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Volume 11
Issue 9
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Metals, Volume 11, Issue 9 (September 2021) – 167 articles

Cover Story (view full-size image): The heat losses in the submerged entry nozzle (SEN) region were investigated using a novel numerical approach by coupling the convection and heat transfer in the liquid slag and melt pool. It was revealed that the parasitic solidification occurred inside an SEN bore with partially or completely absent insulation. SEN clogging was found to promote the solidification of the entrapped melt. The accelerated flow and the impaired superheat inside the mold cavity were detected due to the combined effect of the clogging and parasitic solidification. This study aims to assist in the development of clog-free and energy-saving SEN designs. View this paper
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15 pages, 3910 KiB  
Article
Recovery of Rare Earth Oxides from Flotation Concentrates of Bastnaesite Ore by Ultra-Fine Centrifugal Concentration
by Alex Norgren and Corby Anderson
Metals 2021, 11(9), 1498; https://doi.org/10.3390/met11091498 - 21 Sep 2021
Cited by 3 | Viewed by 2227
Abstract
Historically, the ability to effectively separate carbonate gangue from bastnaesite via flotation has frequently proven to be challenging without sacrificing significant rare earth oxide (REO) grade or recovery. However, in light of the fact that the rare earth bearing minerals often exhibit higher [...] Read more.
Historically, the ability to effectively separate carbonate gangue from bastnaesite via flotation has frequently proven to be challenging without sacrificing significant rare earth oxide (REO) grade or recovery. However, in light of the fact that the rare earth bearing minerals often exhibit higher specific gravities than the carbonate gangue, the possibility exists that the use of gravity separation could be used to achieve such a selective separation. This however is complicated by the fact that, in cases such as this study when the liberation size is finer than 50 microns, most traditional gravity separation methods become increasingly challenging. The aim of this study is to determine the applicability of centrifugal concentrators to beneficiate ultra-fine (UF) bastnaesite and calcite bearing flotation concentrates. By using a UF Falcon, it was possible to achieve initial gravity REO recoveries exceeding 90% while rejecting on the order of 25% to 35% of the total calcium from an assortment of rougher and cleaner flotation concentrates. Additionally, when additional stages of cleaner UF Falcon gravity separation were operated in an open circuit configuration, it was possible, from an original fine feed of 35 microns containing 50.5% REO and 5.5% Ca, to upgrade up to approximately 59% REO and 2.0% calcium. While not the goal of this study, these results also support previous limited data to suggest that UF Falcons are potentially capable of treating a wider range of materials than they were originally designed for, including feeds rich in heavy mineral content. Full article
(This article belongs to the Special Issue Advances in Mineral Processing and Hydrometallurgy—2nd Edition)
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<p>Original TANCO tantalum gravity circuit, recreated for legibility [<a href="#B3-metals-11-01498" class="html-bibr">3</a>].</p> Full article ">Figure 2
<p>Replacement of TANCO gravity circuit with a single Falcon UF 600, recreated for legibility [<a href="#B4-metals-11-01498" class="html-bibr">4</a>].</p> Full article ">Figure 3
<p>Cross sections of bowl of continuous (left) and UF Falcon (right) [<a href="#B5-metals-11-01498" class="html-bibr">5</a>].</p> Full article ">Figure 4
<p>Particle size distributions of MPC (Mountain Pass cleaner flotation concentrate) and other Mountain Pass MLA (Mineral Liberation Analysis) materials.</p> Full article ">Figure 5
<p>False-color image from MLA of the MPC material. Values represent surface area percentages.</p> Full article ">Figure 6
<p>Mineral liberation by individual REE mineral species in MPC material MLA sample.</p> Full article ">Figure 7
<p>Mineral liberation by grouped REE minerals in MPC material MLA sample.</p> Full article ">Figure 8
<p>Cumulative REO grade/recovery for MPC UF Falcon recovery maximization testing.</p> Full article ">Figure 9
<p>Cumulative gravity concentrate Ca grade vs REO recovery for MPC UF Falcon recovery maximization testing.</p> Full article ">Figure 10
<p>Cumulative Ca recovery vs REO recovery for MPC UF Falcon recovery maximization testing.</p> Full article ">Figure 11
<p>Overall flowsheet for MPC UF Falcon REO recovery and grade maximization testing programs. The overlaid values represent the estimated selected cumulative performance.</p> Full article ">Figure 12
<p>Flowsheet of Williams’ UF Falcon test work on enhanced rougher flotation concentrate. Ca, Ba, and Si grades have been converted from the original reported values of CaO, BaO, and SiO<sub>2</sub> from the source documentation [<a href="#B9-metals-11-01498" class="html-bibr">9</a>].</p> Full article ">
16 pages, 11265 KiB  
Article
Deformation Mechanism Investigation on Low Density 18Mn Steels under Different Solid Solution Treatments
by Yong-Tao Huo, Yan-Lin He, Na-Qiong Zhu, Min-Long Ding, Ren-Dong Liu and Yu Zhang
Metals 2021, 11(9), 1497; https://doi.org/10.3390/met11091497 - 21 Sep 2021
Cited by 5 | Viewed by 2219
Abstract
To meet the demand of the 10% weight reduction goal for automotive steel, the microstructure and mechanical properties of Fe-18Mn-Al-C steel with different carbon and aluminum contents were investigated under different solid solution treatments, and the deformation mechanisms of the experimental steels were [...] Read more.
To meet the demand of the 10% weight reduction goal for automotive steel, the microstructure and mechanical properties of Fe-18Mn-Al-C steel with different carbon and aluminum contents were investigated under different solid solution treatments, and the deformation mechanisms of the experimental steels were elucidated. Aided by thermodynamic calculation, transmission electron microscopy (TEM) and in situ scanning electron microscope (SEM) analysis, it was shown that for the 18Mn-1.5Al experimental steel with about 20 mJ/m2 stacking fault energy (SFE), the twinning-induced plasticity (TWIP) effect always dominated in this steel after different solid solution treatments under tensile deformation. With the 7 wt% aluminum addition, the SFE of austenite was affected by temperature and the range of SFE was between 60 and 65 mJ/m2. The existence of ?-ferrite obviously inhibited the TWIP effect. With the increase in the solution treatment temperature, ?-ferrite gradually transformed into the austenite, and the n-value remained low and stable in a large strain range, which were caused by the local hardening during the tensile deformation. Due to the difference in the deformability of the austenite and ?-ferrite structure as well as the inconsistent extension of the slip band, the micro-cracks were easily initiated in the 18Mn-7Al experimental steel; then, it exhibited lower plasticity. Full article
(This article belongs to the Special Issue Strengthening Mechanisms of Metals and Alloys)
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<p>Schematic diagram of the tensile sample for in situ observation (unit: mm).</p> Full article ">Figure 2
<p>SEM (scanning electron microscope) micrographs of steel 1# under different solid solution temperatures: (<b>a</b>) 650 °C, (<b>b</b>) 750 °C, (<b>c</b>) 800 °C, (<b>d</b>) 850 °C, (<b>e</b>) 900 °C and (<b>f</b>) 1050 °C. (δ: δ-ferrite; γ: austenite.)</p> Full article ">Figure 3
<p>SEM micrographs of steel 2# under different solid solution temperatures: (<b>a</b>) 650 °C, (<b>b</b>) 750 °C, (<b>c</b>) 800 °C, (<b>d</b>) 850 °C, (<b>e</b>) 900 °C and (<b>f</b>) 1050 °C.</p> Full article ">Figure 4
<p>X-ray diffraction patterns of steels under different solid solution temperatures: (<b>a</b>) steel 1# and (<b>b</b>) steel 2#.</p> Full article ">Figure 5
<p>TEM (transmission electron microscopy) micrographs of annealing twins microstructure of steels: (<b>a</b>–<b>c</b>) steel 1# and (<b>d</b>–<b>f</b>) steel 2# under the solid solution temperature of 850 °C with (<b>a</b>,<b>d</b>) bright-field images, (<b>b</b>,<b>e</b>) dark-field images and (<b>c</b>,<b>f</b>) corresponding selected area electron diffraction (SAED) patterns.</p> Full article ">Figure 6
<p>Engineering stress–strain curves of steels after solid solution treatment at different temperatures for 6 min: (<b>a</b>) steel 1# and (<b>b</b>) steel 2#.</p> Full article ">Figure 7
<p>Relationship between equilibrium phases and temperature in steels: (<b>a</b>) steel 1# and (<b>b</b>) steel 2#. FCC_A1: austenite. BCC_A2: δ-ferrite. LIQUID: cementite. KAPPA_E21: κ carbide.M23C6: M<sub>23</sub>C<sub>6</sub> carbide.</p> Full article ">Figure 8
<p>Optical microscopy of steel 1# under solid solution treatment at the temperature of 650 °C.</p> Full article ">Figure 9
<p>Instantaneous strain hardening exponent as a function of true strain for steels under solid solution treatment at different temperatures: (<b>a</b>) steel 1# and (<b>b</b>) steel 2#.</p> Full article ">Figure 10
<p>TEM micrographs of deformation twins microstructure of experimental steels: (<b>a</b>–<b>c</b>) steel 1# and (<b>d</b>–<b>f</b>) steel 2# with (<b>a</b>,<b>d</b>) bright field images, (<b>b</b>,<b>e</b>) dark field images and (<b>c</b>,<b>f</b>) corresponding SAED patterns.</p> Full article ">Figure 11
<p>(<b>a</b>) Load-displacement relationship curve of steel 1# and (<b>b–h</b>) SEM analysis of microstructure under different displacements. <math display="inline"><semantics> <mrow> <msub> <mi>L</mi> <mi>γ</mi> </msub> </mrow> </semantics></math> and <math display="inline"><semantics> <mrow> <msub> <mi>L</mi> <mi>δ</mi> </msub> </mrow> </semantics></math> are the size of austenite and δ-ferrite, respectively.</p> Full article ">Figure 12
<p>Relationship between the deformation and the tensile displacement of the two phases.</p> Full article ">Figure 13
<p>Slip bands and microcrack morphology transition of steel 1# with different tensile displacements.</p> Full article ">Figure 14
<p>(<b>a</b>) Load-displacement relationship curve of steel 2# and (<b>b</b>–<b>g</b>) the SEM analysis of microstructure under different displacements.</p> Full article ">
13 pages, 2622 KiB  
Article
Study on Adding Ammonium Hydrogen Fluoride to Improve Manganese Leaching Efficiency of Ammonia Leaching Low-Grade Rhodochrosite
by Peng Yang, Xiaoping Liang, Chengbo Wu, Tengfei Cui and Yu Wang
Metals 2021, 11(9), 1496; https://doi.org/10.3390/met11091496 - 21 Sep 2021
Cited by 4 | Viewed by 2396
Abstract
The ammonia leaching method for treating low-grade rhodochrosite has the advantages of a good impurity removal effect and low environmental pollution. In this paper, aiming at the low leaching efficiency of low-grade rhodochrosite treated by the ammonia leaching method, studies on enhancing the [...] Read more.
The ammonia leaching method for treating low-grade rhodochrosite has the advantages of a good impurity removal effect and low environmental pollution. In this paper, aiming at the low leaching efficiency of low-grade rhodochrosite treated by the ammonia leaching method, studies on enhancing the leaching efficiency of manganese by using ammonium hydrogen fluoride as an additive are carried out. The effects of different ammonia concentrations, leaching temperatures, leaching times, liquid-solid ratios, stirring rates, and the addition of ammonium hydrogen fluoride on the leaching efficiency of manganese with and without ammonium hydrogen fluoride as an additive were comparatively studied, and the parameters of ammonia concentration, ammonia leaching temperature, and ammonium hydrogen fluoride dosage were optimized in the experimental study. The results indicated that ammonium hydrogen fluoride as an additive in the treatment of low-grade rhodochrosite by the ammonia leaching method could effectively increase the leaching efficiency of manganese, and the optimal process parameters were obtained. Meanwhile, the addition of ammonium hydrogen fluoride didn’t affect the quality of the steamed ammonia product. Full article
(This article belongs to the Special Issue Separation and Leaching for Metals Recovery 2021)
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<p>The XRD map of calcined manganese ore roasting sample.</p> Full article ">Figure 2
<p>The technological process.</p> Full article ">Figure 3
<p>Effects of different factors on manganese leaching efficiency with or without annonium hydrogen fluoride. (<b>a</b>) Leaching temperature of 30 °C, liquid–solid ratio of 6:1, stirring rate of 500 r/min, leaching 1.5 h, ammonium hydrogen fluoride dosage of 8% (and 0%); (<b>b</b>) ammonia concentration of 14 mol/L, liquid–solid ratio of 6:1, stirring rate of 500 r/min, leaching time of 1.5 h, amount of ammonium hydrogen fluoride of 8% (and 0%). (<b>c</b>) Leaching temperature of 30 °C, ammonia concentration of 14 mol/L, liquid–solid ratio of 6:1, stirring rate of 500 r/min, ammonium hydrogen fluoride dosage of 8% (and 0%); (<b>d</b>) leaching temperature of 30 °C, ammonia concentration of 14 mol/L, stirring rate of 500 r/min, leaching time of 1 h, amount of ammonium hydrogen fluoride of 8% (and 0%). (<b>e</b>) Leaching temperature of 30 °C, ammonia concentration of 14 mol/L, liquid–solid ratio of 6:1, leaching time of 1 h, amount of ammonium hydrogen fluoride of 8% (and 0%).</p> Full article ">Figure 4
<p>Leaching temperature of 30 °C ammonia concentration of 14 mol/L, liquid–solid ratio of 6:1, stirring rate of 400 r/min, and leaching time of 1 h, and the influence of the addition amount of ammonium hydrogen fluoride on the leaching efficiency of Mn.</p> Full article ">Figure 5
<p>(<b>a</b>) Without ammonium fluoride in XRD diagram of slag leaching; (<b>b</b>) with ammonium fluoride in XRD diagram of slag leaching.</p> Full article ">Figure 6
<p>XRD of manganese carbonate in ammonia vapor product.</p> Full article ">
18 pages, 7828 KiB  
Article
Interface Formation and Bonding Mechanisms of Laser Welding of PMMA Plastic and 304 Austenitic Stainless Steel
by Yijie Huang, Xiangdong Gao, Bo Ma and Yanxi Zhang
Metals 2021, 11(9), 1495; https://doi.org/10.3390/met11091495 - 21 Sep 2021
Cited by 20 | Viewed by 3460
Abstract
Laser welding experiments involving amorphous thermoplastic polymer (PMMA) and 304 austenitic stainless steel plates were conducted to explore the influence of laser welding process parameters on plastic–metal joints. A high-speed camera was applied to record the dynamics of the molten pool and the [...] Read more.
Laser welding experiments involving amorphous thermoplastic polymer (PMMA) and 304 austenitic stainless steel plates were conducted to explore the influence of laser welding process parameters on plastic–metal joints. A high-speed camera was applied to record the dynamics of the molten pool and the formation of bubbles to reveal the bonding mechanisms of the hybrid joints. The influence of process parameters on the joints was analyzed using temperature measurements performed with thermocouples. The microstructure morphology of joints was observed using SEM. The mechanical characterization of the hybrid joints was carried out to understand the effect of the welding conditions on the weld morphology, flaws and shear stress. Different interface temperatures resulted in two types of bubbles and led to different weld morphology characteristics. A stable hybrid joint with the best shear stress was produced with a laser line energy of 20.16 J/mm2, a temperature of 305 °C and small bubbles. The shear stress of the effective joint under the maximum mechanical resistance was 4.17 MPa. The chemical bonds (M-O, M-C) and mechanical anchoring that formed on the steel’s surface contributed to the joint bonding. Range analysis provided guidance for identifying the impact of individual factors in the shear stress for the laser welding of plastic–metal. Full article
(This article belongs to the Special Issue Laser Welding Technology)
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<p>(<b>a</b>) Experimental apparatus used for laser joining PMMA and steel and (<b>b</b>) schematic diagram of the laser light path. Adapted with permission from ref. [<a href="#B31-metals-11-01495" class="html-bibr">31</a>]. Copyright 2021 Optics and Laser Technology.</p> Full article ">Figure 2
<p>High-speed camera images of the bubble formation and movement. (<b>a</b>) Formed molten pool and bubbles, (<b>b</b>) Growth of the molten layer and bubbles fluctuated, (<b>c</b>) Solidification process.</p> Full article ">Figure 3
<p>Macroscopic weld surface morphologies of 25 hybrid joints samples: (<b>a</b>) 304 austenitic stainless sheet side and (<b>b</b>) PMMA side.</p> Full article ">Figure 4
<p>OM images of the bubbles’ morphologies: (<b>a</b>) type I bubbles and (<b>b</b>) type II bubbles.</p> Full article ">Figure 5
<p>Variation in the maximum temperature of the PMMA–steel joint with varying line energy.</p> Full article ">Figure 6
<p>Optical micrographs of the surface morphologies and cross-sections of typical joints of NO.5, NO.17 and NO.7.</p> Full article ">Figure 7
<p>Relationship between the temperature and time of the typical specimens.</p> Full article ">Figure 8
<p>Relationship between the melted depth and laser line energy for typical samples (NO.5, NO.7, and NO.17).</p> Full article ">Figure 9
<p>(<b>a</b>) Cross-sectional SEM image of the joint interface for the NO.5 joint, (<b>b</b>) an enlarged view of the boxed region in (<b>a</b>,<b>c</b>) an enlarged view of the boxed region in (<b>b</b>).</p> Full article ">Figure 10
<p>(<b>a</b>) SEM image and (<b>b</b>) EDS mapping results of the fractured surfaces of the typical sample.</p> Full article ">Figure 11
<p>Example of joint NO.11 to display the ImageJ measurement in the discolored zone: (<b>a</b>) image of joint NO.11 before the ImageJ analysis and (<b>b</b>) image of joint NO.11 after the ImageJ analysis.</p> Full article ">Figure 12
<p>Orthogonal experimental design results of the 25 conditions: measured total joint area and discolored zone.</p> Full article ">Figure 13
<p>Orthogonal experimental design result of the 25 conditions: measured shear force and calculated tensile shear stress.</p> Full article ">
17 pages, 4569 KiB  
Article
Electrochemical Study of Nd and Pr Co-Deposition onto Mo and W from Molten Oxyfluorides
by Vesna S. Cvetković, Dominic Feldhaus, Nataša M. Vukićević, Tanja S. Barudžija, Bernd Friedrich and Jovan N. Jovićević
Metals 2021, 11(9), 1494; https://doi.org/10.3390/met11091494 - 21 Sep 2021
Cited by 7 | Viewed by 3442
Abstract
Electrodeposition processes of neodymium and praseodymium in molten NdF3 + PrF3 + LiF + 1 wt.%Pr6O11 + 1 wt.%Nd2O3 and NdF3 + PrF3 + LiF + 2 wt.%Pr6O11 + 2 [...] Read more.
Electrodeposition processes of neodymium and praseodymium in molten NdF3 + PrF3 + LiF + 1 wt.%Pr6O11 + 1 wt.%Nd2O3 and NdF3 + PrF3 + LiF + 2 wt.%Pr6O11 + 2 wt.%Nd2O3 electrolytes at 1323 K were investigated. Cyclic voltammetry, square wave voltammetry, and open circuit potentiometry were applied to study the electrochemical reduction of Nd(III) and Pr(III) ions on Mo and W cathodes. It was established that a critical condition for Nd and Pr co-deposition in oxyfluoride electrolytes was a constant praseodymium deposition overpotential of ??0.100 V, which was shown to result in co-deposition current densities approaching 6 mAcm?2. Analysis of the results obtained by applied electrochemical techniques showed that praseodymium deposition proceeds as a one-step process involving exchange of three electrons (Pr(III)?Pr(0)) and that neodymium deposition is a two-step process: the first involves one electron exchange (Nd(III)?Nd(II)), and the second involves an exchange of two electrons (Nd(II)?Nd(0)). X-ray diffraction analyses confirmed the formation of metallic Nd and Pr on the working substrate. Keeping the anodic potential to the glassy carbon working anode low results in very low levels of carbon oxides, fluorine and fluorocarbon gas emissions, which should qualify the studied system as an environmentally friendly option for rare earth metal deposition. The newly reported data for Nd and Pr metals co-deposition provide valuable information for the recycling of neodymium-iron-boron magnets. Full article
(This article belongs to the Special Issue Advances in Understanding Metal Electrolysis Processes)
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<p>Electrochemical cell setup.</p> Full article ">Figure 2
<p>Cyclic voltammograms recorded on W working electrode in NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 1 wt.%Pr<sub>6</sub>O<sub>11</sub> + 1 wt.%Nd<sub>2</sub>O<sub>3</sub> electrolyte, E<sub>I</sub> = −0.050 V vs. W to different cathodic end potential E<sub>C</sub>, obtained with sweep rate of 50 mV·s<sup>−1</sup>.</p> Full article ">Figure 3
<p>CVs recorded on: (<b>a</b>) Mo working electrode starting from: E<sub>I</sub> = −0.050 V to E<sub>C</sub> = −0.900 V vs. W; (<b>b</b>) W working electrode starting from: E<sub>I</sub> = −0.050 V to E<sub>C</sub> = −0.950 V vs. W. The voltammograms were obtained with different sweep rates in molten NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 1 wt.%Pr<sub>6</sub>O<sub>11</sub> + 1 wt.%Nd<sub>2</sub>O<sub>3</sub> electrolyte, temperature 1323 K.</p> Full article ">Figure 4
<p>Plots of anodic peak A<sub>1</sub>: (<b>a</b>) current densities vs. square root of scan rate calculated from <a href="#metals-11-01494-f003" class="html-fig">Figure 3</a>a; (<b>b</b>) logarithm of the peak current density vs. peak potential from the same anodic peak in <a href="#metals-11-01494-f003" class="html-fig">Figure 3</a>a; (<b>c</b>) plots of peaks A<sub>3</sub>/C<sub>3</sub>, anodic and cathodic peak current densities vs. square root of scan rates derived from <a href="#metals-11-01494-f003" class="html-fig">Figure 3</a>b.</p> Full article ">Figure 5
<p>Voltammograms obtained in NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 1 wt.%Pr<sub>6</sub>O<sub>11</sub> + 1 wt.%Nd<sub>2</sub>O<sub>3</sub> melt with different holding times (τ = 5 and 60 s) at negative end potential of the cycle on the Mo working electrode; sweep rate 100 mV·s<sup>−1</sup>, temperature 1323 K.</p> Full article ">Figure 6
<p>Voltammograms obtained in NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 1 wt.%Pr<sub>6</sub>O<sub>11</sub> + 1 wt.%Nd<sub>2</sub>O<sub>3</sub> melt with different holding times (τ = 60, 180 and 300 s) at negative end potential of the cycle: (<b>a</b>) Mo working electrode, (<b>b</b>) W working electrode, sweep rate 100 mV·s<sup>−1</sup>.</p> Full article ">Figure 7
<p>(<b>a</b>) SWV obtained with different frequencies on Mo WE; pulse height: 25 mV; potential step: 1 mV; electrolyte NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 1 wt.%Pr<sub>6</sub>O<sub>11</sub> + 1 wt.%Nd<sub>2</sub>O<sub>3</sub>; (<b>b</b>,<b>c</b>) variation of the peak current density vs. the square root of the frequency in <a href="#metals-11-01494-f007" class="html-fig">Figure 7</a>a for peaks II and III.</p> Full article ">Figure 8
<p>CV recorded on Mo working electrode in molten NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 2 wt.%Pr<sub>6</sub>O<sub>11</sub> + 2 wt.% Nd<sub>2</sub>O<sub>3</sub> electrolyte, from E<sub>I</sub> = −0.050 V to E<sub>C</sub> = −0.900 V vs. W; obtained with different scan rates.</p> Full article ">Figure 9
<p>Voltammograms obtained on Mo WE in NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 2 wt.%Pr<sub>6</sub>O<sub>11</sub> + 2 wt.%Nd<sub>2</sub>O<sub>3</sub> melt with different holding times (τ = 30, 60 or 180 s) at negative end potential of the cycle of −0.900 V vs. W; sweep rate 100 mV·s<sup>−1</sup>.</p> Full article ">Figure 10
<p>Open circuit potential transient curves obtained on Mo working electrode in molten NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 2 wt.%Pr<sub>6</sub>O<sub>11</sub> + 2 wt.%Nd<sub>2</sub>O<sub>3</sub> electrolyte, after short polarization of 180 s at −0.900 V vs. W.</p> Full article ">Figure 11
<p>(<b>a</b>) SWV obtained on Mo WE in NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 2 wt.%Pr<sub>6</sub>O<sub>11</sub> + 2 wt.%Nd<sub>2</sub>O<sub>3</sub> electrolyte with different frequency, pulse amplitude = 25 mV, potential step = 1 mV; (<b>b</b>) Variation of the peak current density versus the square root of the frequency in <a href="#metals-11-01494-f011" class="html-fig">Figure 11</a>a for the peak III.</p> Full article ">Figure 12
<p>(<b>a</b>) Appearance of the surface of Mo cathode after deposition at −0.900 V vs. W for 60 min, from NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 2 wt.%Pr<sub>6</sub>O<sub>11</sub> + 2 wt.%Nd<sub>2</sub>O<sub>3</sub> electrolyte, at 1323 K; (<b>b</b>) XRD pattern of the Mo working electrode from <a href="#metals-11-01494-f012" class="html-fig">Figure 12</a>a.</p> Full article ">Figure 13
<p>(<b>a</b>) Appearance of the surface of the Mo cathode after deposition at −0.900 V vs. W for 90 min, from molten NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 2 wt.% Pr<sub>6</sub>O<sub>11</sub> + 2 wt.%Nd<sub>2</sub>O<sub>3</sub> electrolyte, at 1323 K. (the solidified electrolyte was removed from the electrode); (<b>b</b>) XRD pattern of the Mo working electrode from <a href="#metals-11-01494-f013" class="html-fig">Figure 13</a>a.</p> Full article ">Figure 14
<p>Off-gas generated on the GC anode recorded in situ with an FTIR spectrometer during: (<b>a</b>) electrochemical experiments; (<b>b</b>) potentiostatic deposition at −0.900 V vs. W; insert are measurements for CF<sub>4</sub> off-gas; working electrode Mo, in NdF<sub>3</sub> + PrF<sub>3</sub> + LiF + 2 wt.%Pr<sub>6</sub>O<sub>11</sub> + 2 wt.%Nd<sub>2</sub>O<sub>3</sub> electrolyte, T = 1323 K.</p> Full article ">
14 pages, 7242 KiB  
Article
Effects of Process Control Agent Amount, Milling Time, and Annealing Heat Treatment on the Microstructure of AlCrCuFeNi High-Entropy Alloy Synthesized through Mechanical Alloying
by Negar Yazdani, Mohammad Reza Toroghinejad, Ali Shabani and Pasquale Cavaliere
Metals 2021, 11(9), 1493; https://doi.org/10.3390/met11091493 - 20 Sep 2021
Cited by 35 | Viewed by 3376
Abstract
This study was conducted to investigate the characteristics of the AlCrCuFeNi high-entropy alloy (HEA) synthesized through mechanical alloying (MA). In addition, effects of Process Control Agent (PCA) amount and milling time were investigated using X-ray diffraction analysis (XRD), scanning electron microscopy (SEM), and [...] Read more.
This study was conducted to investigate the characteristics of the AlCrCuFeNi high-entropy alloy (HEA) synthesized through mechanical alloying (MA). In addition, effects of Process Control Agent (PCA) amount and milling time were investigated using X-ray diffraction analysis (XRD), scanning electron microscopy (SEM), and energy dispersive X-ray spectroscopy (EDS). The results indicated that the synthesized AlCrCuFeNi alloy is a dual phase (FCC + BCC) HEA and the formation of the phases is strongly affected by the PCA amount. A high amount of PCA postponed the alloying process and prevented solid solution formation. Furthermore, with an increase in the PCA amount, lattice strain decreased, crystallite size increased, and the morphology of the mechanically alloyed particles changed from spherical to a plate-like shape. Additionally, investigation of thermal properties and annealing behavior at different temperatures revealed no phase transformation up to 400 °C; however, the amount of the phases changed. By increasing the temperature to 600 °C, a sigma phase (?) and a B2-ordered solid solution formed; moreover, at 800 °C, the FCC phase decomposed into two different FCC phases. Full article
(This article belongs to the Special Issue Spark Plasma Sintering on Metals and Alloys)
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Figure 1
<p>XRD patterns of AlCrCuFeNi alloy after 60 h of milling with different amounts of Process Control Agent—PCA (1, 2, 3, and 4 wt. %).</p> Full article ">Figure 2
<p>Variation in average crystallite size (nm) and lattice strain (%) of mechanically alloyed AlCrCuFeNi HEA after 60 h of milling as a function of PCA amount.</p> Full article ">Figure 3
<p>SEM images of mechanically alloyed AlCrCuFeNi HEA after 60 h of milling in the presence of (<b>a</b>,<b>b</b>) 1, (<b>c</b>,<b>d</b>) 2, (<b>e</b>,<b>f</b>) 3, and (<b>g</b>,<b>h</b>) 4 wt. % of PCA amount.</p> Full article ">Figure 3 Cont.
<p>SEM images of mechanically alloyed AlCrCuFeNi HEA after 60 h of milling in the presence of (<b>a</b>,<b>b</b>) 1, (<b>c</b>,<b>d</b>) 2, (<b>e</b>,<b>f</b>) 3, and (<b>g</b>,<b>h</b>) 4 wt. % of PCA amount.</p> Full article ">Figure 4
<p>XRD patterns of the AlCrCuFeNi HEA after different time periods of mechanical alloying.</p> Full article ">Figure 5
<p>Average crystallite size (nm) and lattice strain (%) of AlCrCuFeNi HEA as a function of milling time.</p> Full article ">Figure 6
<p>SEM images of the AlCrCuFeNi HEA after (<b>a</b>,<b>b</b>) 10 h, (<b>c</b>,<b>d</b>) 30 h, (<b>e</b>,<b>f</b>) 50 h, and (<b>g</b>,<b>h</b>) 60 h of milling.</p> Full article ">Figure 6 Cont.
<p>SEM images of the AlCrCuFeNi HEA after (<b>a</b>,<b>b</b>) 10 h, (<b>c</b>,<b>d</b>) 30 h, (<b>e</b>,<b>f</b>) 50 h, and (<b>g</b>,<b>h</b>) 60 h of milling.</p> Full article ">Figure 7
<p>The EDS spectrum (<b>a</b>) and quantitative analysis (<b>b</b>) of AlCrCuFeNi alloy after 60 h of milling.</p> Full article ">Figure 8
<p>Differential Scanning Calorimetry (DSC) curve of AlCrCuFeNi HEA.</p> Full article ">Figure 9
<p>XRD patterns of annealed AlCrCuFeNi HEA powders at different temperatures.</p> Full article ">
24 pages, 9338 KiB  
Article
Investigating Nanoindentation Creep Behavior of Pulsed-TIG Welded Inconel 718 and Commercially Pure Titanium Using a Vanadium Interlayer
by Tauheed Shehbaz, Fahd Nawaz Khan, Massab Junaid and Julfikar Haider
Metals 2021, 11(9), 1492; https://doi.org/10.3390/met11091492 - 20 Sep 2021
Cited by 8 | Viewed by 3137
Abstract
In a dissimilar welded joint between Ni base alloys and titanium, creep failure is a potential concern as it could threaten to undermine the integrity of the joint. In this research, the mechanical heterogeneity of a Pulsed TIG welded joint between commercially pure [...] Read more.
In a dissimilar welded joint between Ni base alloys and titanium, creep failure is a potential concern as it could threaten to undermine the integrity of the joint. In this research, the mechanical heterogeneity of a Pulsed TIG welded joint between commercially pure titanium (CpTi) and Inconel 718 (IN718) with a vanadium (V) interlayer was studied through a nanoindentation technique with respect to hardness, elastic modulus, and ambient temperature creep deformation across all regions (fusion zones and interfaces, mainly composed of a dendritic morphology). According to the experimental results, a nanohardness of approximately 10 GPa was observed at the V/IN718 interface, which was almost 70% higher than that at the V/CpTi interface. This happened due to the formation of intermetallic compounds (IMCs) (e.g., Ti2Ni, NiV3, NiTi) and a (Ti, V) solid solution at the V/IN718 and V/CpTi interfaces, respectively. In addition, nanohardness at the V/IN718 interface was inhomogeneous as compared to that at the V/CpTi interface. Creep deformation behavior at the IN718 side was relatively higher than that at different regions on the CpTi side. The decreased plastic deformation or creep effect of the IMCs could be attributed to their higher hardness value. Compared to the base metals (CpTi and IN718), the IMCs exhibited a strain hardening effect. The calculated values of the creep stress exponent were found in the range of 1.51–3.52 and 2.52–4.15 in the V/CpTi and V/IN718 interfaces, respectively. Furthermore, the results indicated that the creep mechanism could have been due to diffusional creep and dislocation climb. Full article
(This article belongs to the Special Issue High-Productivity Welding of Metals and Alloys)
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Figure 1
<p>(<b>a</b>) Schematic diagram. (<b>b</b>) Experimental set-up of TIG welding process.</p> Full article ">Figure 2
<p>(<b>a</b>) SEM micrographs of the CpTi/V/IN718 joint with EDS line scan. Microstructures in the common melt zone of all three materials in the arc facing side: (<b>b</b>) Inconel 718 side; (<b>c</b>) region right above the V interlayer; and (<b>d</b>) the CpTi side.</p> Full article ">Figure 3
<p>(<b>a</b>) Interfacial microstructures of V/IN718. Inset in (<b>a</b>) is a magnified image of the region indicated with the arrow. Elemental mapping of the region shown in (<b>a</b>) for (<b>b</b>) Cr, (<b>c</b>) V, (<b>d</b>) Fe, and (<b>e</b>) Ni.</p> Full article ">Figure 3 Cont.
<p>(<b>a</b>) Interfacial microstructures of V/IN718. Inset in (<b>a</b>) is a magnified image of the region indicated with the arrow. Elemental mapping of the region shown in (<b>a</b>) for (<b>b</b>) Cr, (<b>c</b>) V, (<b>d</b>) Fe, and (<b>e</b>) Ni.</p> Full article ">Figure 4
<p>(<b>a</b>) Interfacial microstructures of V/CpTi interface. (<b>b</b>,<b>c</b>) Elemental mapping of highlighted region in (<b>a</b>).</p> Full article ">Figure 5
<p>XRD spectra of the CpTi/V/IN718 joint.</p> Full article ">Figure 6
<p>(<b>a</b>) SEM Image showing indentation on V/CpTi side. (<b>b</b>) Nano-indentation results of load vs. depth curves for CpTi and vanadium interlayer joint. Inset of (<b>b</b>) shows the pop-ins that appeared in the FZ3 and BM (CpTi) curves.</p> Full article ">Figure 7
<p>(<b>a</b>) SEM images showing nano indents on the V/IN718 side. (<b>b</b>) Nano-indentation results of load vs. depth curves for (N718) and the vanadium interlayer joint.</p> Full article ">Figure 8
<p>Locations selected for nanohardness line scans across the CpTi/V/IN718 joint.</p> Full article ">Figure 9
<p>(<b>a</b>) Nanohardness line scan at the V/IN718 Interface (location (A) in <a href="#metals-11-01492-f008" class="html-fig">Figure 8</a>)). (<b>b</b>) Corresponding hardness and elastic modulus profiles along the line scan.</p> Full article ">Figure 10
<p>(<b>a</b>) Nanohardness line scan at V/CpTi Interface (location (B) in <a href="#metals-11-01492-f008" class="html-fig">Figure 8</a>). (<b>b</b>) Corresponding hardness and elastic modulus profiles along the line scan.</p> Full article ">Figure 11
<p>(<b>a</b>) Nanohardness line scan at V/FZ3 Interface (location (C) in <a href="#metals-11-01492-f008" class="html-fig">Figure 8</a>). (<b>b</b>) Corresponding hardness and elastic modulus profiles along the line scan.</p> Full article ">Figure 12
<p>(<b>a</b>) Nanohardness line scan at V/FZ1 Interface (location (D) in <a href="#metals-11-01492-f008" class="html-fig">Figure 8</a>). (<b>b</b>) Corresponding hardness and elastic modulus profiles along the line scan.</p> Full article ">Figure 13
<p>(<b>a</b>) Nanohardness line scan at FZ3/FZ4 Interface (location (E) in <a href="#metals-11-01492-f008" class="html-fig">Figure 8</a>). (<b>b</b>) CorreScheme 718. interface has a relatively higher nanohardness, approximately 30% greater than FZ2 (which has the highest hardness after the V/IN718 interface) and 70% greater than the IN718 base metal.</p> Full article ">Figure 14
<p>Nanohardness maps. (<b>a</b>) The V/IN718 interface. (<b>b</b>) The V/CpTi interface.</p> Full article ">Figure 15
<p>(<b>a</b>) The load–displacement (P–h) curves of different weld regions of V and CpTi obtained under 200 mN at 200 sec holding time. (<b>b</b>) The displacement–time (h–t) curves obtained in the holding regimes of different regions of V and the CpTi side.</p> Full article ">Figure 16
<p>Experimental data and fitted curves for: (<b>a</b>) BM of the CpTi side; (<b>b</b>) FZ3; (<b>c</b>) FZ4; (<b>d</b>) the FZ4/FZ3 interface; (<b>e</b>) the V/FZ3 interface; (<b>f</b>) the V/CpTi interface; and (<b>g</b>) the V interlayer.</p> Full article ">Figure 16 Cont.
<p>Experimental data and fitted curves for: (<b>a</b>) BM of the CpTi side; (<b>b</b>) FZ3; (<b>c</b>) FZ4; (<b>d</b>) the FZ4/FZ3 interface; (<b>e</b>) the V/FZ3 interface; (<b>f</b>) the V/CpTi interface; and (<b>g</b>) the V interlayer.</p> Full article ">Figure 17
<p>(<b>a</b>) ln (creep strain rate)-ln(stress) plots of different regions of V and CpTi side. (<b>b</b>) Creep stress exponent of different regions in V and the CpTi side.</p> Full article ">Figure 18
<p>(<b>a</b>) The load–displacement (P–h) curves of different regions of V and IN718 obtained under 200 mN at 200 sec holding time. (<b>b</b>) The displacement–time (h–t) curves obtained in the holding regimes of different regions of V and the IN718 side.</p> Full article ">Figure 19
<p>Experimental data and fitted depth–time curves for: (<b>a</b>) BM/IN718; (<b>b</b>) FZ1; (<b>c</b>) FZ2; (<b>d</b>) V/INC 718; and (<b>e</b>) the V interlayer.</p> Full article ">Figure 20
<p>(<b>a</b>) The ln (creep strain rate)-ln (creep stress) curves obtained under the peak load in different regions of IN718. (<b>b</b>) Creep stress exponent of different regions in V and the IN718 side.</p> Full article ">
9 pages, 2485 KiB  
Article
Preparation of Ag0 Nanoparticles by EDM Method as Catalysts for Oxygen Reduction
by Jia Guo, Xiaoming Mu, Shihao Song, Yanwei Ren, Kai Wang and Zunming Lu
Metals 2021, 11(9), 1491; https://doi.org/10.3390/met11091491 - 20 Sep 2021
Cited by 5 | Viewed by 2480
Abstract
At present, platinum-based catalysts are the best cathode catalysts, but due to their high prices, they are difficult to use widely. Under alkaline conditions, silver is a better low-cost substitute. Here, a physical preparation method—electrical discharge machining (EDM)—is used to prepare Ag0 [...] Read more.
At present, platinum-based catalysts are the best cathode catalysts, but due to their high prices, they are difficult to use widely. Under alkaline conditions, silver is a better low-cost substitute. Here, a physical preparation method—electrical discharge machining (EDM)—is used to prepare Ag0 nanoparticles. The method is simple and has a high yield. The diameter of prepared nanoparticles is about 30 nm and the nanoparticle surface is rich in defects. These defects enhance the adsorption of O2. In addition, defects can cause tensile strain on the silver catalyst, causing the d-band center of silver to move upward. The defects and the upward shift of the d-band center jointly improve the adsorption energy and catalytic performance of Ag0. This work provides a new method for the engineering construction of surface defects and the preparation of metal catalysts. Full article
(This article belongs to the Special Issue Research on Metal Nanoparticles)
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<p>Schematic diagram of preparation and synthesis of Ag<sup>0</sup> nanoparticles.</p> Full article ">Figure 2
<p>XRD patterns of Ag and NPS-Ag<sup>0</sup> samples.</p> Full article ">Figure 3
<p>(<b>a</b>) SEM and (<b>b</b>) TEM images of the NPS-Ag<sup>0</sup>, high-resolution TEM images (<b>c</b>,<b>d</b>) of the NPS-Ag<sup>0</sup>. The stacking faults, amorphous region, grain boundary, and twin defects are indicated.</p> Full article ">Figure 4
<p>(<b>a</b>) ORR polarization curves of catalysts in O<sub>2</sub>-saturated 0.1 M KOH solution at scan rate of 5 mV·s<sup>−1</sup> and rotation rate of 1600 rpm. (<b>b</b>) Koutecky–Levich plots at 0.5 V electrode potential (<b>c</b>) Tafel plots from catalysts. (<b>d</b>) Comparison of catalytic activities for ORR of each sample given as the j<sub>k,mass</sub> and j<sub>k,specific</sub>.</p> Full article ">Figure 5
<p>Chronoamperometric response of the NPS-Ag<sup>0</sup>, Ag and Pt/C at 0.5 V (percentage of the current retained-operation time).</p> Full article ">Figure 6
<p>(<b>a</b>) XPS spectra of Ag. (<b>b</b>) XPS spectra of NPS-Ag<sup>0</sup>. (<b>c</b>) Valence band spectra (VBS) of the Ag and NPS-Ag<sup>0</sup>.</p> Full article ">
10 pages, 3741 KiB  
Article
Preparation of Battery-Grade Lithium Carbonate with Lithium-Containing Desorption Solution
by Zheng-Guo Xu and Shu-Ying Sun
Metals 2021, 11(9), 1490; https://doi.org/10.3390/met11091490 - 19 Sep 2021
Cited by 11 | Viewed by 6648
Abstract
In this study, a process for preparing battery-grade lithium carbonate with lithium-rich solution obtained from the low lithium leaching solution of fly ash by adsorption method was proposed. A carbonization-decomposition process was carried out to remove impurities such as iron and aluminum. First, [...] Read more.
In this study, a process for preparing battery-grade lithium carbonate with lithium-rich solution obtained from the low lithium leaching solution of fly ash by adsorption method was proposed. A carbonization-decomposition process was carried out to remove impurities such as iron and aluminum. First, primary Li2CO3 was treated by CO2 to get the more soluble bicarbonates. The decomposition of LiHCO3 produced insoluble Li2CO3 at 90 °C And Li2CO3 was smashed by air stream pulverization. The final precipitation yielded a high purity (99.6%) and homogeneous Li2CO3. Some factors affecting production efficiency were investigated. The results showed that a liquid-solid ratio of 25:1, a carbonization temperature of 25 °C, an air velocity of 2 L/min, and a stirring speed of 400 rpm; a decomposition temperature of 90 °C and a stirring speed of 400 rpm, a molar ratio of EDTA to Ca 2:1; an air pressure of 0.3 MPa and hot water washing precipitate (L/S mass ratio 2:1) promoted ions removal. Full article
(This article belongs to the Special Issue Recent Advances in Leaching and Extractive Metallurgy)
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<p>Flowsheet developed for the recovery of battery-grade Li<sub>2</sub>CO<sub>3</sub> from the lithium-containing desorption solution.</p> Full article ">Figure 2
<p>Washing process of water and ethanol: (<b>a</b>) lithium loss; (<b>b</b>) sodium content.</p> Full article ">Figure 3
<p>Effect of L/S (<b>a</b>) temperature (<b>b</b>) flow rate (<b>c</b>) time (<b>d</b>) on the concentration of impurities [liquid-solid ratio = 20, temperature = 25 °C, flow rate of CO<sub>2</sub> = 3 L/min].</p> Full article ">Figure 4
<p>Effect of liquid-solid ratio on carbonization reaction [temperature = 25 °C, flow rate of CO<sub>2</sub> = 3 L/min].</p> Full article ">Figure 5
<p>Effect of temperature on carbonization reaction [liquid-solid ratio = 20, flow rate of CO<sub>2</sub> = 3 L/min].</p> Full article ">Figure 6
<p>Effect of flow rate on carbonization reaction [liquid-solid ratio = 20, temperature = 25 °C].</p> Full article ">Figure 7
<p>Variation of pH as a function of time.</p> Full article ">Figure 8
<p>Effect of temperature on thermal decomposition crystallization [agitation speed = 300 rpm].</p> Full article ">Figure 9
<p>Sem-analysis of products under different temperatures.</p> Full article ">Figure 10
<p>Effect of agitation speed on thermal decomposition crystallization [temperature = 90 °C].</p> Full article ">Figure 11
<p>Effect of pressure on particle size. (<b>a</b>) average particle size (<b>b</b>) particle size distribution.</p> Full article ">
11 pages, 3702 KiB  
Article
Dynamic Mechanical Properties of Ti–Al3Ti–Al Laminated Composites: Experimental and Numerical Investigation
by Jian Ma, Meini Yuan, Lirong Zheng, Zeyuan Wei and Kai Wang
Metals 2021, 11(9), 1489; https://doi.org/10.3390/met11091489 - 19 Sep 2021
Cited by 4 | Viewed by 2447
Abstract
The Ti–Al3Ti–Al laminated composites with different Al contents were prepared by vacuum hot pressing sintering technology. The effects of Al content on the dynamic mechanical properties of the composites were studied using the combination of Split Hopkinson Pressure Bar experiment and finite element [...] Read more.
The Ti–Al3Ti–Al laminated composites with different Al contents were prepared by vacuum hot pressing sintering technology. The effects of Al content on the dynamic mechanical properties of the composites were studied using the combination of Split Hopkinson Pressure Bar experiment and finite element analysis. The results showed that different Al content changes the fracture mode of the composites. The laminated composites without Al have higher brittleness and lower fracture strain. The Ti–Al3Ti–Al laminated composites containing 10–15%Al have better dynamic mechanical properties than those without Al, but the subsequent increase of Al content is not conducive to the improvement of strength. However, when the Al content in the specimen reaches 30%, the dynamic mechanical properties of the composites decrease, multi-crack phenomenon and relatively large strain occur, and the Al extruded from the layers fills the crack. Full article
(This article belongs to the Special Issue Research on Microstructure and Mechanical Properties of Metallic Materials and Metal-Matrix Composites)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>Schematic of a Split Hopkinson pressure bar apparatus.</p> Full article ">Figure 2
<p>(<b>a</b>) The finite element mesh SHPB experiment system (local), (<b>b</b>) Finite element model of Ti–Al3Ti–Al specimen, (<b>c</b>) Metallographic diagram of Ti–Al3Ti–Al specimen.</p> Full article ">Figure 3
<p>Dynamic compressive stress–strain curves of Ti–Al3Ti–Al laminated composites with different Al contents.</p> Full article ">Figure 4
<p>The specimen after fracture and the fracture morphology observed under electron microscope: (<b>a</b>) Without Al, (<b>b</b>) containing about 15% Al, (<b>c</b>) containing about 30% Al.</p> Full article ">Figure 5
<p>The simulated dynamic stress-strain curves of Ti–Al3Ti–Al laminated composites with different Al contents.</p> Full article ">Figure 6
<p>Energy absorption–time curves of Ti–Al3Ti–Al laminated composites with different Al contents: (<b>a</b>) the whole; (<b>b</b>) Ti layers.</p> Full article ">Figure 7
<p>The deformation degree of Ti layer in Ti–Al3Ti–Al laminated composites with different Al contents: (<b>a</b>) 0% Al, (<b>b</b>) 10% Al, (<b>c</b>) 20% Al, (<b>d</b>) 30%.</p> Full article ">Figure 8
<p>Stress nephogram of longitudinal section of Ti–Al3Ti–Al laminated composites with different Al contents: (<b>a</b>) 0% Al, (<b>b</b>) 10% Al (<b>c</b>), 20% Al, (<b>d</b>) 30%.</p> Full article ">
15 pages, 4579 KiB  
Article
Effect of the Thermomechanical Treatment on the Corrosion of UNSM Processed Inconel 718: An Electrochemical Study
by Ulises Martin, Jacob Ress and David M. Bastidas
Metals 2021, 11(9), 1488; https://doi.org/10.3390/met11091488 - 18 Sep 2021
Cited by 2 | Viewed by 2957
Abstract
In this work, the influence of thermal (TT), mechanical, and thermomechanical (TMT) treatments using the ultrasonic nanocrystal surface modification (UNSM) on the corrosion protection properties of Inconel 718 was studied, correlating the changes in the electrochemical properties with the promoted microstructure. The UNSM [...] Read more.
In this work, the influence of thermal (TT), mechanical, and thermomechanical (TMT) treatments using the ultrasonic nanocrystal surface modification (UNSM) on the corrosion protection properties of Inconel 718 was studied, correlating the changes in the electrochemical properties with the promoted microstructure. The UNSM treatment had a grain refinement effect on the top surface, reducing the grain size from 11.5 to 7.4 µm for the first 10 µm in depth. The high grain boundary density, due to the grain refinement, enabled a faster growth of the passive film. The impedance showed a decrease in the charge transfer resistance by three orders of magnitude, from 106 to 103 ? cm2 for as-received to 1000 °C, as the TT temperature crossed the solvus of the ??/?? and approached the solvus of the ?-phase. The UNSM treatment lowered the pitting corrosion susceptibility, increasing the charge transfer resistance and decreasing the effective capacitance of the double layer, leading to the thickest passive film with 6.8 nm. Full article
(This article belongs to the Special Issue Corrosion and Inhibition Processes)
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<p>Time–temperature–transformation (TTT) diagram of Inconel 718 [<a href="#B21-metals-11-01488" class="html-bibr">21</a>].</p> Full article ">Figure 2
<p>Optical microscopy images showing the cross-section microstructure of Inconel 718: (<b>a</b>) bulk of the AR ×10, and (<b>b</b>) UNSM sample showing the three regions ×50.</p> Full article ">Figure 3
<p>Vickers hardness as a function of the depth for UNSM and TMT samples.</p> Full article ">Figure 4
<p><span class="html-italic">E</span><sub>corr</sub> and <span class="html-italic">i</span><sub>corr</sub> monitoring for the different treatments of Inconel 718 in 3.5 wt.% NaCl.</p> Full article ">Figure 5
<p>EIS analysis of Inconel 718 in 3.5 wt.% NaCl: (<b>a</b>) Nyquist plots for the different treatments, with an inset at higher magnification for the lower frequency data values, and (<b>b</b>) EEC with two time constants.</p> Full article ">Figure 6
<p>SEM images showing the cross-section microstructure of Inconel 718 × 600: (<b>a</b>) AR sample, (<b>b</b>) TT1 sample, (<b>c</b>) TT2 sample, (<b>d</b>) TT3 sample, (<b>e</b>) UNSM sample, and (<b>f</b>) TMT sample.</p> Full article ">Figure 7
<p>SEM micrograph for the EDX analysis of TT3 sample with the selected areas for the elemental analysis ×10,000.</p> Full article ">Figure 8
<p>EDX analysis of TT3 sample: (<b>a</b>) elemental composition of the NbC, (<b>b</b>) elemental composition of the δ-phase, and (<b>c</b>) elemental composition of the γ-phase.</p> Full article ">Figure 9
<p><span class="html-italic">R</span><sub>ct</sub> comparison between LPR and EIS techniques of Inconel 718 samples in 3.5 wt.% NaCl.</p> Full article ">Figure 10
<p><span class="html-italic">C</span><sub>eff,dl</sub> and thickness of the passive film (<span class="html-italic">d</span>) of Inconel 718 samples in 3.5 wt.% NaCl.</p> Full article ">Figure 11
<p>IFM images of the top surface of Inconel 718 after the exposure to the 3.5 wt.% NaCl ×100: (<b>a</b>) AR sample, (<b>b</b>) TT1 sample, (<b>c</b>) TT2 sample, (<b>d</b>) TT3 sample, (<b>e</b>) UNSM sample, and (<b>f</b>) TMT sample.</p> Full article ">
12 pages, 5915 KiB  
Article
Manipulation of Microstructure and Mechanical Properties in N-Doped CoCrFeMnNi High-Entropy Alloys
by Jing Zhang, Kook Noh Yoon, Min Seok Kim, Heh Sang Ahn, Ji Young Kim, Wook Ha Ryu and Eun Soo Park
Metals 2021, 11(9), 1487; https://doi.org/10.3390/met11091487 - 18 Sep 2021
Cited by 14 | Viewed by 3140
Abstract
Herein, we carefully investigate the effect of nitrogen doping in the equiatomic CoCrFeMnNi high-entropy alloy (HEA) on the microstructure evolution and mechanical properties. After homogenization (1100 °C for 20 h), cold-rolling (reduction ratio of 60%) and subsequent annealing (800 °C for 1 h), [...] Read more.
Herein, we carefully investigate the effect of nitrogen doping in the equiatomic CoCrFeMnNi high-entropy alloy (HEA) on the microstructure evolution and mechanical properties. After homogenization (1100 °C for 20 h), cold-rolling (reduction ratio of 60%) and subsequent annealing (800 °C for 1 h), a unique complex heterogeneous microstructure consisting of fine recrystallized grains, large non-recrystallized grains, and nanoscale Cr2N precipitates, were obtained in nitrogen-doped (0.3 wt.%) CoCrFeMnNi HEA. The yield strength and ultimate tensile strength can be significantly improved in nitrogen-doped (0.3 wt.%) CoCrFeMnNi HEA with a complex heterogeneous microstructure, which shows more than two times higher than those compared to CoCrFeMnNi HEA under the identical process condition. It is achieved by the simultaneous operation of various strengthening mechanisms from the complex heterogeneous microstructure. Although it still has not solved the problem of ductility reduction, as the strength increases because the microstructure optimization is not yet complete, it is expected that precise control of the unique complex heterogeneous structure in nitrogen-doped CoCrFeMnNi HEA can open a new era in overcoming the strength–ductility trade-off, one of the oldest dilemmas of structural materials. Full article
(This article belongs to the Special Issue High-Entropy and Complex Concentrated Alloys: A New Generation of Alloys)
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<p>Microstructure and XRD patterns of as-annealed N0 and N0.3 HEAs: (<b>a</b>) BSE image of N0 HEA; (<b>b</b>) BSE image of N0.3 HEA; (<b>c</b>) XRD patterns of N0 and N0.3 HEAs obtained in the range of 30° to 100°; (<b>d</b>) XRD patterns of N0 and N0.3 HEAs obtained in the range of 40° to 55°.</p> Full article ">Figure 2
<p>EBSD results of as-annealed N0.3 HEA: (<b>a</b>) phase map; (<b>b</b>) grain map; (<b>c</b>) grain size distribution map; (<b>d</b>) KAM map; (<b>e</b>) Z-IPF map (color-coded with respect to the <span class="html-italic">z</span>-axis orientation); (<b>f</b>) color legend for IPF map (color-coded with respect to the <span class="html-italic">x</span>-axis orientation).</p> Full article ">Figure 3
<p>EBSD results of as-rolled N0.3 HEA: (<b>a</b>) phase map, (<b>b</b>) KAM map, (<b>c</b>) Z-IPF map, and (<b>d</b>) color legends for IPF map (color-coded with respect to the <span class="html-italic">x</span>- and <span class="html-italic">z</span>-axis orientation).</p> Full article ">Figure 4
<p>High magnification SEM images and EDS line scan of as-annealed N0.3 HEA: (<b>a</b>) BSE image; (<b>b</b>) magnified BSE image of b in (<b>a</b>); (<b>c</b>) EDS line scan of the precipitate c in (<b>b</b>).</p> Full article ">Figure 5
<p>Kikuchi orientation relationship between matrix and the precipitate in as-annealed N0.3 HEA: (<b>a</b>) BSE image; (<b>b</b>) phase map; (<b>c</b>) phase map of c in (<b>b</b>); (<b>d</b>) Kikuchi orientation relationship in the {111}<sub>M</sub>‖{0001}<sub>P</sub> plane; (<b>e</b>) Kikuchi orientation relationship in the &lt;110&gt;<sub>M</sub>‖<math display="inline"><semantics> <mrow> <mo>&lt;</mo> <mn>1</mn> <mover> <mn>1</mn> <mo>−</mo> </mover> <mn>00</mn> <msub> <mo>&gt;</mo> <mi mathvariant="normal">P</mi> </msub> </mrow> </semantics></math> direction.</p> Full article ">Figure 6
<p>(<b>a</b>) Engineering stress–strain curves and (<b>b</b>,<b>c</b>) fractography of an-annealed N0 and N0.3 HEAs.</p> Full article ">
31 pages, 10317 KiB  
Review
Dissimilar Non-Ferrous Metal Welding: An Insight on Experimental and Numerical Analysis
by Jeyaganesh Devaraj, Aiman Ziout and Jaber E. Abu Qudeiri
Metals 2021, 11(9), 1486; https://doi.org/10.3390/met11091486 - 18 Sep 2021
Cited by 20 | Viewed by 7071
Abstract
In recent years Gas Metal Arc Welding (GMAW) technology has expanded its functionalities in various areas which have further motivated its usage in several emerging manufacturing industries. There are several issues and challenges associated with this technology, especially in dissimilar metal welding (DMW). [...] Read more.
In recent years Gas Metal Arc Welding (GMAW) technology has expanded its functionalities in various areas which have further motivated its usage in several emerging manufacturing industries. There are several issues and challenges associated with this technology, especially in dissimilar metal welding (DMW). One of the predominant challenges is selecting appropriate welding parameters which influence the efficiency of this technology. To explore several modern advancements in this expertise, this paper has done an exclusive survey on various standards of GMAW and its variants for selecting suitable parameters for welding dissimilar nonferrous metals. This review summarizes various experimental and numerical results along with related illustrations to highlight the feasibility of welding dissimilar nonferrous metals using traditional GMAW and investigations on advanced GMAW processes such as cold metal transfer (CMT) and pulsed GMAW (P-GMAW). Simulation and modeling of nonferrous DMW have identified several research gaps and modeling problems. Researchers and manufacturers can use this review as a guideline to choose appropriate welding parameters to implement GMAW and its variants for non-ferrous dissimilar welding. It found that by controlling the heat input and effective post-heat treatments, adequate joint properties can be achieved. Automated large -scale manufacturing will widen the utilization scope of GMAW and avoid some costly methods such as laser welding, ultrasonic welding, and friction stir welding etc. Full article
(This article belongs to the Special Issue Numerical Simulation of Metals Welding Process)
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<p>Plan of the current review and the respective phases.</p> Full article ">Figure 2
<p>Architecture of the review.</p> Full article ">Figure 3
<p>Welding simulation standard procedure.</p> Full article ">Figure 4
<p>(<b>a</b>) Schematic of the rectangular heat source model; (<b>b</b>) double ellipsoidal heat source model, where <span class="html-italic">a</span>—length; <span class="html-italic">b</span>—half width; <span class="html-italic">c</span>—depth of penetration; <span class="html-italic">a<sub>f</sub></span>—front length of molten pool; <span class="html-italic">a<sub>r</sub></span>—rear length of molten pool.</p> Full article ">Figure 5
<p>Quality–speed diagram of welding models.</p> Full article ">Figure 6
<p>Layout of SYSWELD modeling adopted conventionally.</p> Full article ">Figure 7
<p>Representation of welding arrangement (<b>a</b>) Weld specimen dimension, (<b>b</b>) Clamping arrangement [<a href="#B47-metals-11-01486" class="html-bibr">47</a>] (from open access journal).</p> Full article ">Figure 8
<p>Color-coded temperature distribution in DMW; rebuilt from [<a href="#B48-metals-11-01486" class="html-bibr">48</a>] with from Elsevier.</p> Full article ">Figure 9
<p>Meshing and layers of weld bead in a welded joint between two dissimilar steels, Modified from [<a href="#B49-metals-11-01486" class="html-bibr">49</a>] with permission from Elsevier.</p> Full article ">Figure 10
<p>Microstructures of weld metal near the Ti-Al interface under different heat inputs: (<b>a</b>) 1.82–1.91 kJ/cm, (<b>b</b>) 1.90–1.99 kJ/cm, and (<b>c</b>) 2.05–2.14 kJ/cm [<a href="#B58-metals-11-01486" class="html-bibr">58</a>]. Figure reused with the permission from Taylor &amp; Francis Group.</p> Full article ">Figure 11
<p>Microstructures of descriptive points in the heat-affected zone of butt-welded titanium Ti6Al4V and aluminum 5A05Al metals [<a href="#B59-metals-11-01486" class="html-bibr">59</a>], (<b>a</b>) Top point, (<b>b</b>) top 1/3th point (<b>c</b>) bottom 1/3th point &amp; (<b>d</b>) root of the interface respectively, reused with permission from Springer Nature.</p> Full article ">Figure 12
<p>Schematic of the welding &amp; brazing CMT technique; Figure reused from [<a href="#B60-metals-11-01486" class="html-bibr">60</a>] with permission from open access Springer.</p> Full article ">Figure 13
<p>Macrographs of welded joints. (<b>I</b>)—0.5 mm, (<b>II</b>)—0.85 mm, (<b>III</b>)—1.20 mm, (<b>a</b>–<b>c</b>) increasing the wire-feed speed from 5–9 m/min [<a href="#B60-metals-11-01486" class="html-bibr">60</a>] (reused with open access journal Springer).</p> Full article ">Figure 14
<p>Hardness variation from the base metal toward the top surface in joints welded at various wire-feed rates [<a href="#B65-metals-11-01486" class="html-bibr">65</a>]. Reused with permission from Elsevier.</p> Full article ">Figure 15
<p>Macroscopic cross section of Joint made from the wire diameter 1 mm and feed rate of 5.0 m/min [<a href="#B69-metals-11-01486" class="html-bibr">69</a>].</p> Full article ">Figure 16
<p>Higher magnification of titanium-weld interface at ZONE A, B, &amp; C in <a href="#metals-11-01486-f015" class="html-fig">Figure 15</a> respectively [<a href="#B69-metals-11-01486" class="html-bibr">69</a>]. Picture is reused with permission from Elsevier.</p> Full article ">Figure 17
<p>Tensile shear strengths of joints formed under different welding parameters; reused from [<a href="#B69-metals-11-01486" class="html-bibr">69</a>] with permission of Elsevier.</p> Full article ">Figure 18
<p>Fracture surfaces of an Al-Cu joint in HAZ (I<sub>main</sub> = 45 A, I<sub>bypass</sub> = 25 A); (<b>a</b>) optical micrograph, (<b>b</b>) SEM image, and (<b>c</b>) magnified view of the fracture surface [<a href="#B73-metals-11-01486" class="html-bibr">73</a>] (With permission from Elsevier).</p> Full article ">Figure 19
<p>Fracture surfaces of an Al–Cu joint around the fusion line ((I<sub>main</sub> = 25 A, I<sub>bypass</sub> = 25 A); (<b>a</b>) optical micrograph, (<b>b</b>) SEM image, and (<b>c</b>) magnified view of the fracture surface [<a href="#B73-metals-11-01486" class="html-bibr">73</a>] (With permission from Elsevier).</p> Full article ">Figure 20
<p>(<b>a</b>) Microstructure of the weld/Cu interface after EMF-assisted CMT welding with a coil current of 1.6 A and frequency of 0 Hz at ×250. (<b>b</b>) Magnification of part I at ×750 (<b>c</b>) Magnification of part II at ×750 [<a href="#B74-metals-11-01486" class="html-bibr">74</a>] (With permission from Taylor &amp; Francis Group).</p> Full article ">Figure 21
<p>Micrographs of Co–Ni joints welded by TIG and EBW, where B- Bead width, h- Reinforcement, B’- Bead width on bottom; h’- Reinforcement on bottom. [<a href="#B79-metals-11-01486" class="html-bibr">79</a>] from open access journal.</p> Full article ">Figure 22
<p>Schematic representation of Al &amp; Mg joint.</p> Full article ">Figure 23
<p>Optical micrograph of the fusion zone at the Mg side of a Mg–Al welded joint; Reused from [<a href="#B77-metals-11-01486" class="html-bibr">77</a>] with permission from Elsevier.</p> Full article ">Figure 24
<p>Microhardness distribution along the weld joint, reused from [<a href="#B77-metals-11-01486" class="html-bibr">77</a>] with permission from Elsevier.</p> Full article ">Figure 25
<p>Microstructure of the welding interface in an Mg–Ti weld (<b>left</b>) and hardness distribution along the brazing interface (<b>right</b>) (top sheet = Mg; bottom sheet = Ti) [<a href="#B72-metals-11-01486" class="html-bibr">72</a>]. Picture reused with permission from Elsevier.</p> Full article ">
9 pages, 5143 KiB  
Article
Natural Aging Effect of Al-20Zn-3Cu Alloy on Mechanical Properties and Its Relation to Microstructural Change
by Min-Jung Park, Hyeongsub So, Kyou-Hyun Kim, Jai-Won Byeon and Leeseung Kang
Metals 2021, 11(9), 1485; https://doi.org/10.3390/met11091485 - 18 Sep 2021
Cited by 4 | Viewed by 2041
Abstract
We investigate the effect of the natural age-hardening response of the Al-20Zn-3Cu alloy with natural aging times up to 12 months. The ultimate tensile strength of the Al-20Zn-3Cu alloy is drastically enhanced from 308 to 320 MPa after 2 months and from 320 [...] Read more.
We investigate the effect of the natural age-hardening response of the Al-20Zn-3Cu alloy with natural aging times up to 12 months. The ultimate tensile strength of the Al-20Zn-3Cu alloy is drastically enhanced from 308 to 320 MPa after 2 months and from 320 to 346 MPa after 9 months. Then, natural age hardening becomes saturated after 9 months. A microstructural investigation reveals that the natural age-hardening mechanism is mainly induced by the diffusion of the Zn element. First, a rapid decrease in the volume fraction of the eutectoid lamellae (?-Al+?-Zn) is observed at the early stage of natural aging, leading to an increase in the tensile strength. This originates from the relatively high diffusivity of Zn due to its low melting temperature. Then, the diffusion of Zn into the Al matrix induces clusters of solute atoms that enhance the growth rate of the nanoprecipitates formed in the Al matrix. As a consequence, the tensile strength of the natural-aged Al-20Zn-3Cu alloy increases drastically after 9 months, whereas the ductility is significantly degraded. Full article
(This article belongs to the Special Issue High-Strength Alloys)
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<p>Fracture strength and elongation of Al-20Zn-3Cu with respect to the natural aging time for the as-cast sample and those at 2 months, 9 months, and 12 months.</p> Full article ">Figure 2
<p>XRD profiles of the Al-20Zn-3Cu alloy at different natural aging times.</p> Full article ">Figure 3
<p>Typical BSE images of Al-20Zn-3Cu: (<b>a</b>) as-cast, (<b>b</b>) 2 months, (<b>c</b>) 9 months, and (<b>d</b>) 12 months. The insets show the decreased volume fraction of the eutectoid lamellar structure consisting of the α-Al and the η-Zn phase.</p> Full article ">Figure 4
<p>Elemental mapping results of Al-20Zn-3Cu: (<b>a</b>) as-cast, (<b>b</b>) 2 months, (<b>c</b>) 9 months, and (<b>d</b>) 12 months. Area I and II respectively indicate the pure Cu and the eutectoid lamellar structure.</p> Full article ">Figure 5
<p>Typical Z-contrast images respectively recorded from (<b>a</b>) as-cast, (<b>b</b>) 2 months, and (<b>c</b>) 9 months natural-aged Al-20Zn-3Cu alloys. Corresponding electron diffraction patterns are presented in the insets.</p> Full article ">
13 pages, 4479 KiB  
Article
Systematic Development of Eutectic High Entropy Alloys by Thermodynamic Modeling and Experimentation: An Example of the CoCrFeNi-Mo System
by Muhammad Mukarram, M. Awais Munir, Mohammad Mujahid and Khurram Yaqoob
Metals 2021, 11(9), 1484; https://doi.org/10.3390/met11091484 - 18 Sep 2021
Cited by 14 | Viewed by 3273
Abstract
Face centered cubic (FCC) high-entropy alloys (HEA) exhibit excellent ductility while body centered cubic (BCC) HEAs are characterized by high strength. Development of fine two-phase eutectic microstructure (consisting of a tough phase such as fcc and a hard phase such as bcc/intermetallic) can [...] Read more.
Face centered cubic (FCC) high-entropy alloys (HEA) exhibit excellent ductility while body centered cubic (BCC) HEAs are characterized by high strength. Development of fine two-phase eutectic microstructure (consisting of a tough phase such as fcc and a hard phase such as bcc/intermetallic) can help in obtaining an extraordinary combination of strength and ductility in HEAs. Designing eutectic high entropy alloys is an extremely difficult task for which different empirical and non-empirical methods have been previously tried. In the present study, the possibility of developing a eutectic microstructure by the addition of Mo to CoCrFeNi was evaluated by calculation of the pseudo-binary phase diagram. Experimental results validated the presence of eutectic reaction in the calculated phase diagrams; however, small changes in the calculated phase diagrams were proposed. It has been shown that calculated pseudo-binary phase diagrams can provide a very good starting point for the development of eutectic HEAs and help in exponentially reducing the amount of experimental effort that may be required otherwise. Eutectic mixture consisting of FCC (A2) phase and intermetallic phases (? and ?) was successfully obtained by the addition of Mo to the CoCrFeNi system. The development of the eutectic microstructure showed a profound effect on the mechanical properties. Hardness of the samples increased from 150 HV for CoCrFeNiMo0.1 to 425.5 HV for CoCrFeNiMo1.0, whereas yield strength increased from around 218 MPa for CoCrFeNiMo0.1 to around 1100 MPa for CoCrFeNiMo1.0. Full article
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>Pseudo-binary phase diagram of CoCrFeNi-Mo system.</p> Full article ">Figure 2
<p>XRD patterns of developed HEAs: (<b>a</b>) CoCrFeNi-Mo<sub>0.1</sub>, (<b>b</b>) CoCrFeNi-Mo<sub>0.25</sub>, (<b>c</b>) CoCrFeNi-Mo<sub>0.5</sub>, (<b>d</b>) CoCrFeNi-Mo<sub>1.0</sub>.</p> Full article ">Figure 3
<p>Scanning electron micrographs of developed HEAs: (<b>a</b>) CoCrFeNi-Mo<sub>0.1</sub>, (<b>b</b>) CoCrFeNi-Mo<sub>0.25</sub>, (<b>c</b>) CoCrFeNi-Mo<sub>0.5</sub>, (<b>d</b>) CoCrFeNi-Mo<sub>1.0</sub>, (<b>e</b>) CoCrFeNi-Mo<sub>1.0</sub> (higher magnification).</p> Full article ">Figure 3 Cont.
<p>Scanning electron micrographs of developed HEAs: (<b>a</b>) CoCrFeNi-Mo<sub>0.1</sub>, (<b>b</b>) CoCrFeNi-Mo<sub>0.25</sub>, (<b>c</b>) CoCrFeNi-Mo<sub>0.5</sub>, (<b>d</b>) CoCrFeNi-Mo<sub>1.0</sub>, (<b>e</b>) CoCrFeNi-Mo<sub>1.0</sub> (higher magnification).</p> Full article ">Figure 4
<p>Modified pseudo-binary phase diagram of CoCrFeNi-Mo system.</p> Full article ">Figure 5
<p>(<b>a</b>) Variation of hardness as a function of amount of Mo in the CoCrFeNiMo<sub>x</sub> HEA; (<b>b</b>) compression testing results for the studied CoCrFeNiMo<sub>x</sub> alloys; (<b>c</b>) variation of yield strength as a function of amount of Mo in the CoCrFeNiMo<sub>x</sub> HEA.</p> Full article ">
15 pages, 9451 KiB  
Article
Study on the Relationship between High Temperature Mechanical Properties and Precipitates Evolution of 7085 Al Alloy after Long Time Thermal Exposures
by Jinxin Zang, Pan Dai, Yanqing Yang, Shuai Liu, Bin Huang, Jigang Ru and Xian Luo
Metals 2021, 11(9), 1483; https://doi.org/10.3390/met11091483 - 18 Sep 2021
Cited by 6 | Viewed by 2783
Abstract
The requirement for 7085 Al alloy as large airframe parts has been increasing due to its low quenching sensitivity and high strength. However, the relationship between high temperature mechanical properties and the evolution of precipitates in hot environments is still unclear. In this [...] Read more.
The requirement for 7085 Al alloy as large airframe parts has been increasing due to its low quenching sensitivity and high strength. However, the relationship between high temperature mechanical properties and the evolution of precipitates in hot environments is still unclear. In this work, thermal exposure followed by tensile tests were conducted on the 7085 Al alloy at various temperatures (100 °C, 125 °C, 150 °C and 175 °C). Variations of hardness, electrical conductivity and tensile properties were investigated. The evolution of the nano scale precipitates was also quantitatively characterized by transmission electron microscopy (TEM). The results show that the hardness and electrical conductivity of the alloy are more sensitive to the temperature than to the time. The strength decreases continuously with the increase of temperature due to the transformation from ?? to ? phase during the process. Furthermore, the main ? phase in the alloy transformed from V3 and V4 to V1 and V2 variants when the temperature was 125 °C. Additionally, with increasing the temperature, the average precipitate radius increased, meanwhile the volume fraction and number density of the precipitates decreased. The strengthening effect of nano scale precipitates on tensile properties of the alloy was calculated and analyzed. Full article
(This article belongs to the Special Issue Research on Microstructure and Mechanical Properties of Metallic Materials and Metal-Matrix Composites)
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<p>Effects of thermal exposure and tensile test temperature on hardness (<b>a</b>) and electrical conductivity (<b>b</b>) of 7085-T74 Al alloy.</p> Full article ">Figure 2
<p>Effects of thermal exposure and tensile test time on hardness (<b>a</b>) and electrical conductivity (<b>b</b>) of 7085-T74 Al alloy.</p> Full article ">Figure 3
<p>(<b>a</b>) Mechanical properties of the 7085 Al alloy after heat treatment at different temperatures, (<b>b</b>) tensile strength, (<b>c</b>) elongation and reduction of area.</p> Full article ">Figure 4
<p>Bright-field TEM images of 7085 Al alloy plate after different heat treatment processes: (<b>a</b>) sample I (25 °C), (<b>b</b>) sample II (100 °C), (<b>c</b>) sample III (125 °C), (<b>d</b>) sample IV (150 °C), and (<b>e</b>) sample V (175 °C).</p> Full article ">Figure 5
<p>(<b>a</b>) Schematic representation of the (001)<sub>Al</sub> pattern and SAED patterns along (001)<sub>Al</sub> zone axis of the 7085 alloy after various heat treatments: (<b>b</b>) sample I (25 °C), (<b>c</b>) sample II (100 °C), (<b>d</b>) sample III (125 °C), (<b>e</b>) sample IV (150 °C), and (<b>f</b>) sample V (175 °C).</p> Full article ">Figure 6
<p>HRTEM images of the precipitates viewed along &lt;110&gt;<sub>Al</sub> orientation in 7085 alloy after various heat treatments: (<b>a</b>) sample I (25 °C), (<b>b</b>) sample II (100 °C), (<b>c</b>) sample III (125 °C), (<b>d</b>) sample IV (150 °C), (<b>e</b>) sample V (175 °C), (<b>f</b>) corresponding FFT pattern of η phase in (<b>e</b>), (<b>g</b>) is from the yellow dashed frame in (<b>i</b>), (<b>h</b>) is the corresponding inverse FFT image of (<b>e</b>), and (<b>i</b>) is from the white solid frame in (<b>h</b>).</p> Full article ">Figure 7
<p>Schematic illustrations showing four types of η variants <span class="html-italic">V</span>1-4 on {111}<sub>Al</sub>. (<b>a</b>) The three-dimensional view, and (<b>b</b>) the projection from [110]<sub>Al</sub> direction.</p> Full article ">Figure 8
<p>Precipitate radius distributions for the 7085 alloy after various heat treatments: (<b>a</b>) sample I (25 °C), (<b>b</b>) sample II (100 °C), (<b>c</b>) sample III (125 °C), (<b>d</b>) sample IV (150 °C), and (<b>e</b>) sample V (175 °C).</p> Full article ">Figure 9
<p>Evolutions of the precipitate parameters of the samples after various heat treatments: (<b>a</b>) precipitate volume fraction, and (<b>b</b>) precipitate number density.</p> Full article ">Figure 10
<p>Distributions of the thickness versus radius of the precipitates in the 7085 alloy after various heat treatments: (<b>a</b>) sample I (25 °C), (<b>b</b>) sample II (100 °C), (<b>c</b>) sample III (125 °C), (<b>d</b>) sample IV (150 °C), and (<b>e</b>) sample V (175 °C). <span class="html-italic">The orange points</span> correspond to the average thickness and radius of the precipitates.</p> Full article ">Figure 11
<p>Comparison of the average aspect ratio–temperature curve with the hardness-temperature curve.</p> Full article ">Figure 12
<p>(<b>a</b>) HRTEM image of η′ precipitate in sample II (100 °C), (<b>b</b>,<b>c</b>) corresponding inverse FFT patterns of (<b>a</b>), (<b>d</b>) HRTEM image of η precipitate in sample II (100 °C), (<b>e</b>,<b>f</b>) corresponding inverse FFT patterns of (<b>d</b>). Dislocations with different Burgers vectors are labeled by different colored ‘T’ symbols.</p> Full article ">Figure 13
<p>Increment in yield strength of the samples after thermal exposure and tensile tests at different temperatures.</p> Full article ">
17 pages, 10039 KiB  
Article
Toward a Simplified Arc Impingement Model in a Direct-Current Electric Arc Furnace
by Mohamad Al-Nasser, Abdellah Kharicha, Hadi Barati, Christoph Pichler, Gernot Hackl, Markus Gruber, Anton Ishmurzin, Christian Redl, Menghuai Wu and Andreas Ludwig
Metals 2021, 11(9), 1482; https://doi.org/10.3390/met11091482 - 17 Sep 2021
Cited by 7 | Viewed by 3396
Abstract
A 2D axisymmetric two-phase model was developed to study the effect of an arc impingement on the liquid metal inside an electric arc furnace. In addition to the arc flow dynamics, the model covered the heat transfer and magneto hydrodynamics of the arc [...] Read more.
A 2D axisymmetric two-phase model was developed to study the effect of an arc impingement on the liquid metal inside an electric arc furnace. In addition to the arc flow dynamics, the model covered the heat transfer and magneto hydrodynamics of the arc and the liquid metal. Through a parametric study, three different parameters were considered to predict the most important factors affecting the arc and overall behaviour of the process: the arc gap, the density of the gas, and the total electric current. Understanding the effect of these parameters can show the key factors altering the arc dynamics. The study showed that the total applied current was the most important parameter that influenced the impingement depth and mixing of the liquid metal. The depth of the impingement and strength of the mixing of the liquid bath were directly proportional to the current applied in the furnace. The initial arc gap distance was found to be crucial for sustaining a continuous and stable arc. The value of the gas density was very important for the velocity profile; however, it had no significant effect on the impingement depth. This showed that a constant density could be used instead of a varying gas density with temperature to increase the computational efficiency. The study assessed the effects of the aforementioned factors on the arc impingement depth, velocity magnitude, and arc stability. The conclusions acquired and challenges are also presented. Full article
(This article belongs to the Special Issue Modeling and Simulation of Metallurgical Processes in Ironmaking and Steelmaking)
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<p>Electric arc furnace schematic.</p> Full article ">Figure 2
<p>Geometry of the domain: <span class="html-italic">R</span> = 0.8, <span class="html-italic">H</span> = 1 m. The electrode length was 0.04 and the radius was 0.008; the arc gap (<span class="html-italic">G</span>) was 0.25, and the metal bath depth was initially 0.71 m.</p> Full article ">Figure 3
<p>Electric arc during the impingement of the liquid metal: (1) actual arc length; (2) arc gap.</p> Full article ">Figure 4
<p>Velocity magnitude as a function of the mesh size.</p> Full article ">Figure 5
<p>Impingement depth as a function of the mesh size.</p> Full article ">Figure 6
<p>Simulation domain. The yellow frame shows the zoomed domain for a better visual depiction.</p> Full article ">Figure 7
<p>Arc impingement arc gap of 2.5 × 10<sup>−1</sup> m.</p> Full article ">Figure 8
<p>Liquid metal impingement by an arc jet for a 2.0 × 10<sup>−1</sup> m arc gap.</p> Full article ">Figure 9
<p>Liquid metal impingement by an arc jet for a 3.0 × 10<sup>−1</sup> m gap distance.</p> Full article ">Figure 10
<p>Liquid metal impingement by an arc jet for a gas density of 1.0 × 10<sup>−1</sup> kg/m<sup>3</sup>.</p> Full article ">Figure 11
<p>Liquid metal impingement by an arc jet for a gas density of 1 × 10<sup>−2</sup> kg/m<sup>3</sup>.</p> Full article ">Figure 12
<p>The zoomed domain is enlarged (yellow frame) due to a deeper arc impingement.</p> Full article ">Figure 13
<p>Liquid metal impingement by an arc jet for a total current of 2.0 × 10<sup>1</sup> kA.</p> Full article ">Figure 14
<p>Liquid metal impingement by an arc jet for a total current of 3.0 × 10<sup>1</sup> kA.</p> Full article ">
18 pages, 4775 KiB  
Article
Tube Expansion by Single Point Incremental Forming: An Experimental and Numerical Investigation
by Carlos Suntaxi, Gabriel Centeno, M. Beatriz Silva, Carpóforo Vallellano and Paulo A. F. Martins
Metals 2021, 11(9), 1481; https://doi.org/10.3390/met11091481 - 17 Sep 2021
Cited by 4 | Viewed by 2429
Abstract
In this paper, we revisit the formability of tube expansion by single point incremental forming to account for the material strain hardening and the non-proportional loading paths that were not taken into consideration in a previously published analytical model of the process built [...] Read more.
In this paper, we revisit the formability of tube expansion by single point incremental forming to account for the material strain hardening and the non-proportional loading paths that were not taken into consideration in a previously published analytical model of the process built upon a rigid perfectly plastic material. The objective is to provide a new insight on the reason why the critical strains at failure of tube expansion by single point incremental forming are far superior to those of conventional tube expansion by rigid tapered conical punches. For this purpose, we replaced the stress triaxiality ratio that is responsible for the accumulation of damage and cracking by tension in monotonic, proportional loading paths, by integral forms of the stress triaxiality ratio that are more adequate for the non-proportional paths resulting from the loading and unloading cycles of incremental tube expansion. Experimental and numerical simulation results plotted in the effective strain vs. stress triaxiality space confirm the validity of the new damage accumulation approach for handling the non-proportional loading paths that oscillate cyclically from shearing to biaxial stretching, as the single point hemispherical tool approaches, contacts and moves away from a specific location of the incrementally expanded tube surface. Full article
(This article belongs to the Special Issue Tube and Sheet Metal Forming Processes and Applications)
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<p>Flow curve of the aluminum AA6063T6 tubes (adapted from [<a href="#B9-metals-11-01481" class="html-bibr">9</a>]).</p> Full article ">Figure 2
<p>Forming limit curve (FLC) and fracture forming limit (FFL) line of the aluminum AA6063-T6 tube in principal strain space. The red line represents the experimental strain loading path of conventional tube expansion with a rigid tapered conical punch having a semi-angle of 15° (adapted from [<a href="#B12-metals-11-01481" class="html-bibr">12</a>]).</p> Full article ">Figure 3
<p>Forming limit curve (FLC) and fracture forming limit (FFL) line of the aluminum AA6063-T6 tube in the effective strain vs. stress triaxiality space, obtained from analytical transformation assuming material isotropy, linear strain paths and plane stress loading conditions.</p> Full article ">Figure 4
<p>Schematic representation of incremental tube expansion showing (<b>a</b>) the single point hemispherical tool path and (<b>b</b>) the multi-stage forming strategy. A photograph of a specimen after eight forming stages is included in (<b>c</b>) together with a detail showing the final cracked surface.</p> Full article ">Figure 5
<p>(<b>a</b>) Experimental determination of the in-plane strains by CGA using the automatic measuring system ARGUS<sup>®</sup>, and (<b>b</b>) representation of these strains in principal strain space.</p> Full article ">Figure 6
<p>Finite element modeling of conventional tube expansion with a rigid tapered conical punch having a semi-angle of 15°. (<b>a</b>) Initial mesh with a detail of mesh refinement and (<b>b</b>) predicted distribution of effective strain at the end of the process.</p> Full article ">Figure 7
<p>Finite element modeling of incremental tube expansion. (<b>a</b>) Initial mesh and (<b>b</b>) final mesh after eight forming stages.</p> Full article ">Figure 8
<p>Experimental vs. finite element predicted in-plane strains for conventional tube expansion with a rigid tapered conical punch (red markers) and incremental tube expansion with a single point hemispherical tool after eight forming stages (black markers).</p> Full article ">Figure 9
<p>(<b>a</b>) Detail of a tube section after the incremental expansion and (<b>b</b>) evolution of the tube wall thickness with the longitudinal distance to the upper tube end.</p> Full article ">Figure 10
<p>Finite element predicted in-plane strains of point A during the eight forming stages of incremental tube expansion.</p> Full article ">Figure 11
<p>(<b>a</b>) Linear, proportional strain path in principal strain space, (<b>b</b>) non-proportional strain path discretized through a series of piecewise linear strain paths in principal strain space, (<b>c</b>) representation of the strain paths (<b>a</b>,<b>b</b>) in the effective strain vs. stress triaxiality space.</p> Full article ">Figure 12
<p>Schematic representation of the non-proportional, cyclic path of incremental tube expansion experienced by an arbitrary tube location with a plot of the <math display="inline"> <semantics> <mrow> <mover accent="true"> <mi>ε</mi> <mo>¯</mo> </mover> <mo>=</mo> <mi>f</mi> <mrow> <mo>(</mo> <mi>η</mi> <mo>)</mo> </mrow> </mrow> </semantics> </math> evolutions based on the three different integral forms of stress triaxiality: (<b>a</b>) envelope stress triaxiality <math display="inline"> <semantics> <mrow> <msub> <mi>η</mi> <mrow> <mi>e</mi> <mi>n</mi> <mi>v</mi> </mrow> </msub> </mrow> </semantics> </math>, (<b>b</b>) average positive stress triaxiality <math display="inline"> <semantics> <mrow> <msub> <mover accent="true"> <mi>η</mi> <mo>¯</mo> </mover> <mrow> <mi>p</mi> <mi>o</mi> <mi>s</mi> </mrow> </msub> </mrow> </semantics> </math> and (<b>c</b>) average stress triaxiality <math display="inline"> <semantics> <mover accent="true"> <mi>η</mi> <mo>¯</mo> </mover> </semantics> </math>.</p> Full article ">Figure 13
<p>Finite element computed non-proportional, cyclic path of point A (<a href="#metals-11-01481-f009" class="html-fig">Figure 9</a>) with several <math display="inline"> <semantics> <mrow> <mover accent="true"> <mi>ε</mi> <mo>¯</mo> </mover> <mo>=</mo> <mi>f</mi> <mrow> <mo>(</mo> <mi>η</mi> <mo>)</mo> </mrow> </mrow> </semantics> </math> evolutions obtained from different assumptions and integral forms of stress-triaxiality.</p> Full article ">Figure 14
<p>Finite element computed evolutions of the loading paths experienced by a point A located 1.5 mm away from the upper tube end in incremental and conventional tube expansion processes. Note: the red and blue triangular markers correspond to the experimentally determined “gauge length” strains at fracture.</p> Full article ">
14 pages, 2957 KiB  
Article
Effect of the Basicity on Mineralogical Phases and Micro-Structure of Dephosphorization Slag in the New Double Slag Converter Steelmaking Process
by Wenkui Yang, Jian Yang, Runhao Zhang, Han Sun and Yunlong Qiu
Metals 2021, 11(9), 1480; https://doi.org/10.3390/met11091480 - 17 Sep 2021
Cited by 6 | Viewed by 2150
Abstract
In the present work, the effect of the basicity at the lower range from 0.98 to 2.13 on dephosphorization of hot metal at 1623 K was studied through high-temperature laboratorial experiments. With the increase of the basicity from 0.98 to 2.13, the P [...] Read more.
In the present work, the effect of the basicity at the lower range from 0.98 to 2.13 on dephosphorization of hot metal at 1623 K was studied through high-temperature laboratorial experiments. With the increase of the basicity from 0.98 to 2.13, the P and C contents in hot metal rapidly decrease and increase at first, and then gradually decrease and increase, respectively. From the scanning electron microscopy-energy dispersive X-ray spectroscopy (SEM-EDS) and X-ray diffraction (XRD) results, with the increase of the basicity, the phase containing the high P content changes from the matrix phase into the phosphorus (P)-rich phase. Under the present experimental conditions, the P-rich phase can only be precipitated from the liquid slag when the basicity is higher than 1.55, which is a benefit to the dephosphorization. As the Raman intensity of the P-O-Ca structure unit in the P-rich phase is significantly higher than that of the P-O-Si structure unit, most of the phosphorus in the P-rich phase exists in the P-O-Ca structure unit and a small amount of phosphorus exists in the P-O-Si structure unit. With the increase of the basicity of the dephosphorization slag, the activity coefficient of P2O5, ?(P2O5) , in the liquid phase decreases, while the basicity in the liquid phase increases. Full article
(This article belongs to the Special Issue Oxygen Steelmaking Process)
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<p>Effect of basicity on the contents of elements in the hot metal after dephosphorization.</p> Full article ">Figure 2
<p>Effect of basicity on the removal ratios of the elements in the hot metal after dephosphorization.</p> Full article ">Figure 3
<p>Effect of the basicity of the dephosphorization slag on the P<sub>2</sub>O<sub>5</sub>, MnO, T.Fe, and Al<sub>2</sub>O<sub>3</sub> contents in slag.</p> Full article ">Figure 4
<p>(<b>a</b>) Compositions of dephosphorization slags located in the CaO-SiO<sub>2</sub>-FeO ternary phase diagram; (<b>b</b>) compositions of the corresponding P-rich phases located in the CaO-SiO<sub>2</sub>-P<sub>2</sub>O<sub>5</sub> ternary phase diagram.</p> Full article ">Figure 5
<p>SEM images of dephosphorization slag with magnification of 200 times at different basicities.</p> Full article ">Figure 6
<p>SEM images of the dephosphorization slag with magnification of 1000 times at different basicities.</p> Full article ">Figure 7
<p>XRD patterns of the dephosphorization slag with different basicity.</p> Full article ">Figure 8
<p>Raman spectra analysis of Fe-rich phase in R155 slag, matrix phase C in R213 slag, and P-rich phases in R155 slag and R213 slag.</p> Full article ">Figure 9
<p>Effect of the basicity of the dephosphorization slag on the <math display="inline"><semantics> <mrow> <msub> <mi>γ</mi> <mrow> <msub> <mrow> <mrow> <mo>(</mo> <mi mathvariant="normal">P</mi> </mrow> </mrow> <mn>2</mn> </msub> <msub> <mi mathvariant="normal">O</mi> <mn>5</mn> </msub> <mo>)</mo> </mrow> </msub> </mrow> </semantics></math> and basicity in the liquid phase.</p> Full article ">
19 pages, 4763 KiB  
Article
Prediction of Cracking Susceptibility of Commercial Aluminum Alloys during Solidification
by Fan Zhang, Songmao Liang, Chuan Zhang, Shuanglin Chen, Duchao Lv, Weisheng Cao and Sindo Kou
Metals 2021, 11(9), 1479; https://doi.org/10.3390/met11091479 - 17 Sep 2021
Cited by 10 | Viewed by 4702
Abstract
Cracking during solidification is a complex phenomenon which has been investigated from various angles for decades using both experimental and theoretical methods. In this paper, cracking susceptibility was investigated by a simulation method for three series of aluminum alloys: AA2xxx, AA6xxx, and AA7xxx [...] Read more.
Cracking during solidification is a complex phenomenon which has been investigated from various angles for decades using both experimental and theoretical methods. In this paper, cracking susceptibility was investigated by a simulation method for three series of aluminum alloys: AA2xxx, AA6xxx, and AA7xxx alloys. The simulation tool was developed using the CALPHAD method and is readily applicable to multicomponent alloy systems. For each series of alloys, cracking susceptible index values were calculated for more than 1000 alloy compositions by high-throughput calculation. Cracking susceptible maps were then constructed for these three series of aluminum alloys using the simulated results. The effects of major and minor alloying elements were clearly demonstrated by these index maps. The cooling rate effect was also studied, and it was concluded that back diffusion in the solid can significantly improve the cracking susceptibility. Full article
(This article belongs to the Special Issue Calphad Tools for the Metallurgy of Solidification)
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<p>Al-Cu-Mg cracking susceptibility maps (<b>a</b>) by improved Scheil model considering back diffusion in the solid under the cooling rate of 20 °C/s and (<b>b</b>) experimentally tested data [<a href="#B54-metals-11-01479" class="html-bibr">54</a>] (Reprinted with permission from ref. [<a href="#B24-metals-11-01479" class="html-bibr">24</a>]. Copyright 2017 Elsevier).</p> Full article ">Figure 2
<p>Calculated cracking susceptibility index (CSI) maps for AA2xxx alloys: (<b>a</b>) Al-0.6Fe-0.5Mn<b>-0.2Si</b>-xCu-yMg alloys; (<b>b</b>) Al-0.6Fe-0.5Mn<b>-0.6Si</b>-xCu-yMg alloys; (<b>c</b>) Al-0.6Fe-0.5Mn<b>-1.2Si</b>-xCu-yMg alloys; (<b>d</b>) Al-1.2Si-xCu-yMg alloys. Simulation is performed by improved Scheil model considering back diffusion in the solid under the cooling rate of 20 °C/s.</p> Full article ">Figure 3
<p>Calculated <span class="html-italic">CSI</span> values in the binary systems: (<b>a</b>) Al-Si alloys and (<b>b</b>) Al-Mg alloys.</p> Full article ">Figure 4
<p>Al-Mg-Si cracking susceptibility maps: (<b>a</b>) by improved Scheil model considering back diffusion in the solid under the cooling rate of 20 °C/s and (<b>b</b>) experimentally tested data [<a href="#B55-metals-11-01479" class="html-bibr">55</a>] (Reprinted with permission from ref. [<a href="#B24-metals-11-01479" class="html-bibr">24</a>]. Copyright 2017 Elsevier).</p> Full article ">Figure 5
<p>Calculated cracking susceptibility index (<span class="html-italic">CSI</span>) maps for AA6xxx alloys: (<b>a</b>) Al-0.6Fe-0.5Mn-0.2Zn-<b>0.2Cu</b>-xMg-ySi; (<b>b</b>) Al-0.6Fe-0.5Mn-0.2Zn-<b>0.5Cu</b>-xMg-ySi; (<b>c</b>) Al-0.6Fe-0.5Mn-0.2Zn-<b>1.2Cu</b>-xMg-ySi; (<b>d</b>) Al-0.2Zn-1.2Cu-xSi-yMg; (<b>e</b>) Al-1.2Cu-xSi-yMg. Simulation is performed by improved Scheil model considering back diffusion in the solid under the cooling rate of 20 °C/s.</p> Full article ">Figure 6
<p>Al-Mg-Zn cracking susceptibility maps: (<b>a</b>) by improved Scheil model with the cooling rate of 20 °C/s and (<b>b</b>) experimentally tested data [<a href="#B56-metals-11-01479" class="html-bibr">56</a>] (Reprinted with permission from ref. [<a href="#B4-metals-11-01479" class="html-bibr">4</a>]. Copyright 2004 Elsevier).</p> Full article ">Figure 7
<p>Al-Mg-Zn-0.5Cu cracking susceptibility maps: (<b>a</b>) by improved Scheil model considering back diffusion in the solid under the cooling rate of 20 °C/s and (<b>b</b>) experimentally tested data [<a href="#B56-metals-11-01479" class="html-bibr">56</a>] (Reprinted with permission from ref. [<a href="#B4-metals-11-01479" class="html-bibr">4</a>]. Copyright 2004 Elsevier).</p> Full article ">Figure 8
<p>Calculated cracking susceptibility index (<span class="html-italic">CSI</span>) maps for AA7xxx alloys: (<b>a</b>) Al-0.2Si-0.3Fe-0.2Mn-<b>0.5Cu</b>-xZn-yMg; (<b>b</b>) Al-0.2Si-0.3Fe-0.2Mn-<b>1.5Cu</b>-xZn-yMg; (<b>c</b>) Al-0.2Si-0.3Fe-0.2Mn-<b>2.2Cu</b>-xZn-yMg. Simulation is performed by improved Scheil model considering back diffusion in the solid under the cooling rate of 20 °C/s.</p> Full article ">Figure 9
<p>Al-Mg-Zn crack susceptibility maps: (<b>a</b>) by Scheil model, (<b>b</b>) by improved Scheil model with the cooling rate of 100 °C/s, (<b>c</b>) by improved Scheil model with the cooling rate of 10 °C/s, and (<b>d</b>) by lever rule.</p> Full article ">Figure 10
<p>Solidification paths <math display="inline"><semantics> <mrow> <mi>T</mi> <mo> </mo> <mi>v</mi> <mi>s</mi> <mo> </mo> <msup> <mrow> <mrow> <mo>(</mo> <mrow> <msub> <mi>f</mi> <mi>s</mi> </msub> </mrow> <mo>)</mo> </mrow> </mrow> <mrow> <mn>1</mn> <mo>/</mo> <mn>2</mn> </mrow> </msup> </mrow> </semantics></math> by Scheil model for (<b>a</b>) Al-Mg-Si alloys along <math display="inline"><semantics> <mrow> <mi>w</mi> <mrow> <mo>(</mo> <mrow> <mi>Mg</mi> </mrow> <mo>)</mo> </mrow> <mo>/</mo> <mi>w</mi> <mrow> <mo>(</mo> <mrow> <mi>Si</mi> </mrow> <mo>)</mo> </mrow> <mo>=</mo> <mn>1.73</mn> </mrow> </semantics></math> and (<b>b</b>) Al-Mg-Zn alloys along <math display="inline"><semantics> <mrow> <mi>w</mi> <mrow> <mo>(</mo> <mrow> <mi>Zn</mi> </mrow> <mo>)</mo> </mrow> <mo>/</mo> <mi>w</mi> <mrow> <mo>(</mo> <mrow> <mi>Mg</mi> </mrow> <mo>)</mo> </mrow> <mo>=</mo> <mn>5.4</mn> </mrow> </semantics></math>.</p> Full article ">Figure 11
<p>Calculated <span class="html-italic">CSI</span> for a series of Al-4Mg-xCu-<span class="html-italic">n</span>Ti (<span class="html-italic">n</span> = 0 and 0.2 wt%) alloys.</p> Full article ">
15 pages, 6286 KiB  
Article
Excellent Wear Resistance of a High-Speed Train Brake Disc Steel with High Hardening Ratcheting Strain Zone
by Lei Yang, Tingwei Zhou, Zhenlin Xu, Yizhu He, Xuewen Hu and Hai Zhao
Metals 2021, 11(9), 1478; https://doi.org/10.3390/met11091478 - 17 Sep 2021
Cited by 3 | Viewed by 2450
Abstract
Wear resistance of brake discs has a significant effect on the safety of high-speed trains. In this work, the rolling–sliding wear resistance of a low-carbon martensitic brake disc steel was investigated. The microstructure evolution and mechanical properties from the worn surface to the [...] Read more.
Wear resistance of brake discs has a significant effect on the safety of high-speed trains. In this work, the rolling–sliding wear resistance of a low-carbon martensitic brake disc steel was investigated. The microstructure evolution and mechanical properties from the worn surface to the matrix were analyzed to promote parameters for the application of this steel. The results indicated that a ratcheting strain zone was formed at the surface of the brake disc steel under asymmetric cycling load and presented the morphology of the plastic flow line. An equation between the shear strain and the depth from the surface was established, which reflected the gradient distribution of the shear strain. The martensite lath refined into the nano/sub-micron grain and strip on the topmost surface. The micron scratch test results exhibited that the hardening rate and fracture toughness of the ratcheting strain zone varied continually along with depth, and the highest hardening rate occurred on the worn surface due to the increase of dislocation density and grain refinement. Additionally, the brake disc steel obtained better wear resistance than that of other wear-resistant materials used for railways due to steel having the highest hardening rate in the ratcheting strain zone. Full article
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<p>Schematic diagram of the production process of the test steel.</p> Full article ">Figure 2
<p>Schematic diagram of the rolling–sliding wear tester.</p> Full article ">Figure 3
<p>SEM image of the wear surface of test steel: (<b>a</b>) debris and (<b>b</b>) peeling pits and cracks.</p> Full article ">Figure 4
<p>(<b>a</b>) OM, (<b>b</b>) FESEM and (<b>c</b>) TEM image of the matrix microstructure of the test steel.</p> Full article ">Figure 5
<p>(<b>a</b>,<b>b</b>) Various forms of structural response to cyclic loading, (<b>c</b>) OM image of the longitudinal section of wear sample, and (<b>d</b>) schematic diagram of specimen cutting and the RS zone.</p> Full article ">Figure 6
<p>FESEM images of plastic flow lines in the RS zone.</p> Full article ">Figure 7
<p>Curves of (<b>a</b>) plastic flow line displacement field and (<b>b</b>) shear strain along with the depth.</p> Full article ">Figure 8
<p>FESEM image of (<b>a</b>) martensite lath characteristics, and (<b>b</b>) the topmost surface microstructure of the RS zone.</p> Full article ">Figure 9
<p>X-ray diffraction image of test steel: (<b>a</b>) RS zone and matrix; (<b>b</b>) (110) diffraction peak; (<b>c</b>) (200) diffraction peak.</p> Full article ">Figure 10
<p>((Δ<span class="html-italic">K</span>)<sup>2</sup> − <span class="html-italic">α</span>)/<span class="html-italic">K</span><sup>2</sup> vs. <span class="html-italic">H</span><sup>2</sup> plot for martensite peaks (<b>a</b>) matrix and (<b>b</b>)RS zone assuming Gaussian and Lorentzian peak broadening, respectively.</p> Full article ">Figure 11
<p>(<b>a</b>) Scratch morphology and (<b>b</b>) cross-section of positions A and B.</p> Full article ">Figure 12
<p>Curves of (<b>a</b>) scratch width and (<b>b</b>) scratch depth with depth from the surface.</p> Full article ">Figure 13
<p>Hardening rate and fracture toughness along with <span class="html-italic">D<sub>S</sub></span> distribution: (<b>a</b>) hardening rate and (<b>b</b>) fracture toughness.</p> Full article ">Figure 14
<p>(<b>a</b>) Wear resistance and (<b>b</b>) hardening rate of materials in railway reported.</p> Full article ">
13 pages, 3387 KiB  
Article
Microstructure and Mechanical Properties of the ((CoCrFeNi)95Nb5)100?xMox High-Entropy Alloy Coating Fabricated under Different Laser Power
by Wenrui Wang, Qi Sun, Dingzhi Wang, Junsong Hou, Wu Qi, Dongyue Li and Lu Xie
Metals 2021, 11(9), 1477; https://doi.org/10.3390/met11091477 - 17 Sep 2021
Cited by 7 | Viewed by 2365
Abstract
In this paper, the ((CoCrFeNi)95Nb5)100?xMox (x = 1, 1.5 and 2) high-entropy alloy (HEA) coatings were fabricated on the substrate of 45# steel by laser cladding process under different laser beam power. The influence [...] Read more.
In this paper, the ((CoCrFeNi)95Nb5)100?xMox (x = 1, 1.5 and 2) high-entropy alloy (HEA) coatings were fabricated on the substrate of 45# steel by laser cladding process under different laser beam power. The influence of laser beam power and molybdenum element content on the microstructure and microhardness of the HEA coatings was investigated. Results show that the HEA coatings were composed of face-centered cubic (FCC) phase and Laves phase, had low porosity, and bonded well to the substrate. The Mo1 coating is composed of cellular dendritic structures and columnar dendritic structures. With the increase of molybdenum element content, the columnar dendritic structures disappeared, the grains are refined, and the arrangement of grains is more compact. The volume fraction of the interdendritic phase under the laser beam power of 800 W was small and irregular. After the laser beam power was increased to 1000 W, the volume fraction of the interdendritic phase was increased. Under the laser beam power of 1200 W, the volume fraction of the interdendritic phase was small again. Therefore, the coatings fabricated under the laser beam power of 1000 W had a larger volume fraction of the interdendritic phase and higher microhardness. With the increase in molybdenum content, the grain changed from columnar dendrite to cellular dendrite, and the microhardness of the coating increased. The characteristics of the laser cladding process, the formation of Laves phase, and the fine grain strengthening lead to high microhardness of the coatings. Full article
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<p>Synchronous laser cladding equipment.</p> Full article ">Figure 2
<p>Solid work model of specimen with coating.</p> Full article ">Figure 3
<p>Schematic diagram of the laser cladding process.</p> Full article ">Figure 4
<p>(<b>a</b>) XRD patterns of the ((CoCrFeNi)<sub>95</sub>Nb<sub>5</sub>)<sub>100−<span class="html-italic">x</span></sub>Mo<sub><span class="html-italic">x</span></sub> (<span class="html-italic">x</span> = 1, 1.5, 2) HEA powder; (<b>b</b>–<b>d</b>) SEM images of the ((CoCrFeNi)<sub>95</sub>Nb<sub>5</sub>)<sub>100−<span class="html-italic">x</span></sub>Mo<sub><span class="html-italic">x</span></sub> (<span class="html-italic">x</span> = 1, 1.5, 2) HEA powder, where (<b>b</b>) <span class="html-italic">x</span> = 1, (<b>c</b>) <span class="html-italic">x</span> = 1.5, (<b>d</b>) <span class="html-italic">x</span> = 2.</p> Full article ">Figure 5
<p>(<b>a</b>) XRD patterns of the ((CoCrFeNi)<sub>95</sub>Nb<sub>5</sub>)<sub>100−<span class="html-italic">x</span></sub>Mo<sub><span class="html-italic">x</span></sub> (<span class="html-italic">x</span> = 1, 1.5, 2) HEA coatings, (<b>b</b>–<b>d</b>) SEM images of the ((CoCrFeNi)<sub>95</sub>Nb<sub>5</sub>)<sub>100−<span class="html-italic">x</span></sub>Mo<sub><span class="html-italic">x</span></sub> (<span class="html-italic">x</span> = 1, 1.5, 2) HEA coatings.</p> Full article ">Figure 6
<p>(<b>a</b>) XRD patterns of the ((CoCrFeNi)<sub>95</sub>Nb<sub>5</sub>)<sub>99</sub>Mo<sub>1</sub> HEA coatings prepared under different laser beam power, (<b>b</b>–<b>d</b>) SEM images of the((CoCrFeNi)<sub>95</sub>Nb<sub>5</sub>)<sub>99</sub>Mo<sub>1</sub> HEA coatings prepared under different laser beam power.</p> Full article ">Figure 7
<p>(<b>a</b>) SEM image of longitudinal section of the ((CoCrFeNi)<sub>95</sub>Nb<sub>5</sub>)<sub>99</sub>Mo<sub>1</sub> coating prepared under the laser beam power of 800 W, (<b>b</b>) distributions of alloying elements by EDS line scanning.</p> Full article ">Figure 7 Cont.
<p>(<b>a</b>) SEM image of longitudinal section of the ((CoCrFeNi)<sub>95</sub>Nb<sub>5</sub>)<sub>99</sub>Mo<sub>1</sub> coating prepared under the laser beam power of 800 W, (<b>b</b>) distributions of alloying elements by EDS line scanning.</p> Full article ">Figure 8
<p>Microhardness of the ((CoCrFeNi)<sub>95</sub>Nb<sub>5</sub>)<sub>100−<span class="html-italic">x</span></sub>Mo<sub><span class="html-italic">x</span></sub> (<span class="html-italic">x</span> = 1, 1.5, 2) HEA coating.</p> Full article ">
16 pages, 6749 KiB  
Article
A Comparison of Laboratory Coal Testing with the Blast Furnace Process and Coal Injection
by Julian Steer, Mark Greenslade and Richard Marsh
Metals 2021, 11(9), 1476; https://doi.org/10.3390/met11091476 - 17 Sep 2021
Cited by 9 | Viewed by 2574 | Correction
Abstract
The injection of coal through tuyeres into a blast furnace is widely adopted throughout the industry to reduce the amount of coke used and to improve the efficiency of the iron making process. Coals are selected depending on their availability, cost, and the [...] Read more.
The injection of coal through tuyeres into a blast furnace is widely adopted throughout the industry to reduce the amount of coke used and to improve the efficiency of the iron making process. Coals are selected depending on their availability, cost, and the physical and chemical properties determined by tests, such as the volatile matter content, fixed carbon, and ash content. This paper describes research comparing the laboratory measured properties of injection coals that were used over a two-month production period compared to the process variables and measurements of the blast furnace during that study period. In addition to the standard tests, a drop tube furnace (DTF) was used to compare the burnout of coals and the char properties against the production data using a range of statistical techniques. Linear regression modelling indicated that the coal type was the most important predictor of the coal rate but that the properties measured using laboratory tests of those coals were a minor feature in the model. However, comparisons of the Spearman’s correlations between different variables indicated that the reverse Boudouard reactivity of the chars, prepared in the DTF from the coals, did appear to be related to some extent to the coal and coke rates on production. It appears that the constant process adjustments made by the process control systems on the furnace make it difficult to identify strong correlations with the laboratory data and that the frequency of coal sampling and the coal blend variability are likely to contribute to this difficulty. Full article
(This article belongs to the Special Issue Advances in Ironmaking and Steelmaking Processes)
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<p>The range of the measured volatile matter contents for oven-dried coals and coal blends used during the study periods (◦ represents outliers).</p> Full article ">Figure 2
<p>The ranges of the measured ash contents for the oven-dried coals and coal blends used during the study periods.</p> Full article ">Figure 3
<p>The ranges of the measured ash contents for the oven-dried coals and coal blends used during the study periods (◦ represents outliers).</p> Full article ">Figure 4
<p>Schematic of drop tube furnace.</p> Full article ">Figure 5
<p>Scatter plot of the relationship between oxygen flow and the iron-production rate.</p> Full article ">Figure 6
<p>Scatter plot of the relationship between the gasification reactivity of the injectant coal char and the coke rate during blast furnace production.</p> Full article ">Figure 7
<p>Scatter plot of the relationship between the injectant coal and the coke rate during blast furnace production.</p> Full article ">Figure 8
<p>Scatter plot of the relationship between the gasification reactivity of the injectant coal and the coal injection rate during blast furnace production.</p> Full article ">Figure 9
<p>The range of the coal injection rates for each of the coals/coal blends used during the trial period of blast furnace production.</p> Full article ">Figure 10
<p>Scatter plot of the DTF burnouts of the injection coals versus their char gasification reactivities.</p> Full article ">Figure 11
<p>Scatter plot of the volatile matter content of the injectant coals versus their char gasification reactivities.</p> Full article ">Figure 12
<p>Box plots of the range of coal injection rates for each of the coals/coal blends used during blast furnace production (◦ represents outliers and * extreme outliers).</p> Full article ">Figure 13
<p>Box plots of the range of coke addition rates for the time periods corresponding to the addition of different coals/coal blends during blast furnace production (◦ = high/low potential outliers and * = high/low extreme values).</p> Full article ">Figure 14
<p>Box plots of the range of char gasification reactivity for each of the coals/coal blends used during blast furnace production (◦ = high/low potential outliers and * = high/low extreme values).</p> Full article ">Figure 15
<p>Range of coal burnout for each of the coals/coal blends used during observed blast furnace production (◦ = high/low potential outliers).</p> Full article ">Figure 16
<p>Relationship between the measured and predicted injection coal rate in the multiple linear regression model.</p> Full article ">Figure 16 Cont.
<p>Relationship between the measured and predicted injection coal rate in the multiple linear regression model.</p> Full article ">Figure 17
<p>Relationship between the measured and predicted coke rate in the multiple linear regression model.</p> Full article ">Figure 17 Cont.
<p>Relationship between the measured and predicted coke rate in the multiple linear regression model.</p> Full article ">
8 pages, 1289 KiB  
Article
Thermoelectric Power in Ce Systems with Unstable Valence
by Tomasz Toliński
Metals 2021, 11(9), 1475; https://doi.org/10.3390/met11091475 - 17 Sep 2021
Cited by 1 | Viewed by 1785
Abstract
In this paper, we report on a few exemplary tests of the applicability of analysis based on the interconfiguration fluctuation model (ICF) for a description of the temperature dependence of the thermoelectric power, S(T). The examples include a series of [...] Read more.
In this paper, we report on a few exemplary tests of the applicability of analysis based on the interconfiguration fluctuation model (ICF) for a description of the temperature dependence of the thermoelectric power, S(T). The examples include a series of alloys: CeNi2(Si1?yGey)2, Ce(Ni1?xCux)2Si2, and the fluctuating valence (FV) compound CeNi4Ga. The two series develop from CeNi2Si2 being the FV system, where the f states occupation increases progressively with the Ge or Cu substitution. We find here that the ICF model parameters are of similar magnitude both for the analysis of the temperature dependence of the magnetic susceptibility and thermoelectric power. The ICF-type model appears to be a powerful tool for the analysis of S(T) dependences in Ce-based FV compounds and alloys. Full article
(This article belongs to the Special Issue Synthesis, Characterization and Simulation of Novel Magnetic Intermetallic Compounds)
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<p>Magnetic susceptibility of CeNi<sub>2</sub>(Si<sub>1−y</sub>Ge<sub>y</sub>)<sub>2</sub>. The ICF model is applied for y = 0.13 and 0.63 (measured up to 1000 K), Equation (2).</p> Full article ">Figure 2
<p>Thermoelectric power as a function of temperature fitted with Equation (5). (<b>a</b>) y = 0.0, (<b>b</b>) y = 0.13, (<b>c</b>) y = 0.63, (<b>d</b>) y = 0.88. The obtained values of the parameters are gathered in <a href="#metals-11-01475-t001" class="html-table">Table 1</a>. <span class="html-italic">S</span>(<span class="html-italic">T</span>) measurements were described in Ref. [<a href="#B16-metals-11-01475" class="html-bibr">16</a>].</p> Full article ">Figure 3
<p>Occupancy of the <span class="html-italic">f</span> states as a function of temperature derived from the thermoelectric power <span class="html-italic">S</span>(<span class="html-italic">T</span>) for CeNi<sub>2</sub>(Si<sub>1−y</sub>Ge<sub>y</sub>)<sub>2</sub>.</p> Full article ">Figure 4
<p>Thermoelectric power of Ce(Ni<sub>1−x</sub>Cu<sub>x</sub>)<sub>2</sub>Si<sub>2</sub> as a function of temperature fitted with Equation (5). (<b>a</b>) x = 0.25 and (<b>b</b>) x = 0.88. See the text for the values of the parameters.</p> Full article ">Figure 5
<p>Occupancy of the <span class="html-italic">f</span> states as a function of temperature derived from the thermoelectric power <span class="html-italic">S</span>(<span class="html-italic">T</span>) for Ce(Ni<sub>1−x</sub>Cu<sub>x</sub>)<sub>2</sub>Si<sub>2</sub>.</p> Full article ">Figure 6
<p>Left axis: Thermoelectric power <span class="html-italic">S</span>(<span class="html-italic">T</span>) of CeNi<sub>4</sub>Ga fitted with Equation (5). See the text for the values of the parameters. Right axis: Occupancy of the <span class="html-italic">f</span> states as a function of temperature.</p> Full article ">
13 pages, 4916 KiB  
Article
Experimental Investigation of Material Properties in FSW Dissimilar Aluminum-Steel Lap Joints
by Michelangelo Mortello, Matteo Pedemonte, Nicola Contuzzi and Giuseppe Casalino
Metals 2021, 11(9), 1474; https://doi.org/10.3390/met11091474 - 16 Sep 2021
Cited by 13 | Viewed by 2310
Abstract
The friction stir lap welding of AA5083 H111 aluminum alloy and S355J2 grade DH36 structural steel was investigated. A polycrystalline cubic boron nitride with tungsten and rhenium additives tool was used. According to visual inspection, radiographic examination, and tensile test, it was observed [...] Read more.
The friction stir lap welding of AA5083 H111 aluminum alloy and S355J2 grade DH36 structural steel was investigated. A polycrystalline cubic boron nitride with tungsten and rhenium additives tool was used. According to visual inspection, radiographic examination, and tensile test, it was observed that the best results were obtained for rotation speeds of about 700–800 rpm, with a feed speed ranging between 1.3 and 1.9 mm/s. From the fatigue tests, it is possible to state that there was a preferential propagation of cracks in the part of the aluminum alloy base material. Furthermore, a different response to fatigue stress for samples extracted from the same weld at different positions was observed, which introduces an overall variability in weld behavior along the welding direction. The specimens obtained in the second part of the weld endured a larger number of cycles before reaching failure, which can be related to progressively varying thermal conditions, dissipation behavior, and better metal coupling as the tool travels along the welding line. Full article
(This article belongs to the Special Issue Advances in Processing and Mechanical Behavior in Lightweight Metals and Alloys)
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<p>Schematic setup of the process in cross sectional (<b>a</b>) and axonometry (<b>b</b>) view.</p> Full article ">Figure 2
<p>H. Loitz-Robotik Machine set up.</p> Full article ">Figure 3
<p>Megastir Q70 (PCBN/WRe) tool.</p> Full article ">Figure 4
<p>Sample 994.</p> Full article ">Figure 5
<p>Sample 995.</p> Full article ">Figure 6
<p>Sample 999.</p> Full article ">Figure 7
<p>Sample 987 macro- and micro-section.</p> Full article ">Figure 8
<p>X-ray test for samples 953 (<b>a</b>) and 996 (<b>b</b>).</p> Full article ">Figure 9
<p>X-ray test for sample 957.</p> Full article ">Figure 10
<p>Tensile and fatigue tests samples.</p> Full article ">Figure 11
<p>Shear force versus intermetallic thickness.</p> Full article ">Figure 12
<p>Fracture zone for welding 957 specimen (<b>a</b>) and welding 3 specimen (<b>b</b>).</p> Full article ">Figure 13
<p>Number of cycles at break for sample 957.</p> Full article ">Figure 14
<p>Number of cycles at break for sample 996.</p> Full article ">Figure 15
<p>Fracture in the HAZ on the aluminum side for welding 9 specimen C.</p> Full article ">
12 pages, 3974 KiB  
Article
Minimizing the Negative Effects of Coolant Channels on the Torsional and Torsional-Axial Stiffness of Drills
by Amir Parsian, Mahdi Eynian, Martin Magnevall and Tomas Beno
Metals 2021, 11(9), 1473; https://doi.org/10.3390/met11091473 - 16 Sep 2021
Cited by 2 | Viewed by 2235
Abstract
Coolant channels allow internal coolant delivery to the cutting region and significantly improve drilling, but these channels also reduce the torsional and torsional-axial stiffness of the drills. Such a reduction in stiffness can degrade the quality of the drilled holes. The evacuation of [...] Read more.
Coolant channels allow internal coolant delivery to the cutting region and significantly improve drilling, but these channels also reduce the torsional and torsional-axial stiffness of the drills. Such a reduction in stiffness can degrade the quality of the drilled holes. The evacuation of cutting chips and the delivery of the cutting fluid put strict geometrical restrictions on the cross-section design of the drill. This necessitates careful selection and optimization of features such as the geometry of the coolant channels. This paper presents a new method that uses Prandtl’s stress function to predict the torsional and torsional-axial stiffness values. Using this method drills with one central channel are compared to those with two eccentric coolant channels, which shows that with the same cross-section area, the reduction of axial and torsional-axial stiffness is notably smaller for the design with two eccentric channels compared to a single central channel. The stress function method is further used to select the appropriate location of the eccentric coolant channels to minimize the loss of torsional and torsional-axial stiffness. These results are verified by comparison to the results of three-dimensional finite element analyses. Full article
(This article belongs to the Special Issue Modeling and Simulation of Metal Processing)
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<p>A drill cross-section, <span class="html-italic">A</span>, with the boundary <math display="inline"><semantics> <mi>Γ</mi> </semantics></math>. (<b>a</b>) without hole, (<b>b</b>) with a hole, <span class="html-italic">B</span>.</p> Full article ">Figure 2
<p>Typical drill cross-sections: (<b>a</b>) solid drill (<b>b</b>) with a single coolant channel (<b>c</b>) with two coolant channels. Drills with coolant channels have total hole areas equal to 17 mm<sup>2</sup>. (Length parameters are in mm.)</p> Full article ">Figure 3
<p>Prandtl’s stress function <math display="inline"><semantics> <mi>ψ</mi> </semantics></math> for the cross-section shown in <a href="#metals-11-01473-f002" class="html-fig">Figure 2</a>a; in local minima points, the shear stress components (derivatives of Prandtl’s stress function) are zero.</p> Full article ">Figure 4
<p>Prandtl’s stress function (<math display="inline"><semantics> <mi>ψ</mi> </semantics></math>), for drills with (<b>a</b>) a single central coolant channel and (<b>b</b>) twin eccentric channels placed at the minima shown in <a href="#metals-11-01473-f003" class="html-fig">Figure 3</a>.</p> Full article ">Figure 5
<p>Tangential displacement with 3D FEA for the (<b>a</b>) solid drill (max: 9.21 μm), (<b>b</b>) drill with a central channel (max: 10.21 μm), (<b>c</b>) two-channel drill (max: 9.31 μm).</p> Full article ">Figure 6
<p>Torsional stiffness versus total pre-twist angle for the solid drill, according to the 3D FEM model.</p> Full article ">Figure 7
<p>Axial displacement in meters calculated by 3D FEM for the (<b>a</b>) solid drill (max: 1.43 μm), (<b>b</b>) drill with a central channel (max: 1.82 μm), (<b>c</b>) two-channel drill (max: 1.46 μm).</p> Full article ">Figure 8
<p>Axial deformation (warping) of the top surface as calculated by the 3D FEM (<b>a</b>) solid drill (max: 1.43 μm), (<b>b</b>) drill with a central channel (max: 1.82 μm), (<b>c</b>) two-channel drill (max: 1.46 μm).</p> Full article ">Figure 9
<p>Torsional stiffness versus total cross-section area of the coolant channels.</p> Full article ">Figure 10
<p>Reduction in torsional and torsional-axial stiffness versus the radial position of the upper coolant channel.</p> Full article ">
8 pages, 1175 KiB  
Article
Role of Heat Treatment on Atomic Order and Ordering Domains in Ni45Co5Mn36.6In13.4 Ribbons
by Yan Feng, Xueman Wan, Xiaohai Bian, Yanling Ai and Haibo Wang
Metals 2021, 11(9), 1472; https://doi.org/10.3390/met11091472 - 16 Sep 2021
Cited by 3 | Viewed by 1752
Abstract
The effects of cooling rate and annealed temperature on the state of atomic order and microstructure of L21 domains of Ni45Co5Mn36.6In13.4 ribbons are investigated comprehensively. The state of atomic order is quantitatively studied by in [...] Read more.
The effects of cooling rate and annealed temperature on the state of atomic order and microstructure of L21 domains of Ni45Co5Mn36.6In13.4 ribbons are investigated comprehensively. The state of atomic order is quantitatively studied by in situ X-ray diffraction (XRD), and the microstructure of ordered domains is revealed by transmission electron microscopy (TEM). As-spun ribbons show B2 structure of low atomic order, exhibiting the dispersive L21 domains’ morphology. By applying heat treatment around the order–disorder transition temperature followed by furnace cooling or quenching into water, respectively, we found the strong dependence of ordered domains on cooling rates. Furnace cooling samples show L21 domains with small sized antiphase boundary, revealing a high degree of atomic order, while quenching hinders the formation of ordered domains. Annealing above the order–disorder transition temperature followed by quenching preserves the disordered atomic state with the mixture of L21 structure in B2 matrix. Full article
(This article belongs to the Section Entropic Alloys and Meta-Metals)
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<p>(<b>a</b>) (111) reflections of in situ heating powder XRD (X-ray diffraction) analysis. The inset is the (111) reflection of long-time annealed Ni<sub>45</sub>Co<sub>5</sub>Mn<sub>36.6</sub>In<sub>13.4</sub> ribbons. (<b>b</b>) DSC (Differential Scanning Calorimetry) heating curves of Ni<sub>45</sub>Co<sub>5</sub>Mn<sub>36.6</sub>In<sub>13.4</sub> ribbons and the transition temperatures were determined by tangent method on the endothermic peak. (<b>c</b>) (111) reflections of XRD analysis of as-spun and QW (ice-water quenching) Ni<sub>45</sub>Co<sub>5</sub>Mn<sub>36.6</sub>In<sub>13.4</sub> ribbons and (<b>d</b>) that of as-spun and FC (furnace cooling) Ni<sub>45</sub>Co<sub>5</sub>Mn<sub>36.6</sub>In<sub>13.4</sub> ribbons.</p> Full article ">Figure 2
<p>The dark-field images using (111) reflection taken from Ni<sub>45</sub>Co<sub>5</sub>Mn<sub>36.6</sub>In<sub>13.4</sub> ribbons for (<b>a</b>) as-spun, quenched at (<b>b</b>) 700 K, (<b>c</b>) 800 K, (<b>d</b>) 900 K, (<b>e</b>) 1000 K and (<b>f</b>) 1100 K; and annealed at (<b>g</b>) 700 K, (<b>h</b>) 800 K, (<b>i</b>) 900 K, (<b>j</b>) 1000 K and (<b>k</b>) 1100 K, respectively. Inset of each of image is the corresponding SAD patterns along (<math display="inline"><semantics> <mrow> <mover accent="true"> <mn>1</mn> <mo>¯</mo> </mover> <mn>10</mn> </mrow> </semantics></math>) zone axis of the bright domains in the dark field image.</p> Full article ">
15 pages, 3209 KiB  
Article
Corrosion Behavior of AlFeCrCoNiZrx High-Entropy Alloys in 0.5 M Sulfuric Acid Solution
by Yuhong Yao, Yaohua Jin, Wei Gao, Xiaoyu Liang, Jian Chen and Shidong Zhu
Metals 2021, 11(9), 1471; https://doi.org/10.3390/met11091471 - 16 Sep 2021
Cited by 17 | Viewed by 3000
Abstract
AlCoCrFeNiZrx (x = 0, 0.1, 0.2, 0.3, and 0.5) high-entropy alloys (HEAs) were prepared by a non-consumable vacuum arc melting technology, and the microstructure and corrosion behavior were investigated by XRD, SEM, immersion tests, and electrochemical measurements. The results indicate that [...] Read more.
AlCoCrFeNiZrx (x = 0, 0.1, 0.2, 0.3, and 0.5) high-entropy alloys (HEAs) were prepared by a non-consumable vacuum arc melting technology, and the microstructure and corrosion behavior were investigated by XRD, SEM, immersion tests, and electrochemical measurements. The results indicate that galvanic corrosion of the AlCoCrFeNiZrx alloys occurred in 0.5 M H2SO4 solution, and only 0.1 mol of the added Zr could greatly improve the corrosion resistance of the alloys. The corrosion properties of the AlCoCrFeNiZrx HEAs had similar change tendencies with the increase in the Zr content in the immersion tests, potentiodynamic polarization measurements, and electrochemical impedance analysis, that is, the corrosion resistance of the AlCoCrFeNiZrx alloys in a 0.5 M H2SO4 solution first increased and then decreased with the increase in the Zr content. The Zr0.1 alloys were found to have the best selective corrosion and general corrosion resistance with the smallest corrosion rate, whereas the Zr0.3 alloys presented the worst selective corrosion and general corrosion resistance with the highest corrosion rate from both the immersion tests and the potentiodynamic polarization measurements. Full article
(This article belongs to the Special Issue Wear and Corrosion Behavior of High-Entropy Alloy)
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<p>Volume loss (<b>a</b>) and the average corrosion rates (<b>b</b>) of as-cast AlCoCrFeNiZr<span class="html-italic">x</span> alloys in a 0.5 M H<sub>2</sub>SO<sub>4</sub> solution for the 0–192 h immersion tests at 25 °C.</p> Full article ">Figure 2
<p>Corroded surface morphologies (<b>a1</b>, <b>b1</b>,<b>c1</b>,<b>d1</b>,<b>e1</b>) and typical EDS spectra in region A (<b>a2</b>,<b>b2</b>,<b>c2</b>,<b>d2</b>,<b>e2</b>) and region B (<b>a3</b>,<b>b3</b>,<b>c3</b>,<b>d3</b>,<b>e3</b>) of as-cast AlCoCrFeNiZr<span class="html-italic">x</span> alloys in a 0.5 M H<sub>2</sub>SO<sub>4</sub> solution for 0–192 h immersion tests at 25 °C: (<b>a</b>) <span class="html-italic">x</span> = 0, (<b>b</b>) <span class="html-italic">x</span> = 0.1, (<b>c</b>) <span class="html-italic">x</span> = 0.2, (<b>d</b>) <span class="html-italic">x</span> = 0.3, and (<b>e</b>) <span class="html-italic">x</span> = 0.5.</p> Full article ">Figure 3
<p>XRD patterns and SEM micrographs of as-cast AlCoCrFeNiZr<span class="html-italic">x</span> (<span class="html-italic">x</span> = 0, 0.1, 0.2, 0.3, and 0.5) alloys: (<b>a</b>) XRD patterns, (<b>b</b>) <span class="html-italic">x</span> = 0, (<b>c</b>) <span class="html-italic">x</span> = 0.1, (<b>d</b>) <span class="html-italic">x</span> = 0.2, (<b>e</b>) <span class="html-italic">x</span> = 0.3, and (<b>f</b>) <span class="html-italic">x</span> = 0.5.</p> Full article ">Figure 4
<p>Metallograph (<b>a</b>) and TEM bright-field image (<b>b</b>) as well as corresponding selected area electron diffraction patterns (<b>c</b>) of the Zr0 alloys.</p> Full article ">Figure 5
<p>Potentiodynamic polarization curves of AlCoCrFeNiZr<span class="html-italic">x</span> (x = 0, 0.1, 0.2, 0.3, and 0.5) alloys in a 0.5 M H2SO4 solution.</p> Full article ">Figure 6
<p>The Nyquist curves (<span class="html-italic">Z</span>im versus <span class="html-italic">Z</span>re) of AlCoCrFeNiZr<span class="html-italic">x</span> (<span class="html-italic">x</span> = 0, 0.1, 0.2, 0.3, and 0.5) alloys in a 0.5 M H<sub>2</sub>SO<sub>4</sub> solution.</p> Full article ">Figure 7
<p>EIS equivalent circuits of AlCoCrFeNiZr<span class="html-italic">x</span> (<span class="html-italic">x</span> = 0, 0.1, 0.2, 0.3, and 0.5) alloys in a 0.5 M H<sub>2</sub>SO<sub>4</sub> solution: (<b>a</b>) <span class="html-italic">x</span> = 0, 0.1, 0.3, and 0.5; (<b>b</b>) <span class="html-italic">x</span> = 0.2.</p> Full article ">
11 pages, 1361 KiB  
Article
Investigation of the Possibility of Tailoring the Chemical Com-Position of the NiTi Alloy by Selective Laser Melting
by Evgenii Borisov, Kirill Starikov, Anatoly Popovich and Tatiana Tihonovskaya
Metals 2021, 11(9), 1470; https://doi.org/10.3390/met11091470 - 16 Sep 2021
Cited by 8 | Viewed by 2234
Abstract
In this work a study of the selective laser melting process of two NiTi alloys of equiatomic, and rich Ni composition were conducted. A study of the influence of the technological parameters on the alloy density was carried out. Values of technological parameters [...] Read more.
In this work a study of the selective laser melting process of two NiTi alloys of equiatomic, and rich Ni composition were conducted. A study of the influence of the technological parameters on the alloy density was carried out. Values of technological parameters were obtained to ensure production of samples with the lowest number of defects. When using process parameters with the same energy density but different values of the constituent technological parameters, the amount of nickel carried away by evaporation changed insignificantly. An increase in the energy density led to an increase in the amount of nickel carried away, causing final samples with lower Ni content. When using multiple laser processing in the low-energy parameter set, it was possible to achieve a decrease in the nickel content in the alloy, similar to that with single high-energy processing. DSC studies showed a significant increase in transformation temperatures upon repeated laser processing due to the higher evaporation of nickel. The use of double laser treatment gave a decrease in the final density of the sample compared to a single treatment, but its value is still higher than when using a single treatment with a higher energy density. Full article
(This article belongs to the Special Issue Advanced Manufacturing of Novel Metallic Related Materials)
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<p>A sample prepared by selective laser melting from NiTi 50.7 alloy powder. (<b>a</b>) with VED = 173.6 J/mm<sup>3</sup>, parameter set 2 (<a href="#metals-11-01470-t002" class="html-table">Table 2</a>); (<b>b</b>) with VED = 130.2 (×2) J/mm<sup>3</sup>, parameter set 4_2 (<a href="#metals-11-01470-t003" class="html-table">Table 3</a>).</p> Full article ">Figure 2
<p>Dependence of the Ni content in compact samples on the bulk energy density during the SLM process. Black—NiTi 50.0, gray—NiTi 50.7.</p> Full article ">Figure 3
<p>Dependence of the Ni content in compact samples on the number of laser treatments. Black dots—50.0, gray dots—50.7.</p> Full article ">Figure 4
<p>DSC curves for samples made from NiTi 50.7 powder with parameter sets 2 with VED = 173.6 J/mm<sup>3</sup> (<b>a</b>), parameter set 4 with VED = 130.2 J/mm<sup>3</sup> (<b>b</b>), parameter set 4_2 with VED = 130.2 × 2 J/mm<sup>3</sup> (<b>c</b>) according to <a href="#metals-11-01470-t002" class="html-table">Table 2</a> and <a href="#metals-11-01470-t003" class="html-table">Table 3</a>.</p> Full article ">Figure 4 Cont.
<p>DSC curves for samples made from NiTi 50.7 powder with parameter sets 2 with VED = 173.6 J/mm<sup>3</sup> (<b>a</b>), parameter set 4 with VED = 130.2 J/mm<sup>3</sup> (<b>b</b>), parameter set 4_2 with VED = 130.2 × 2 J/mm<sup>3</sup> (<b>c</b>) according to <a href="#metals-11-01470-t002" class="html-table">Table 2</a> and <a href="#metals-11-01470-t003" class="html-table">Table 3</a>.</p> Full article ">
3 pages, 170 KiB  
Editorial
Sustainable Steel Industry: Energy and Resource Efficiency, Low-Emissions and Carbon-Lean Production
by Valentina Colla and Teresa Annunziata Branca
Metals 2021, 11(9), 1469; https://doi.org/10.3390/met11091469 - 16 Sep 2021
Cited by 4 | Viewed by 3299
Abstract
The three pillars of sustainability represented by the environment, economy and society in the steel manufacturing industry are directly connected to the efficient and effective management of resources, such as energy, raw materials, by-products and water [...] Full article
(This article belongs to the Special Issue Sustainable Steel Industry: Energy and Resource Efficiency, Low-Emissions and Carbon-Lean Production)
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