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Volume 14
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Metals, Volume 14, Issue 8 (August 2024) – 84 articles

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16 pages, 4869 KiB  
Article
An Efficient and Stable MXene-Immobilized, Cobalt-Based Catalyst for Hydrogen Evolution Reaction
by Wei Guo, Buxiang Wang and Qing Shu
Metals 2024, 14(8), 922; https://doi.org/10.3390/met14080922 - 14 Aug 2024
Abstract
Hydrogen (H2) is considered to be the best carbon-free energy carrier that can replace fossil fuels because of its high energy density and the advantages of not producing greenhouse gases and air pollutants. As a green and sustainable method for hydrogen [...] Read more.
Hydrogen (H2) is considered to be the best carbon-free energy carrier that can replace fossil fuels because of its high energy density and the advantages of not producing greenhouse gases and air pollutants. As a green and sustainable method for hydrogen production, the electrochemical hydrogen evolution reaction (HER) has received widespread attention. Currently, it is a great challenge to prepare economically stable electrocatalysts for the HER using non-precious metals. In this study, a Co/Co3O4/Ti3C2Tx catalyst was synthesized by supporting Co/Co3O4 with Ti3C2Tx. The results show that Co/Co3O4/Ti3C2Tx has excellent HER activity and durability in 1 mol L−1 KOH, and the overpotential and Tafel slope at 10 mA·cm−2 were 87 mV and 61.90 mV dec−1, respectively. The excellent HER activity and stability of Co/Co3O4/Ti3C2Tx can be explained as follows: Ti3C2Tx provides a stable skeleton and a large number of attachment sites for Co/Co3O4, thus exposing more active sites; the unique two-dimensional structure of Ti3C2Tx provides an efficient conductive network for rapid electron transfer between the electrolyte and the catalyst during electrocatalysis; Co3O4 makes the Co/Co3O4/Ti3C2Tx catalyst more hydrophilic, which can accelerate the release rate of bubbles; Co/Co3O4 can accelerate the adsorption and deionization of H2O to synthesize H2. This study provides a new approach for the design and preparation of low-cost and high-performance HER catalysts. Full article
(This article belongs to the Special Issue Deposition of Metals and Their Application in Catalytic Processes of Energy Interest (2nd Edition))
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<p>Schematic diagram of Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span> preparation.</p> Full article ">Figure 2
<p>FT-IR spectrum of Ti<sub>3</sub>C<sub>2.</sub></p> Full article ">Figure 3
<p>SEM images of Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span> and Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>. (<b>a</b>) Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>; (<b>b</b>) Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>.</p> Full article ">Figure 4
<p>Elemental analysis diagram of Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>. (<b>a</b>) SEM diagram of Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>; (<b>b</b>) EDS mapping of Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span> and (<b>c</b>) Co element, (<b>d</b>) O element, (<b>e</b>) C element and (<b>f</b>) Ti element; (<b>g</b>) atomic ratios of Co, O, C and Ti elements.</p> Full article ">Figure 5
<p>TEM images of (<b>a</b>) Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span> and (<b>b</b>) Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>.</p> Full article ">Figure 6
<p>XRD patterns of Ti<sub>3</sub>AlC<sub>2</sub>, Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span> and Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>. (<b>a</b>) Ti<sub>3</sub>AlC<sub>2</sub> and Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>; (<b>b</b>) Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>.</p> Full article ">Figure 7
<p>XPS diagram of Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>. (<b>a</b>) XPS full spectrum, (<b>b</b>) C1s spectrum, (<b>c</b>) Co2p spectrum and (<b>d</b>) O1s spectrum.</p> Full article ">Figure 7 Cont.
<p>XPS diagram of Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>. (<b>a</b>) XPS full spectrum, (<b>b</b>) C1s spectrum, (<b>c</b>) Co2p spectrum and (<b>d</b>) O1s spectrum.</p> Full article ">Figure 8
<p>Contact angles of Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>, Co and Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>. (<b>a</b>) Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>, (<b>b</b>) Co and (<b>c</b>) Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span><sub>.</sub></p> Full article ">Figure 9
<p>N<sub>2</sub> adsorption–desorption curves and pore size distributions of Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>, Co and Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>. (<b>a</b>) N<sub>2</sub> adsorption–desorption curve; (<b>b</b>) pore size distribution (0–600 nm).</p> Full article ">Figure 10
<p>LSV diagram of (<b>a</b>) Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>, Co and Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>; (<b>b</b>) Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span> (stability test) in 0.5 mol L<sup>−1</sup> H<sub>2</sub>SO<sub>4</sub> electrolyte.</p> Full article ">Figure 11
<p>(<b>a</b>): LSV diagram of Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>, Co and Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span> in 1 mol L<sup>−1</sup> KOH electrolyte; CV results of (<b>b</b>) Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>, (<b>c</b>) Co and (<b>d</b>) Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span> under different scanning rates (10, 20, 30, 40, 50, 60, 70, 80, 90 and 100 mV s<sup>−1</sup>); (<b>e</b>) C<sub>dl</sub> value and (<b>f</b>) Tafel diagram of Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>, Co and Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span>; (<b>g</b>) LSV diagram of Co/Co<sub>3</sub>O<sub>4</sub>/Ti<sub>3</sub>C<sub>2</sub>T<span class="html-italic"><sub>x</sub></span> (stability test) in 1 mol L<sup>−1</sup> KOH electrolyte.</p> Full article ">
13 pages, 18656 KiB  
Article
Evolutions on Microstructure and Impact Toughness of G115 Steel after Long-Term Aging at 700 °C
by Jianming Yu, Shaohai Ma, Kui Liang, Kai Yan, Xisheng Yang and Shuli Zhang
Metals 2024, 14(8), 921; https://doi.org/10.3390/met14080921 - 14 Aug 2024
Abstract
The microstructure and impact toughness evolution of G115 steel after long-term (ranging from 500 h to 10,000 h) aging at 700 °C were investigated in this study. The results showed that the microstructure of the G115 steel evolved from a finer-grained matrix with [...] Read more.
The microstructure and impact toughness evolution of G115 steel after long-term (ranging from 500 h to 10,000 h) aging at 700 °C were investigated in this study. The results showed that the microstructure of the G115 steel evolved from a finer-grained matrix with minor precipitates to a coarse-grained matrix with more precipitate with aging time, presenting a decrease in the local deformation degree in the matrix. The impact toughness of the steel decreased with aging time, presenting the largest decline at the initial aging times. The decrease in impact toughness was attributed to the coarsening of precipitates (M23C6 and Laves phase) in the steel matrix. The stable impact toughness during the whole aging process (from 500 h to 10,000 h) should be related to the comprehensive effects, including the precipitation of the Laves phase, the increase in high-angle grain boundaries, and the softening of the metal matrix. Full article
(This article belongs to the Special Issue Corrosion of Metals: Behaviors and Mechanisms)
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<p>(<b>a</b>) Schematic diagram of the impact specimen (unit: mm). (<b>b</b>) The ZBC2452-CD Charpy impact tester.</p> Full article ">Figure 2
<p>Metallographic morphology of G115 steel with different aging times at 700 °C: (<b>a</b>) BM, (<b>b</b>) 500 h, (<b>c</b>) 1000 h, (<b>d</b>) 3000 h, (<b>e</b>) 5000 h, (<b>f</b>) 10,000 h; PAGs are marked in images.</p> Full article ">Figure 3
<p>Average grain size of PAG in G115 steel with different aging times at 700 °C.</p> Full article ">Figure 4
<p>SEM morphologies of G115 steel with different aging times at 700 °C: (<b>a</b>) BM, (<b>b</b>) 500 h, (<b>c</b>) 1000 h, (<b>d</b>) 3000 h, (<b>e</b>) 5000 h, (<b>f</b>) 10,000 h.</p> Full article ">Figure 5
<p>Magnified back-scattered electron morphologies of G115 steel with different aging times at 700 °C: (<b>a</b>) BM, (<b>b</b>) 500 h, (<b>c</b>) 1000 h, (<b>d</b>) 3000 h, (<b>e</b>) 5000 h, (<b>f</b>) 10,000 h.</p> Full article ">Figure 6
<p>EBSD maps of G115 steel with different aging times at 700 °C: (<b>a</b>) BM, (<b>b</b>) 3000 h, (<b>c</b>) 5000 h, (<b>d</b>) 10,000 h, (<b>1</b>) inverse pole figures and (<b>2</b>) grain orientation spread maps.</p> Full article ">Figure 7
<p>Grain boundary distribution maps of the G115 steel with different aging times at 700 °C: (<b>a</b>) BM, (<b>b</b>) 3000 h, (<b>c</b>) 5000 h, (<b>d</b>) 10,000 h.</p> Full article ">Figure 8
<p>Kernel average misorientation maps of the G115 steel with different aging times at 700 °C: (<b>a</b>) BM, (<b>b</b>) 500 h, (<b>c</b>) 1000 h, (<b>d</b>) 3000 d, (<b>e</b>) 5000 h, (<b>f</b>) 10,000 h.</p> Full article ">Figure 9
<p>TEM images of G115 steel with different aging times at 700 °C: (<b>a</b>) BM, (<b>b</b>) 3000 h, (<b>c</b>) 5000 h, (<b>d</b>) 10,000 h; EDS results were obtained at the local area.</p> Full article ">Figure 10
<p>Impact energy evolution of the G115 steel after long-term aging at 700 °C.</p> Full article ">Figure 11
<p>SEM micrographs of the impact fracture for G115 steel with different aging times at 700 °C: ((<b>a</b>) radiation area, (<b>b</b>) fiber region) BM, ((<b>c</b>) radiation area, (<b>d</b>) fiber region) 500 h, ((<b>e</b>) radiation area, (<b>f</b>) fiber region) 3000 h, ((<b>g</b>) radiation area, (<b>h</b>) fiber region) 5000 h, and ((<b>i</b>) radiation area, (<b>j</b>) fiber region)) 10,000 h.</p> Full article ">
14 pages, 11167 KiB  
Article
Effect of Temperature and Strain on Bonding of Similar AA3105 Aluminum Alloys by the Roll Bonding Process
by Mauro Carta, Pasquale Buonadonna, Barbara Reggiani, Lorenzo Donati, Francesco Aymerich and Mohamad El Mehtedi
Metals 2024, 14(8), 920; https://doi.org/10.3390/met14080920 - 14 Aug 2024
Abstract
Accumulative roll bonding (ARB) is a severe plastic deformation process that enables the production of materials with ultrafine microstructures and enhances the characteristics of the base material, particularly in metal matrix composites. The primary objective of this study is to experimentally investigate the [...] Read more.
Accumulative roll bonding (ARB) is a severe plastic deformation process that enables the production of materials with ultrafine microstructures and enhances the characteristics of the base material, particularly in metal matrix composites. The primary objective of this study is to experimentally investigate the bonding strength in AA3105 strips that underwent the roll bonding process, with a specific focus on examining the influence of temperature and reduction rate on bonding. Three temperature levels (200 °C, 300 °C, and 400 °C) and three thickness reduction levels (35%, 50%, and 65%) were considered. The T-peel test was carried out to assess the bonding quality. It was employed to determine the peak force required to separate the two bonded strips. Additionally, ANOVA analysis was performed to develop a regression equation for analyzing peak force. Optical microscopy was used to evaluate the interface bonding quality in the longitudinal section. The results indicate that the bonding strength increases with both temperature and percentage reduction. Full article
(This article belongs to the Special Issue Metal Rolling and Heat Treatment Processing)
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<p>Schematic diagram of the ARB process [<a href="#B3-metals-14-00920" class="html-bibr">3</a>].</p> Full article ">Figure 2
<p>(<b>a</b>) Specimen with the two sheets bent to 90° in order to perform the T-peel test; (<b>b</b>) T-RB sample during the T-peel test.</p> Full article ">Figure 3
<p>T-peel specimens after testing with different deformation modes: (<b>a</b>) sample with elastic deformation indicating inferior bonding; (<b>b</b>) slight plastic deformation with moderate bonding strength; (<b>c</b>) significant plastic deformation and ductile behavior suggesting a strong bond between strips.</p> Full article ">Figure 4
<p>Examples of load normalized width vs. displacement of the T-peel tests for every condition: (<b>a</b>) 200 °C (<b>b</b>) 300 °C; (<b>c</b>) 400 °C.</p> Full article ">Figure 5
<p>Different representations of experimental data: (<b>a</b>) diagram of F/w and (<b>b</b>) 3D representation of F/w as a function of temperature and reduction rate.</p> Full article ">Figure 6
<p>Surface plot for peak force per width as a function of temperature and reduction percentage.</p> Full article ">Figure 7
<p>Comparison between experimental and predicted results.</p> Full article ">Figure 8
<p>Micrographs by OM at 200 °C and reductions of (<b>a</b>) 35%, (<b>b</b>) 50%, and (<b>c</b>) 65%; at 300 °C and reductions of (<b>d</b>) 35%, (<b>e</b>) 50%, and (<b>f</b>) 65%; at 400 °C and reductions of (<b>g</b>) 35%, (<b>h</b>) 50%, and (<b>i</b>) 65%.</p> Full article ">Figure 9
<p>Micrographs by OM with polarized light at 200 °C and reductions of (<b>a</b>) 35%, (<b>b</b>) 50%, and (<b>c</b>) 65%.</p> Full article ">Figure 10
<p>Micrographs by OM with polarized light at 300 °C and reductions of (<b>a</b>) 35%, (<b>b</b>) 50%, and (<b>c</b>) 65%.</p> Full article ">Figure 11
<p>Micrographs by OM with polarized light at 400 °C and reductions of (<b>a</b>) 35%, (<b>b</b>) 50%, and (<b>c</b>) 65%.</p> Full article ">Figure 12
<p>SEM image of 400 °C sample (50% reduction), 3000×.</p> Full article ">
13 pages, 6168 KiB  
Article
Cooling Air Velocity on Iron Ore Pellet Performance Based on Experiments and Simulations
by Liming Ma, Jianliang Zhang, Zhengjian Liu, Qiuye Cai, Liangyuan Hao, Shaofeng Lu, Huiqing Jiang and Yaozu Wang
Metals 2024, 14(8), 919; https://doi.org/10.3390/met14080919 - 14 Aug 2024
Abstract
During the pellet cooling process, cooling air velocity is crucial for optimizing the cooling rate, evaluating the utilization rate of cooling heat energy, and improving pellet performance. As the simulated cooling air velocity increased, the gas temperature at the cooling endpoint decreased from [...] Read more.
During the pellet cooling process, cooling air velocity is crucial for optimizing the cooling rate, evaluating the utilization rate of cooling heat energy, and improving pellet performance. As the simulated cooling air velocity increased, the gas temperature at the cooling endpoint decreased from 87 °C to 51 °C, and the solid temperature decreased from 149 °C to 103 °C. The total enthalpy of the recovered gas initially reduced and then increased while the heat recovery rate gradually increased. During the experiment, the inhomogeneity of pellet quality gradually increased with the rise in cooling air velocity. The effect of cooling air velocity on pellet properties is primarily reflected in the formation of cracks and low-melting liquid phases (FeO and fayalite). As the cooling air velocity increases, the softening onset temperature of the pellet decreases significantly. The melting zone decreases from 193 °C to 105 °C, and the permeability of the adhesive zone increases. Full article
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<p>Schematic diagram of the cooling experiment for the iron ore pellets ((<b>a</b>) green pellet preparation; (<b>b</b>) drying of green pellets; (<b>c</b>) preheating and roasting; (<b>d</b>) cooling of iron ore pellets).</p> Full article ">Figure 2
<p>Schematic diagram of cooling simulation of the iron ore pellets ((<b>a</b>) schematic diagram of pellet cooling equipment; (<b>b</b>) modelling of iron ore pellet cooling; (<b>c</b>) iron ore pellet cooling baseline model; (<b>d</b>) accuracy calibration of cooling models of iron ore pellet).</p> Full article ">Figure 3
<p>Influence of cooling air velocity on pellet performance ((<b>a</b>) cooling time; (<b>b</b>) pellet strength).</p> Full article ">Figure 4
<p>Light microscopic analysis of pellets at different cooling air velocities ((<b>a</b>–<b>d</b>) 2 m<sup>3</sup>/h; (<b>e</b>–<b>h</b>) 4 m<sup>3</sup>/h; (<b>i</b>–<b>l</b>) 6 m<sup>3</sup>/h).</p> Full article ">Figure 5
<p>SEM-EDS analysis positions of pellets at different cooling air velocities ((<b>a</b>–<b>c</b>) 2 m<sup>3</sup>/h; (<b>d</b>–<b>f</b>) 4 m<sup>3</sup>/h; (<b>g</b>–<b>i</b>) 6 m<sup>3</sup>/h).</p> Full article ">Figure 6
<p>Effect of cooling gas on gas–solid phase temperature ((<b>a</b>) gas temperature; (<b>b</b>) temperature of pellet).</p> Full article ">Figure 7
<p>Effect of cooling air velocity on cooling rate and enthalpy ((<b>a</b>) cooling rate; (<b>b</b>) enthalpy).</p> Full article ">Figure 8
<p>Effect of cooling air velocity on heat recovery.</p> Full article ">Figure 9
<p>Softening–melting characteristics of pellets at different cooling air velocities.</p> Full article ">
20 pages, 17160 KiB  
Article
Molecular Dynamics Study on the Mechanical Behaviors of Nanotwinned Titanium
by Bingxin Wu, Kaikai Jin and Yin Yao
Metals 2024, 14(8), 918; https://doi.org/10.3390/met14080918 - 14 Aug 2024
Abstract
Titanium and titanium alloys have been widely applied in the manufacture of aircraft engines and aircraft skins, the mechanical properties of which have a crucial influence on the safety and lifespan of aircrafts. Based on nanotwinned titanium models with different twin boundary spacings, [...] Read more.
Titanium and titanium alloys have been widely applied in the manufacture of aircraft engines and aircraft skins, the mechanical properties of which have a crucial influence on the safety and lifespan of aircrafts. Based on nanotwinned titanium models with different twin boundary spacings, the impacts of different loadings and twin boundary spacings on the plastic deformation of titanium were studied in this paper. It was found that due to the different contained twin boundaries, the different types of nanotwinned titanium possessed different dislocation nucleation abilities on the twin boundaries, different types of dislocation–twin interactions occurred, and significant differences were observed in the mechanical properties and plastic deformation mechanisms. For the {101-2} twin, basal plane dislocations were likely to nucleate on the twin boundary. The plastic deformation mechanism of the material under tensile loading was dominated by partial dislocation slip on the basal plane and face-centered cubic phase transitions, and the yield strength of the titanium increased with decreasing twin boundary spacing. However, under compression loading, the plastic deformation mechanism of the material was dominated by a combination of partial dislocation slip on the basal plane and twin boundary migration. For the {101-1} twin under tensile loading, the plastic deformation mechanism of the material was dominated by partial dislocation slip on the basal plane and crack nucleation and propagation, while under compression loading, the plastic deformation mechanism of the material was dominated by partial dislocation slip on the basal plane and twin boundary migration. For the {1124} twin, the interaction of its twin boundary and dislocation could produce secondary twins. Under tensile loading, the plastic deformation mechanism of the material was dominated by dislocation–twin and twin–twin interactions, while under compression loading, the plastic deformation mechanism of the material was dominated by partial dislocation slip on the basal plane, and the product of the dislocation–twin interactions was basal dislocation. All these results are of guiding value for the optimal design of microstructures in titanium, which should be helpful for achieving strong and tough metallic materials for aircraft manufacturing. Full article
(This article belongs to the Special Issue Deformation of Metals and Alloys: Theory, Simulations and Experiments)
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<p>Potential energy per atom for {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>2</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>011</mn> </mrow> </semantics></math>&gt;, {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt;, and {<math display="inline"><semantics> <mrow> <mn>11</mn> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>4</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>22</mn> <mover accent="true"> <mrow> <mn>4</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; twin boundaries.</p> Full article ">Figure 2
<p>Atomistic configurations of {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>2</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>011</mn> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings: (<b>a</b>) 13.8 nm; (<b>b</b>) 11.1 nm; (<b>c</b>) 6.9 nm.</p> Full article ">Figure 3
<p>Stress–strain curves of {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>2</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>011</mn> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under tensile loadings.</p> Full article ">Figure 4
<p>Atomistic configurations of {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>2</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>011</mn> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under tensile loadings: (<b>a</b>) 13.8 nm; (<b>b</b>) 11.1 nm; (<b>c</b>) 6.9 nm.</p> Full article ">Figure 5
<p>(<b>a</b>) Basal partial dislocation emitted from (<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>2</mn> </mrow> </semantics></math>) TB; (<b>b</b>) creation of numerous basal/prismatic interfaces.</p> Full article ">Figure 6
<p>Stress–strain curves of {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>2</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>011</mn> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under compressive loadings.</p> Full article ">Figure 7
<p>Atomistic configurations of {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>2</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>011</mn> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under compressive loadings: (<b>a</b>) 27.6 nm; (<b>b</b>) 11.1 nm; (<b>c</b>) 6.9 nm.</p> Full article ">Figure 8
<p>The phenomenon of detwinning due to (<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>2</mn> </mrow> </semantics></math>) TB migration: (<b>a</b>) 11.1 nm; (<b>b</b>) 6.9 nm.</p> Full article ">Figure 9
<p>Stress–strain curves of {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under tensile loadings.</p> Full article ">Figure 10
<p>Atomistic configurations of {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under tensile loadings: (<b>a</b>) 13.5 nm; (<b>b</b>) 10.8 nm.</p> Full article ">Figure 11
<p>(<b>a</b>) A (<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>) pyramidal partial dislocation; (<b>b</b>) a prismatic partial dislocation.</p> Full article ">Figure 12
<p>Stress–strain curves of {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under compressive loadings.</p> Full article ">Figure 13
<p>Atomistic configurations of {<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under compressive loadings: (<b>a</b>) 13.5 nm; (<b>b</b>) 6.7 nm.</p> Full article ">Figure 14
<p>(<b>a</b>) A (<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>) pyramidal slip; (<b>b</b>) the migration of (<math display="inline"><semantics> <mrow> <mn>10</mn> <mover accent="true"> <mrow> <mn>1</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>) TB due to twinning dislocation slip.</p> Full article ">Figure 15
<p>Stress–strain curves of {<math display="inline"><semantics> <mrow> <mn>11</mn> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>4</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>22</mn> <mover accent="true"> <mrow> <mn>4</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under tensile loadings.</p> Full article ">Figure 16
<p>Atomistic configurations of {<math display="inline"><semantics> <mrow> <mn>11</mn> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>4</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>22</mn> <mover accent="true"> <mrow> <mn>4</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; nanotwinned titanium with twin boundary spacings of 17 nm under tensile loading.</p> Full article ">Figure 17
<p>(<b>a</b>,<b>b</b>) The (<math display="inline"><semantics> <mrow> <mn>11</mn> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>) twin cutting through basal SF boundaries, and (<b>c</b>,<b>d</b>) (<math display="inline"><semantics> <mrow> <mn>11</mn> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>1</mn> </mrow> </semantics></math>) TB migration leading to detwinning.</p> Full article ">Figure 18
<p>Atomistic configurations of {<math display="inline"><semantics> <mrow> <mn>11</mn> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>4</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>22</mn> <mover accent="true"> <mrow> <mn>4</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; nanotwinned titanium with twin boundary spacings of 7.3 nm under tensile loading.</p> Full article ">Figure 19
<p>Stress–strain curves of {<math display="inline"><semantics> <mrow> <mn>11</mn> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>4</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>22</mn> <mover accent="true"> <mrow> <mn>4</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under compressive loadings.</p> Full article ">Figure 20
<p>Atomistic configurations of the {<math display="inline"><semantics> <mrow> <mn>11</mn> <mover accent="true"> <mrow> <mn>2</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mn>4</mn> </mrow> </semantics></math>}&lt;<math display="inline"><semantics> <mrow> <mn>22</mn> <mover accent="true"> <mrow> <mn>4</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> <mover accent="true"> <mrow> <mn>3</mn> </mrow> <mo mathvariant="normal">-</mo> </mover> </mrow> </semantics></math>&gt; nanotwinned titanium with different twin boundary spacings under compressive loadings, (<b>a</b>) 17 nm; (<b>b</b>) 7.3 nm.</p> Full article ">
18 pages, 7849 KiB  
Article
Exploring Mechanical Properties Using the Hydraulic Bulge Test and Uniaxial Tensile Test with Micro-Samples for Metals
by Jian Bao, Haoran Ding, Zhiquan Zuo and Jian Peng
Metals 2024, 14(8), 917; https://doi.org/10.3390/met14080917 - 13 Aug 2024
Viewed by 253
Abstract
The hydraulic bulge test with micro-samples is expected to be useful in the damage assessment of long-service-period metals to understand the degeneration of their mechanical properties. Since the hydraulic bulge test has a different stress state from the classical uniaxial tensile test, we [...] Read more.
The hydraulic bulge test with micro-samples is expected to be useful in the damage assessment of long-service-period metals to understand the degeneration of their mechanical properties. Since the hydraulic bulge test has a different stress state from the classical uniaxial tensile test, we need to understand their correlation and differences. In this study, the hydraulic bulge test and the uniaxial tensile test are employed to analyze the mechanical properties of three typical metals used in pressure vessels: 316L, 16MnDR, and Q345R. By utilizing Kruglov’s vertex thickness and Panknin’s curvature radius equivalent, the pressure–displacement curves from the hydraulic bulge test are converted into biaxial stress–strain curves. Based on the equivalent plastic energy model, the biaxial stress–strain curves are converted into uniaxial stress–strain curves with an error less than 10% in the strain hardening stage, achieving the unified characterization of mechanical properties under different stress states. Moreover, the hydraulic bulge test provides a more extensive strain hardening stage, and the fracture strains are 9–16.5% larger than those of uniaxial tensile test. This paper provides a reference for using the hydraulic bulge test with micro-samples in studying the mechanical properties and presents the advantages of this novel test method. Full article
(This article belongs to the Special Issue Fatigue, Creep Behavior and Fracture Mechanics of Metals)
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<p>HBT device, sample, and typical result: (<b>a</b>) schematic diagram of the HBT device; (<b>b</b>) HBT device and pressure–displacement measurement system; (<b>c</b>) micro-sample; (<b>d</b>) pressure–displacement curve.</p> Full article ">Figure 2
<p>UTT: (<b>a</b>) micro-sample dimensions; (<b>b</b>) full-field strain measurement by DIC.</p> Full article ">Figure 3
<p>HBT pressure–displacement curves and fracture morphologies of 316L, 16MnDR, and Q345R.</p> Full article ">Figure 4
<p>Yield pressure determination method.</p> Full article ">Figure 5
<p>Stress–strain curves of 316L, 16MnDR, and Q345R: (<b>a</b>) engineering stress–strain curves; (<b>b</b>) true stress–strain curves.</p> Full article ">Figure 6
<p>Comparison of mechanical parameters of the UTT and HBT: (<b>a</b>) yield parameters; (<b>b</b>) ultimate tensile parameters.</p> Full article ">Figure 7
<p>Comparison of fracture morphologies between the HBT and UTT tests: (<b>a1</b>–<b>a3</b>) 316L-HBT; (<b>b1</b>–<b>b3</b>) 316L-UTT; (<b>c1</b>–<b>c3</b>) 16MnDR-HBT; (<b>d1</b>–<b>d3</b>) 16MnDR-UTT; (<b>e1</b>–<b>e3</b>) Q345R-HBT; (<b>f1</b>–<b>f3</b>) Q345R-UTT.</p> Full article ">Figure 8
<p>Structural model of the HBT.</p> Full article ">Figure 9
<p>Comparison of the biaxial stress–strain curve provided by the HBT and the uniaxial stress–strain curve provided by the UTT: (<b>a</b>) 316L; (<b>b</b>) 16MnDR; (<b>c</b>) Q345R.</p> Full article ">Figure 10
<p>Stress–plastic work curves obtained from the HBT and UTT.</p> Full article ">Figure 11
<p>The true stress–plastic strain curves obtained from the HBT and UTT: (<b>a</b>) 316L; (<b>b</b>) 16MnDR; (<b>c</b>) Q345R.</p> Full article ">
19 pages, 5372 KiB  
Article
Enhanced Generative Adversarial Networks for Isa Furnace Matte Grade Prediction under Limited Data
by Huaibo Ma, Zhuorui Li, Bo Shu, Bin Yu and Jun Ma
Metals 2024, 14(8), 916; https://doi.org/10.3390/met14080916 - 13 Aug 2024
Viewed by 230
Abstract
Due to the scarcity of modeling samples and the low prediction accuracy of the matte grade prediction model in the copper melting process, a new prediction method is proposed. This method is based on enhanced generative adversarial networks (EGANs) and random forests (RFs). [...] Read more.
Due to the scarcity of modeling samples and the low prediction accuracy of the matte grade prediction model in the copper melting process, a new prediction method is proposed. This method is based on enhanced generative adversarial networks (EGANs) and random forests (RFs). Firstly, the maximum relevance minimum redundancy (MRMR) algorithm is utilized to screen the key influencing factors of matte grade and remove redundant information. Secondly, the GAN data augmentation model containing different activation functions is constructed. And, the generated data fusion criterion based on the root mean squared error (RMSE) and the coefficient of determination (R2) is designed, which can tap into the global character distributions of the copper melting data to improve the quality of the generated data. Finally, a matte grade prediction model based on RF is constructed, and the industrial data collected from the copper smelting process are used to verify the effectiveness of the model. The experimental results show that the proposed method can obtain high-quality generated data, and the prediction accuracy is better than other models. The R2 is improved by at least 2.68%, and other indicators such as RMSE, mean absolute error (MAE), and mean absolute percentage error (MAPE) are significantly improved. Full article
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<p>Generative adversarial network structure.</p> Full article ">Figure 2
<p>RF model prediction flowchart.</p> Full article ">Figure 3
<p>Overall framework of the matte grade prediction model, fusing the generative adversarial network and random forest.</p> Full article ">Figure 4
<p>Isa furnace melting process flow diagram.</p> Full article ">Figure 5
<p>The box plot screens out the outliers in the data.</p> Full article ">Figure 6
<p>PPC between input variables and matte grade: (<b>a</b>) PPC between input variables and matte grade in <span class="html-italic">D</span><sub>1</sub>, <span class="html-italic">D</span><sub>2</sub>, <span class="html-italic">D</span><sub>3</sub>, and <span class="html-italic">D</span>; (<b>b</b>) PCC between input variables and matte grade in <span class="html-italic">D<sub>f</sub></span><sub>1</sub>, <span class="html-italic">D<sub>f</sub></span><sub>2</sub>, <span class="html-italic">D<sub>f</sub></span><sub>3</sub>, <span class="html-italic">D<sub>f</sub></span><sub>4</sub>, <span class="html-italic">D</span><sub>1</sub>, <span class="html-italic">D</span><sub>2</sub>, <span class="html-italic">D</span><sub>3</sub>, and <span class="html-italic">D</span>.</p> Full article ">Figure 7
<p>The frequency distribution histogram of the real data <span class="html-italic">D</span> and the generated data <span class="html-italic">D<sub>f</sub></span><sub>4</sub>.</p> Full article ">Figure 8
<p>The normal distribution curve of the generated data <span class="html-italic">D<sub>f</sub></span><sub>4</sub> and the real data <span class="html-italic">D</span>.</p> Full article ">Figure 9
<p>Comparison of generated data with real data: (<b>a</b>) frequency distribution histogram; (<b>b</b>) normal distribution curve; (<b>c</b>) PCA dimension reduction.</p> Full article ">Figure 10
<p>Comparison of ablation experiments: (<b>a</b>) variable selection experiment comparison; (<b>b</b>) data enhancement experiment comparison; (<b>c</b>) data fusion experiment comparison.</p> Full article ">Figure 11
<p>Visualization of the prediction results: (<b>a</b>) map of the distribution of the prediction error of matte grade; (<b>b</b>) scatter plot of the distribution of the actual and predicted values of matte grade.</p> Full article ">Figure 12
<p>Parameters importance. (<b>a</b>) SHAP values of different features for different samples. (<b>b</b>) Feature importance ranking by SHAP.</p> Full article ">Figure 13
<p>Contribution of each feature to the local samples. (<b>a</b>) Test sample 1. (<b>b</b>) Test sample 2.</p> Full article ">
13 pages, 3533 KiB  
Article
Nodular Graphite Dissolution and Nucleus Observation: High-Temperature Dynamics of Ductile Iron Recycling
by I. Adhiwiguna, N. Nobakht and R. Deike
Metals 2024, 14(8), 915; https://doi.org/10.3390/met14080915 - 13 Aug 2024
Viewed by 195
Abstract
This investigation examines the dynamic behavior of the nodular graphite structure in ductile cast iron at elevated temperatures during the recycling process. It comprises a systematic analysis of the impact of high temperature on the change in chemical composition, followed by a set [...] Read more.
This investigation examines the dynamic behavior of the nodular graphite structure in ductile cast iron at elevated temperatures during the recycling process. It comprises a systematic analysis of the impact of high temperature on the change in chemical composition, followed by a set of examinations of the nodular graphite structure dissolution mechanism at the early phase of the remelting process. The results indicate that prolonged holding at higher temperatures affects the carbon or silicon concentration due to oxidation, which correlates with the operating temperature and the dynamic concentration proportion of those two main alloying elements. It is also substantiated that the dissolution of nodular graphite, the only carbon source during the ductile cast iron remelting process, does not occur primarily in the liquid state but has already started during the solid phase because of austenitization. This dissolution is governed mainly by a surface reaction, as indicated by the residual graphite structure with preserved nonmetallic nuclei. Hence, this approach also provides an alternative method for observing the nodular graphite core by intentionally partially dissolving the graphite structure. Full article
(This article belongs to the Special Issue Casting Alloy Design and Characterization—2nd Edition)
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<p>The average change in the carbon content during the high-temperature holding of (<b>a</b>) HS-CI and (<b>b</b>) LS-CI—additional sample LS-CI 1400 °C* was melted in an alumina-based crucible.</p> Full article ">Figure 2
<p>The average change in the carbon and silicon contents in (<b>a</b>) HS-CI and (<b>b</b>) LS-CI during holding at a temperature of 1300 °C.</p> Full article ">Figure 3
<p>The average change in the carbon and silicon contents in (<b>a</b>) HS-CI and (<b>b</b>) LS-CI during holding at a temperature of 1500 °C.</p> Full article ">Figure 4
<p>Microstructure of nital-etched LS-CI (<b>a</b>) before and (<b>b</b>) after ET-Exp (1200 °C for 10 min)—transforming (1) nodular graphite in (2) pearlite and (3) ferrite matrix system into (4) residual graphite in (5) ledeburite and (6) martensite matrix system.</p> Full article ">Figure 5
<p>Microstructure development LS-CI after being held at 1200 °C followed by water quenching for (<b>a1</b>) 5, (<b>b1</b>) 10, and (<b>c1</b>) 15 minutes in as-polished condition as well as for (<b>a2</b>) 5, (<b>b2</b>) 10, and (<b>c2</b>) 15 minutes in nital-etched condition.</p> Full article ">Figure 6
<p>Microstructure of the nital-etched LS-CI after 15 minutes of holding at 1200 °C followed by water quenching indicating (<b>a</b>) residual graphite in red circle detailed in (<b>b</b>): (1) graphite, (2) martensite, and (3) ledeburite.</p> Full article ">Figure 7
<p>Microstructure of the nital-etched LS-CI after 15 minutes of holding at (<b>a</b>) 1100 °C, (<b>b</b>) 1200 °C, and (<b>c</b>) 1300 °C followed by water quenching—detailed red arrow G: residual graphite.</p> Full article ">Figure 8
<p>(<b>a</b>) Rest nodular graphite and (<b>b</b>) detail of its nonmetallic nucleus in the LS-CI sample.</p> Full article ">Figure 9
<p>FactSage calculation results for the oxidation of HS-CI and LS-CI at explored temperatures.</p> Full article ">
18 pages, 14710 KiB  
Article
Full Density Powder Metallurgical Cold Work Tool Steel through Nitrogen Sintering and Capsule-Free Hot Isostatic Pressing
by Anok Babu Nagaram, Giulio Maistro, Erik Adolfsson, Yu Cao, Eduard Hryha and Lars Nyborg
Metals 2024, 14(8), 914; https://doi.org/10.3390/met14080914 - 12 Aug 2024
Viewed by 218
Abstract
Vanadis 4E (V4E) is a powder metallurgical cold work tool steel predominantly used in application with demand for wear resistance, high hardness, and toughness. It is of interest to have a processing route that enables full density starting from clean gas-atomized powder allowing [...] Read more.
Vanadis 4E (V4E) is a powder metallurgical cold work tool steel predominantly used in application with demand for wear resistance, high hardness, and toughness. It is of interest to have a processing route that enables full density starting from clean gas-atomized powder allowing component shaping capabilities. This study presents a process involving freeze granulation of powder to facilitate compaction by means of cold isostatic pressing, followed by sintering to allow for capsule-free hot isostatic pressing (HIP) and subsequent heat treatments of fully densified specimens. The sintering stage has been studied in particular, and it is shown how sintering in pure nitrogen at 1150 °C results in predominantly closed porosity, while sintering at 1200 °C gives near full density. Microstructural investigation shows that vanadium-rich carbonitride (MX) is formed as a result of the nitrogen uptake during sintering, with coarser appearance for the higher temperature. Nearly complete densification, approximately 7.80 ± 0.01 g/cm3, was achieved after sintering at 1200 °C, and after sintering at 1150 °C, followed by capsule-free HIP, hardening, and tempering. Irrespective of processing once the MX is formed, the nitrogen is locked into this phase and the austenite is stabilised, which means any tempering tends to result in a mixture of austenite and tempered martensite, the former being predominate during the sequential tempering, whereas martensite formation during cooling from austenitization temperatures becomes limited. Full article
(This article belongs to the Special Issue Powder Metallurgy of Metallic Materials)
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<p>Schematic illustration of freeze granulation (FG) method.</p> Full article ">Figure 2
<p>(<b>a</b>) Uniaxial pressing of granules, (<b>b</b>) isostatic pressing of the pre-pressed sample, and (<b>c</b>) samples in sealed vacuum bags after CIP.</p> Full article ">Figure 3
<p>Relative green density of V4E after compaction at different CIP pressures of 100 MPa, 200 MPa, and 300 MPa (<b>a</b>). Debinding curves for V4E showing relative weight as a function of time for two different final debinding temperatures (<b>b</b>).</p> Full article ">Figure 4
<p>Open and closed porosity of V4E samples after sintering at 1150 °C and 1200 °C, measured using the Archimedes density method.</p> Full article ">Figure 5
<p>Optical micrographs of V4E material sintered at 1150 °C (<b>a</b>) and 1200 °C (<b>b</b>).</p> Full article ">Figure 6
<p>Optical micrographs of V4E material after sintering at (<b>a</b>) 1150 °C and (<b>b</b>) 1200 °C, followed by capsule-free HIP, hardening, and tempering, adapted from Ref. [<a href="#B26-metals-14-00914" class="html-bibr">26</a>].</p> Full article ">Figure 7
<p>SEM micrographs of V4E samples after sintering at (<b>a</b>) 1150 °C, and (<b>b</b>) 1200 °C, and (<b>c</b>,<b>d</b>) after subsequent capsule-free HIP at 1140 °C, hardening, and tempering.</p> Full article ">Figure 8
<p>SEM images and EDS mapping of V4E material after sintering at (<b>top row</b>) 1150 °C and (<b>bottom row</b>) 1200 °C.</p> Full article ">Figure 9
<p>SEM images and EDS mapping of V4E material after sintering at 1200 °C.</p> Full article ">Figure 10
<p>X-ray Diffraction pattern (<b>left</b>) and EBSD mapping (<b>right</b>) of V4E material after sintering at 1200 °C, followed by capsule-free HIP, hardening, and tempering.</p> Full article ">Figure 11
<p>Dilatometry curves of V4E PM tool steel material at sintering temperatures of 1150 °C and 1200 °C.</p> Full article ">Figure 12
<p>(<b>a</b>) DSC signal vs. temperature for V4E PM tool steel sample (<b>b</b>) Thermodynamic simulation of V4E using ThermoCalc software (2024a Version).</p> Full article ">Figure 13
<p>Thermodynamic calculations studies of V4E PM tool steel (<b>a</b>) at different pressures of nitrogen (<b>b</b>), with varying amounts of N in the sample.</p> Full article ">Figure 14
<p>JMatPro simulation (Version 12.4) of V4E PM cold work tool steel with 0.7% nitrogen.</p> Full article ">Figure 15
<p>Martensite transition of V4E cold work tool steel with and without nitrogen predicted by JMatPro simulation (Version 12.4) for hardening temperature of 1100 °C (green line).</p> Full article ">
13 pages, 27460 KiB  
Article
Comparative Study of Heat Transfer Simulation and Effects of Different Scrap Steel Preheating Methods
by Pengcheng Xiao, Yuxin Jin, Liguang Zhu, Chao Wang and Rong Zhu
Metals 2024, 14(8), 913; https://doi.org/10.3390/met14080913 - 12 Aug 2024
Viewed by 237
Abstract
The materials charged into a converter comprise molten iron and scrap steel. Adjusting the ratio by increasing scrap steel and decreasing molten iron is a steelmaking raw material strategy designed specifically for China’s unique circumstances, with the goal of lowering carbon emissions. To [...] Read more.
The materials charged into a converter comprise molten iron and scrap steel. Adjusting the ratio by increasing scrap steel and decreasing molten iron is a steelmaking raw material strategy designed specifically for China’s unique circumstances, with the goal of lowering carbon emissions. To maintain the converter tapping temperature, scrap must be preheated to provide additional heat. Current scrap preheating predominantly utilizes horizontal tunnel furnaces, resulting in high energy consumption and low efficiency. To address these issues, a three-stage shaft furnace for scrap preheating was designed, and Fluent software was used to compare and study the preheating efficiency of the new three-stage furnace against the traditional horizontal furnace under various operational conditions. Initially, a three-dimensional transient multi-field coupling model was developed for two scrap preheating scenarios, examining the effects of both furnaces on scrap surface and core temperatures across varying preheating durations and gas velocities. Simulation results indicate that, under identical gas heat consumption conditions, scrap achieves markedly higher final temperatures in the shaft furnace compared to the horizontal furnace, with scrap surface and core temperatures increasing notably with extended preheating times and higher gas velocities, albeit with a gradual decrease in heating rate as the scrap temperature rises. At a gas velocity of 9 m/s and a preheating time of 600 s, the shaft furnace achieves the highest waste heat utilization rate for scrap, with scrap averaging 325 °C higher than in the horizontal furnace, absorbing an additional 202 MJ of heat per ton. In the horizontal preheating furnace, scrap steel exhibits a heat absorption efficiency of 35%, whereas in the vertical furnace, this efficiency increases notably to 63%. In the vertical furnace, the waste heat recovery rate of scrap steel reaches 57%. Full article
(This article belongs to the Special Issue Advanced Tundish Metallurgy and Clean Steel Technology—Second Edition)
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<p>Horizontal preheating furnace.</p> Full article ">Figure 2
<p>Two types of preheating furnace models. (<b>a</b>) Vertical triple-stage continuous furnace. (<b>b</b>) Continuous horizontal furnace.</p> Full article ">Figure 3
<p>Comparison of simulation results from the horizontal furnace model with actual temperature measurements. (<b>a</b>) The temperature inside the horizontal furnace model. (<b>b</b>) The temperature of scrap steel inside the horizontal furnace model. (<b>c</b>) Measured temperature.</p> Full article ">Figure 4
<p>Preheating effectiveness of scrap at different times in two types of preheating furnaces. (<b>a</b>) The temperature distribution following a 600 s preheating period in the horizontal furnace. (<b>b</b>) The temperature distribution following a 900 s preheating period in the horizontal furnace. (<b>c</b>) The temperature distribution following a 1200 s preheating period in the horizontal furnace. (<b>d</b>) The temperature distribution following a 600 s preheating period in the vertical furnace. (<b>e</b>) The temperature distribution following a 900 s preheating period in the vertical furnace. (<b>f</b>) The temperature distribution following a 1200 s preheating period in the vertical furnace.</p> Full article ">Figure 5
<p>Two different types of preheating furnaces with varying initial velocities of combustible gases and their impact on the preheating of scrap. (<b>a</b>) The temperature field with a gas velocity of 5 m/s in the horizontal furnace. (<b>b</b>) The temperature field with a gas velocity of 7 m/s in the horizontal furnace. (<b>c</b>) The temperature field with a gas velocity of 9 m/s in the horizontal furnace. (<b>d</b>) The temperature field with a gas velocity of 5 m/s in the vertical furnace. (<b>e</b>) The temperature field with a gas velocity of 7 m/s in the vertical furnace. (<b>f</b>) The temperature field with a gas velocity of 9 m/s in the vertical furnace.</p> Full article ">Figure 6
<p>Preheating effect of scrap cores under different preheating times in horizontal preheating furnace. (<b>a</b>) Temperature field of scrap cores preheated for 600 s in a horizontal furnace. (<b>b</b>) Temperature field of scrap cores preheated for 900 s in a horizontal furnace. (<b>c</b>) Temperature field of scrap cores preheated for 600 s in a horizontal furnace.</p> Full article ">Figure 7
<p>Preheating effect of scrap cores under different preheating times in horizontal preheating furnace. (<b>a</b>) Temperature field of scrap cores preheated for 600 s in a vertical furnace. (<b>b</b>) Temperature field of scrap cores preheated for 900 s in a vertical furnace. (<b>c</b>) Temperature field of scrap cores preheated for 1200 s in a vertical furnace.</p> Full article ">Figure 8
<p>Temperature of scrap at different times in two types of preheating furnaces at the same gas velocity. (<b>a</b>) Temperature of scrap preheated for 600 s in a horizontal furnace. (<b>b</b>) Temperature of scrap preheated for 900 s in a horizontal furnace. (<b>c</b>) Temperature of scrap preheated for 1200 s in a horizontal furnace. (<b>d</b>) Temperature of scrap preheated for 600 s in a vertical furnace. (<b>e</b>) Temperature of scrap preheated for 900 s in a vertical furnace. (<b>f</b>) Temperature of scrap preheated for 1200 s in a vertical furnace.</p> Full article ">Figure 9
<p>Temperature of scrap at different gas velocities in two types of preheating furnaces at the same time. (<b>a</b>) Temperature of scrap in the horizontal furnace when gas velocity is 5 m/s. (<b>b</b>) Temperature of scrap in the horizontal furnace when gas velocity is 7 m/s. (<b>c</b>) Temperature of scrap in the horizontal furnace when gas velocity is 9 m/s. (<b>d</b>) Temperature of scrap in the vertical furnace when gas velocity is 5 m/s. (<b>e</b>) Temperature of scrap in the vertical furnace when gas velocity is 7 m/s. (<b>f</b>) Temperature of scrap in the vertical furnace when gas velocity is 9 m/s.</p> Full article ">
14 pages, 5049 KiB  
Article
Parametric Study of Vanadium Extraction Process from Refining Tailings of Crude Titanium Tetrachloride
by Weitong Du, Tingfeng Yao, Haiming Cheng, Dianchun Ju and Zhuo Chen
Metals 2024, 14(8), 912; https://doi.org/10.3390/met14080912 - 12 Aug 2024
Viewed by 311
Abstract
The recovery of vanadium from titanium tetrachloride tail residue is a resource-efficient and environment-friendly method for treating hazardous vanadium-containing solid waste. In this study, to maximize the recovery rate of vanadium in the vanadium extraction process, the independent calcination and leaching factors were [...] Read more.
The recovery of vanadium from titanium tetrachloride tail residue is a resource-efficient and environment-friendly method for treating hazardous vanadium-containing solid waste. In this study, to maximize the recovery rate of vanadium in the vanadium extraction process, the independent calcination and leaching factors were optimized using response surface methodology, in terms of calcination temperature (750–950 °C), calcination time (60–180 min), leaching liquid–solid ratio (5–25 mL/g), and leaching time (30–150 min). The results revealed that the calcination temperature was the most effective parameter for vanadium recovery, while the liquid–solid ratio was the least effective factor. Additionally, the optimal conditions were identified as a calcination temperature of 937 °C, a calcination time of 150 min, a leaching solid-to-liquid ratio of 17.4 mL/g, and a leaching time of 150 min. The maximum predicted recovery rate of vanadium by the model regression equation reached 93.1% and showed high credibility consistent with the experimental recovery rate of 93%. Full article
(This article belongs to the Section Extractive Metallurgy)
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<p>A schematic diagram of the vanadium extraction process from the refining tailings of crude titanium tetrachloride.</p> Full article ">Figure 2
<p>Particle size distribution of raw tailing.</p> Full article ">Figure 3
<p>XRD pattern of refined vanadium-containing tailings from crude titanium tetrachloride.</p> Full article ">Figure 4
<p>XRD pattern of samples at various temperatures.</p> Full article ">Figure 5
<p>Predicted values versus experimental values of vanadium recovery rate.</p> Full article ">Figure 6
<p>Effect of leaching time and calcination time interaction on vanadium recovery rate.</p> Full article ">Figure 7
<p>Effect of calcination time and calcination temperature interaction on vanadium recovery rate.</p> Full article ">Figure 8
<p>Effect of liquid–solid ratio and calcination temperature interaction on vanadium recovery rate.</p> Full article ">Figure 9
<p>The morphology and XRD patterns of (<b>a</b>) NH<sub>4</sub>VO<sub>3</sub> and (<b>b</b>) V<sub>2</sub>O<sub>5</sub> products.</p> Full article ">
15 pages, 16997 KiB  
Article
Active and Passive Filling Stir Repairing of AISI 304 Alloy
by Vincenzo Lunetto, Dario Basile, Valentino Razza and Pasquale Russo Spena
Metals 2024, 14(8), 911; https://doi.org/10.3390/met14080911 - 11 Aug 2024
Viewed by 407
Abstract
This study investigates active filling friction stir repair (AF-FSR) and passive filling friction stir repair (PF-FSR) for repairing AISI 304 stainless steel sheets, focusing on addressing the challenges posed by high melting point metals. The research involved repairing overlapping 2 mm thick sheets [...] Read more.
This study investigates active filling friction stir repair (AF-FSR) and passive filling friction stir repair (PF-FSR) for repairing AISI 304 stainless steel sheets, focusing on addressing the challenges posed by high melting point metals. The research involved repairing overlapping 2 mm thick sheets with pre-drilled holes of 2, 4, and 6 mm diameters, simulating broken components. Various process parameters, including rotational speed, dwell time, and the use of metal fillers, were tested to evaluate their impact on repair quality. The results demonstrated that PF-FSR provided superior mechanical strength to AF-FSR, particularly for larger pre-hole diameters. PF-FSR achieved higher shear tension strength due to better defect filling and reduced void formation, with shear tension strengths exceeding 25 kN for larger pre-holes and lower variability in strength measurements. AF-FSR was less effective for larger pre-holes, resulting in significant voids and reduced strength. Microstructural analysis revealed that PF-FSR facilitated more efficient material mixing and filling, minimizing unrepaired regions. However, excessive rotational speeds and dwell times in PF-FSR led to deformation and flash formation, highlighting the need for optimal parameter selection. Although further studies are needed, this study confirms the feasibility of FSR techniques for repairing small defects in AISI 304 steels, offering valuable insights for sustainable manufacturing practices in industries such as automotive and aerospace, where efficient and reliable repair methods are critical. Full article
(This article belongs to the Special Issue Advances in Welding and Mechanical Joining of Metals)
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<p>The implementation of the “6R methodology” in industrial practice (adapted from [<a href="#B2-metals-14-00911" class="html-bibr">2</a>]).</p> Full article ">Figure 2
<p>Schematic illustration of the overlapped coupons subjected to AF-FSR, PF-FSR, and P-FSSW. The red area highlights the pre-hole in the upper sheets for AF-FSR and PF-FSR. Such samples were also used as shear tension samples.</p> Full article ">Figure 3
<p>Schematic illustration of (<b>a</b>) AF-FSR and (<b>b</b>) PF-FSR processes. Example for a pre-hole diameter of 6 mm. Arrows show the roto-translation motion of the tool pin.</p> Full article ">Figure 4
<p>(<b>a</b>) Equipment and (<b>b</b>) tool geometry used for the repair tests.</p> Full article ">Figure 5
<p>P-FSSW, AF-FSR, and PF-FSR processes: vertical force [kN] and spindle torque [Nm].</p> Full article ">Figure 6
<p>Cross-section of repaired samples obtained through AF-FSR (1000 rpm and 10 s dwell time) with different pre-hole diameters: (<b>a</b>) 2 mm (D2); (<b>b</b>) 4 mm (D4); (<b>c</b>) 6 mm (D6). Image (<b>d</b>) refers to the typical cross-section of a P-FSSW joint (no pre-hole, D0). The white dotted lines refer to the upper sheets before AF- or P-FSSW. The green-colored circle in the image (<b>a</b>) highlight the location of the microstructure reported in Figure 10c.</p> Full article ">Figure 7
<p>Cross-section of repaired samples obtained through AF- and PF-FSR (1000 rpm and 15 s dwell time) with different pre-hole diameters: (<b>a</b>) AF-FSR 4 mm (D4); (<b>b</b>) AF-FSR 6 mm (D6); (<b>c</b>) PF-FSR 4 mm (D4); (<b>d</b>) PF-FSR 6 mm (D6).</p> Full article ">Figure 7 Cont.
<p>Cross-section of repaired samples obtained through AF- and PF-FSR (1000 rpm and 15 s dwell time) with different pre-hole diameters: (<b>a</b>) AF-FSR 4 mm (D4); (<b>b</b>) AF-FSR 6 mm (D6); (<b>c</b>) PF-FSR 4 mm (D4); (<b>d</b>) PF-FSR 6 mm (D6).</p> Full article ">Figure 8
<p>Cross-section of PF-FSR samples with a 6 mm pre-hole obtained with a dwell time of 20 s at varying tool rotational speeds: (<b>a</b>) 1000 rpm and (<b>b</b>) 1500 rpm.</p> Full article ">Figure 9
<p>Schematic material flows beneath the tool during (<b>a</b>) AF-FSR and (<b>b</b>) PF-FSR. The SZ and the TMAZ regions have been drawn for a D6 case from the cross-section of <a href="#metals-14-00911-f006" class="html-fig">Figure 6</a>c and <a href="#metals-14-00911-f007" class="html-fig">Figure 7</a>d.</p> Full article ">Figure 10
<p>Typical microstructures of the (<b>a</b>) BM and (<b>b</b>) SZ of the repaired AF- and PF-FSR regions. In addition, (<b>c</b>) an example of inhomogeneous metal flow lines (enclosed in the red dotted lines) found in some AF-FSR samples, as found in the location of the metallographic image highlighted by the green-colored circle in <a href="#metals-14-00911-f006" class="html-fig">Figure 6</a>a.</p> Full article ">Figure 11
<p>Schematic material flows beneath the tool during AF-FSR divided into four subsequent steps: (<b>a</b>–<b>d</b>) The final shape of SZ and TMAZ were drawn for a D2 case from the cross-section of <a href="#metals-14-00911-f006" class="html-fig">Figure 6</a>a and marked off by black dotted lines.</p> Full article ">Figure 12
<p>Shear tension strength and elongation at the peak load results for the tested samples. If not specified, the tool rotational speed is 1000 rpm.</p> Full article ">Figure 13
<p>Fracture surfaces of the repaired sheet stacks via AF-FSR and PF-FSR with pre-hole diameters of 4 and 6 mm.</p> Full article ">
13 pages, 4316 KiB  
Article
Influence of Top Slag Containing TiO2 and VOx on Hot Metal Pre-Desulfurization
by Biwen Yang, Bo Song, Liang Chen, Honghong Sun, Derek O. Northwood, Kristian E. Waters and Hao Ma
Metals 2024, 14(8), 910; https://doi.org/10.3390/met14080910 - 11 Aug 2024
Viewed by 310
Abstract
The desulfurization capacity of top slag in the process of pre-desulfurization of hot metal containing vanadium and titanium was researched. The top slag system of CaO-SiO2-MgO-Al2O3-TiO2-VOx that was formed by blast furnace slag and [...] Read more.
The desulfurization capacity of top slag in the process of pre-desulfurization of hot metal containing vanadium and titanium was researched. The top slag system of CaO-SiO2-MgO-Al2O3-TiO2-VOx that was formed by blast furnace slag and a CaO desulfurization agent reduced the sulfur in hot metal from 0.08 wt.% to 0.02 wt.%. It was found that the resulfurization of the slag happened in the later periods of the desulfurization process. The vanadium–titanium oxides were both acidic in the desulfurization slag. TiO2 and VOx reacted with the basic oxides to form CaTiO3 and MgV2O4 at 1623 K, which reduced free CaO and was not conducive to top slag desulfurization. The results of calculation showed that the top slag desulfurization accounted for 15% of the total desulfurization. Using the ionic and molecule coexistence theory of slag structure, it is shown that the desulfurization efficiency could be enhanced by adjusting both the amount of desulfurization agent and the composition of the blast furnace slag before pre-desulfurization. Full article
(This article belongs to the Special Issue Modeling Thermodynamic Systems and Optimizing Metallurgical Processes)
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<p>The schematic of high-temperature furnace. 1. Plug, 2. water outlet, 3. corundum tube, 4. MoSi<sub>2</sub> heating elements (U-shape), 5. sleeve, 6. cooling water; 7. water inlet, 8. gas inlet, 9. firebrick, 10. corundum crucible, 11. hot metal, 12. crucible, 13. thermal insulation material, 14. firebrick, 15. thermocouple; 16. gas outlet.</p> Full article ">Figure 2
<p>Diagram of crucible and the placement of experimental material. (<b>a</b>) The effect of CaO-based desulfurizer dosage on top slag desulfurization. (<b>b</b>) The influence of TiO<sub>2</sub> and VOx on top slag desulfurization.</p> Full article ">Figure 3
<p>XRD patterns of top slag with different V and Ti oxide contents.</p> Full article ">Figure 4
<p>Effect of BF slag amount on top slag desulfurization. (<b>a</b>) Increment of the main components of top slag; (<b>b</b>) sulfur content in molten iron and <span class="html-italic">L<sub>S</sub></span> after experiment.</p> Full article ">Figure 5
<p>Scheme of embedding the desulfurizing agent into BF slag.</p> Full article ">Figure 6
<p>The relationship between the “activity” (<span class="html-italic">N<sub>i</sub></span>) and the moles number (<span class="html-italic">n<sub>i</sub></span>) in the slag. (<b>a</b>) CaO; (<b>b</b>) SiO<sub>2</sub>; (<b>c</b>) Al<sub>2</sub>O<sub>3</sub>; (<b>d</b>) MgO; (<b>e</b>) TiO<sub>2</sub>; (<b>f</b>) V<sub>2</sub>O<sub>3</sub>; (<b>g</b>) FeO; (<b>h</b>) MgO·V<sub>2</sub>O<sub>3</sub>; (<b>i</b>) CaO·TiO<sub>2</sub>.</p> Full article ">Figure 7
<p>Relationship between the main component activity of slag and content of (<b>a</b>) TiO<sub>2,</sub> (<b>b</b>) V<sub>2</sub>O<sub>3</sub>.</p> Full article ">Figure 8
<p>Relationship between <span class="html-italic">L<sub>S</sub></span> and the content of (<b>a</b>) V<sub>2</sub>O<sub>3</sub> and (<b>b</b>) TiO<sub>2</sub>.</p> Full article ">Figure A1
<p>The process schematic of automatic slag melting point tester: 1. observation instrument; 2. amplifier; 3. sample delivery equipment; 4. high-temperature heating furnace; 5. luminescence meter; 6. argon gas; 7. temperature controller; 8. console; 9. heating element; 10. Al<sub>2</sub>O<sub>3</sub> gasket; 11. sample; 12. thermocouple.</p> Full article ">
14 pages, 3496 KiB  
Article
Analysis of Longitudinal Cracking and Mold Flux Optimization in High-Speed Continuous Casting of Hyper-Peritectic Steel Thin Slabs
by Zhipeng Yuan, Liguang Zhu, Xingjuan Wang and Kaixuan Zhang
Metals 2024, 14(8), 909; https://doi.org/10.3390/met14080909 - 11 Aug 2024
Viewed by 333
Abstract
Longitudinal crack defects are a frequent occurrence on the surface of thin slabs during the high-speed continuous casting process. Therefore, this study undertakes a detailed analysis of the solidification characteristics of hyper-peritectic steel thin slabs. By establishing a three-dimensional heat transfer numerical model [...] Read more.
Longitudinal crack defects are a frequent occurrence on the surface of thin slabs during the high-speed continuous casting process. Therefore, this study undertakes a detailed analysis of the solidification characteristics of hyper-peritectic steel thin slabs. By establishing a three-dimensional heat transfer numerical model of the thin slab, the formation mechanism of longitudinal cracks caused by uneven growth of the initial shell is determined. Based on the mechanism of longitudinal crack formation, by adjusting the performance parameters of the mold flux, the contradiction between the heat transfer control and lubrication improvement of the mold flux is fully coordinated, further reducing the incidence of longitudinal cracks on the surface of the casting thin slab. The results show that, using the optimized mold flux, the basicity increases from 1.60 to 1.68, the F- mass fraction increases from 10.67% to 11.22%, the Na2O mass fraction increases from 4.35% to 5.28%, the Li2O mass fraction increases from 0.68% to 0.75%, and the carbon mass fraction reduces from 10.86% to 10.47%. The crystallization performance and rheological properties of the mold flux significantly improve, reducing the heat transfer performance while ensuring the lubrication ability of the molten slag. After optimizing the mold flux, a surface detection system was used to statistically analyze the longitudinal cracks on the surface of the casting thin slab. The proportion of longitudinal cracks (crack length/steel coil length, where each coil produced is about 32 m long) on the surface of the thin slab decreases from 0.056% to 0.031%, and the surface quality of the thin slab significantly improves. Full article
(This article belongs to the Special Issue Green and Intelligent Steelmaking Technologies with Low Carbon Emissions)
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<p>(<b>a</b>) Phase diagram of solidification; (<b>b</b>) solidification two-phase zone transition diagram.</p> Full article ">Figure 2
<p>(<b>a</b>) The surface temperatures of blank shells at different casting speeds; (<b>b</b>) the location of solid phase transformation on the surface of thin slab shells at different casting speeds.</p> Full article ">Figure 3
<p>Thickness of thin slab shell at different casting speeds.</p> Full article ">Figure 4
<p>Longitudinal crack morphology of hyper-peritectic steel continuous casting thin slab.</p> Full article ">Figure 5
<p>Schematic diagram of SHTT-II crystallization temperature tester.</p> Full article ">Figure 6
<p>Schematic diagram of viscosity experimental setup.</p> Full article ">Figure 7
<p>Temperature control curve of melting point and melting rate tester.</p> Full article ">Figure 8
<p>(<b>a</b>) Mold heat flow density; (<b>b</b>) liquid slag layer thickness and mold flux consumption statistics.</p> Full article ">Figure 9
<p>(<b>a</b>) The overall surface condition of the casting thin slab after using the B-type mold flux; (<b>b</b>) local conditions of the surface of the casting thin slab after acid washing treatment.</p> Full article ">
16 pages, 6646 KiB  
Article
Detrimental Effects of βo-Phase on Practical Properties of TiAl Alloys
by Toshimitsu Tetsui and Kazuhiro Mizuta
Metals 2024, 14(8), 908; https://doi.org/10.3390/met14080908 - 9 Aug 2024
Viewed by 240
Abstract
The TNM alloy, a βo-phase-containing TiAl alloy, has been withdrawn from use as a last-stage turbine blade in commercial jet engines as it suffered frequent impact fractures in service, raising doubts regarding the necessity of the βo-phase in practical [...] Read more.
The TNM alloy, a βo-phase-containing TiAl alloy, has been withdrawn from use as a last-stage turbine blade in commercial jet engines as it suffered frequent impact fractures in service, raising doubts regarding the necessity of the βo-phase in practical TiAl alloys. Here, we evaluate the practical properties required for jet engine blades for various TiAl alloys and investigate the effects of the βo-phase thereupon. First, we explore the influence of the βo-phase content on the impact resistance and machinability for forged Ti–43.5Al–xCr and cast Ti–46.0Al–xCr alloys; the properties deteriorate significantly at increasing βo-phase contents. Subsequently, two practical TiAl alloys—TNM alloy and TiAl4822—were prepared with and without the βo-phase by varying the heat treatment temperature for the former and the Cr concentration for the latter. In addition to impact resistance and machinability, the creep strength is significantly reduced by the presence of the βo-phase. Overall, these findings suggest that the βo-phase is an undesirable phase in practical TiAl alloys, especially those used for jet engine blades, because, although the disordered β-phase is soft at high temperatures, it changes to significantly more brittle and harder βo-phase after cooling. Full article
(This article belongs to the Section Metal Casting, Forming and Heat Treatment)
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<p>Appearance of Ti–43.5Al–xCr (at. %) alloys hot-forged after heating at 1330 °C.</p> Full article ">Figure 2
<p>Backscattered electron images showing microstructures of forged Ti–43.5Al–xCr ternary alloys subjected to the 1280 °C/5 h/furnace cooling heat treatment protocol after hot-forged at 1330 °C: x = (<b>a</b>) 2.0Cr, (<b>b</b>) 2.5Cr, (<b>c</b>) 3.0Cr, (<b>d</b>) 3.5Cr, and (<b>e</b>) 4.0Cr.</p> Full article ">Figure 3
<p>Backscattered electron images showing microstructures of cast Ti–46.0Al–xCr ternary alloys subjected to hot isostatic pressing (HIP) using the 1200 °C/4 h/186 MPa protocol: x = (<b>a</b>) 2.0Cr, (<b>b</b>) 2.5Cr, (<b>c</b>) 3.0Cr, (<b>d</b>) 3.5Cr, and (<b>e</b>) 4.0Cr.</p> Full article ">Figure 4
<p>Relationship between Cr concentration and ratio of α<sub>2</sub>- and β<sub>o</sub>-phase in (<b>a</b>) forged Ti–43.5Al–xCr alloys and (<b>b</b>) cast Ti–46.0Al–xCr alloys.</p> Full article ">Figure 5
<p>Relationship between the area ratio of α<sub>2</sub>- and β<sub>o</sub>-phase and the mean absorbed energy obtained in the Charpy impact test at RT and 700 °C for (<b>a</b>) forged Ti–43.5Al–xCr alloys and (<b>b</b>) cast Ti–46.0Al–xCr alloys.</p> Full article ">Figure 6
<p>Results of machining tests conducted on forged Ti–43.5Al–xCr and cast Ti–46.0Al–xCr alloys, showing the relationship between Cr concentration and tool weight loss at each Cr incorporation stage.</p> Full article ">Figure 7
<p>Backscattered electron images showing microstructures of modified TNM alloys hot-forged at 1330 °C and then subjected to the first heat treatment stage at (<b>a</b>) 1177, (<b>b</b>) 1207, and (<b>c</b>) 1237 °C for 3 h followed by air cooling and subjected to the second heat treatment stage at 850 °C for 6 h followed by furnace cooling.</p> Full article ">Figure 8
<p>Backscattered electron images showing microstructures of modified TiAl4822 alloys after HIP at 1200 °C/4 h/186 MPa: (<b>a</b>) Ti–47.0Al–2.0Nb–1.79Cr, (<b>b</b>) Ti–47.0Al–2.0Nb–1.91Cr, and (<b>c</b>) Ti–47.0Al–2.0Nb–2.14Cr.</p> Full article ">Figure 9
<p>Results of machining tests conducted on the modified TNM and TiAl4822 alloys, showing the relationship between the (<b>a</b>) first-stage heat treatment temperature and tool weight loss for modified TNM alloys, and (<b>b</b>) Cr content and tool weight loss for modified TiAl4822 alloys.</p> Full article ">Figure 10
<p>Creep curves of (<b>a</b>) modified TNM alloys and (<b>b</b>) modified TiAl4822 alloys subjected to a creep test performed at 750 °C/225 MPa.</p> Full article ">Figure 11
<p>Microstructures near ruptured position in creep specimens, showing the locations of creep voids: (<b>a</b>) modified TNM alloy subjected to the first-stage heat treatment at 1177 °C and (<b>b</b>) modified TiAl4822 alloy with 2.14Cr.</p> Full article ">Figure 12
<p>Results obtained from tests conducted to analyze the microstructure, Vickers hardness (HV), and brittleness at room temperature for alloys aimed at a single phase corresponding to each phase in the Ti–Al–Cr ternary alloy system at 1200 °C.</p> Full article ">
11 pages, 16463 KiB  
Article
Influence of Inclusion Parameter and Depth on the Rotating Bending Fatigue Behavior of Bearing Steel
by Lijun Xu, Zhonghua Zhan and Shulan Zhang
Metals 2024, 14(8), 907; https://doi.org/10.3390/met14080907 - 9 Aug 2024
Viewed by 288
Abstract
Inclusions are an important parameter affecting the fatigue life of materials. In this paper, the type, size, and quantity of inclusions in bearing steel were quantitatively analyzed using scanning electron microscopy and automatic scanning electron microscopy with an X-ray energy dispersive spectroscopy function. [...] Read more.
Inclusions are an important parameter affecting the fatigue life of materials. In this paper, the type, size, and quantity of inclusions in bearing steel were quantitatively analyzed using scanning electron microscopy and automatic scanning electron microscopy with an X-ray energy dispersive spectroscopy function. The effects of the inclusion parameters and positions on the rotating bending fatigue properties were analyzed using the rotating bending fatigue test. The results proved that for samples 1 and 2, the inclusions were mainly sulfides, Ti-containing inclusions, and their composite inclusions. For samples 3 and 4, the inclusions were mainly oxides or sulfide–oxide complexes. The number and maximum size of inclusions in sample 2 were relatively small. This was mainly due to the difference in the content of Al, S, and Ca elements in the different samples. The inclusion distance to the surface and the maximum inclusion size had a larger influence on the rotating bending fatigue life in comparison to the inclusion type. Moreover, nitride–oxides had a more detrimental effect on the rotating bending fatigue life as compared to the sulfide–oxide complex inclusions. A model was established on the basis of the inclusion size, depth, and stress by using the Python software. The simulation demonstrated that using five parameters fit well with the experiment results. Full article
(This article belongs to the Special Issue Inclusion Metallurgy (2nd Edition))
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<p>Specimen geometry and size for the rotating bending fatigue test.</p> Full article ">Figure 2
<p>Inclusion images of samples 1#, 2#, 3#, and 4#: (<b>a</b>) sulfides in sample 1#; (<b>b</b>) oxides in sample 2#; (<b>c</b>) Ti carbides or nitrides in sample 3#; (<b>d</b>) complexes of sulfides and oxides in sample 4#.</p> Full article ">Figure 3
<p>The number of cycles to failure obtained from the rotating bending fatigue tests for samples under different stress: (<b>a</b>) sample 1#; (<b>b</b>) sample 2#; (<b>c</b>) sample 3#; (<b>d</b>) sample 4#.</p> Full article ">Figure 4
<p>SEM fracture images showing fatigue cracks initiated at inclusions and EDS results of inclusions of samples 2# and 4# subjected to cyclic stress of 1400 MPa: (<b>a</b>) sample 2#; (<b>b</b>) magnification of Figure (<b>a</b>); (<b>c</b>) EDS result of (<b>b</b>) inclusion of (<b>e</b>); (<b>d</b>) sample 4#; (<b>e</b>) magnification of (<b>d</b>); (<b>f</b>) EDS result of (<b>e</b>) inclusion.</p> Full article ">Figure 5
<p>SEM fracture images showing fatigue fracture morphology: (<b>a</b>) sample 2#; (<b>b</b>) sample 4#.</p> Full article ">Figure 6
<p>The relationship among number of cycles to failure, inclusion size, and depth for the samples subjected to different stresses: (<b>a</b>) 1300 MPa; (<b>b</b>) 1400 MPa.</p> Full article ">Figure 7
<p>The relationship among number of cycles to failure, inclusion type, and depth for the four kinds of samples subjected to stress of 1600 MPa.</p> Full article ">Figure 8
<p>Simulation results for the round and MSE.</p> Full article ">
20 pages, 2991 KiB  
Article
Analysis of the Feeding Behavior in a Bottom-Blown Lead-Smelting Furnace
by Kena Sun, Xiaowu Jie, Yonglu Zhang, Wei Gao, Derek O. Northwood, Kristian E. Waters and Hao Ma
Metals 2024, 14(8), 906; https://doi.org/10.3390/met14080906 - 9 Aug 2024
Viewed by 293
Abstract
Computational fluid dynamics (CFD) software was used to simulate the feeding behavior in a bottom-blown lead-smelting furnace. The results show that when the particle size is less than 30 μm, 20% of the particles are suspended in the gas phase and do not [...] Read more.
Computational fluid dynamics (CFD) software was used to simulate the feeding behavior in a bottom-blown lead-smelting furnace. The results show that when the particle size is less than 30 μm, 20% of the particles are suspended in the gas phase and do not enter the melt pool for smelting, thus resulting in material loss. When the particle size exceeds 75 μm, the particles settle in the metal layer. When the particle size is 40–60 μm, the particles are distributed in the slag and metal phases, and the material is uniformly distributed in the molten pool; additionally, the average velocity of the particles exceeds 1.4 m/s, the average temperature exceeds 960 K, and the particles exhibit better behavior within this range, thus rendering it the optimal range of particle sizes for feeding. Full article
(This article belongs to the Special Issue Modeling Thermodynamic Systems and Optimizing Metallurgical Processes)
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<p>Bottom-blown lead furnace structural model, 1—oxygen lances, 2—feeds, 3—exhaust gas outlet.</p> Full article ">Figure 2
<p>Measurement of the slag’s physical property parameters: (<b>a</b>) viscosity, (<b>b</b>) density, (<b>c</b>) surface tension.</p> Full article ">Figure 3
<p>Velocity distribution of particles with different sizes.</p> Full article ">Figure 4
<p>Distributions of the average and maximum velocity of the particles.</p> Full article ">Figure 5
<p>Temperature distribution of particles with different sizes.</p> Full article ">Figure 6
<p>Distributions of the average and maximum temperatures of the particles.</p> Full article ">Figure A1
<p>Bubble morphology diagram of numerical calculation and water model experiment, (<b>a</b>) Numerical calculation of bubble morphology at the outlet of the spray gun; (<b>b</b>) Bubble morphology at the outlet of the spray gun in water model experiment, Adapted from Ref. [<a href="#B34-metals-14-00906" class="html-bibr">34</a>].</p> Full article ">
14 pages, 3601 KiB  
Article
Analyzing Impact of Processing Parameters and Material Properties on Symmetry of Wire-Arc Directed Energy Deposit Beads
by Stephen Price, Kiran Judd, Matthew Gleason, Kyle Tsaknopoulos, Danielle L. Cote and Rodica Neamtu
Metals 2024, 14(8), 905; https://doi.org/10.3390/met14080905 - 9 Aug 2024
Viewed by 407
Abstract
Wire arc-directed energy deposit (wire-arc DED) enables the efficient manufacturing of large-scale metal parts. Many factors can impact overall part quality during manufacturing, including processing parameters such as feed rate, travel speed, and various material properties. Previous works have sought to use machine [...] Read more.
Wire arc-directed energy deposit (wire-arc DED) enables the efficient manufacturing of large-scale metal parts. Many factors can impact overall part quality during manufacturing, including processing parameters such as feed rate, travel speed, and various material properties. Previous works have sought to use machine learning to evaluate and predict these impacts, but they have primarily focused on the width and height of single-layer deposits. Building upon these studies, this work offers a novel technique to characterize and evaluate the asymmetry of deposited beads to better understand the impact these parameters have. Specifically, leveraging time-series analysis techniques, the surface profiles of beads can be compared and analyzed to identify the degree of asymmetry. Additionally, this work relates these factors to the extent to which substrates warp during the depositing of material. With a better understanding of these impacts, manufacturing processes can be optimized for improved quality and reduced waste. These findings highlight that, while material selection and processing parameters do not strongly correlate with bead asymmetry, beads are deposited with varying degrees of asymmetry, requiring further analysis to identify the source. In contrast, substrate warping is significantly influenced by the thermal properties of the materials used. Of the properties analyzed, heat capacity, thermal diffusivity and thermal conductivity were found to be most relevant to substrate warping. Additionally, while to a lesser extent, material properties of the wire were found to be similarly correlated to warping as their substrate counterparts. These insights can inform the optimization of manufacturing processes, leading to improved part quality and reduced material waste. This study also underscores the need for further research into the interplay between processing conditions and material characteristics in wire-arc DED. Full article
(This article belongs to the Special Issue Editorial Board Members’ Collection Series: Additive Manufacturing Technology)
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Graphical abstract
Full article ">Figure 1
<p>Validating edge cases of calculating asymmetry with a minimum and maximum asymmetry.</p> Full article ">Figure 2
<p>Overview of the process to calculate substrate curvature as a result of warping.</p> Full article ">Figure 3
<p>Distribution of the number of beads deposited per substrate.</p> Full article ">Figure 4
<p>Front and back of 4140 steel substrate with beads deposited on the top and bottom, creating a cubic warp.</p> Full article ">Figure 5
<p>Side profile of cubic warping of 4140 steel substrate.</p> Full article ">Figure 6
<p>Highlighting differing levels of symmetry in deposited beads’ surface profiles.</p> Full article ">Figure 7
<p>Correlation matrix relating processing parameters, bead dimensions, and bead location to asymmetry.</p> Full article ">Figure 8
<p>Correlation matrix relating material to asymmetry.</p> Full article ">
12 pages, 4369 KiB  
Article
Catalytic Activity Evaluation of the Molten Salt-Modified Novel Ni Electrodes for Urea Electrooxidation in Alkaline Solutions
by Dawid Kutyła, Michihisa Fukumoto, Hiroki Takahashi, Marek Wojnicki and Piotr Żabiński
Metals 2024, 14(8), 904; https://doi.org/10.3390/met14080904 - 9 Aug 2024
Viewed by 393
Abstract
The presented paper characterized the molten salt-modified Ni electrode with excellent catalytic activity towards alkaline urea electrooxidation reaction. The electrodes were modified by electrodeposition of Al from molten salt electrolytes containing NaCl-KCl-AlF3 at a temperature of 750 °C and applied potential of [...] Read more.
The presented paper characterized the molten salt-modified Ni electrode with excellent catalytic activity towards alkaline urea electrooxidation reaction. The electrodes were modified by electrodeposition of Al from molten salt electrolytes containing NaCl-KCl-AlF3 at a temperature of 750 °C and applied potential of −1.9 V. The porous surface was obtained by anodic polarization with a potential of −0.4 V until the anodic current was equal to 0 mAcm−2. The prepared deposits’ structure, surface morphology, and composition were analyzed using scanning electron microscopy (SEM) and X-ray diffraction (XRD). Anodic polarization was applied to assess the electrocatalytic activity and elucidate the urea electrooxidation mechanism in 1 M KOH + 0.33 M urea solution. The nanocrystalline structure, fine grain size, and microcracks on the surface of the studied electrodes contributed to their notably high electrochemically active surface area (ECSA). The cyclic voltammetry in the non-Faradaic regions of the samples shows that molten salt modification can increase the double layer capacitance of bare Ni plates by around ten times, from 0.29 mFcm−2 to 2.16 mFcm−2. Polarization of the electrodes in urea-containing KOH solution with potential of +1.52 V shows a significant difference in catalytic performance. For the bare nickel sample, the registered current density from the urea electrooxidation reaction was around +1 mAcm−2, and for the molten salt-modified one, it was +38 mAcm−2, which indicates the fact that the molten salt surface treatment can be a promising tool in tailoring the electrochemical properties of materials. Full article
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<p>Graphical representation of nickel electrode molten salt treatment via the deposition/dissolution process of Al.</p> Full article ">Figure 2
<p>Current vs. time plot of deposition/dissolution of Al during molten salt treatment of Ni electrode. Deposition time: 3600 s, applied potential: −1.9 V. Dissolution time: 5400 s, applied potential: −0.5 V. Temperature of molten salt treatment: 900 °C.</p> Full article ">Figure 3
<p>SEM images of the Ni porous electrode obtained by molten salt deposition/dissolution of Al. Figure captions and magnifications: (<b>A</b>)—500×, (<b>B</b>)—1000×, (<b>C</b>)—2000×, (<b>D</b>)—4000×.</p> Full article ">Figure 4
<p>X-ray diffraction patterns of the unmodified Ni and Ni porous electrode obtained by molten salt deposition/dissolution of Al. Scan speed: 0.25 deg/min.</p> Full article ">Figure 5
<p>The cyclic voltammetry scans with different sweep rates registered for the unmodified Ni (<b>left</b>) and Ni porous (<b>right</b>) electrodes in 1 M NaOH electrolyte.</p> Full article ">Figure 6
<p>The cyclic voltammetry scans with different sweep rates were registered for the unmodified Ni and Ni porous electrode in 1 M NaOH electrolyte.</p> Full article ">Figure 7
<p>The linear voltammetry scans registered for the unmodified Ni and Ni porous electrode in 1 M NaOH with and without the 0.33 M urea in electrolyte.</p> Full article ">Figure 8
<p>The chronoamperometric scans registered for the unmodified Ni and Ni porous electrodes in 1 M NaOH with and without the 0.33 M urea in electrolyte.</p> Full article ">
15 pages, 10315 KiB  
Article
Corrosion Behaviors of Weathering Steels in the Actual Marine Atmospheric Zone and Immersion Zone
by Ying Yang, Tianzi Lin, Guohui Wang, Yubo Wang, Minghui Shao, Fandi Meng and Fuhui Wang
Metals 2024, 14(8), 903; https://doi.org/10.3390/met14080903 - 9 Aug 2024
Viewed by 320
Abstract
The corrosion behaviors of three bridge steels in a real tropical marine environment for 2 years were studied. One weathering steel (WS) was designed with higher levels of nickel, copper, and molybdenum compared to the other. These two kinds of WSs and one [...] Read more.
The corrosion behaviors of three bridge steels in a real tropical marine environment for 2 years were studied. One weathering steel (WS) was designed with higher levels of nickel, copper, and molybdenum compared to the other. These two kinds of WSs and one kind of ordinary high-strength low-alloy steel (Q345qe) were compared under two conditions (marine atmospheric zone and marine immersion zone at Sanya Marine Environmental Test Station). The morphology, corrosion rate, and corrosion product analysis of the steels were performed through SEM, XPS, FTIR and other characterization methods. The results demonstrated that weathering steels facilitate the densification of the corrosion product layer due to the addition of alloying elements Cr, Ni, and Cu, promoting rust nucleation and enhancing the compactness of the protective layer. However, in an immersion environment, the extensive erosion by chloride ions renders the benefits of WS ineffective. Full article
(This article belongs to the Special Issue Corrosion of Metals: Behaviors and Mechanisms)
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<p>Climatic data during the exposure period at Sanya station: (<b>a</b>) the average temperature for each month; (<b>b</b>) the average relative humidity for each month. Photos of the corrosion test equipment and samples in the (<b>c</b>) marine atmospheric zone and (<b>d</b>) marine immersion zone. (<b>e</b>) Schematic diagram of the size of the samples.</p> Full article ">Figure 2
<p>Metallographic structure of bridge steels: (<b>a</b>) #1 steel; (<b>b</b>) #2 steel; and (<b>c</b>) #3 steel.</p> Full article ">Figure 3
<p>Mass changes and corrosion rates of the three steels at different time points for (<b>a</b>,<b>b</b>) atmospheric zone and (<b>c</b>,<b>d</b>) immersion zone.</p> Full article ">Figure 4
<p>Corrosion depths of the three steels after two years for (<b>a</b>) atmospheric zone and (<b>b</b>) immersion zone.</p> Full article ">Figure 5
<p>Macroscopic corrosion morphology of three kinds of steels at different time points for (<b>a</b>) atmospheric zone and (<b>b</b>) immersion zone.</p> Full article ">Figure 6
<p>Morphology of substrates after complete removal of rust for (<b>a</b>) atmospheric zone and (<b>b</b>) immersion zone.</p> Full article ">Figure 6 Cont.
<p>Morphology of substrates after complete removal of rust for (<b>a</b>) atmospheric zone and (<b>b</b>) immersion zone.</p> Full article ">Figure 7
<p>Microscopic corrosion morphology of 1# steel in atmospheric zone at different times: (<b>a</b>,<b>b</b>) 0.5 years and its high magnification micrograph; (<b>c</b>,<b>d</b>) 1 year and its high magnification micrograph with the feature of clustered corrosion products (red plus sign); (<b>e</b>,<b>f</b>) 1.5 years and its high magnification micrograph; (<b>g</b>,<b>h</b>) 2 years and its high magnification micrograph.</p> Full article ">Figure 8
<p>Microscopic corrosion morphology of 2# steel in atmospheric zone at different times: (<b>a</b>,<b>b</b>) 0.5 years and its high magnification micrograph; (<b>c</b>,<b>d</b>) 1 year and its high magnification micrograph with the feature of rod clustered corrosion products (red plus sign); (<b>e</b>,<b>f</b>) 1.5 years and its high magnification micrograph; (<b>g</b>,<b>h</b>) 2 years and its high magnification micrograph.</p> Full article ">Figure 9
<p>Microscopic corrosion morphology of 3# steel in atmospheric zone at different times: (<b>a</b>,<b>b</b>) 0.5 years and its high magnification micrograph; (<b>c</b>,<b>d</b>) 1 year and its high magnification micrograph with the feature of needle-like corrosion products (red plus sign); (<b>e</b>,<b>f</b>) 1.5 years and its high magnification micrograph; (<b>g</b>,<b>h</b>) 2 years and its high magnification micrograph.</p> Full article ">Figure 10
<p>Microscopic corrosion morphology of 1# steel in immersion zone at different times: (<b>a</b>,<b>b</b>) 0.5 years and its high magnification micrograph; (<b>c</b>,<b>d</b>) 1 year and its high magnification micrograph with the feature of spherical corrosion products (red plus sign); (<b>e</b>,<b>f</b>) 1.5 years and its high magnification micrograph; (<b>g</b>,<b>h</b>) 2 years and its high magnification micrograph.</p> Full article ">Figure 11
<p>Microscopic corrosion morphology of 2# steel in immersion zone at different times: (<b>a</b>,<b>b</b>) 0.5 years and its high magnification micrograph; (<b>c</b>,<b>d</b>) 1 year and its high magnification micrograph with the feature of fine needles corrosion products (red plus sign); (<b>e</b>,<b>f</b>) 1.5 years and its high magnification micrograph; (<b>g</b>,<b>h</b>) 2 years and its high magnification micrograph.</p> Full article ">Figure 12
<p>Microscopic corrosion morphology of 3# steel in immersion zone at different times: (<b>a</b>,<b>b</b>) 0.5 years and its high magnification micrograph; (<b>c</b>,<b>d</b>) 1 year and its high magnification micrograph with the feature of blocky and strip-like corrosion products (red plus sign); (<b>e</b>,<b>f</b>) 1.5 years and its high magnification micrograph; (<b>g</b>,<b>h</b>) 2 years and its high magnification micrograph.</p> Full article ">Figure 13
<p>XRD results of the corrosion products of steels exposed to the marine atmospheric environment for 0.5 years.</p> Full article ">Figure 14
<p>FTIR results of different steels in the atmospheric zone after 2 years.</p> Full article ">Figure 15
<p>XRD results of the corrosion products of steels immersed in the marine immersion zone for 0.5 years.</p> Full article ">Figure 16
<p>FTIR results of different steels in the immersion zone after 2 years.</p> Full article ">
15 pages, 8475 KiB  
Article
Effect of Si Addition in NiCrAl Coating on Corrosion in Molten Nitrate Salt
by Thamrongsin Siripongsakul, Patchaporn Kettrakul, Noparat Kanjanaprayut and Piyorose Promdirek
Metals 2024, 14(8), 902; https://doi.org/10.3390/met14080902 - 8 Aug 2024
Viewed by 318
Abstract
The materials used in concentrating solar power (CSP) systems are becoming of interest because of the high energy efficiency of energy storage. Molten salts can be used as both heat-storage media and heat-transfer fluid in a CSP system. In molten salts, steel alloys [...] Read more.
The materials used in concentrating solar power (CSP) systems are becoming of interest because of the high energy efficiency of energy storage. Molten salts can be used as both heat-storage media and heat-transfer fluid in a CSP system. In molten salts, steel alloys used in vessels and pipelines are highly vulnerable to hot corrosion. To protect steel alloys, applying a coating is an excellent strategy to extend the life of the alloy. NiCrAl coatings are well-suited for high-temperature environments. The purpose of this study was to investigate the corrosion behavior of NiCrAl with Si addition coatings on AISI304 in molten salt. NiCrAl coatings with and without Si addition were deposited using the high-velocity oxygen fuel (HVOF) technique. The corrosion test was performed using an immersion test in modified solar salt with 0.5% NaCl at 400–600 °C. Gravimetric methods evaluate the weight change for immersion tests. At 400 °C, an increased amount of weight gain due to the oxidation reaction and molten salt infiltration was observed. At 600 °C, the corrosion reaction was more dominant, and apparent oxidation was decreased; however, oxidation products NiO and sodium aluminum silicate were detected. Si addition supports the formation of the protective oxide sodium aluminum silicate, which inhibits molten salt oxidation reaction and molten salt infiltration. Full article
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<p>Cross-sections of the coatings taken using an optical microscope: (<b>a</b>) NiCrAl coating, (<b>b</b>) 2.7% Si + NiCrAl coating.</p> Full article ">Figure 2
<p>Mass change of coated samples after immersion in modified molten nitrate salt mixtures at (<b>a</b>) 400 °C, (<b>b</b>) 500 °C, and (<b>c</b>) 600 °C and mass change rate (<b>d</b>) 400 °C, (<b>e</b>) 500 °C, and (<b>f</b>) 600 °C for 0, 25, 50, and 100 h.</p> Full article ">Figure 3
<p>Cross-section morphologies of (<b>a</b>) a NiCrAl-coated sample, (<b>b</b>) a NiCrAl + 2.7% Si-coated sample in the molten salt at 400 °C for 25 h, (<b>c</b>) a NiCrAl-coated sample, (<b>d</b>) a NiCrAl + 2.7% Si-coated sample in the molten salt at 600 °C for 100 h, (<b>e</b>,<b>f</b>) magnified cross-section morphologies of (<b>c</b>) and (<b>d</b>), respectively.</p> Full article ">Figure 4
<p>EDS mapping of (<b>a</b>) a NiCrAl-coated sample, (<b>b</b>) a NiCrAl + 2.7% Si-coated sample in the molten salt at 400 °C for 25 h, (<b>c</b>) a NiCrAl-coated sample, (<b>d</b>) a NiCrAl + 2.7% Si-coated sample in the molten salt at 600 °C for 100 h.</p> Full article ">Figure 4 Cont.
<p>EDS mapping of (<b>a</b>) a NiCrAl-coated sample, (<b>b</b>) a NiCrAl + 2.7% Si-coated sample in the molten salt at 400 °C for 25 h, (<b>c</b>) a NiCrAl-coated sample, (<b>d</b>) a NiCrAl + 2.7% Si-coated sample in the molten salt at 600 °C for 100 h.</p> Full article ">Figure 5
<p>XRD pattern of (<b>a</b>) NiCrAl and (<b>b</b>) NiCrAl + 2.7% Si after immersion in modified molten solar salt at 400 °C for 25 h.</p> Full article ">Figure 6
<p>XRD pattern of (<b>a</b>) NiCrAl and (<b>b</b>) NiCrAl + 2.7% Si after immersion in modified molten solar salt at 600 °C for 100 h.</p> Full article ">
13 pages, 12851 KiB  
Article
Synergistic Effect of Phase Transformation and Stress-Induced Twinning on the Antibacterial Property and Elastic Modulus of Ti-13Nb-13Zr-7Ag
by Diangeng Cai, Jiayu Wang, Lei Yang, Xiaocen Xu and Erlin Zhang
Metals 2024, 14(8), 901; https://doi.org/10.3390/met14080901 - 8 Aug 2024
Viewed by 290
Abstract
Ti-13Nb-13Zr-7Ag (TNZ-7Ag) has a great potential for biomedical application due to its low elastic modulus and excellent antibacterial properties. However, it is difficult to balance low elastic modulus and high antibacterial properties. In this study, the TNZ-7Ag was treated by a pre-deformation and [...] Read more.
Ti-13Nb-13Zr-7Ag (TNZ-7Ag) has a great potential for biomedical application due to its low elastic modulus and excellent antibacterial properties. However, it is difficult to balance low elastic modulus and high antibacterial properties. In this study, the TNZ-7Ag was treated by a pre-deformation and aging treatment to avoid this problem. The results proved that the stress-induced twinning caused a large range of Ti2Ag particle agglomeration, and in turn enhanced the antibacterial performance of Ag-containing titanium alloys. And the twinned martensite formed during the pre-deformation promoted the precipitation of Ti2Ag phase and inhibited the growth of α phase. As a result, TNZ-7Ag with both low elastic modulus and strong antibacterial properties was obtained by the treatments. All results demonstrated that pre-deformation based on the synergistic effect of aging treatment was an effective strategy to develop novel biomedical titanium alloys. Low-elastic-modulus antibacterial titanium alloys, which can be used for the development of novel biomedical titanium alloys, were prepared. Full article
(This article belongs to the Special Issue Design, Phase Transformation and Mechanical Properties of Titanium Alloy)
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<p>XRD patterns of TNZ-7Ag with different deformation. (<b>a</b>) XRD patterns (the T4 condition listed in red), (<b>b</b>) the zoomed-in region from 50–65°.</p> Full article ">Figure 2
<p>SEM microstructure of TNZ-7Ag with different deformation. (<b>a</b>) 20%, (<b>b</b>) 30%, (<b>c</b>) 40%, (<b>d</b>) 50%.</p> Full article ">Figure 3
<p>TEM microstructure of TNZ-7Ag with 50% deformation. (<b>a</b>) TNZ-7Ag with 50% deformation, (<b>b</b>) bright-field, (<b>c</b>) dark-field images and (<b>d</b>) SAED of twins and (<b>e</b>) the image and (<b>f</b>) SAED of β and β′ phase.</p> Full article ">Figure 4
<p>Element mapping and EDS results of β-phase separation. (<b>a</b>) STEM, (<b>b</b>) Ag, (<b>c</b>) Nb, (<b>d</b>) Zr, (<b>e</b>) Ti, (<b>f</b>) EDS results.</p> Full article ">Figure 5
<p>XRD patterns of TNZ-7Ag with 50% pre-deformation and different aging durations. (<b>a</b>) XRD, (<b>b</b>) the percentage of β phase.</p> Full article ">Figure 6
<p>The SEM microstructure of TNZ-7Ag with different pre-deformation + aging treatments. (<b>a</b>,<b>e</b>) 50%/0.5 h, (<b>b</b>,<b>f</b>) 50%/1 h, (<b>c</b>,<b>g</b>) 50%/1.5 h, (<b>d</b>,<b>h</b>) 50%/2 h.</p> Full article ">Figure 7
<p>TEM microstructure of TNZ-7Ag with 50%-deformation + 600 °C aging treatments. (<b>a</b>) 50%/0.5 h, (<b>b</b>) 50%/1.5 h, (<b>c</b>) 50%/2 h, (<b>d</b>,<b>e</b>) high-resolution microstructure of 50%/2 h, (<b>f</b>,<b>g</b>) Fourier transform map.</p> Full article ">Figure 8
<p>The elastic modulus of TNZ-7Ag with 50% pre-deformation and aging for 0.5–2 h, TNZ-7Ag without pre-deformation and Ti-13Nb-13Zr (data from ref. [<a href="#B11-metals-14-00901" class="html-bibr">11</a>]) listed in red.</p> Full article ">Figure 9
<p>Tafel curves and data results of TNZ-7Ag with 50% pre-deformation and aged at 600 °C for 0.5–2 h. (<b>a</b>) Tafel curves, (<b>b</b>) corrosion current density.</p> Full article ">Figure 10
<p>Ag ion concentration and release rates from TNZ-7Ag alloys with 50% pre-deformation and aged for 0.5–2 h after 1 day, 3 days, and 7 days of immersion. (<b>a</b>) Ag ion concentration, (<b>b</b>) Ag ion release rate.</p> Full article ">Figure 11
<p>Bacterial colonies on TNZ-7Ag with 50% deformation + 600 °C aging for 0.5–2 h and the antibacterial rate. (<b>a</b>) cp Ti, (<b>b</b>) 0.5 h, (<b>c</b>) 1 h, (<b>d</b>) 1.5 h, (<b>e</b>) 2 h, (<b>f</b>) antibacterial rate (* <span class="html-italic">p</span> &gt; 0.05, n = 3).</p> Full article ">Figure 12
<p>The OD value and relative growth rate (RGR) of MC3T3-E1 with the extracts of different samples after 1 day, 3 days, and 7 days of culture. (<b>a</b>) OD values, (<b>b</b>) RGR (* <span class="html-italic">p</span> &gt; 0.05, n = 3).</p> Full article ">
14 pages, 8890 KiB  
Article
Formation of Metastable Solid Solutions in Bi-Ge Films during Low-Temperature Treatment
by Sergiy Bogatyrenko, Pavlo Kryshtal, Adam Gruszczyński and Aleksandr Kryshtal
Metals 2024, 14(8), 900; https://doi.org/10.3390/met14080900 - 8 Aug 2024
Viewed by 312
Abstract
We investigated the mechanism and kinetics of the formation of metastable BiGe solid phases during the amorphous-to-crystalline transformation of Ge films in contact with Bi. Ge/Bi/Ge sandwich films with a Bi film between amorphous Ge films, which were fabricated by sequential deposition of [...] Read more.
We investigated the mechanism and kinetics of the formation of metastable BiGe solid phases during the amorphous-to-crystalline transformation of Ge films in contact with Bi. Ge/Bi/Ge sandwich films with a Bi film between amorphous Ge films, which were fabricated by sequential deposition of the components in a vacuum, were used in this study. The total thickness and composition of the sandwich films varied in the range from 30 to 400 nm and from 22 to 48 wt% Bi, respectively. Electron diffraction, high-resolution (S)TEM imaging, EDX, and EEL spectroscopy were used for in situ and ex situ characterization of the morphology, composition, and structure of Ge/Bi/Ge films in the temperature range of 20–271 °C. We proved the formation of polycrystalline Ge films containing up to 28 wt% Bi during low-temperature treatment. The interaction process was activated at ≈150 °C, resulting in the crystallization of Ge with the simultaneous formation of a quasi-homogeneous supersaturated solid solution throughout the entire volume of the film at ≈210 °C. We showed that the formation of crystalline Ge films with an extended solid solubility of Bi depended mostly on the overall composition of the tri-layer film. The role of metal-induced crystallization of the amorphous germanium in the formation of the supersaturated solid phases is discussed. Full article
(This article belongs to the Special Issue Advances in Nanostructured Metallic Materials)
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Graphical abstract
Full article ">Figure 1
<p>The temperature dependencies of the Q-factor of a quartz resonator loaded by Bi-Ge three-layer films with different Ge film thicknesses during the heating–cooling cycle. Filled dots correspond to heating, and empty dots correspond to cooling.</p> Full article ">Figure 2
<p>HAADF STEM images of the morphology of <span class="html-italic">a</span>-Ge(12.5 nm)/Bi(5 nm)/<span class="html-italic">a</span>-Ge(12.5 nm) layered film at 20 °C (<b>a</b>), 100 °C (<b>b</b>), 150 °C (<b>c</b>), 180 °C (<b>d</b>), 195 °C (<b>e</b>), and 205 °C (<b>f</b>) during heating. Insets show the corresponding SAED patterns. The scale bar is the same for all images.</p> Full article ">Figure 3
<p>Radially averaged intensity profiles of the diffraction maxima of <span class="html-italic">a</span>-Ge(12.5 nm)/Bi(5 nm)/<span class="html-italic">a</span>-Ge(12.5 nm) layered film at different temperatures.</p> Full article ">Figure 4
<p>HAADF-STEM images of the same area of <span class="html-italic">a</span>-Ge(12.5 nm)/Bi(5 nm)/<span class="html-italic">a</span>-Ge(12.5 nm) layered film at a temperature of 150 °C after 30 (<b>a</b>) and 120 (<b>b</b>) minutes of annealing.</p> Full article ">Figure 5
<p>HAADF STEM images and corresponding EDX chemical element maps of <span class="html-italic">a</span>-Ge(12.5 nm)/Bi(5 nm)/<span class="html-italic">a</span>-Ge(12.5 nm) layered film at 150 °C (<b>a</b>–<b>c</b>), 205 °C (<b>d</b>–<b>f</b>), and 20 °C (<b>g</b>–<b>i</b>) (after cooling from 205 °C).</p> Full article ">Figure 6
<p>(<b>a</b>) HAADF-STEM image of <span class="html-italic">a</span>-Ge(12.5 nm)/Bi(5 nm)/<span class="html-italic">a</span>-Ge(12.5 nm) layered film at 150 °C. (<b>b</b>) and (<b>c</b>) show the fast Fourier transformations from the areas marked by orange and blue squares, respectively.</p> Full article ">Figure 7
<p>HAADF-STEM image (<b>a</b>) and SAED pattern (inset) of the BiGe metastable phase at 210 °C. (<b>b</b>) shows the overall EDX spectrum of the film.</p> Full article ">Figure 8
<p>HAADF-STEM images of lamellas of Ge/Bi/Ge layered system (<b>a</b>–<b>c</b>) at different temperatures and corresponding EDX chemical element maps (<b>d</b>–<b>f</b>). (<b>g</b>–<b>i</b>) show concentration profiles along yellow arrows in figure (<b>d</b>–<b>f</b>). Inset in (<b>b</b>) shows typical FFT from the solid solution region.</p> Full article ">Figure 9
<p>(<b>a</b>) HAADF-STEM image and (<b>b</b>) false-color map of the volume plasmon peak energy from the area marked by a yellow rectangle in (<b>a</b>) of Ge/Bi/Ge film annealed at 150 °C. (<b>c</b>) Low-loss EELS spectrum of the BiGe alloy obtained from the area marked by a red square in (<b>a</b>). (<b>d</b>) is a zoomed-in view of the plasmon peak of the BiGe alloy in (<b>c</b>) overlayed with the plasmon peaks of <span class="html-italic">a</span>-Ge (blue square in (<b>a</b>)) and Bi (green square in (<b>a</b>)).</p> Full article ">
7 pages, 3209 KiB  
Communication
Phase Mapping Using a Combination of Multi-Functional Scanning Electron Microscopy Detectors and Imaging Modes
by Gang Liu, Yonghua Zhao and Shuai Wang
Metals 2024, 14(8), 899; https://doi.org/10.3390/met14080899 - 7 Aug 2024
Viewed by 314
Abstract
Microstructure degradation and phase transformations are critical concerns in nickel-based superalloys during thermal exposure. Understanding the phase transformation mechanism requires the detailed mapping of the distribution of each phase at different degradation stages and in various precipitation sizes. However, differentiating between phases in [...] Read more.
Microstructure degradation and phase transformations are critical concerns in nickel-based superalloys during thermal exposure. Understanding the phase transformation mechanism requires the detailed mapping of the distribution of each phase at different degradation stages and in various precipitation sizes. However, differentiating between phases in large areas, typically on the scale of millimeters and often relying on scanning electron microscopy (SEM) techniques, has traditionally been a challenging task. In this study, we present a novel and efficient phase mapping method that leverages multiple imaging detectors and modes in SEM. This approach allows for the relatively rapid and explicit differentiation and mapping of the distribution of various phases, including MC, M23C6, γ′, and η phases, as demonstrated in a typical superalloy subjected to aging experiments at 800 °C. Full article
(This article belongs to the Special Issue State-of-Art: Metals Failure Analysis)
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Figure 1

Figure 1
<p>Example of bulk MC carbides on the grain boundary (GB) after heat treatment (before the aging experiment).(<b>a</b>) MC on the GB; (<b>b</b>) MC on the GBs junction.</p> Full article ">Figure 2
<p>Example 1 of phase mapping via image combination from different detectors and imaging modes. “ETD”, “ABS”, and “T1” are different detectors. “Standard” and “Optiplan” are different imaging modes. The specific meaning of these detectors and modes is explained in the manuscript. (<b>a</b>–<b>d</b>) Images from different combinations; (<b>e</b>) schematic of the phase mapping. The red arrow in (<b>c</b>) indicates the plate-like η phases.</p> Full article ">Figure 3
<p>Another example of phase mapping via image combination from different detectors and imaging modes. (<b>a</b>) “T1 + Optiplan”; (<b>b</b>) “ETD + Optiplan”; (<b>c</b>) “ABS + Standard”; (<b>d</b>) schematic of phase mapping; (<b>e</b>) SEM-EDS elemental mapping; (<b>f</b>) EPMA elemental mapping.</p> Full article ">Figure 4
<p>Phase identification in <a href="#metals-14-00899-f003" class="html-fig">Figure 3</a> via FIB lift-out and TEM characterization: (<b>a</b>) lift-out area as indicated by the box area; (<b>b</b>) image of the FIB specimen; (<b>c</b>) schematic of the phase mapping in (<b>b</b>); (<b>d</b>) TEM-EDS mapping of the interested area in (<b>b</b>); (<b>e</b>–<b>g</b>) images of MC carbides and representative elemental mapping in the box area of (<b>b</b>); (<b>h</b>) dark field image of dot-like MC carbides; (<b>i</b>–<b>k</b>) diffraction patterns of MC, η, and M<sub>23</sub>C<sub>6</sub>, respectively.</p> Full article ">
12 pages, 6747 KiB  
Article
Optimizing Tempering Parameters to Enhance Precipitation Behavior and Impact Toughness in High-Nickel Steel
by Guojin Sun and Qi Wang
Metals 2024, 14(8), 898; https://doi.org/10.3390/met14080898 - 7 Aug 2024
Viewed by 315
Abstract
This study explores the effects of tempering on the precipitation behavior and impact toughness of high-nickel steel. The specimens underwent double quenching at 870 °C and 770 °C, followed by tempering at various temperatures. Advanced characterization techniques including optical microscopy (OM), scanning electron [...] Read more.
This study explores the effects of tempering on the precipitation behavior and impact toughness of high-nickel steel. The specimens underwent double quenching at 870 °C and 770 °C, followed by tempering at various temperatures. Advanced characterization techniques including optical microscopy (OM), scanning electron microscopy (SEM), and transmission electron microscopy (TEM) were used to elucidate precipitation phenomena. Additionally, electron backscatter diffraction (EBSD) was employed to assess the misorientation distribution after tempering. Charpy impact tests were performed on specimens tempered at different temperatures to evaluate their toughness. The findings reveal that with increasing tempering temperature, the fraction of low-angle grain boundaries decreases, which correlates positively with enhanced impact toughness. The results demonstrate that tempering at 580 °C optimizes the material’s microstructure, achieving an impact toughness value of approximately 163 J. Full article
(This article belongs to the Special Issue Research on Mechanical Properties, Deformation, and Microstructures of Metallic Materials)
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Figure 1
<p>Dilatometry measurement of phase transformation temperatures.</p> Full article ">Figure 2
<p>Standard Charpy impact specimen (unit: mm).</p> Full article ">Figure 3
<p>SEM images of steel at various tempering temperatures for 2h. (<b>a</b>) As-quenched; (<b>b</b>) 500 °C; (<b>c</b>) 580 °C; (<b>d</b>) 620 °C.</p> Full article ">Figure 4
<p>Misorientation distribution at different tempering temperatures. (<b>a</b>) 500 °C; (<b>b</b>) 580 °C; (<b>c</b>) 620 °C.</p> Full article ">Figure 5
<p>Misorientation distribution within grains at different tempering temperatures (<b>a</b>,<b>b</b>) at 500 °C and (<b>c</b>,<b>d</b>) at 620 °C.</p> Full article ">Figure 5 Cont.
<p>Misorientation distribution within grains at different tempering temperatures (<b>a</b>,<b>b</b>) at 500 °C and (<b>c</b>,<b>d</b>) at 620 °C.</p> Full article ">Figure 6
<p>Precipitation distribution at various tempering temperatures: (<b>a</b>) 500 °C; (<b>b</b>) 580 °C; (<b>c</b>) 620 °C.</p> Full article ">Figure 7
<p>TEM images of steel at various tempering temperatures: (<b>a</b>) 500 °C; (<b>b</b>) 580 °C; (<b>c</b>) 620 °C.</p> Full article ">Figure 8
<p>Precipitation size distribution curves at different tempering temperatures: (<b>a</b>) 500 °C; (<b>b</b>) 580 °C; (<b>c</b>) 620 °C.</p> Full article ">Figure 9
<p>EDX results for the matrix and precipitates: (<b>a</b>) matrix; (<b>b</b>) precipitates.</p> Full article ">Figure 10
<p>Impact toughness after tempering.</p> Full article ">
46 pages, 15485 KiB  
Review
A Comprehensive Review of Laser Powder Bed Fusion in Jewelry: Technologies, Materials, and Post-Processing with Future Perspective
by Geethapriyan Thangamani, Stefano Felicioni, Elisa Padovano, Sara Biamino, Mariangela Lombardi, Daniele Ugues, Paolo Fino and Federica Bondioli
Metals 2024, 14(8), 897; https://doi.org/10.3390/met14080897 - 6 Aug 2024
Viewed by 599
Abstract
In recent years, additive manufacturing (AM) has played a significant role in various fashion industries, especially the textile and jewelry manufacturing sectors. This review article delves deeply into the wide range of methods and materials used to make intricately designed jewelry fabrication using [...] Read more.
In recent years, additive manufacturing (AM) has played a significant role in various fashion industries, especially the textile and jewelry manufacturing sectors. This review article delves deeply into the wide range of methods and materials used to make intricately designed jewelry fabrication using the additive manufacturing (AM) process. The Laser Powder Bed Fusion (L-PBF) process is examined for its suitability in achieving complex design and structural integrity in jewelry fabrication even with respect to powder metallurgy methods. Moreover, the review explores the use of precious materials, such as gold, silver, copper, platinum, and their alloys in additive manufacturing. Processing precious materials is challenging due to their high reflectivity and thermal conductivity, which results in poor densification and mechanical properties. To address this issue, the review article proposes three different strategies: (i) adding alloying elements, (ii) coating powder particles, and (iii) using low-wavelength lasers (green or blue). Finally, this review examines crucial post-processing techniques to improve surface quality, robustness, and attractiveness. To conclude, this review emphasizes the potential of combining additive manufacturing (AM) with traditional craftsmanship for creating jewelry, exploring the potential future directions and developments in the field of additive manufacturing (AM) for jewelry fabrication. Full article
(This article belongs to the Section Additive Manufacturing)
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<p>Support structures designed using CAD software for “the Ojo”, Reprinted with permission from ref. [<a href="#B10-metals-14-00897" class="html-bibr">10</a>]. 2016 Springer Nature.</p> Full article ">Figure 2
<p>(<b>a</b>) Laser beam on thinly spread Au powder using the L-PBF process, (<b>b</b>) <b>“</b>the Ojo” pendant built from the Au powder bed, (<b>c</b>) support remotion by pliers, (<b>d</b>) <b>t</b>hree different pendants before mechanical finishing and polishing, Reprinted from ref. [<a href="#B8-metals-14-00897" class="html-bibr">8</a>].</p> Full article ">Figure 3
<p>(<b>a</b>) “The Ojo” hand polished after 1 h stream finishing, (<b>b</b>) “the Ojo” series after mechanical and hand polishing and further stone setting, Reprinted with permission from ref. [<a href="#B10-metals-14-00897" class="html-bibr">10</a>]. 2016 Springer Nature.</p> Full article ">Figure 4
<p>Current and future trends in AM of the jewelry industry.</p> Full article ">Figure 5
<p>Innovative technique for jewelry production: powder metallurgy and advanced manufacturing methods.</p> Full article ">Figure 6
<p>SEM image of commercial copper powder fabricated by gas atomization processes with lower magnification (<b>a</b>) and higher magnification (<b>b</b>).</p> Full article ">Figure 7
<p>(<b>a</b>) Press and sinter of Tanishq pure gold coins by Titan, (<b>b</b>) MIM-produced 18-karat gold 3N pieces, Reprinted from ref. [<a href="#B19-metals-14-00897" class="html-bibr">19</a>].</p> Full article ">Figure 8
<p>(<b>a</b>,<b>b</b>) Multi-color rings, Reprinted from ref. [<a href="#B21-metals-14-00897" class="html-bibr">21</a>].</p> Full article ">Figure 9
<p>(<b>a</b>) Comparison of different materials reflectivity, Reprinted from ref. [<a href="#B23-metals-14-00897" class="html-bibr">23</a>] and (<b>b</b>) heat absorption and reflection of powder bed by the laser beam, Reprinted from ref. [<a href="#B22-metals-14-00897" class="html-bibr">22</a>].</p> Full article ">Figure 10
<p>SEM images of 18 K YG (<b>a</b>–<b>d</b>) obtained by L-PBF with different laser power and (<b>e</b>–<b>h</b>) jewelry parts, Reprinted from ref. [<a href="#B22-metals-14-00897" class="html-bibr">22</a>].</p> Full article ">Figure 11
<p>SEM images of AM 99.9% silver samples at 30 μm layer thickness with a scan speed and hatch distance of (<b>a</b>) 800 mm/s and 0.14 mm, and (<b>b</b>) 600 mm/s and 0.15 mm, Reprinted with permission from ref. [<a href="#B27-metals-14-00897" class="html-bibr">27</a>]. 2020 Elsevier.</p> Full article ">Figure 12
<p>CAD geometry of the Ag ring (<b>a</b>,<b>b</b>) and the product obtained by L-PBF (<b>c</b>), used under CC BY 4.0 [<a href="#B28-metals-14-00897" class="html-bibr">28</a>].</p> Full article ">Figure 13
<p>(<b>a</b>) Additive manufacturing of platinum ring; (<b>b</b>) defects of the ring shank on one side, Reprinted from ref. [<a href="#B32-metals-14-00897" class="html-bibr">32</a>].</p> Full article ">Figure 14
<p>SEM image of the left side (<b>a</b>) and right side (<b>b</b>) of the ring shank surface (red color arrow in <a href="#metals-14-00897-f013" class="html-fig">Figure 13</a>a), Reprinted from ref. [<a href="#B32-metals-14-00897" class="html-bibr">32</a>].</p> Full article ">Figure 15
<p>SEM images of the right ring shank surface area (<b>a</b>,<b>b</b>), Reprinted from ref. [<a href="#B32-metals-14-00897" class="html-bibr">32</a>].</p> Full article ">Figure 16
<p>(<b>a</b>) L-PBF with 1070 nm wavelength, Reprinted from ref. [<a href="#B34-metals-14-00897" class="html-bibr">34</a>], (<b>b</b>) damage of the optical mirror due to laser back reflection from the copper substrate and specific defects, Reprinted with permission from ref. [<a href="#B36-metals-14-00897" class="html-bibr">36</a>], 2019 Elsevier (<b>c</b>) micro-balling effect, Reprinted from ref. [<a href="#B38-metals-14-00897" class="html-bibr">38</a>], (<b>d</b>) balling effect, Reprinted from ref. [<a href="#B38-metals-14-00897" class="html-bibr">38</a>], (<b>e</b>) delamination, Reprinted from ref. [<a href="#B38-metals-14-00897" class="html-bibr">38</a>], (<b>f</b>) elevated edge, Reprinted from ref. [<a href="#B38-metals-14-00897" class="html-bibr">38</a>].</p> Full article ">Figure 17
<p>Schematic representation of a strategy to improve the laser absorptivity in additive manufacturing processes for jewelry application.</p> Full article ">Figure 18
<p>L-PBF of Au alloys: (<b>a</b>) white-gold and (<b>b</b>) red gold, Reprinted from ref. [<a href="#B25-metals-14-00897" class="html-bibr">25</a>].</p> Full article ">Figure 19
<p>Gold alloy ring obtained by (<b>a</b>) standard (76.2 W, 250 mm/s, layer thickness 30 μm, 50% of overlapping) and (<b>b</b>) optimized parameters (65 W, 330 mm/s, layer thickness 20 μm, and 75% of overlapping); (<b>c</b>–<b>f</b>) SEM images of pave obtained using the two sets of parameter, Reprinted from ref. [<a href="#B47-metals-14-00897" class="html-bibr">47</a>].</p> Full article ">Figure 20
<p>Effect of particle size distribution on the resolution and quality of the obtained surfaces: (<b>a</b>) and (<b>c</b>) PSD &lt; 63 μm; (<b>b</b>) and (<b>d</b>) PSD &lt; 15 μm, Reprinted from ref. [<a href="#B47-metals-14-00897" class="html-bibr">47</a>].</p> Full article ">Figure 21
<p>L-PBF of 925 Ag: (<b>a</b>–<b>d</b>) microstructure investigation with different powder compositions, and (<b>e</b>–<b>h</b>) jewelry parts, Reprinted from ref. [<a href="#B22-metals-14-00897" class="html-bibr">22</a>].</p> Full article ">Figure 22
<p>SEM images of L-PBF (<b>a</b>) Ag and (<b>b</b>) 925 Ag in the build direction (BD), Reprinted from ref. [<a href="#B30-metals-14-00897" class="html-bibr">30</a>].</p> Full article ">Figure 23
<p>L-PBF samples of Ag–Cu alloy with different La<sub>2</sub>O<sub>3</sub> content (<b>a</b>) and SEM microstructure obtained on (<b>b</b>) 0, (<b>c</b>) 0.4, (<b>d</b>) 0.8, and (<b>e</b>) 1.2% La<sub>2</sub>O<sub>3</sub> samples, Reprinted from ref. [<a href="#B55-metals-14-00897" class="html-bibr">55</a>].</p> Full article ">Figure 24
<p>Influence of the addition of solid suppressing elements in ternary 95Pt-1.5Cu-3.5Ru-X alloys, (X representing gallium, indium, tin, germanium or zinc) (<b>a</b>) and impact of germanium (<b>b</b>) and gallium addition (<b>c</b>) on the SST depth, Reprinted from ref. [<a href="#B32-metals-14-00897" class="html-bibr">32</a>].</p> Full article ">Figure 25
<p>Comparison of reflectivity of pure copper and its alloy powders: effect of (<b>a</b>) different elements, Reprinted with permission from ref. [<a href="#B61-metals-14-00897" class="html-bibr">61</a>], 2020 Science China Press, different amounts of (<b>b</b>) CNTs, used under CC BY 3.0 [<a href="#B62-metals-14-00897" class="html-bibr">62</a>] (<b>c</b>) C, Reprinted from ref. [<a href="#B63-metals-14-00897" class="html-bibr">63</a>], and (<b>d</b>) Cr and C, Reprinted with permission from ref. [<a href="#B64-metals-14-00897" class="html-bibr">64</a>]. 2019 John Wiley and Sons.</p> Full article ">Figure 26
<p>SEM micrograph of L-PBF samples of (<b>a</b>) pure copper, (<b>b</b>) Cu-0.5%CNTs, (<b>c</b>) Cu-1%CNTs, and (<b>d</b>) Cu-1.5%CNTs, used under CC BY 3.0 [<a href="#B62-metals-14-00897" class="html-bibr">62</a>].</p> Full article ">Figure 27
<p>SEM micrographs of Cu + C0.1 samples. The white arrows in (<b>a</b>) indicate the laser scan direction (<b>b</b>) surface cracks (<b>c</b>) the orientation of the crack parallel to BD (<b>d</b>) cracks filled with carbon segregation, Reprinted from ref. [<a href="#B63-metals-14-00897" class="html-bibr">63</a>].</p> Full article ">Figure 28
<p>SEM images of (<b>a</b>–<b>c</b>) Cu alloy and (<b>d</b>–<b>f</b>) CuCr0.3 underlining the epitaxial growth of the columnar grains parallel to the build direction and precipitates and carbon segregation at the grain boundaries, Reprinted with permission from ref. [<a href="#B64-metals-14-00897" class="html-bibr">64</a>]. 2019 John Wiley and Sons.</p> Full article ">Figure 29
<p>Effect of scan speed on the microstructure of tin-bronze samples (<b>a</b>–<b>d</b>). Non-optimized (<b>e</b>) and optimized (<b>f</b>) surface roughness in the as-built condition and after surface finishing step (<b>g</b>,<b>h</b>) jewelry parts, Reprinted from ref. [<a href="#B22-metals-14-00897" class="html-bibr">22</a>].</p> Full article ">Figure 30
<p>(<b>a</b>) Schematic representation of Ni-coated copper powders, (<b>b</b>) SEM image, and (<b>c</b>) EDS analysis, Reprinted from refs. [<a href="#B69-metals-14-00897" class="html-bibr">69</a>,<a href="#B70-metals-14-00897" class="html-bibr">70</a>].</p> Full article ">Figure 31
<p>Optical absorption of pure copper and surface-modified copper powders with different wavelengths. (<b>a</b>) Ni-coated copper powders, Reprinted from refs. [<a href="#B69-metals-14-00897" class="html-bibr">69</a>,<a href="#B70-metals-14-00897" class="html-bibr">70</a>], (<b>b</b>) copper powder treated in a nitrogen environment, Reprinted with permission from ref. [<a href="#B72-metals-14-00897" class="html-bibr">72</a>]. 2020 Elsevier.</p> Full article ">Figure 32
<p>SEM image of uncoated (<b>a</b>,<b>b</b>) and coated (<b>c</b>,<b>d</b>) Cu samples obtained by L-PBF with different VED, Reprinted from refs. [<a href="#B69-metals-14-00897" class="html-bibr">69</a>,<a href="#B70-metals-14-00897" class="html-bibr">70</a>].</p> Full article ">Figure 33
<p>Optical image of the top surface and cross section of L-PBF produced parts: (<b>a</b>,<b>b</b>) virgin CuCr1; (<b>c</b>,<b>d</b>) surface-modified CuCr1, Reprinted with permission from ref. [<a href="#B72-metals-14-00897" class="html-bibr">72</a>]. 2020 Elsevier.</p> Full article ">Figure 34
<p>Comparison of reflectivity behavior of pure and coated copper powders at 1064 nm, Reprinted from ref. [<a href="#B74-metals-14-00897" class="html-bibr">74</a>].</p> Full article ">Figure 35
<p>Laser absorptivity of various high reflective metals from 200 to 1400 nm. The wavelength ranges of (<b>a</b>) IR: infrared laser, (<b>b</b>) GL: green laser, and (<b>c</b>) BL: blue laser are shown, Reprinted from ref. [<a href="#B78-metals-14-00897" class="html-bibr">78</a>].</p> Full article ">Figure 36
<p>Comparison of laser-based additive manufacturing of copper cubic structure (four layers) with a green and red laser (laser source: TRUMPF product), Reprinted from ref. [<a href="#B81-metals-14-00897" class="html-bibr">81</a>].</p> Full article ">Figure 37
<p>L-PBF of pure copper: microstructural analysis of samples obtained by (<b>a</b>–<b>c</b>) infrared laser and (<b>d</b>–<b>f</b>) green lasers, Reprinted with permission from ref. [<a href="#B84-metals-14-00897" class="html-bibr">84</a>]. 2023 Springer Nature.</p> Full article ">Figure 38
<p>Post-processing methods used for jewelry fabrication.</p> Full article ">Figure 39
<p>Time comparison of different polishing processes: Vibratory Bowl Finishers, Barreling, Centrifugal Tumbling (CF), Drag Finishing (DF), and Stream Finishing (SF), Reprinted from ref. [<a href="#B94-metals-14-00897" class="html-bibr">94</a>].</p> Full article ">Figure 40
<p>Post-processing of L-PBF jewelry parts made of tin-bronze alloy. As-built samples (<b>a</b>,<b>b</b>) and after dry (<b>c</b>,<b>d</b>) and wet blasting with corundum (<b>e</b>,f), Reprinted from ref. [<a href="#B22-metals-14-00897" class="html-bibr">22</a>].</p> Full article ">Figure 41
<p>Laser polishing of a pendant using optimized laser parameters (<b>a</b>,<b>b</b>), Reprinted from ref. [<a href="#B94-metals-14-00897" class="html-bibr">94</a>].</p> Full article ">Figure 42
<p>Electro-mechanical polished sample (<b>a</b>) before and (<b>b</b>) after a 15-min treatment, Reprinted from ref. [<a href="#B94-metals-14-00897" class="html-bibr">94</a>].</p> Full article ">Figure 43
<p>Surface polishing of yellow-gold jewelry followed by various steps: (<b>a</b>) sandblasted, (<b>b</b>) electro-polished, and (<b>c</b>) hand-polished, Reprinted from ref. [<a href="#B23-metals-14-00897" class="html-bibr">23</a>].</p> Full article ">Figure 44
<p>AM of jewelry production and integration with TM.</p> Full article ">Figure 45
<p>Future perspective of jewelry and furniture production by AM.</p> Full article ">
14 pages, 10852 KiB  
Article
Effects of Partially Replacing Mo with Nb on the Microstructure and Properties of High-Strength Low-Alloy Steel during Reverse Austenization
by Liang Luo, Jiajun Zhang, Hao Fu, Fuhu Chen, Jianchun Qin and Yimin Li
Metals 2024, 14(8), 896; https://doi.org/10.3390/met14080896 - 6 Aug 2024
Viewed by 326
Abstract
This study investigated the effects of partially replacing expensive Mo with cheaper Nb on the microstructure and properties of high-strength low-alloy (HSLA) steel during reverse austenisation. The mechanical properties of the steel in the hot-rolled state were lower with a partial replacement of [...] Read more.
This study investigated the effects of partially replacing expensive Mo with cheaper Nb on the microstructure and properties of high-strength low-alloy (HSLA) steel during reverse austenisation. The mechanical properties of the steel in the hot-rolled state were lower with a partial replacement of Mo by Nb. However, after pre-tempering and reheating and quenching, the strength increased greatly while the ductility and toughness did not decrease much. Thus, the negative effects of replacing Mo with Nb were mostly alleviated, and a good balance between strength, ductility and toughness was achieved. After heat treatment, the mass percentage of precipitates increased substantially, which helped to pin grain boundaries during austenisation. The percent of high-angle grain boundaries greatly increased while the average effective grain size decreased, which improved grain refinement. The results showed that combining a partial replacement of Mo by Nb with heat treatment allows the microstructure and mechanical properties of HSLA steel to be effectively controlled while improving the balance between cost and performance. These findings provide valuable insights into the preparation and design of steels with similar microstructures. 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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<p>Hot-rolling and heat treatment processes for the steels.</p> Full article ">Figure 2
<p>Mechanical properties of hot-rolled and heat-treated steel samples: (<b>a</b>) tensile strength and yield strength; (<b>b</b>) elongation and impact absorption energy.</p> Full article ">Figure 3
<p>SEM images showing the microstructures of the steel samples: (<b>a</b>) GQ0, (<b>b</b>) GQ1, (<b>c</b>) GQ0-HT and (<b>d</b>) GQ1-HT.</p> Full article ">Figure 4
<p>TEM images showing the microstructures of the steel samples after heat treatment: (<b>a</b>–<b>f</b>) GQ0-HT and (<b>g</b>–<b>l</b>) GQ1-HT. The first column shows bright-field images, the second column displays local magnifications of the first column and the third column exhibits the corresponding high-angle annular dark-field (HAADF) images of the first column.</p> Full article ">Figure 5
<p>TEM images showing the precipitates of GQ0-HT: (<b>a</b>,<b>c</b>) bright-field images and (<b>b</b>,<b>d</b>) HAADF images.</p> Full article ">Figure 6
<p>TEM images showing the precipitates of GQ1-HT: (<b>a</b>,<b>c</b>) bright-field images and (<b>b</b>,<b>d</b>) HAADF images.</p> Full article ">Figure 7
<p>Inverse pole figure (IPF) maps of the steel samples: (<b>a</b>) GQ0, (<b>b</b>) GQ1, (<b>c</b>) GQ0-HT and (<b>d</b>) GQ1-HT.</p> Full article ">Figure 8
<p>Effective grain size distributions of the steel samples: (<b>a</b>) GQ0, (<b>b</b>) GQ1, (<b>c</b>) GQ0-HT and (<b>d</b>) GQ1-HT.</p> Full article ">Figure 9
<p>EBSD grain boundary maps of the steel samples: (<b>a</b>) GQ0, (<b>b</b>) GQ1, (<b>c</b>) GQ0-HT and (<b>d</b>) GQ1-HT (red lines indicate neighbouring grain orientation differences of 2° to 15°; black lines indicate orientation differences greater than 15°).</p> Full article ">Figure 10
<p>Distributions of grain boundary misorientations for the steel samples: (<b>a</b>) GQ0, (<b>b</b>) GQ1, (<b>c</b>) GQ0-HT and (<b>d</b>) GQ1-HT.</p> Full article ">Figure 11
<p>Precipitate size distributions of the steel samples.</p> Full article ">
11 pages, 3710 KiB  
Article
Effect of Cr3C2 Addition on Microstructure and Mechanical Properties of WC-CoNiFe Cemented Carbides
by Jinbo Wu, Daping Ren, Bo Xie, Rongyue He, Zhanji Geng, Zichun Zhang, Yang Liu, Dong Wang, Yanghui Zhu and Wei Zhang
Metals 2024, 14(8), 895; https://doi.org/10.3390/met14080895 - 6 Aug 2024
Viewed by 379
Abstract
In traditional cemented carbides, Co is mainly used as a binder. Recently, replacing Co with medium- to high-entropy alloys has shown significant improvements in hardness, fracture toughness, high-temperature oxidation resistance, and corrosion resistance, making it a research focus globally. Both the typical grain [...] Read more.
In traditional cemented carbides, Co is mainly used as a binder. Recently, replacing Co with medium- to high-entropy alloys has shown significant improvements in hardness, fracture toughness, high-temperature oxidation resistance, and corrosion resistance, making it a research focus globally. Both the typical grain refiner Cr3C2 and medium- to high-entropy alloy binders affect the WC grain size in cemented carbides. This study investigates the synergistic grain refinement mechanism of Cr3C2 and medium- to high-entropy alloy binders on WC grains and their impact on the microstructure and mechanical properties of cemented carbides. The results show that increasing the Cr3C2 addition refined WC grains in the WC-CoNiFe alloy, increased coercivity, and enhanced hardness, with transverse rupture strength first increasing and then decreasing. The alloy achieved optimal performance at 0.6 wt.%Cr3C2, with a hardness of 91.25 HRA and transverse rupture strength of 3883.2 MPa. Full article
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<p>XRD pattern of WC−CoNiFe carbide with different Cr<sub>3</sub>C<sub>2</sub> content.</p> Full article ">Figure 2
<p>Microstructure of WC-CoNiFe with different Cr<sub>3</sub>C<sub>2</sub> content: (<b>a</b>) 0 wt.%; (<b>b</b>) 0.2 wt.%; (<b>c</b>) 0.6 wt.%; (<b>d</b>) 0.9 wt.%; (<b>e</b>) 2 wt.%.</p> Full article ">Figure 3
<p>WC−CoNiFe cobalt magnetism and coercive magnetism with different content of Cr<sub>3</sub>C<sub>2</sub>.</p> Full article ">Figure 4
<p>Transverse rupture strength and hardness of WC−CoNiFe with different Cr<sub>3</sub>C<sub>2</sub> content.</p> Full article ">Figure 5
<p>Transverse fracture morphology of alloys with different Cr<sub>3</sub>C<sub>2</sub> content: (<b>a</b>) 0 wt. %; (<b>b</b>) 0.2 wt. %; (<b>c</b>) 0.6 wt. %; (<b>d</b>) 0.9 wt. %; (<b>e</b>) 2 wt. %.</p> Full article ">
17 pages, 25787 KiB  
Article
Machine-Learning-Assisted Design of Novel TiZrNbVAl Refractory High-Entropy Alloys with Enhanced Ductility
by Xinyi Zhao, Zihang Wei, Junfeng Zhao, Yandong Jia, Shuo Cao, Dan Wang and Yucheng Lei
Metals 2024, 14(8), 894; https://doi.org/10.3390/met14080894 - 5 Aug 2024
Viewed by 417
Abstract
Refractory high-entropy alloys (RHEAs) typically exhibit excellent high-temperature strength but limited ductility. In this study, a comprehensive machine learning strategy with integrated material knowledge is proposed to predict the elongation of TiZrNbVAl RHEAs. By referring to the ductility theories, a set of cost-effective [...] Read more.
Refractory high-entropy alloys (RHEAs) typically exhibit excellent high-temperature strength but limited ductility. In this study, a comprehensive machine learning strategy with integrated material knowledge is proposed to predict the elongation of TiZrNbVAl RHEAs. By referring to the ductility theories, a set of cost-effective material features is developed with various mathematical forms of thermodynamic parameters. These features are proven to effectively incorporate material knowledge into ML modeling. They also offer potential alternatives to those obtained from costly first-principles calculations. Based on Pearson correlation coefficients, the linear relationships between pairwise features were compared, and the seven key features with the greatest impact on the model were selected for ML modeling. Regression tasks were performed to predict the ductility of TiZrNbVAl, and the CatBoost gradient boosting algorithm exhibiting the best performance was eventually selected. The established optimized model achieves high predictive accuracies exceeding 0.8. These key features were further analyzed using interpretable ML methods to elucidate their influences on various ductility mechanisms. According to the ML results, different compositions of TiZrNbVAl with excellent tensile properties were prepared. The experimental results indicate that Ti44Zr24Nb17V5Al10 and Ti44Zr26Nb8V13Al9 both exhibited ultimate tensile strengths of approximately 1180 MPa and elongations higher than 21%. They verified that the ML strategy proposed in this study is an effective approach for predicting the properties of RHEAs. It is a potential method that can replace costly first-principles calculations. Thermodynamic parameters have been shown to effectively predict alloy ductility to a certain extent. Full article
(This article belongs to the Special Issue Numerical and Experimental Advances in Metal Processing)
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<p>Flowchart of TiZrNbVAl RHEA with improved elongation based on ML model.</p> Full article ">Figure 2
<p>Alloying element distribution of 22 TiZrNbVAl components with elongation of more than 10%.</p> Full article ">Figure 3
<p>Pearson correlation map between any two features in the initial 11 features. The colors red and blue are used to represent positive and negative correlations, respectively. The intensity of the color shows the correlation strength.</p> Full article ">Figure 4
<p>Eight distinct models were assessed for their performance using the MAE and MSE as evaluation metrics.</p> Full article ">Figure 5
<p>The performance of the ML model based on the initial dataset.</p> Full article ">Figure 6
<p>(<b>a</b>) Tensile engineering stress–strain curves of Ti<sub>44</sub>Zr<sub>24</sub>Nb<sub>17</sub>V<sub>5</sub>Al<sub>10</sub>, Ti<sub>44</sub>Zr<sub>26</sub>Nb<sub>8</sub>V<sub>13</sub>Al<sub>9</sub>, and Ti<sub>38</sub>Zr<sub>30</sub>Nb<sub>17</sub>V<sub>5</sub>Al<sub>10</sub> at room temperature. (<b>b</b>) The comparison of ultimate tensile strength and elongation of the alloys collected and three RHEAs in this work.</p> Full article ">Figure 7
<p>Tensile fracture morphologies of (<b>a</b>) Ti<sub>44</sub>Zr<sub>24</sub>Nb<sub>17</sub>V<sub>5</sub>Al<sub>10</sub>, (<b>b</b>) Ti<sub>44</sub>Zr<sub>26</sub>Nb<sub>8</sub>V<sub>13</sub>Al<sub>9</sub>, and (<b>c</b>) Ti<sub>38</sub>Zr<sub>30</sub>Nb<sub>17</sub>V<sub>5</sub>Al<sub>10</sub>.</p> Full article ">Figure 8
<p>SHAP waterfall plots for nine randomly selected RHEA samples. The green number is the predicted elongation, and the orange in parenthesis is the experimental elongation.</p> Full article ">Figure 8 Cont.
<p>SHAP waterfall plots for nine randomly selected RHEA samples. The green number is the predicted elongation, and the orange in parenthesis is the experimental elongation.</p> Full article ">Figure 9
<p>The relationship between the SHAP value and (<b>a</b>) <span class="html-italic">VEC</span>, (<b>b</b>) <span class="html-italic">σ<sub>a</sub></span>, (<b>c</b>) <span class="html-italic">σ<sub>T</sub></span>, (<b>d</b>) <span class="html-italic">σ<sub>χ</sub></span>, (<b>e</b>) <span class="html-italic">σ<sub>VEC</sub></span>, (<b>f</b>) Δ<span class="html-italic">H<sub>mix</sub></span>, and (<b>g</b>) Δ<span class="html-italic">S<sub>mix</sub></span> of the nine RHEA samples.</p> Full article ">Figure 10
<p>XRD patterns of Ti<sub>44</sub>Zr<sub>24</sub>Nb<sub>17</sub>V<sub>5</sub>Al<sub>10</sub>, Ti<sub>44</sub>Zr<sub>26</sub>Nb<sub>8</sub>V<sub>13</sub>Al<sub>9</sub>, and Ti<sub>38</sub>Zr<sub>30</sub>Nb<sub>17</sub>V<sub>5</sub>Al<sub>10</sub> RHEAs.</p> Full article ">Figure 11
<p>SEM images–EDS mapping and TEM images–SAED patterns of (<b>a</b>,<b>b</b>) Ti<sub>44</sub>Zr<sub>24</sub>Nb<sub>17</sub>V<sub>5</sub>Al<sub>10</sub>, (<b>c</b>,<b>d</b>) Ti<sub>44</sub>Zr<sub>26</sub>Nb<sub>8</sub>V<sub>13</sub>Al<sub>9</sub>, (<b>e</b>,<b>f</b>) Ti<sub>38</sub>Zr<sub>30</sub>Nb<sub>17</sub>V<sub>5</sub>Al<sub>10</sub>.</p> Full article ">
14 pages, 11569 KiB  
Article
Effect of Swing Amplitude on Microstructure and Properties of TC4 Titanium Alloy in Laser Welding
by Jianhui Liang and Zhanqi Liu
Metals 2024, 14(8), 893; https://doi.org/10.3390/met14080893 - 5 Aug 2024
Viewed by 444
Abstract
The welding of TC4 titanium alloy sheets with a thickness of 1 mm was successfully accomplished by a swinging laser. The microstructure and mechanical properties of the welding seam under different swing amplitudes were studied. In this paper, the microstructure, phase composition, mechanical [...] Read more.
The welding of TC4 titanium alloy sheets with a thickness of 1 mm was successfully accomplished by a swinging laser. The microstructure and mechanical properties of the welding seam under different swing amplitudes were studied. In this paper, the microstructure, phase composition, mechanical properties, and fracture morphology of the weld with swing frequency of 50 Hz and different swing amplitudes (0.2 mm, 1 mm, 2 mm, and 3 mm) were tested and analyzed. The results show that basket-weave microstructures are present in the fusion zone of welds under different oscillation amplitudes, but the morphology of martensite within the basket-weave differs. The weld microstructure is mainly composed of acicular α′ martensite, initial α phase, secondary α phase, and residual β phase. The hardness of the weld is higher than that of the base metal, and the overall hardness decreases from the weld center to the base metal. When the oscillation amplitude A = 1 mm, the weld microstructure has the smallest average grain size, the highest microhardness (388.86 HV), the largest tensile strength (1115.4 MPa), and quasi-cleavage fracture occurs. At an oscillation amplitude of A = 2 mm, the tensile specimen achieves the maximum elongation of 14%, with ductile fracture as the dominant mechanism. Full article
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<p>Microstructure of 1 mm thick TC4 titanium alloy.</p> Full article ">Figure 2
<p>Welding assembly drawing.</p> Full article ">Figure 3
<p>Size of the tensile specimen (unit: mm).</p> Full article ">Figure 4
<p>Weld morphology: (<b>a</b>) A = 0.2 mm, face of weld; (<b>b</b>) A = 0.2 mm, back of weld; (<b>c</b>) A = 1 mm, face of weld; (<b>d</b>) A = 1 mm, back of weld; (<b>e</b>) A = 2 mm, face of weld; (<b>f</b>) A = 2 mm, back of weld; (<b>g</b>) A = 3 mm, face of weld; (<b>h</b>) A = 3 mm, back of weld.</p> Full article ">Figure 5
<p>Microstructure of weld seam: (<b>a</b>) A = 0.2 mm; (<b>b</b>) A = 1 mm; (<b>c</b>) A = 2 mm; (<b>d</b>) A = 3 mm.</p> Full article ">Figure 6
<p>Grains in the fusion zone of weld seam: (<b>a</b>) grain size distribution diagram; (<b>b</b>) distribution diagram of maximum and average grain size.</p> Full article ">Figure 7
<p>Microstructure of weld when A = 1 mm: (<b>a</b>) Interface of base material and heat affected zone; (<b>b</b>) Fusion zone; (<b>c</b>) Fusion zone.</p> Full article ">Figure 8
<p>Distribution of elements in different areas of the weld: (<b>a</b>) Interface of base material and heat affected zone; (<b>b</b>) Fusion zone; (<b>c</b>) Fusion zone.</p> Full article ">Figure 9
<p>TC4 titanium alloy phase transition: (<b>a</b>) phase transition at temperature rise; (<b>b</b>) phase transition during cooling.</p> Full article ">Figure 10
<p>XRD comparison.</p> Full article ">Figure 11
<p>Tensile properties: (<b>a</b>) stress–strain curves of tensile samples with different parameters; (<b>b</b>) fracture locations of tensile samples; (<b>c</b>) line graph of tensile strength and elongation.</p> Full article ">Figure 12
<p>Distribution chart of weld joint hardness.</p> Full article ">Figure 13
<p>Fracture section of tensile fracture: (<b>a</b>–<b>c</b>) A = 0.2 mm; (<b>d</b>–<b>f</b>) A = 1 mm; (<b>g</b>–<b>i</b>) A = 2 mm; (<b>j</b>–<b>l</b>) A = 3 mm.</p> Full article ">
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