view (<b>c</b>) of the GB model. The purple balls represent Fe atoms, the green balls are interstitial sites, and the red balls indicate the substitutional sites.</p> Full article ">Figure 2
<p>Solution energies of Si, P, and S located at different sites in the γ-Fe Σ5 (210)[001] GB.</p> Full article ">Figure 3
<p>Local atomic configuration of pure GB (<b>a</b>), Si-doped GB (<b>b</b>), P-doped GB (<b>c</b>), and S-doped GB (<b>d</b>) after structure relaxation. The atomic distance is Å. The purple balls are Fe atoms, the green ball presents Si atom, the orange ball is P atom, and the blue ball is S atom.</p> Full article ">Figure 4
<p>Calculated charge density in the (001) plane for cases of the pure GB (<b>a</b>), Si-doped GB (<b>b</b>), P-doped GB (<b>c</b>), and S-doped GB (<b>d</b>). The unit is e/a.u.<sup>3</sup>. The color change from blue to red indicates the increase of charge density.</p> Full article ">Figure 5
<p>The planar projection of electron density differences for the (001) plane in the pure GB (<b>a</b>), Si-doped GB (<b>b</b>), P-doped GB (<b>c</b>), and S-doped GB (<b>d</b>). The unit is e/a.u.<sup>3</sup>. The red area means a charge depletion, while the blue area represents a charge accumulation.</p> Full article ">Figure 6
<p>The DOS of pure GB (<b>a</b>), Si-doped GB (<b>b</b>), P-doped GB (<b>c</b>), and S-doped GB (<b>d</b>).</p> Full article ">Figure 7
<p>The tensile strain–stress curves of pure GB, Si-doped GB, P-doped GB, and S-doped GB.</p> Full article ">Figure 8
<p>Local atomic configuration of pure GB (<b>a</b>), Si-doped GB (<b>b</b>), P-doped GB (<b>c</b>), and S-doped GB (<b>d</b>) at 30% strain-induced variations.</p> Full article ">Figure 9
<p>Charge density distribution in the (001) planes of the pure GB (<b>a</b>), Si-doped GB (<b>b</b>), P-doped GB (<b>c</b>), and S-doped GB (<b>d</b>) at 0%, 5%, 10%, 15%, and 20% strain-induced variations.</p> Full article ">Figure 10
<p>Atomic arrangement of GB models with four impurity atoms (<b>a</b>–<b>e</b>) and eight impurity atoms (<b>f</b>). The purple balls are Fe atoms, and the red balls are possible substituted sites.</p> Full article ">Figure 11
<p>Solution energies of pure GB, Si-doped GB, P-doped GB, and S-doped GB with different concentrations.</p> Full article ">Figure 12
<p>Charge density of the most stable GB models with 4Si (<b>a</b>), 4P (<b>b</b>), and 4S (<b>c</b>).</p> Full article ">Figure 13
<p>The concentration effects of Si (<b>a</b>), P (<b>b</b>), and S (<b>c</b>) on the GB, and the maximum stress as a function of solute concentration (<b>d</b>).</p> Full article ">Figure 13 Cont.
<p>The concentration effects of Si (<b>a</b>), P (<b>b</b>), and S (<b>c</b>) on the GB, and the maximum stress as a function of solute concentration (<b>d</b>).</p> Full article ">
16 pages, 12284 KiB  
Article
Microstructure Evolution and Strengthening Mechanisms of Mg–Steel Welds Subjected to Multiple Microshot Peening Treatment
by Jianghui Wang and Chuan Xu
Metals 2024, 14(4), 470; https://doi.org/10.3390/met14040470 - 17 Apr 2024
Viewed by 598
Abstract
A surface modification through multiple microshot peening (MSP) was performed on Mg–steel weldment. Application of MSP was found beneficial to the elimination of surface microdefects owing to severe plastic deformation induced by MSP. Moreover, MSP treatment transformed the residual tensile stress of the [...] Read more.
A surface modification through multiple microshot peening (MSP) was performed on Mg–steel weldment. Application of MSP was found beneficial to the elimination of surface microdefects owing to severe plastic deformation induced by MSP. Moreover, MSP treatment transformed the residual tensile stress of the weld surface into residual compressive stress, which was beneficial to inhibit the initiation and propagation of surface microdefects. Strain strengthening and grain refining were introduced into the shot peened joint, resulting in the notable increase in surface hardness and tensile strength. Compared with an untreated joint, the tensile strength of optimized Mg/steel weldment was markedly enhanced and raised 28% to 244 MPa, and fracture ultimately occurred in the Mg alloy base material. Moreover, the refinement of weld grain induced by MSP treatment was beneficial to strengthen the stress corrosion sensitivity of Mg/steel joints, while also promoting the formation of a denser Mg(OH)2 passivation film on the weld surface and enhancing the corrosion resistance of the joints. Full article
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<p>The schematic illustration of shot peening process (<b>a</b>) and strengthening mechanism (<b>b</b>).</p> Full article ">Figure 2
<p>Sampling position of Mg–steel joint: Ⅰ tensile test specimen, Ⅱ specimen for microhardness and microstructure analysis.</p> Full article ">Figure 3
<p>Typical surface microstructure of Mg–steel specimen: (<b>a</b>,<b>c</b>,<b>d</b>) without shot peening (Joint S1); (<b>b</b>,<b>e</b>) with shot peening process (Joint S5).</p> Full article ">Figure 4
<p>Surface roughness measurement results of weld and base metal after MSP treatment: (<b>a</b>) maximum profile height Ry, (<b>b</b>) arithmetic average roughness Ra.</p> Full article ">Figure 5
<p>Depth distribution of residual stress on Mg–steel weld after varied shot peening passes.</p> Full article ">Figure 6
<p>The typical metallographic diagram of (<b>a</b>) base material, (<b>b</b>) cross-section of traditional Mg–steel weld (Joint S1), and (<b>c</b>) cross-section of shot peened Mg–steel weld (Joint S5).</p> Full article ">Figure 7
<p>Calculation results of average crystallite dimension and lattice distortion of weld surface after microshot peening treatment.</p> Full article ">Figure 8
<p>TEM field images of weld surface zone: (<b>a</b>) and (<b>b</b>) Joint S5, (<b>c</b>) Joint S2, (<b>d</b>) Joint S3, (<b>e</b>) Joint S4, and (<b>f</b>) Joint S6.</p> Full article ">Figure 9
<p>The grain refining process of Mg–steel weld under multiple shot peening: (<b>a</b>) low-density dislocations in original grains, (<b>b</b>) dislocation multiplies under projectile impact, (<b>c</b>) dislocation walls and dislocation cells, (<b>d</b>) subgrain boundaries, (<b>e</b>) dynamic recrystallization nucleation, and (<b>f</b>) refined grain.</p> Full article ">Figure 10
<p>Schematic diagram of hardness measurement position (<b>a</b>) and the microhardness variations of weld surface zone with varied peening parameters (<b>b</b>).</p> Full article ">Figure 11
<p>Tensile strength of Mg–steel weldment under MSP treatment.</p> Full article ">Figure 12
<p>Typical fracture features of Mg–steel joint under shot peening treatment: (<b>a</b>–<b>c</b>) fracture locations of Joints S1, S5, and S6, (<b>d</b>–<b>f</b>) fracture appearances of Joints S1, S5, and S6.</p> Full article ">Figure 13
<p>Fracture diagram of Mg–steel joint of S6 (with Almen intensity of 0.25 mm N): (<b>a</b>) fracture process, (<b>b</b>,<b>c</b>) enlarged view of initial weld area and fractured weld zone, respectively.</p> Full article ">Figure 14
<p>Schematic diagram of the influence of grain refinement on the evolution of stress corrosion cracks: (<b>a</b>) untreated joint, (<b>b</b>) Joint S5 (with MSP strength of 0.20 mm N).</p> Full article ">Figure 15
<p>Corrosion morphologies of Mg/steel weldments after 30 min salt bath test (3.5 wt.% NaCl solution): (<b>a</b>) Joint S1 (without MSP), (<b>b</b>) Joint S5 (MSP intensity of 0.20 mm N).</p> Full article ">
14 pages, 24089 KiB  
Article
Effect of Precipitated Particles on Austenite Grain Growth of Al- and Nb-Microalloyed 20MnCr Gear Steel
by Yingqi Zhu, Shitao Fan, Xiuzhen Lian and Na Min
Metals 2024, 14(4), 469; https://doi.org/10.3390/met14040469 - 17 Apr 2024
Cited by 1 | Viewed by 688
Abstract
The paper deals with the effect of the morphology characteristics, grain size, and the volume fraction of AlN- and NbC-precipitated particles on the prior austenite grain growth behavior in the Al- and Nb-microalloying 20MnCr gear steel during pseudo-carburizing heat treatments. The results indicate [...] Read more.
The paper deals with the effect of the morphology characteristics, grain size, and the volume fraction of AlN- and NbC-precipitated particles on the prior austenite grain growth behavior in the Al- and Nb-microalloying 20MnCr gear steel during pseudo-carburizing heat treatments. The results indicate that the Nb addition in 20MnCr gear steel have a better effect on preventing austenite grain growth. The coarsening time after pseudo-carburizing in the Nb-microalloyed 20MnCr steel are improved by about 4 h compared with the Al-microalloyed steel. The precipitated particles coarsen and the number decreases with the pseudo-carburization temperature increasing, resulting in a reduction in the pinning pressure of the precipitated particles on the austenite grain boundaries. When the pseudo-carburization temperature reaches 1150 °C, the precipitated particles no longer have the ability to pin the austenite grain boundaries. In addition, the kinetics model for austenite grain growth under the process of the pinning and coarsening of the precipitated particles was established. Full article
(This article belongs to the Special Issue Microalloying in Ferrous and Non-ferrous Alloys)
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<p>The schematic of the heat treatment process.</p> Full article ">Figure 2
<p>Optical micrographs of prior austenite grain in the C0 steel after pseudo-carburization at different temperature for 0.5 h: (<b>a</b>) 950 °C; (<b>b</b>) 1050 °C; and (<b>c</b>) 1150 °C.</p> Full article ">Figure 3
<p>Optical micrographs of prior austenite grain in the C0 steel after pseudo-carburization at different temperature for 1 h: (<b>a</b>) 950 °C; (<b>b</b>) 1050 °C; and (<b>c</b>) 1150 °C.</p> Full article ">Figure 4
<p>Optical micrographs of prior austenite grain in the C0 steel after pseudo-carburization at different temperature for 5 h: (<b>a</b>) 950 °C; (<b>b</b>) 1050 °C; and (<b>c</b>) 1150 °C.</p> Full article ">Figure 5
<p>Optical micrographs of prior austenite grain in the C1 steel after pseudo-carburization at different temperature for 0.5 h: (<b>a</b>) 950 °C; (<b>b</b>) 1050 °C; and (<b>c</b>) 1150 °C.</p> Full article ">Figure 6
<p>Optical micrographs of prior austenite grain in the C1 steel after pseudo-carburization at different temperature for 1 h: (<b>a</b>) 950 °C; (<b>b</b>) 1050 °C; and (<b>c</b>) 1150 °C.</p> Full article ">Figure 7
<p>Optical micrographs of prior austenite grain in the C1 steel after pseudo-carburization at different temperature for 5 h: (<b>a</b>) 950 °C; (<b>b</b>) 1050 °C; and (<b>c</b>) 1150 °C.</p> Full article ">Figure 8
<p>Grain size of prior austenite grain in different (<b>a</b>,<b>c</b>) temperatures and (<b>b</b>,<b>d</b>) times.</p> Full article ">Figure 9
<p>Determination of precipitated particles in the C0 steel after pseudo-carburization at 950 °C for 0.5 h: (<b>a</b>) TEM micrograph (BF), (<b>b</b>) corresponding selected area electron diffraction (SAED) pattern, and (<b>c</b>) EDS corresponding to the precipitated particle in (<b>a</b>).</p> Full article ">Figure 10
<p>SEM micrograph of the precipitated particles in the C1 steel after pseudo-carburization at 1150 °C for 0.5 h.</p> Full article ">Figure 11
<p>Determination of the precipitated particles in the C1 steel after pseudo-carburization at 1050 °C for 0.5 h: (<b>a</b>) TEM micrograph (BF), (<b>b</b>) NbC precipitate’s SAED pattern, and (<b>c</b>) EDS corresponding to the precipitated particle in (<b>a</b>).</p> Full article ">Figure 12
<p>The size distribution of the NbC-precipitated particles in the C1 steel after pseudo-carburization at different temperatures for different times. The height of bars represents probability density (frequency). (<b>a</b>) at 950 °C for 0.5 h; (<b>b</b>) at 950 °C for 1 h; (<b>c</b>) at 950 °C for 5 h; (<b>d</b>) at 1050 °C for 0.5 h; (<b>e</b>) at 1050 °C for 1 h; (<b>f</b>) at 1050 °C for 5 h; (<b>g</b>) at 1150 °C for 0.5 h; (<b>h</b>) at 1150 °C for 1 h; (<b>i</b>) at 1150 °C for 5 h.</p> Full article ">Figure 13
<p>The volume fractions of the NbC-precipitated particles in the C1 steel.</p> Full article ">Figure 14
<p>Size distribution of the AlN-precipitated particles in the C0 steel after pseudo-carburization at 950 °C for 0.5 h.</p> Full article ">Figure 15
<p>Pinning pressure of the NbC-precipitated particles in the C1 steel at different pseudo-carburizing treatments.</p> Full article ">
14 pages, 16731 KiB  
Article
Corrosion Behavior of FeCrMnxAlCu High-Entropy Alloys in NaOH Solutions
by Yang Yang, Junpeng Cui, Zhipeng Wang and Li Feng
Metals 2024, 14(4), 468; https://doi.org/10.3390/met14040468 - 17 Apr 2024
Viewed by 801
Abstract
FeCrMnxAlCu (x = 2.0, 1.5, 1.0, 0.5, and 0.0) high-entropy alloys (HEAs) were prepared using vacuum arc melting. The phase structure, microstructure, and element distribution of FeCrMnxAlCu (x = 2.0, 1.5, 1.0, 0.5, and 0.0) HEAs were analyzed using [...] Read more.
FeCrMnxAlCu (x = 2.0, 1.5, 1.0, 0.5, and 0.0) high-entropy alloys (HEAs) were prepared using vacuum arc melting. The phase structure, microstructure, and element distribution of FeCrMnxAlCu (x = 2.0, 1.5, 1.0, 0.5, and 0.0) HEAs were analyzed using X-ray diffraction, scanning electron microscopy, and energy-dispersive spectroscopy. The corrosion resistance of the alloy in a NaOH solution was evaluated using a potentiodynamic polarization curve, electrochemical impedance spectroscopy, an immersion test, WLI, and X-ray photoelectron spectroscopy. The results showed that FeCrMnxAlCu HEAs are an FCC+BCC bi-phase mixed structure with typical dendrite and interdendrite structures. The corrosion test demonstrated that the HEAs presented a breakthrough characteristic of activation–passivation–passivation. With an increase in Mn content, the corrosion potential shifted first positively and then negatively, and the corrosion current first decreased and then increased. Among HEAs, the FeCrMn0.5AlCu HEA had the best electrochemical corrosion resistance. After corrosion, both oxide and hydroxide corrosion product films were formed on the surface, which reduced the ion diffusion rate, slowed down the corrosion process, and improved the corrosion resistance. Full article
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<p>The X-ray diffraction pattern of the FeCrMn<sub>x</sub>AlCu HEAs.</p> Full article ">Figure 2
<p>The SEM images of the FeCrMn<sub>x</sub>AlCu HEAs: (<b>a</b>) x = 0, (<b>b</b>) x = 0.5, (<b>c</b>) x = 1.0, (<b>d</b>) x = 1.5, (<b>e</b>) x = 2.0.</p> Full article ">Figure 3
<p>Face scanning EDS map of FeCrMn<sub>1.5</sub>AlCu HEAs.</p> Full article ">Figure 4
<p>Electrochemical test results of FeCrMn<sub>x</sub>AlCu HEAs in 1 M NaOH solution: (<b>a</b>) polarization curves, (<b>b</b>) Nyquist plots, (<b>c</b>) Bode phase angle, (<b>d</b>) Bode mode plots.</p> Full article ">Figure 5
<p>The SEM images of the FeCrMn<sub>x</sub>AlCu HEAs after electrochemical testing: (<b>a</b>) x = 0, (<b>b</b>) x = 0.5, (<b>c</b>) x = 1.0, (<b>d</b>) x = 1.5, (<b>e</b>) x = 2.0.</p> Full article ">Figure 6
<p>The corrosion rate of the FeCrMn<sub>x</sub>AlCu HEAs after immersion.</p> Full article ">Figure 7
<p>The surface SEM images of FeCrMn<sub>x</sub>AlCu HEAs after immersion testing: (<b>a</b>) x = 0, (<b>b</b>) x = 0.5, (<b>c</b>) x = 1.0, (<b>d</b>) x = 1.5, (<b>e</b>) x = 2.0.</p> Full article ">Figure 8
<p>EDS spectrum of FeCrMn<sub>0.5</sub>AlCu HEAs after immersion.</p> Full article ">Figure 9
<p>Three-dimensional images of the surface of FeCrMn<sub>x</sub>AlCu HEAs after immersion: (<b>a</b>) x = 0, (<b>b</b>) x = 0.5, (<b>c</b>) x = 1.0, (<b>d</b>) x = 1.5, (<b>e</b>) x = 2.0.</p> Full article ">Figure 10
<p>High-resolution spectroscopy of FeCrMn<sub>x</sub>AlCu HEAs after corrosion: (<b>a</b>) x = 0, (<b>b</b>) x = 0.5, (<b>c</b>) x = 2.0.</p> Full article ">Figure 10 Cont.
<p>High-resolution spectroscopy of FeCrMn<sub>x</sub>AlCu HEAs after corrosion: (<b>a</b>) x = 0, (<b>b</b>) x = 0.5, (<b>c</b>) x = 2.0.</p> Full article ">Figure 10 Cont.
<p>High-resolution spectroscopy of FeCrMn<sub>x</sub>AlCu HEAs after corrosion: (<b>a</b>) x = 0, (<b>b</b>) x = 0.5, (<b>c</b>) x = 2.0.</p> Full article ">Figure 11
<p>The atomic percentage of each element’s valence state on the surface of FeCrMn<sub>x</sub>AlCu HEAs after corrosion.</p> Full article ">
15 pages, 7779 KiB  
Article
Mechanical Properties of Interfaces between Mg and SiC: An Ab Initio Study
by Zhipeng Yao, Samaneh Nasiri, Mingjun Yang and Michael Zaiser
Metals 2024, 14(4), 467; https://doi.org/10.3390/met14040467 - 16 Apr 2024
Viewed by 836
Abstract
Covalently bonded particles may exhibit extremely high strength, but their performance in the reinforcement of metal alloys crucially depends on the properties of their interfaces with the embedding matrix. Here, density functional theory is used for investigating a range of interface configurations between [...] Read more.
Covalently bonded particles may exhibit extremely high strength, but their performance in the reinforcement of metal alloys crucially depends on the properties of their interfaces with the embedding matrix. Here, density functional theory is used for investigating a range of interface configurations between magnesium and silicon carbide in view of their mechanical properties. Interfaces are analyzed not only in terms of interface energy/work of separation but also in terms of the interfacial shear stresses required to induce interface-parallel displacements. These properties are studied for bilayer systems with different orientations of the Mg and SiC layers and for different terminations of the SiC layer (Si or C atoms located at the interface). The results are discussed in terms of their implication for mechanical behavior of SiC reinforced Mg alloys. Full article
(This article belongs to the Special Issue Multi-Scale Simulation of Metals and Alloys)
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<p>Views of the geometry of the investigated models of C-terminated interfaces; the black-lined boxes indicate the supercells used in the calculations, <span class="html-fig-inline" id="metals-14-00467-i001"><img alt="Metals 14 00467 i001" src="/metals/metals-14-00467/article_deploy/html/images/metals-14-00467-i001.png"/></span> Si atoms, <span class="html-fig-inline" id="metals-14-00467-i002"><img alt="Metals 14 00467 i002" src="/metals/metals-14-00467/article_deploy/html/images/metals-14-00467-i002.png"/></span> C atoms, <span class="html-fig-inline" id="metals-14-00467-i003"><img alt="Metals 14 00467 i003" src="/metals/metals-14-00467/article_deploy/html/images/metals-14-00467-i003.png"/></span> Mg atoms; left: CSi(0001)Mg(0001), center: CSi(111)Mg(0001), right: CSi(111)Mg(10<math display="inline"><semantics> <mover> <mn>1</mn> <mo>¯</mo> </mover> </semantics></math>0).</p> Full article ">Figure 2
<p><b>Left</b>: Top views of the geometry of four variant models for SiC(111)Mg(0001): (<b>I</b>) bridge, (<b>II</b>) hollow, (<b>III</b>) top-C, and (<b>IV</b>) top-Si ); <b>right</b>: top views of the geometry of four variant models for SiC(111)Mg(10<math display="inline"><semantics> <mover> <mn>1</mn> <mo>¯</mo> </mover> </semantics></math>0): (<b>I</b>) bridge, (<b>II</b>) hollow, (<b>III</b>) top-C, and (<b>IV</b>) top-Si; color scheme the same as in <a href="#metals-14-00467-f001" class="html-fig">Figure 1</a>.</p> Full article ">Figure 3
<p>Comparison between interfaces of Mg(0001) with SiC(111) and SiC(0001) polytypes; four interface configurations are considered: bridge, hollow, top-C, and top-Si; top graph: C termination; bottom graph: Si termination.</p> Full article ">Figure 4
<p>Comparison of interface energy curves between Mg(0001) and C-terminated CSi(111) vs. Si-terminated SiC(111); four interface configurations are considered: bridge, hollow, top-C, and top-Si.</p> Full article ">Figure 5
<p>Interface tensile stress curves, comparison of interfaces between Mg(0001) and C-terminated CSi(111) vs. Si-terminated SiC(111); four interface configurations are considered: bridge, hollow, top-C, and top-Si.</p> Full article ">Figure 6
<p>Comparison between interface energy curves for interfaces between CSi(111) and two different orientations of the Mg slab: Mg(0001) and Mg(10<math display="inline"><semantics> <mover> <mn>1</mn> <mo>¯</mo> </mover> </semantics></math>0); four interface configurations are considered: bridge, hollow, top-C, and top-Si.</p> Full article ">Figure 7
<p>Interface tensile stress curves, comparison of interfaces between CSi(111) and two different orientations of the Mg layer: Mg(0001) and Mg(10<math display="inline"><semantics> <mover> <mn>1</mn> <mo>¯</mo> </mover> </semantics></math>0); four interface configurations are considered: bridge, hollow, top-C, and top-Si.</p> Full article ">Figure 8
<p>Interface energy surfaces: (<b>top left</b>), CSi(111)Mg(0001); (<b>top right</b>) SiC(111)Mg(0001); (<b>bottom left</b>), CSi(111)Mg(10<math display="inline"><semantics> <mover> <mn>1</mn> <mo>¯</mo> </mover> </semantics></math>0); (<b>bottom right</b>), Si(111)Mg(10<math display="inline"><semantics> <mover> <mn>1</mn> <mo>¯</mo> </mover> </semantics></math>0); the lines marked P1 and P2 indicate the directions along which energy and IFSS profiles were taken.</p> Full article ">Figure 9
<p>Interface energy and IFSS profiles taken along the lines indicated by P1 and P2 in <a href="#metals-14-00467-f008" class="html-fig">Figure 8</a>: top, CSi(111)Mg(0001); bottom, SiC(111)Mg(0001); P1 = [<math display="inline"><semantics> <mrow> <mn>10</mn> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mn>0</mn> </mrow> </semantics></math>] (Mg)/[<math display="inline"><semantics> <mrow> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mn>2</mn> </mrow> </semantics></math>] (SiC), P2 = [<math display="inline"><semantics> <mrow> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mn>2</mn> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mn>0</mn> </mrow> </semantics></math>] (Mg)/[<math display="inline"><semantics> <mrow> <mn>1</mn> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mn>0</mn> </mrow> </semantics></math>] (SiC).</p> Full article ">Figure 10
<p>Interface energy and IFSS profiles taken along the lines indicated by P1 and P2 in <a href="#metals-14-00467-f008" class="html-fig">Figure 8</a>; top: CSi(111)Mg(10<math display="inline"><semantics> <mover> <mn>1</mn> <mo>¯</mo> </mover> </semantics></math>0), bottom: SiC(111)Mg(10<math display="inline"><semantics> <mover> <mn>1</mn> <mo>¯</mo> </mover> </semantics></math>0); P1 = [0001] (Mg)/[<math display="inline"><semantics> <mrow> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mn>2</mn> </mrow> </semantics></math>] (SiC), P2 = [<math display="inline"><semantics> <mrow> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mn>2</mn> <mover> <mn>1</mn> <mo>¯</mo> </mover> <mn>0</mn> </mrow> </semantics></math>] (Mg)/[<math display="inline"><semantics> <mover> <mn>1</mn> <mo>¯</mo> </mover> </semantics></math>10] (SiC).</p> Full article ">
18 pages, 10455 KiB  
Article
Study on Al Evaporation during AlV55 Melting and Alloy Preparation
by Bin Sun, Heli Wan, Baoqiang Xu, Xianjun Lei and Lanjie Li
Metals 2024, 14(4), 466; https://doi.org/10.3390/met14040466 - 16 Apr 2024
Viewed by 872
Abstract
Vacuum induction melting is a commonly used method for preparing AlV55 alloys. However, this method results in high Al evaporation losses, leading to poor cost control. Moreover, the influence of the process parameters on the alloy composition remains unclear. In this study, the [...] Read more.
Vacuum induction melting is a commonly used method for preparing AlV55 alloys. However, this method results in high Al evaporation losses, leading to poor cost control. Moreover, the influence of the process parameters on the alloy composition remains unclear. In this study, the evaporation pattern of Al in the melting and preparation processes of AlV55 alloys is studied, and measures for controlling the system pressure are proposed to effectively reduce Al evaporation. The experimental results show that smelting under an Ar gas atmosphere of 2000 Pa can reduce the evaporation loss of Al from 11.48% under vacuum conditions (60 Pa) to 0.58% of the amount of raw material added, effectively improving the metal utilization rate and reducing production costs. The influence of various process parameters on the alloy composition and Al volatilization are investigated to enable the effective control of the compositional uniformity and impurity content of the resulting alloys. Under optimal conditions, the impurity contents of C, O, and N are 0.03%, 0.0073%, and 0.013%, respectively; this reduces the amount of Al lost by evaporation compared to conventional methods, and the alloy produced meets commercial standards. Full article
(This article belongs to the Special Issue Remelting and Casting Processes in the Production of Metals and Alloys)
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<p>(<b>a</b>) Photograph of AlV65 alloy scrap; (<b>b</b>) XRD pattern of AlV65.</p> Full article ">Figure 2
<p>Schematic of the AlV55 alloy preparation process.</p> Full article ">Figure 3
<p>Schematic diagram of the reaction furnace system.</p> Full article ">Figure 4
<p>Saturated vapor pressures <span class="html-italic">P</span> of different elements in the alloy as a function of temperature <span class="html-italic">T</span>.</p> Full article ">Figure 5
<p>Phase diagram of Al-V alloy.</p> Full article ">Figure 6
<p>Molecular free path of Al.</p> Full article ">Figure 7
<p>Schematic diagram of the Al evaporation process.</p> Full article ">Figure 8
<p>XRD pattern of AlV55 alloy.</p> Full article ">Figure 9
<p>(<b>a</b>) SEM image and EDS analysis of AlV55 alloy; (<b>b</b>,<b>c</b>) scanning results of (<b>b</b>) Spot 1 and (<b>c</b>) Spot 2 in (<b>a</b>).</p> Full article ">Figure 10
<p>Theoretical evaporation rate of Al.</p> Full article ">Figure 11
<p>Theoretical and measured values of Al evaporation.</p> Full article ">Figure 12
<p>Effect of temperature on the V content and Al evaporation loss of the alloy product.</p> Full article ">Figure 13
<p>SEM images of the alloy product obtained at different temperatures: (<b>a</b>) 1973 K, (<b>b</b>) 2023 K, (<b>c</b>) 2073 K, (<b>d</b>) 2123 K, and (<b>e</b>) 2173 K.</p> Full article ">Figure 14
<p>Effect of holding time on the V content and Al evaporation of the alloy product.</p> Full article ">Figure 15
<p>SEM images of the alloy product obtained at different holding times: (<b>a</b>) 5 min, (<b>b</b>) 10 min, (<b>c</b>) 15 min, (<b>d</b>) 20 min, and (<b>e</b>) 25 min.</p> Full article ">Figure 16
<p>Effect of material ratio on V content.</p> Full article ">Figure 17
<p>Impurity removal effect.</p> Full article ">Figure 18
<p>(<b>a</b>) Photograph of the alloy section with sampling points indicated; (<b>b</b>–<b>f</b>) SEM images of the sampling points: (<b>b</b>) Point 1, (<b>c</b>) Point 2, (<b>d</b>) Point 3, (<b>e</b>) Point 4, and (<b>f</b>) Point 5.</p> Full article ">
15 pages, 2041 KiB  
Article
Optimization of High-Alumina Blast Furnace Slag Based on Exergy Analysis
by Zhen Wang, Haiyan Zheng, Yan Zhang and Liang Ge
Metals 2024, 14(4), 465; https://doi.org/10.3390/met14040465 - 15 Apr 2024
Viewed by 913
Abstract
Raw material with a high Al2O3 content has led to an increase in the Al2O3 content in blast furnace slag, which has affected the normal operation of a blast furnace. The exergy analysis method is an important [...] Read more.
Raw material with a high Al2O3 content has led to an increase in the Al2O3 content in blast furnace slag, which has affected the normal operation of a blast furnace. The exergy analysis method is an important method for studying the energy utilization of high-alumina blast furnace smelting. In this paper, to investigate the impact of slag composition on exergy efficiency and optimize exergy efficiency during the smelting process of high Al2O3 iron ore, a gray box exergy analysis model of blast furnace smelting and an objective function for minimizing the total exergy loss were developed. The results indicated that the blast furnace smelting process had an exergy efficiency (η) of 28.29% for hot metal and slag; the exergy efficiency of the blast furnace did not significantly increase with the increasing w(MgO)/w(Al2O3) and R (w(CaO)/w(SiO2)), but the exergy efficiency of the blast furnace declined with increasing w(Al2O3). The regional optimal solution for the objective function method was 7129.42 MJ with slag compositions of R = 1.295, w(MgO)/w(Al2O3) = 0.545, and w(Al2O3) = 15%. Full article
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<p>Exergy analysis model diagram of blast furnace smelting.</p> Full article ">Figure 2
<p>Exergy flows of the whole blast furnace.</p> Full article ">Figure 3
<p>Effect of <span class="html-italic">R</span> on <span class="html-italic">η</span>.</p> Full article ">Figure 4
<p>Effect of <span class="html-italic">R</span> on <span class="html-italic">E</span><sub>x, ph, slag</sub> and <span class="html-italic">E</span><sub>x, ch, slag</sub>.</p> Full article ">Figure 5
<p>Effect of <span class="html-italic">w</span>(MgO)/<span class="html-italic">w</span>(Al<sub>2</sub>O<sub>3</sub>) on <span class="html-italic">η</span>.</p> Full article ">Figure 6
<p>Effect of <span class="html-italic">w</span>(MgO)/<span class="html-italic">w</span>(Al<sub>2</sub>O<sub>3</sub>) on <span class="html-italic">E</span><sub>x, slag</sub> and <span class="html-italic">E</span><sub>total</sub>.</p> Full article ">Figure 7
<p>Effect of <span class="html-italic">w</span>(Al<sub>2</sub>O<sub>3</sub>) on <span class="html-italic">η</span>.</p> Full article ">Figure 8
<p>Effect of <span class="html-italic">w</span>(Al<sub>2</sub>O<sub>3</sub>) on <span class="html-italic">E</span><sub>x, slag</sub>, <span class="html-italic">E</span><sub>x, gas</sub>, and <span class="html-italic">E</span><sub>xl, ex</sub>: <span class="html-italic">R</span> = 1.20 and <span class="html-italic">w</span>(MgO)/<span class="html-italic">w</span>(Al<sub>2</sub>O<sub>3</sub>) = 0.45.</p> Full article ">Figure 9
<p>The objective function, gray box model, and actual production exergy loss are compared.</p> Full article ">
20 pages, 8697 KiB  
Review
A Comprehensive Review of Fatigue Strength in Pure Copper Metals (DHP, OF, ETP)
by Eduardo Jiménez-Ruiz, Rubén Lostado-Lorza and Carlos Berlanga-Labari
Metals 2024, 14(4), 464; https://doi.org/10.3390/met14040464 - 15 Apr 2024
Viewed by 2056
Abstract
Due to their exceptional electrical and thermal conductivity properties, high-purity copper (Cu-DHP) and copper alloys of similar composition, such as electrolytic tough-pitch (ETP), oxygen-free electronic (OFE) and oxygen-free (OF), have often been used in the manufacture of essential components for the electrical, electronic [...] Read more.
Due to their exceptional electrical and thermal conductivity properties, high-purity copper (Cu-DHP) and copper alloys of similar composition, such as electrolytic tough-pitch (ETP), oxygen-free electronic (OFE) and oxygen-free (OF), have often been used in the manufacture of essential components for the electrical, electronic and power generation industries. Since these components are subject to cyclic loads in service, they can suffer progressive structural damage that causes failure due to fatigue. The purpose of this review is to examine the most relevant aspects of mechanical fatigue in Cu-DHP, ETP, OFE and OF. The impact of many factors on fatigue strength (Se), including the frequency, temperature, chemical environment, grain size, metallurgical condition and load type, were analyzed and discussed. Stress–life (S-N) curves under zero mean stress (σm = 0) were found for high-cycle fatigue (HCF). For non-zero mean stress (σm ≠ 0), stress curves were based on a combination of Gerber, Soderberg and ASME elliptic failure criteria. Stress–life (S-N) curves were also developed to correlate fatigue strength (Se) with stress amplitude (σa), yield strength (Syp) and ultimate strength (Sut). Finally, for low-cycle fatigue (LCF), strain–life (ε-N) curves that establish a relationship between the number of cycles to failure (N) and total strain amplitude (εplastic) were determined. Hence, this review, as well as the proposed curves, provide valuable information to understand fatigue failure for these types of materials. Full article
(This article belongs to the Special Issue Fatigue Behavior in Metallic Materials)
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<p>Fatigue strength (S<sub>e</sub>) data for Cu-DHP. Individual values correspond to pipe and strips of Cu-DHP annealed and cold-worked in several grades and tested in air at room temperature.</p> Full article ">Figure 2
<p>Effect of load type on the fatigue strength (S<sub>e</sub>) of copper of high purity Reprinted with permission from ref. [<a href="#B14-metals-14-00464" class="html-bibr">14</a>] John Wiley &amp; Sons 1985. Tests of smooth bars conducted in air at room temperature under zero mean stress conditions (R = −1).</p> Full article ">Figure 3
<p>Effect of frequency on the fatigue strength (S<sub>e</sub>) of copper of high purity reprinted with permission from ref. [<a href="#B14-metals-14-00464" class="html-bibr">14</a>] John Wiley &amp; Sons 1985. Testing was conducted in air at room temperature in zero mean stress conditions (R = −1), material TP or OF, reverse bending and push-pull, 0.1, 2.1 and 32.3 Hz cold-worked, 30–50, 85 and 10–150 Hz annealed. 30–50 Hz in several environments: lab air, vacuum, damp nitrogen, damp pure air and dry pure air.</p> Full article ">Figure 4
<p>Reference data by frequency (reprinted with permission from refs. [<a href="#B21-metals-14-00464" class="html-bibr">21</a>,<a href="#B22-metals-14-00464" class="html-bibr">22</a>,<a href="#B23-metals-14-00464" class="html-bibr">23</a>,<a href="#B24-metals-14-00464" class="html-bibr">24</a>] Elsevier Science &amp; Technology Journals 1988, 1996, 1979). Test conditions: Fintova: cold-rolled ETP in air at room temperature, push-pull. Marti: OFHC in air at room temperature, push-pull. Stanzl-Tschegg: ETP 99.98 and ETP in air at room temperature, push-pull. Phung: strengthened OFHC in air at room temperature, push-pull.</p> Full article ">Figure 5
<p>Effect of grain size on the fatigue strength (S<sub>e</sub>) of high-purity copper reprinted with permission from ref. [<a href="#B14-metals-14-00464" class="html-bibr">14</a>] John Wiley &amp; Sons 1985. Test conditions. 0–0.05: Annealed TP or OF, in air at room temperature, reverse bending and push-pull. 0.051–0.1: Annealed TP or OF, in air at room temperature, reverse bending and push-pull. 0.101–0.50: Annealed TP or OF, in air at room temperature, push-pull.</p> Full article ">Figure 6
<p>Effect of grain size on the fatigue strength (S<sub>e</sub>) of high-purity copper. (UFG: ultrafine grained.) Test conditions. 0–0.05: Annealed TP or OF, in air at room temperature, reverse bending and push-pull. UFG: OF, in air at room temperature, reverse bending, torsion and push-pull.</p> Full article ">Figure 7
<p>Effect of metallurgical condition on the fatigue strength (S<sub>e</sub>) of copper of high purity reprinted with permission from ref. [<a href="#B14-metals-14-00464" class="html-bibr">14</a>] John Wiley &amp; Sons 1985. Test conditions. 70% cold-worked: Cold-worked 70% OF or TP, in air at room temperature. Annealed: Annealed TP or OF, mainly in air, a few in damp nitrogen at room temperature, reverse bending and push-pull, several frequencies. Cold-worked: Cold-worked TP or OF in air at room temperature, reverse bending and push-pull and several frequencies.</p> Full article ">Figure 8
<p>Fatigue strength (S<sub>e</sub>) data categorized by environment. Murphy reprinted from Ref. [<a href="#B11-metals-14-00464" class="html-bibr">11</a>]. Test conditions. Air RT: Annealed TP in air at room temperature, reverse bending and push-pull, 85 Hz. Damp nitrogen: Annealed TP or OF at room temperature, push-pull, 30–50 Hz. Damp pure air: Annealed TP or OF at room temperature, push-pull, 30–50 Hz. Lab air: Annealed TP or OF at room temperature, push-pull, 30–50 Hz. Vacuum: Annealed TP or OF at room temperature, push-pull, 30–50 Hz.</p> Full article ">Figure 9
<p>Fatigue strength (S<sub>e</sub>) data categorized by temperature (Simon reprinted from Ref. [<a href="#B42-metals-14-00464" class="html-bibr">42</a>]). Test condition. 4: Annealed OF in air/liquid, push-pull. 20: Annealed OF in air/liquid, push-pull. 77: Annealed OF in air/liquid/vacuum/liquid nitrogen, push-pull. 90: Annealed OF in air/liquid, push-pull. 295: Annealed OF in air/liquid/vacuum/liquid nitrogen/dry purified air/damp purified air, push-pull.</p> Full article ">Figure 10
<p>Fatigue strength (S<sub>e</sub>) data categorized by type of pure copper (Murphy reprinted with permission from ref. [<a href="#B14-metals-14-00464" class="html-bibr">14</a>] John Wiley &amp; Sons 1985). Test conditions. Cu-DHP: Cu-DHP at different metallurgical states, in air at room temperature, reverse bending and push-pull. OF: OF at several metallurgical states in air at room temperature, push-pull and reverse bending, 10–150 Hz.</p> Full article ">Figure 11
<p>Fatigue strength (S<sub>e</sub>) data of ETP, OFE, OF and Cu-DHP.</p> Full article ">Figure 12
<p>Fatigue strength (S<sub>e</sub>) data of ETP, OFE, OF and Cu–DHP compiled according to the corresponding authors. Cu–DHP data are depicted using red triangles, while other copper samples are represented by black dots.</p> Full article ">Figure 13
<p>Haigh diagram with Murphy data and a proposed limit of 10<sup>7</sup> cycles. Test carried out on OF in air at room temperature at 85 Hz, push-pull under different mean stress conditions.</p> Full article ">Figure 14
<p>Haigh diagram with Murphy data and a proposed limit of 10<sup>8</sup> cycles. Test conducted on OF, in air at room temperature at 85 Hz, push-pull under different mean stress conditions.</p> Full article ">Figure 15
<p>Plastic strain amplitude (<span class="html-italic">ε<sub>plastic</sub></span>) vs. number of cycles (<span class="html-italic">N</span>) to failure of Cu-DHP, OF and OFHC. Test conditions. Cu-DHP: Cu-DHP in air at room temperature, push-pull. OF: Annealed OF in air/vacuum at room temperature, push-pull, several frequencies and temperatures. OFHC: Annealed OFHC in air at room temperature, push-pull, 0.5 Hz.</p> Full article ">Figure 16
<p>Plastic strain amplitude (<span class="html-italic">ε<sub>plasti</sub></span><sub>c</sub>) vs. number of cycles (<span class="html-italic">N</span>) to failure according to different authors. Test conditions: Coffin: OFHC in air at room temperature. Muhamma: Cu-DHP in air at room temperature, push-pull. Pan: Annealed OFHC in air at room temperature, push-pull, 0.5 Hz. Simon: Annealed OF in air at room temperature, push-pull, several frequencies. Takahasi: OF at several temperatures in vacuum. Yoichi Tamiya: Cu-DHP and other coppers in air at room temperature.</p> Full article ">Figure 17
<p>Plastic strain amplitude (<span class="html-italic">ε<sub>plasti</sub></span><sub>c</sub>) vs. number of cycles (<span class="html-italic">N</span>) to failure of Cu-DHP with different adjustment lines. Test conditions described in previous figures.</p> Full article ">Figure 18
<p>Plastic strain amplitude (<span class="html-italic">ε<sub>plasti</sub></span><sub>c</sub>) vs. number of cycles (<span class="html-italic">N</span>) to failure: the effect of temperature on fatigue strength (S<sub>e</sub>) at a low number of cycles reprinted from Ref. [<a href="#B51-metals-14-00464" class="html-bibr">51</a>]. Test conditions, OF in vacuum at several temperatures.</p> Full article ">Figure 19
<p>Optical microscopy image depicting a longitudinal cross-section of a copper-copper brazing joint. (<b>a</b>) Cu-DHP with average grain size 16–60 mm, (<b>b</b>) Cu-DHP with coarse grain size, (<b>c</b>) brazing material.</p> Full article ">Figure 20
<p>Optical microscopy image depicting a longitudinal cross-section of a copper-copper brazing joint. (<b>a</b>) Cu-DHP with averaged grain size 16–60 mm, (<b>b</b>) Cu-DHP with coarse grain size, (<b>c</b>) brazing material.</p> Full article ">Figure 21
<p>Photo of the copper tube failure studied by McDougal and Stevenson reprinted with permission from ref. [<a href="#B57-metals-14-00464" class="html-bibr">57</a>] Springer Nature BV, 2004.</p> Full article ">Figure 22
<p>Photo of the copper tube failure studied by Stevenson reprinted with permission from ref. [<a href="#B58-metals-14-00464" class="html-bibr">58</a>] Springer Nature BV, 2004.</p> Full article ">Figure 23
<p>Low-cycle fatigue failure in a Cu-DHP tube. Red and blue arrows show stretch marks from fatigue.</p> Full article ">
18 pages, 6263 KiB  
Article
Orientation Relationship of the Intergrowth Al13Fe3 and Al13Fe4 Intermetallics Determined by Single-Crystal X-ray Diffraction
by Yibo Liu, Changzeng Fan, Zhefeng Xu, Ruidong Fu, Bin Wen and Lifeng Zhang
Metals 2024, 14(4), 463; https://doi.org/10.3390/met14040463 - 15 Apr 2024
Viewed by 1037
Abstract
In the Al-Fe binary system, the Al13Fe3 phase as well as the Al13Fe4 phase has similar icosahedral building blocks like those appearing in quasicrystals. Therefore, it is of vital importance to clarify the formation process of these [...] Read more.
In the Al-Fe binary system, the Al13Fe3 phase as well as the Al13Fe4 phase has similar icosahedral building blocks like those appearing in quasicrystals. Therefore, it is of vital importance to clarify the formation process of these two phases. Coexistence of the Al13Fe3 and Al13Fe4 phases was discovered from the educts obtained with a nominal atomic ratio of Al/Fe of 9:2 by high-pressure sintering for the first time. Firstly, single crystal X-ray diffraction (SXRD) combined with a scanning electron microscope (SEM) with energy dispersive X-ray spectroscopy (EDX) measurement capabilities were adopted to determine the detailed crystal structures of both phases, which were sharply refined with regard to Al13Fe3 and Al13Fe4. Secondly, the orientation relationship between Al13Fe3 and Al13Fe4 was directly deduced from the SXRD datasets and the coexistence structure model was consequently constructed. Finally, seven pairs of parallel atomic planes and their unique orientation relations were determined from the reconstructed reciprocal space precession images. In addition, the real space structure model of the intergrowth crystal along with one kind of interfacial atomic structure were constructed from the determined orientation relations between two phases. Full article
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<p>Diffraction spots of the entire sample in the reciprocal space. (<b>a</b>) Projection of three sets of data in reciprocal space in a random direction; (<b>b</b>) projection of the Al<sub>13</sub>Fe<sub>3</sub>-1 phase and Al<sub>13</sub>Fe<sub>4</sub> phase along the a* axis of the Al<sub>13</sub>Fe<sub>4</sub> phase; (<b>c</b>) projection of the Al<sub>13</sub>Fe<sub>3</sub>-1 phase and Al<sub>13</sub>Fe<sub>3</sub>-2 phase along the a* axis of the Al<sub>13</sub>Fe<sub>3</sub>-1 phase; (<b>d</b>) projection of the Al<sub>13</sub>Fe<sub>3</sub>-2 phase and Al<sub>13</sub>Fe<sub>4</sub> phase along the b* axis of the Al<sub>13</sub>Fe<sub>3</sub>-2 phase.</p> Full article ">Figure 2
<p>Reciprocal lattice patterns and crystal structure diagram of the Al<sub>13</sub>Fe<sub>3</sub>-1 phase projected in different directions: (<b>a</b>) a* axis; (<b>b</b>) b* axis; (<b>c</b>) c* axis; (<b>d</b>) c axis.</p> Full article ">Figure 3
<p>Reciprocal lattice patterns and crystal structure diagram of Al<sub>13</sub>Fe<sub>4</sub> phase projected in different directions: (<b>a</b>) a* axis; (<b>b</b>) b* axis; (<b>c</b>) c* axis; (<b>d</b>) b axis.</p> Full article ">Figure 4
<p>Reciprocal lattice patterns and crystal structure diagram of the Al<sub>13</sub>Fe<sub>3</sub>-2 phase projected in different directions: (<b>a</b>) a* axis; (<b>b</b>) b* axis; (<b>c</b>) c* axis; (<b>d</b>) c axis.</p> Full article ">Figure 5
<p>Cluster assembly of the Al<sub>13</sub>Fe<sub>3</sub> phase in a unit cell: (<b>a</b>) the central atoms and isolated atoms of the cluster; (<b>b</b>) cluster assembly model.</p> Full article ">Figure 6
<p>Atomic environment diagram: (<b>a</b>) Al4 atom; (<b>b</b>) Fe1 atom. Symmetry codes: (i) −<span class="html-italic">x</span> + 2/3, −<span class="html-italic">x</span> + <span class="html-italic">y</span> + 1/3, −<span class="html-italic">z</span> + 5/6; (ii) −<span class="html-italic">y</span> + 2/3, <span class="html-italic">x</span> − <span class="html-italic">y</span> + 1/3, <span class="html-italic">z</span> + 1/3; (iii) <span class="html-italic">y</span> − 1/3, −<span class="html-italic">x</span> + <span class="html-italic">y</span> + 1/3, −<span class="html-italic">z</span> + 1/3; (iv) −<span class="html-italic">y</span> + 1, −<span class="html-italic">x</span> + 1, <span class="html-italic">z</span> + 1/2.</p> Full article ">Figure 7
<p>The cluster in the Al13Fe4 phase (The yellow atoms represent the Fe atoms and the blue atoms represent the Al atoms). (<b>a</b>) The inner shell of the Al7(2)(1@12@46) cluster; (<b>b</b>) the outer-shell of the Al7(2)(1@12@46) cluster; (<b>c</b>) the Al15(1)(1@12) cluster; (<b>d</b>) cluster assembly model in the Al13Fe4 unit cell.</p> Full article ">Figure 8
<p>Atomic environment diagram: (<b>a</b>) Al7 atom; (<b>b</b>) Al15 atom. Symmetry codes: (i) <span class="html-italic">x</span>, −<span class="html-italic">y</span>, <span class="html-italic">z</span>; (vii) −<span class="html-italic">x</span> + 1, <span class="html-italic">y</span>, −<span class="html-italic">z</span> + 1; (viii) −<span class="html-italic">x</span> + 1, −<span class="html-italic">y</span>, −<span class="html-italic">z</span> + 1.</p> Full article ">Figure 9
<p>The structural models in real space for the Al<sub>13</sub>Fe<sub>3</sub> and Al<sub>13</sub>Fe<sub>4</sub> phases are presented: the orientation of cell edges (<b>a</b>) and the unit cell (<b>b</b>) (The saffron yellow atoms represent the Fe atoms and the blue atoms represent the Al atoms).</p> Full article ">Figure 10
<p>The precession images of intergrowth crystals are presented as follows: (<b>a</b>) Al<sub>13</sub>Fe<sub>3</sub>-1(0<span class="html-italic">kl</span>), where the red and blue circles depict the crystal planes of the Al<sub>13</sub>Fe<sub>3</sub>-1 phase and Al<sub>13</sub>Fe<sub>4</sub> phase, respectively. (<b>b</b>,<b>c</b>) show the same phase, with circles representing (<b>b</b>) Al<sub>13</sub>Fe<sub>3</sub>-1(<span class="html-italic">h</span>0<span class="html-italic">l</span>) and (<b>c</b>) Al<sub>13</sub>Fe<sub>3</sub>-1(<span class="html-italic">hk</span>0). Similarly, (<b>d</b>) Al<sub>13</sub>Fe<sub>4</sub>(0<span class="html-italic">kl</span>) is shown, where the red and blue circles depict the crystal planes of the Al<sub>13</sub>Fe<sub>4</sub> phase and Al<sub>13</sub>Fe<sub>3</sub>-1 phase, respectively. (<b>e</b>,<b>f</b>) show the same phase, with circles representing (<b>e</b>) Al<sub>13</sub>Fe<sub>4</sub>(<span class="html-italic">h</span>0<span class="html-italic">l</span>) and (<b>f</b>) Al<sub>13</sub>Fe<sub>4</sub>(<span class="html-italic">hk</span>0).</p> Full article ">Figure 11
<p>The precession images of intergrowth crystals are presented as follows: (<b>a</b>) Al<sub>13</sub>Fe<sub>3</sub>-2(0<span class="html-italic">kl</span>), where the green and purple circles depict the crystal planes of the Al<sub>13</sub>Fe<sub>3</sub>-2 phase and Al<sub>13</sub>Fe<sub>4</sub> phase, respectively. (<b>b</b>,<b>c</b>) show the same phase, with circles representing (<b>b</b>) Al<sub>13</sub>Fe<sub>3</sub>-2(<span class="html-italic">h</span>0<span class="html-italic">l</span>) and (<b>c</b>) Al<sub>13</sub>Fe<sub>3</sub>-2(<span class="html-italic">hk</span>0). Similarly, (<b>d</b>) Al<sub>13</sub>Fe<sub>4</sub>(0<span class="html-italic">kl</span>) is shown, where the red and blue circles depict the crystal planes of the Al<sub>13</sub>Fe<sub>4</sub> phase and Al<sub>13</sub>Fe<sub>3</sub>-2 phase, respectively. (<b>e</b>,<b>f</b>) show the same phase, with circles representing (e) Al<sub>13</sub>Fe<sub>4</sub>(<span class="html-italic">h</span>0<span class="html-italic">l</span>) and (<b>f</b>) Al<sub>13</sub>Fe<sub>4</sub>(<span class="html-italic">hk</span>0).</p> Full article ">Figure 12
<p>Stereographic projection of the orientation relationships between Al<sub>13</sub>Fe<sub>3</sub> and Al<sub>13</sub>Fe<sub>4</sub> based on the Al<sub>13</sub>Fe<sub>3</sub> phase: (<b>a</b>) the orientation relationships between crystal planes, (<b>b</b>) the orientation relationships between crystal directions.</p> Full article ">Figure 13
<p>(<b>a</b>): (033) surface of the Al<sub>13</sub>Fe<sub>3</sub> phase; (<b>b</b>): (<math display="inline"><semantics> <mrow> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>) surface of the Al<sub>13</sub>Fe<sub>4</sub> phase; (<b>c</b>): the atomic interface model of Al<sub>13</sub>Fe<sub>3</sub>(033)/Al<sub>13</sub>Fe<sub>4</sub>(<math display="inline"><semantics> <mrow> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>) (The golden atoms represent the Fe atoms and the blue atoms represent the Al atoms).</p> Full article ">
14 pages, 3337 KiB  
Article
Study of Assimilation of Cored Wire into Liquid Steel Baths
by Edgar-Ivan Castro-Cedeno, Julien Jourdan, Jonathan Martens, Jean-Pierre Bellot and Alain Jardy
Metals 2024, 14(4), 462; https://doi.org/10.3390/met14040462 - 15 Apr 2024
Viewed by 1143
Abstract
Cored wire is a widespread technology used for performing additions into liquid metal baths as an alternative to bulk-additions. A laboratory-scale study was performed in which the kinetics of assimilation of cored wire in liquid steel baths were studied. An original dataset of [...] Read more.
Cored wire is a widespread technology used for performing additions into liquid metal baths as an alternative to bulk-additions. A laboratory-scale study was performed in which the kinetics of assimilation of cored wire in liquid steel baths were studied. An original dataset of positions of the wire/melt interface of cored wire as a function of the time and steel bath temperature was produced. The dataset was compared against results of simulations made with a transient 1D (radial) thermal model of the assimilation of cored wire, and demonstrated reasonable agreement. Hence, this paper provides a dataset that can be used as a resource for the validation of future developments in the field of modeling cored wire injection into liquid metal baths. Full article
(This article belongs to the Topic Advanced Processes in Metallurgical Technologies)
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<p>Schematic representation of possible assimilation routes for cored wires immersed into liquid steel.</p> Full article ">Figure 2
<p>Schematic of experimental apparatus used for the experiments of the immersion of cored wire into liquid steel baths.</p> Full article ">Figure 3
<p>Photograph of extraction of wire V from the bath during the immersion experiment. Notice the shell of solidified bath material around the cored wire.</p> Full article ">Figure 4
<p>Cored wires with solidified metal shell after wire immersion experiments at various initial steel bath temperatures: (<b>a</b>) wire I (1818 K); (<b>b</b>) wire II (1823 K); (<b>c</b>) wire III (1823 K); (<b>d</b>) wire IV (1853 K); (<b>e</b>) wire V (1858 K); (<b>f</b>) wire VI (1873 K).</p> Full article ">Figure 5
<p>Schematic of the wire + solidified shell ensemble post-immersion. The immersion times of radial sections throughout the length of the wire can be estimated using Equations (1)–(4), using the measured length of wire before immersion (<span class="html-italic">L<sub>bef</sub></span>), length of wire after immersion (<span class="html-italic">L<sub>aft</sub></span>), length of non-immersed part of the wire (<span class="html-italic">L<sub>nim</sub></span>), and timestamps of camera recordings (<span class="html-italic">t<sub>im</sub></span>, <span class="html-italic">t<sub>hold</sub></span>, <span class="html-italic">t<sub>ext</sub></span>).</p> Full article ">Figure 6
<p>Schematic of the procedure for measuring the diameter of the wire + shell ensemble post-immersion. To obtain estimations of mean diameter, the specimen was rotated for each of the 10 individual measurements taken at fixed wire lengths.</p> Full article ">Figure 7
<p>Comparison of estimations of radius of the wire + shell ensemble (experimental) and computed evolution of the wire/melt interface (simulations): (<b>a</b>) wire <span class="html-italic">I</span> (1818 K); (<b>b</b>) wire <span class="html-italic">II</span> (1823 K), (<b>c</b>) wire <span class="html-italic">III</span> (1823 K); (<b>d</b>) wire <span class="html-italic">IV</span> (1853 K); (<b>e</b>) wire <span class="html-italic">V</span> (1858 K); (<b>f</b>) wire <span class="html-italic">VI</span> (1873 K).</p> Full article ">Figure 8
<p>Deviations between estimations of thickness of the casing + shell ensemble (experimental) and thicknesses from the computed evolution of the wire/melt interface (simulations).</p> Full article ">Figure 9
<p>Extrapolations of wire assimilation times for cored wire used in the lab-scale experiments, as a function of bath temperature, convective heat transfer coefficient, and magnitude of contact thermal resistances.</p> Full article ">
18 pages, 3661 KiB  
Article
Planning Mechanical Behavior of A356 Alloy Wheels by Using Distinct Heat Treatments
by Adriano L. Tonetti, Wislei R. Osório, Ausdinir D. Bortolozo and Giovana S. Padilha
Metals 2024, 14(4), 461; https://doi.org/10.3390/met14040461 - 13 Apr 2024
Viewed by 1155
Abstract
The aim of this investigation concerns evaluating the mechanical strength and microhardness values of A356 alloy samples in distinctive heat treatments, including those commonly applied to automotive wheels. It is recognized that A356 and Al-Si-based alloys exhibit considerable versatility across numerous industrial applications. [...] Read more.
The aim of this investigation concerns evaluating the mechanical strength and microhardness values of A356 alloy samples in distinctive heat treatments, including those commonly applied to automotive wheels. It is recognized that A356 and Al-Si-based alloys exhibit considerable versatility across numerous industrial applications. The mechanical behavior obtained is intimately associated with different operational parameters (e.g., cooling rates, solution treatment, quenching, and artificial aging). In this study, a group of samples are quenched at 30, 60 and 80 °C. Another set is quenched and subsequently aged at three different temperatures, i.e., 180, 200, and 220 °C for 5 h, and mechanical responses are compared. Microstructural characterization, X-ray diffraction (XRD) analysis, tensile testing, and microhardness measurements are carried out. Using the Rietveld data and based on the modified Williamson–Hall method, the microstrains, crystallite size, and dislocation densities are calculated. Based on this, the resulting mechanical strengths from distinctive quenching and aging are understood. It was found that there exists a “quasi-optimal range” of operational parameters involving different A356 alloy treatments, which vary depending on the manufacturing route. Considering A356 alloy wheels, the planning of the powder coat treatment before or after T6 treating provides better mechanical properties and ductility. Full article
(This article belongs to the Topic Alloys and Composites Corrosion and Mechanical Properties)
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<p>Representative temperature vs. time profile for the heating rate (I), solution treating (II), quenching (III), and three distinctive T6 heat treatments for the examined A356 alloy samples (IV). All adopted quenching are carried out in 120 s, and all aging treatments are performed for 180 min (3 h).</p> Full article ">Figure 2
<p>Experimental engineering stress vs. strain curves of A356 alloy after distinct (<b>a</b>) quenching and (<b>b</b>) aging. The as-cast (non-treated) results are also depicted.</p> Full article ">Figure 3
<p>Relation among UTS, YS, and ductility in distinctive (<b>a</b>) quenching and (<b>b</b>) aging temperatures.</p> Full article ">Figure 4
<p>Microhardness obtained after three different quenching and three different aging temperatures.</p> Full article ">Figure 5
<p>Correlation between yield strength (YS) and microhardness and ultimate tensile strength (UTS) and microhardness by using (<b>a</b>) exponential and (<b>b</b>) linear regression equations based on Krishna et al.’ [<a href="#B43-metals-14-00461" class="html-bibr">43</a>] and Cerri and Ghio’s [<a href="#B42-metals-14-00461" class="html-bibr">42</a>] equations.</p> Full article ">Figure 6
<p>Correlation between YS and UTS with microhardness fitted by (<b>a</b>) Cerri and Ghio’s [<a href="#B42-metals-14-00461" class="html-bibr">42</a>] linear equation and (<b>b</b>) YS and UTS with microhardness (in MPa) considering the investigations of Abdulwahab et al. [<a href="#B40-metals-14-00461" class="html-bibr">40</a>] and Rometsch and Schaffer [<a href="#B39-metals-14-00461" class="html-bibr">39</a>].</p> Full article ">Figure 7
<p>XRD patterns of the A356 alloy samples in condition: (<b>a</b>) as-cast, (<b>b</b>) water-quenched at 30 °C only, (<b>c</b>) artificially aged (T6) at 180 °C for 3 h; and (<b>d</b>) the calculated crystallite size and remaining strain determined by using Rietveld data (gray lines).</p> Full article ">Figure 8
<p>(<b>a</b>) Experimental DSC curves obtained by using 10 °C/min of the A356 alloy samples: (<b>a</b>) quenched at 30 °C and artificially aged at 220 °C and (<b>b</b>) the as-cast condition; (<b>b</b>) A typical XRD pattern corresponding with the as-cast A356 alloy sample, evidencing (inset) remaining Mg<sub>2</sub>Si particles.</p> Full article ">Figure 9
<p>Typical SEM micrographs of the A356 alloy samples of the as-cast samples (<b>a</b>), and quenched at (<b>b</b>) 30 °C, (<b>c</b>) 60 °C, and (<b>d</b>) 80 °C, evidencing Fe-rich particles and Mg<sub>2</sub>Si particles (as-cast sample).</p> Full article ">Figure 10
<p>(<b>a</b>) A typical micrograph of the fractured surface of the specimens examined, SEM micrographs of the samples quenched at (<b>b</b>) 30 °C and (<b>c</b>) 60 °C; and the samples after aging at (<b>d</b>) 180 °C, (<b>e</b>) 200 °C, and (<b>f</b>) 220 °C. The white arrows indicate the cleavage planes and yellow and red dashed circles represent tendon fossa (dimples sites) [<a href="#B50-metals-14-00461" class="html-bibr">50</a>] and tearing edges, respectively. Yellow arrows depict the constituted porosity/voids.</p> Full article ">
16 pages, 5550 KiB  
Article
Study on Recovery of Lithium from Lithium-Containing Aluminum Electrolyte
by Rui Ji, Xi Cui, Wenzheng Zhang, Shichao Wang, Mingliang Yang and Tao Qu
Metals 2024, 14(4), 460; https://doi.org/10.3390/met14040460 - 13 Apr 2024
Viewed by 978
Abstract
The current process of recovering lithium from wasted aluminum electrolyte mostly entails extracting lithium from lithium-containing aluminum electrolyte by acid leaching and dissipation. Aiming at the disadvantages of the existing treatment method, such as the long process flow, environmental pollution, poor working environment, [...] Read more.
The current process of recovering lithium from wasted aluminum electrolyte mostly entails extracting lithium from lithium-containing aluminum electrolyte by acid leaching and dissipation. Aiming at the disadvantages of the existing treatment method, such as the long process flow, environmental pollution, poor working environment, etc., we propose a new technology to extract lithium from the wasted aluminum electrolyte and systematically investigate the effects of raw material particle size, holding time, temperature and other factors on the recovery of lithium. The results show that the better process conditions for the recovery of lithium are as follows: the raw material particle size is 75~150 μm, the additive is CaCl2, the mass ratio of calcium chloride to lithium-containing aluminum electrolyte is 3:5, the reaction temperature is 1473 K, and the holding time is 3 h. After the product of the reaction is crushed and leaching is carried out by using deionized water (pH = 6.8), the temperature of the leaching is 368 K, the leaching time is 3 h, and the solid–liquid ratio is 1/3, and the leaching rate of Li can be up to 75.1%. In addition, the purity of the recovered AlF3 is more than 92.7%. This process realizes the comprehensive and efficient use of lithium-containing aluminum electrolyte and provides a new idea for the development of lithium extraction technology from lithium-containing aluminum electrolyte. Full article
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<p>XRD pattern of lithium-containing aluminum electrolyte.</p> Full article ">Figure 2
<p>SEM-EDS analysis: (<b>a</b>) SEM image of Li-containing aluminum electrolyte, (<b>b</b>) element distribution, (<b>c</b>) elemental analysis.</p> Full article ">Figure 3
<p>Structure of Na<sub>2</sub>LiAlF<sub>6</sub> substance.</p> Full article ">Figure 4
<p>Variation in Gibbs free energy with reaction temperature at 1 atm, (<b>a</b>) Li<sub>3</sub>AlF<sub>6</sub>-related reactions, (<b>b</b>) Na<sub>3</sub>AlF<sub>6</sub>-related reactions, (<b>c</b>) reactions (11)–(14).</p> Full article ">Figure 5
<p>High-temperature melting furnace: (<b>a</b>) description of the experimental equipment, (<b>b</b>) physical drawing of the experimental equipment.</p> Full article ">Figure 6
<p>XRD images of chlorinated roasting products at different temperatures.</p> Full article ">Figure 7
<p>XRD images of condensate fractions from 1473 K and 1573 K CaCl<sub>2</sub> roasting under 1 atm conditions.</p> Full article ">Figure 8
<p>Li leaching rate of chlorinated roasting products at different temperatures.</p> Full article ">Figure 9
<p>Effect of different chlorination roasting temperature and holding time on the leaching rate of Li.</p> Full article ">Figure 10
<p>Effect of different raw material particle size dimensions and different CaCl<sub>2</sub> percentages on the leaching rate of Li.</p> Full article ">Figure 11
<p>Effect of leaching conditions on leaching rate of Li.</p> Full article ">Figure 12
<p>Physical phase analysis of leaching slag.</p> Full article ">
15 pages, 12435 KiB  
Article
Additive Manufacturing for Rapid Sand Casting: Mechanical and Microstructural Investigation of Aluminum Alloy Automotive Prototypes
by Silvia Cecchel and Giovanna Cornacchia
Metals 2024, 14(4), 459; https://doi.org/10.3390/met14040459 - 13 Apr 2024
Viewed by 1121
Abstract
The automotive industry is undergoing a rapid evolution to meet today’s challenges; therefore, continuous innovation and product development are needed. Validation tests on prototypes play a crucial role in moving new components into industrial production. There is also a pressing need for faster [...] Read more.
The automotive industry is undergoing a rapid evolution to meet today’s challenges; therefore, continuous innovation and product development are needed. Validation tests on prototypes play a crucial role in moving new components into industrial production. There is also a pressing need for faster prototyping processes. In this context, rapid sand casting (RSC), based on additive manufacturing technology, offers a promising solution for a quick production of sand molds. While this technology is already employed in the industry, the need to deepen the general understanding of its impact on the casting properties is still a relevant item. In this study, different geometries of automotive prototypes made of aluminum EN AC 42100-T6 alloy were experimentally analyzed. Microstructural examinations, tensile tests, and fractography and porosity analyses were conducted. The findings demonstrate the considerable potential of RSC, giving, in general, high mechanical properties. A comparative analysis with prototypes produced through traditional sand casting revealed similar results, with RSC exhibiting superior yield strength and stress at brake. However, both technologies revealed a reduced elongation percentage, as expected. Future efforts will focus on standardizing the RSC process to enhance ductility levels. Full article
(This article belongs to the Special Issue Light Alloy and Its Application (2nd Edition))
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<p>Component geometry and overall dimensions.</p> Full article ">Figure 2
<p>Examples of SDAS measurements. Red line represents L length.</p> Full article ">Figure 3
<p>Tensile sample location and nomenclature. The areas of extraction are identified by the highlighted red and green specimens, named from A to D.</p> Full article ">Figure 4
<p>X-ray image of the analyzed components. Porosities are indicated by red arrows. Tensile specimens location is identified by red and green colors.</p> Full article ">Figure 5
<p>Macrostructures of EN AC 42100-T6 RSC components, with porosity ratio percentage, porosity average (avg) size, and SDAS values for each section from the shoulders of the tensile specimens analyzed.</p> Full article ">Figure 6
<p>Microstructure at different magnifications of EN AC 42100-T6 RSC samples.</p> Full article ">Figure 7
<p>Microstructure at different magnifications of EN AC 42100-T6 samples produced using SC and RSC technology.</p> Full article ">Figure 8
<p>SEM images and EDS analysis of EN AC 42100-T6 samples produced using SC (<b>a</b>) and RSC technology (<b>b</b>,<b>c</b>).</p> Full article ">Figure 9
<p>Tensile test curve for EN AC 42100-T6 for the different RSC components analyzed. Only the curve with the highest yield strength for each component was reported for comparison purposes.</p> Full article ">Figure 10
<p>SEM on the fracture surfaces of the tensile samples made of EN AC 42100-T6 using RSC, at different magnifications. Red arrows indicate the location of the subsequent magnifications. (<b>a</b>) highest elongation percentage and (<b>b</b>) lowest elongation percentage.</p> Full article ">Figure 11
<p>SEM on the fracture surfaces of the tensile samples made of EN AC 42100-T6 using sand casting, at different magnifications. Red arrows indicate the location of the subsequent magnifications. (<b>a</b>) highest elongation percentage and (<b>b</b>) lowest elongation percentage.</p> Full article ">
25 pages, 19567 KiB  
Article
Evaluation of Energy Utilization Efficiency and Optimal Energy Matching Model of EAF Steelmaking Based on Association Rule Mining
by Lingzhi Yang, Zhihui Li, Hang Hu, Yuchi Zou, Zeng Feng, Weizhen Chen, Feng Chen, Shuai Wang and Yufeng Guo
Metals 2024, 14(4), 458; https://doi.org/10.3390/met14040458 - 12 Apr 2024
Cited by 2 | Viewed by 902
Abstract
In the iron and steel industry, evaluating the energy utilization efficiency (EUE) and determining the optimal energy matching mode play an important role in addressing increasing energy depletion and environmental problems. Electric Arc Furnace (EAF) steelmaking is a typical short crude steel production [...] Read more.
In the iron and steel industry, evaluating the energy utilization efficiency (EUE) and determining the optimal energy matching mode play an important role in addressing increasing energy depletion and environmental problems. Electric Arc Furnace (EAF) steelmaking is a typical short crude steel production route, which is characterized by an energy-intensive fast smelting rhythm and diversified raw charge structure. In this paper, the energy model of the EAF steelmaking process is established to conduct an energy analysis and EUE evaluation. An association rule mining (ARM) strategy for guiding the EAF production process based on data cleaning, feature selection, and an association rule (AR) algorithm was proposed, and the effectiveness of this strategy was verified. The unsupervised algorithm Auto-Encoder (AE) was adopted to detect and eliminate abnormal data, complete data cleaning, and ensure data quality and accuracy. The AE model performs best when the number of nodes in the hidden layer is 18. The feature selection determines 10 factors such as the hot metal (HM) ratio and HM temperature as important data features to simplify the model structure. According to different ratios and temperatures of the HM, combined with k-means clustering and an AR algorithm, the optimal operation process for the EUE in the EAF steelmaking under different smelting modes is proposed. The results indicated that under the conditions of a low HM ratio and low HM temperature, the EUE is best when the power consumption in the second stage ranges between 4853 kWh and 7520 kWh, the oxygen consumption in the second stage ranges between 1816 m3 and 1961 m3, and the natural gas consumption ranges between 156 m3 and 196 m3. Conversely, under the conditions of a high HM ratio and high HM temperature, the EUE tends to decrease, and the EUE is best when the furnace wall oxygen consumption ranges between 4732 m3 and 5670 m3, and the oxygen consumption in the second stage ranges between 1561 m3 and 1871 m3. By comparison, under different smelting modes, the smelting scheme obtained by the ARM has an obvious effect on the improvement of the EUE. With a high EUE, the improvement of the A2B1 smelting mode is the most obvious, from 24.7% to 53%. This study is expected to provide technical ideas for energy conservation and emission reduction in the EAF steelmaking process in the future. Full article
(This article belongs to the Special Issue Green and Intelligent Steelmaking Technologies with Low Carbon Emissions)
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<p>Energy balance diagram of EAF steelmaking process.</p> Full article ">Figure 2
<p>Research framework of this study.</p> Full article ">Figure 3
<p>Association rules mining flow diagram.</p> Full article ">Figure 4
<p>The statistical interface of energy data.</p> Full article ">Figure 5
<p>The energy composition and flow of the furnace number ‘20306983’.</p> Full article ">Figure 6
<p>Correlation heat map between variables.</p> Full article ">Figure 7
<p>Feature selection results for the extreme gradient boosting (XGBoost) model.</p> Full article ">Figure 8
<p>The k-means clustering algorithm discretization result. Note that ELW, EL2, EL4, etc. in the figure represent the clustering and discretization results of different features. Each symbol represents power consumption, power consumption in the second stage, power consumption in the fourth stage, furnace wall oxygen consumption, furnace wall oxygen consumption in the first stage, furnace wall oxygen consumption in the second stage, carbon powder weight, natural gas consumption, and EUE. The line of the different colors represents the labels of each cluster, and the numbers below represent the boundaries of each grouping.</p> Full article ">Figure 9
<p>Optimal energy matching model of EAF based on ARM.</p> Full article ">Figure 10
<p>EUE stacked bar chart under different smelting modes.</p> Full article ">
18 pages, 2936 KiB  
Article
Numerical Simulation of Cathode Nodule Local Effects
by Xiaoyu Wang, Chun Li and Jun Tie
Metals 2024, 14(4), 457; https://doi.org/10.3390/met14040457 - 12 Apr 2024
Viewed by 810
Abstract
As one of the main factors decreasing current efficiency and product quality, the growth of nodules deserves attention in the copper electrorefining process. Three-dimensional (3D) Finite Element Method models combining tertiary current distribution and fluid flow were established in this study to investigate [...] Read more.
As one of the main factors decreasing current efficiency and product quality, the growth of nodules deserves attention in the copper electrorefining process. Three-dimensional (3D) Finite Element Method models combining tertiary current distribution and fluid flow were established in this study to investigate the details of the growth of columnar nodules, including the electrolyte flow around the nodule and its effects. Compared with an inert nodule, a significant impact of the electrochemical reaction of an active nodule has been observed on the fluid flow, which may be one of the reasons for the formation of small nodule clusters on the cathode. Furthermore, the local current density is not even on the nodule surface under the comprehensive influence of local electrolyte flow, local overvoltage, and the angle with the anode surface. Thus, the head of an active nodule grows faster than the root, and the upper parts grow faster than the lower parts, leading to asymmetric growth of the nodules. Full article
(This article belongs to the Special Issue Modeling and Simulation of Multiphase Transport Phenomena in Process Metallurgy)
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Graphical abstract
Full article ">Figure 1
<p>The geometric model and mesh partition used in this paper (<b>a</b>,<b>b</b>) are the three-dimensional geometric models of non-nodules and nodules of φ5 × 23 mm, respectively; (<b>c</b>) is the mesh partition under the condition of nodules, and the color legend shows the grid unit size at the nodules. (<b>d</b>) is the boundary layer mesh at the cathode surface, and the anode is consistent.</p> Full article ">Figure 2
<p>The electrolyte flow velocities in electrolytic cells. (<b>a</b>) represents a condition of no nodules, (<b>b</b>) represents an electrolytic cell with inert nodules, (<b>c</b>) represents a real nodule electrolytic cell with a length of 23 mm, and (<b>d</b>) represents a real nodule electrolytic cell with a length of 15 mm. (<b>a</b>–<b>d</b>) are composed of cross-sectional views (YZ cross-section) and longitudinal views (XZ cross-section) to demonstrate the internal fluid dynamics within the respective electrolytic cell configurations. A–D cross-sections are shown in 3D on the right.</p> Full article ">Figure 3
<p>The average of the z-direction velocity vectors at 20 mm, (<b>a</b>) above nodule (red lines in <a href="#metals-14-00457-f002" class="html-fig">Figure 2</a>), (<b>b</b>) below nodule (orange lines in <a href="#metals-14-00457-f002" class="html-fig">Figure 2</a>); 0 mm, 25 mm, 50 mm, and 75 mm represent cross-sections (XZ cross-sections) at distances of 0 mm, 25 mm, 50 mm, and 75 mm from the center of the nodule, below nodule, and under a different nodule. A–D represent the cross-section of the 3D diagram on the right side of <a href="#metals-14-00457-f002" class="html-fig">Figure 2</a>.</p> Full article ">Figure 3 Cont.
<p>The average of the z-direction velocity vectors at 20 mm, (<b>a</b>) above nodule (red lines in <a href="#metals-14-00457-f002" class="html-fig">Figure 2</a>), (<b>b</b>) below nodule (orange lines in <a href="#metals-14-00457-f002" class="html-fig">Figure 2</a>); 0 mm, 25 mm, 50 mm, and 75 mm represent cross-sections (XZ cross-sections) at distances of 0 mm, 25 mm, 50 mm, and 75 mm from the center of the nodule, below nodule, and under a different nodule. A–D represent the cross-section of the 3D diagram on the right side of <a href="#metals-14-00457-f002" class="html-fig">Figure 2</a>.</p> Full article ">Figure 4
<p>Distribution of current density of the surface of the cathode; (<b>a</b>) the distribution of the surface of nodule current density and the angle θ to the anode surface along the nodule, and the arrows represent current density vectors, (<b>b</b>) the distribution of current density along the head and root of the nodule, local 1 includes the U1 and D1 locals; local 2 includes the U2 and D2 locals.</p> Full article ">Figure 5
<p>The variation in local overvoltage and copper ion flux with the distance from the cathode surface to the head and root edges of the nodule. (<b>a</b>) Local overvoltage, (<b>b</b>) copper ion flux; local 1 in the figure includes locals U1 and D1; local 2 includes locals U2 and D2; local 3 includes locals U3 and D3.</p> Full article ">Figure 6
<p>Local flow rate of electrolyte near the cross-section of the center of a nodule magnification view, the arrows represent the direction of electrolyte flow.</p> Full article ">Figure 7
<p>Copper ion concentration distribution near the nodule. (<b>a</b>) Nodule center cross section (XZ section), (<b>b</b>) the variation trend of the nodule at the upper and lower edges with respect to the cathode surface distance; local 1 includes the U1 and D1 locals, local 2 includes the U2 and D2 locals, and local 3 includes the U3 and D3 locals.</p> Full article ">Figure 8
<p>Contour streamlines of copper ion concentration and electrolyte velocity plots for various locals of YZ cross-section for the nodule size φ5 × 23 mm for locals 1, 2, and 3; (<b>a</b>,<b>d</b>) are for the cross-section of 2 mm from the cathode surface; (<b>b</b>,<b>e</b>) are for the cross-section of 10 mm from the cathode surface; (<b>c</b>,<b>f</b>) are for the cross-section of 22.5 mm from the cathode surface.</p> Full article ">
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