[go: up one dir, main page]
Next Issue
Volume 12, December
Previous Issue
Volume 12, October
 
 
Search for Articles:
Title / Keyword
Author / Affiliation / Email
Journal
Article Type
 
 
Advanced Search
 
Section
Special Issue
Volume
Issue
Number
Page
 
4.9
2.6
Journals
Metals
Volume 12
Issue 11
share announcement

Metals, Volume 12, Issue 11 (November 2022) – 222 articles

Cover Story (view full-size image): The fatigue behavior of austenitic stable and metastable stainless steels with different surface morphologies was investigated regarding the VHCF regime using an ultrasonic fatigue testing system. The AISI 904L is stable against deformation-induced phase formation, while the AISI 347 is in the metastable state and shows martensite formation. The specimens from stable austenite failed in the HCF and VHCF regimes. In contrast, the metastable austenite achieved true fatigue limits and failed only in the HCF regime. Due to surface modification, an increase in the fatigue strength of metastable AISI 347 was achieved. View this paper
  • Issues are regarded as officially published after their release is announced to the table of contents alert mailing list.
  • You may sign up for e-mail alerts to receive table of contents of newly released issues.
  • PDF is the official format for papers published in both, html and pdf forms. To view the papers in pdf format, click on the "PDF Full-text" link, and use the free Adobe Reader to open them.
Order results
Result details
Section
Show export options expand_more Show export options expand_less
Select all
Export citation of selected articles as:
38 pages, 11767 KiB  
Article
Optimal Design of Three-Stress Accelerated Degradation Test Plan for Motorized Spindle with Poor Prior Information
by Hongxun Zhao, Zhaojun Yang, Chuanhai Chen, Zhifeng Liu, Wei Luo and Chunlei Hua
Metals 2022, 12(11), 1996; https://doi.org/10.3390/met12111996 - 21 Nov 2022
Cited by 1 | Viewed by 1355
Abstract
Accurate optimal design for the test plan with limited prior information is impossible since the optimal design method of a three-stress accelerated degradation test plan for a motorized spindle is based on the determination of model parameters. In order to optimize the test [...] Read more.
Accurate optimal design for the test plan with limited prior information is impossible since the optimal design method of a three-stress accelerated degradation test plan for a motorized spindle is based on the determination of model parameters. In order to optimize the test plan with poor prior information, a “dynamic” optimal design method is proposed in this article. Firstly, a three-stress accelerated degradation model with a stress coupling term is established based on the correlation of the degradation rate of the motorized spindle, and the parameters in the model are regarded as variables to represent the deviation between the prior information and the true value of the motorized spindle when the prior information is poor. Then, based on the information theory and the sequential design method, an optimal design method of the three-stress accelerated degradation test plan of the motorized spindle with the information entropy as the objective function is proposed to realize the “dynamic” optimization of the test plan. Finally, the usability of the proposed method is verified by taking a Chinese model spindle as an example, and the validity of the method is verified by checking the model accuracy of the accelerated degradation model of the motorized spindle after the test. Full article
(This article belongs to the Special Issue Quality Prediction and Control Technology Design for Intelligent Manufacturing)
►▼ Show Figures

Figure 1

Figure 1
<p>The performance degradation track of motorized spindle.</p> Full article ">Figure 2
<p>Relationship of model parameters.</p> Full article ">Figure 3
<p>General flow of sequential design for ADT plan of motorized spindle.</p> Full article ">Figure 4
<p>Optimization process of three-stress ADT plan of motorized spindle with poor prior information.</p> Full article ">Figure 5
<p>The Actual test equipment of motorized spindle.</p> Full article ">Figure 6
<p>Iterative process of the optimal solution of pre-test optimization.</p> Full article ">Figure 6 Cont.
<p>Iterative process of the optimal solution of pre-test optimization.</p> Full article ">Figure 7
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 7).</p> Full article ">Figure 7 Cont.
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 7).</p> Full article ">Figure 8
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 6).</p> Full article ">Figure 8 Cont.
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 6).</p> Full article ">Figure 9
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 5).</p> Full article ">Figure 9 Cont.
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 5).</p> Full article ">Figure 10
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 4).</p> Full article ">Figure 10 Cont.
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 4).</p> Full article ">Figure 11
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 3).</p> Full article ">Figure 11 Cont.
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 3).</p> Full article ">Figure 12
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 2).</p> Full article ">Figure 12 Cont.
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 2).</p> Full article ">Figure 13
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 1).</p> Full article ">Figure 13 Cont.
<p>The update process of probability density of each model parameter with the test (<span class="html-italic">l</span> = 1).</p> Full article ">Figure 14
<p>The update process of the posterior distribution of each model parameter.</p> Full article ">Figure 14 Cont.
<p>The update process of the posterior distribution of each model parameter.</p> Full article ">Figure 15
<p>The Quantile-Quantile plot of model and standard normal distribution.</p> Full article ">
13 pages, 4660 KiB  
Article
The Growth of Intermetallic Compounds and Its Effect on Bonding Properties of Cu/Al Clad Plates by CFR
by Long Li, Guangping Deng, Weiguo Zhai, Sha Li, Xiangyu Gao and Tao Wang
Metals 2022, 12(11), 1995; https://doi.org/10.3390/met12111995 - 21 Nov 2022
Cited by 4 | Viewed by 1480
Abstract
Cu/Al clad plates prepared using a corrugated + flat rolling (CFR) technique were annealed at 300–450 °C for 10–240 min. Furthermore, the interfacial diffusion behavior and the bonding properties of the Cu/Al clad plates were studied in detail. The results demonstrated that, at [...] Read more.
Cu/Al clad plates prepared using a corrugated + flat rolling (CFR) technique were annealed at 300–450 °C for 10–240 min. Furthermore, the interfacial diffusion behavior and the bonding properties of the Cu/Al clad plates were studied in detail. The results demonstrated that, at the initial stage of the annealing process, the development of the first IMCs layer was restrained by the high atomic concentration gradient in the new bonding interface zone, and the second intermetallic compounds (IMCs) layer preferentially formed in the new bonding interface zone, leading to a slight increase in the growth activation energy of the clad plates. In addition, the atoms’ diffusion behavior at the peak and trough interfaces was not significantly affected by the matrix microstructure, and there was no discernible difference in the growth activation energy at these two positions. Ultimately, it was shown that the maximum average peel strength at the peak and trough interfaces reached 53.07 N/mm and 41.23 N/mm, respectively, when annealing at 350 °C for 10 min. Full article
(This article belongs to the Special Issue Process and Numerical Simulation of Oxygen Steelmaking)
►▼ Show Figures

Figure 1

Figure 1
<p>The schematic diagram of CFR technique.</p> Full article ">Figure 2
<p>The positions and specifications of the peel test samples.</p> Full article ">Figure 3
<p>SEM images of the Cu/Al clad plates: (<b>a</b>,<b>c</b>) at peak, (<b>b</b>,<b>d</b>) at trough, (<b>a</b>,<b>b</b>) by CCRB, (<b>c</b>,<b>d</b>) by FRB; (<b>e</b>) and (<b>f</b>) are the EDX line scanning results along the line in (<b>a</b>) and (<b>b</b>), respectively.</p> Full article ">Figure 4
<p>Interface SEM images with different annealing temperature: (<b>a</b>,<b>e</b>) 300 °C, (<b>b</b>,<b>f</b>) 350 °C, (<b>c</b>,<b>g</b>) 400 °C, (<b>d</b>,<b>h</b>) 450 °C; (<b>a</b>–<b>d</b>) peak, (<b>e</b>–<b>h</b>) trough, annealing time: 60 min.</p> Full article ">Figure 5
<p>Function of intermetallic sub-layer thickness and annealing time.</p> Full article ">Figure 6
<p>Arrhenius plot for the growth rate of the intermetallic layer: (<b>a</b>) CuAl<sub>2</sub>, (<b>b</b>) Cu<sub>9</sub>Al<sub>4</sub>, (<b>c</b>) CuAl.</p> Full article ">Figure 7
<p>Interface SEM images with different annealing times: (<b>a</b>,<b>b</b>) 10 min, (<b>c</b>,<b>d</b>) 20 min, (<b>e</b>,<b>f</b>) 40 min, (<b>g</b>,<b>h</b>) 60 min; (<b>a</b>,<b>c</b>,<b>e</b>,<b>g</b>) peak, (<b>b</b>,<b>d</b>,<b>f</b>,<b>h</b>) trough.</p> Full article ">Figure 8
<p>Plots of peel strength versus the displacement at (<b>a</b>) peak and (<b>b</b>) trough; (<b>c</b>) statistical diagram of average peel strength.</p> Full article ">Figure 9
<p>SEM images of peeling section at peak: (<b>a</b>,<b>b</b>) 10 min, (<b>c</b>,<b>d</b>) 20 min, (<b>e</b>,<b>f</b>) 40 min, (<b>g</b>,<b>h</b>) 60 min; (<b>a</b>,<b>c</b>,<b>e</b>,<b>g</b>) Cu side, (<b>b</b>,<b>d</b>,<b>f</b>,<b>h</b>) Al side.</p> Full article ">Figure 10
<p>SEM images of peeling section at trough: (<b>a</b>,<b>b</b>) 10 min, (<b>c</b>,<b>d</b>) 20 min, (<b>e</b>,<b>f</b>) 40 min, (<b>g</b>,<b>h</b>) 60 min; (<b>a</b>,<b>c</b>,<b>e</b>,<b>g</b>) Cu side, (<b>b</b>,<b>d</b>,<b>f</b>,<b>h</b>) Al side.</p> Full article ">
8 pages, 1746 KiB  
Article
Inclusion Characteristics in Steel with CeO2 Nanoparticle Addition
by Hui Kong, Xiang Cheng, Shoulin Huang and Yue Qiu
Metals 2022, 12(11), 1994; https://doi.org/10.3390/met12111994 - 21 Nov 2022
Cited by 3 | Viewed by 1367
Abstract
The application of Ce oxides in oxide metallurgy has received extensive attention, but until now, the direct adding of CeO2 into molten steel to generate Ce oxides has not occurred. In this paper, a mixture of CeO2 and Si nanoparticles were [...] Read more.
The application of Ce oxides in oxide metallurgy has received extensive attention, but until now, the direct adding of CeO2 into molten steel to generate Ce oxides has not occurred. In this paper, a mixture of CeO2 and Si nanoparticles were added into molten steel. The resultant formation of micrometer scale Ce-bearing oxides confirmed its adding validity. This behavior may be interpreted as the reactivity between CeO2 and [Al], and the improved wettability between CeO2 and molten steel with the assistance of Si powder. Thus, when the quantity of CeO2 is kept constant, its added yield should increase when increasing the added quantity of Si. This was verified by the larger percentage of Ce-bearing oxides of the total oxides and the greater average content of Ce in Ce-bearing oxides after normalization. Moreover, compared with the blank sample, statistical results indicated that the oxides in CeO2-modified samples were refined, and their dispersion homogeneity was enhanced. This comparison indicates the effectiveness of the external adding method in oxide metallurgy. Full article
(This article belongs to the Topic Advanced Processes in Metallurgical Technologies)
►▼ Show Figures

Figure 1

Figure 1
<p>SEM micrograph and EDS mapping images of various elements for typical oxide in 1#.</p> Full article ">Figure 2
<p>SEM micrograph and EDS mapping images of various elements for typical oxide in 2#.</p> Full article ">Figure 3
<p>SEM micrograph and EDS mapping images of various elements for typical oxide in 3#.</p> Full article ">Figure 4
<p>The size distribution of oxides for samples.</p> Full article ">
16 pages, 6381 KiB  
Article
Effect of Shoulder Fillet Radius on Welds in Bobbin Tool Friction Stir Welding of A1050
by Huilin Miao, Takuya Miura, Wei Jiang, Masato Okada and Masaaki Otsu
Metals 2022, 12(11), 1993; https://doi.org/10.3390/met12111993 - 21 Nov 2022
Cited by 2 | Viewed by 1614
Abstract
In this study, five bobbin tools with different shoulder fillet radii were employed for the bobbin tool friction stir welding (BT-FSW) of A1050-O sheets to systematically evaluate the effects of shoulder fillet radius on the welding defect formation, flash formation, weld thickness, grain [...] Read more.
In this study, five bobbin tools with different shoulder fillet radii were employed for the bobbin tool friction stir welding (BT-FSW) of A1050-O sheets to systematically evaluate the effects of shoulder fillet radius on the welding defect formation, flash formation, weld thickness, grain size of the stir zone, and tensile properties. The quality classifications of the joints’ appearance were summarized as process windows, and the appropriate welding condition range for each shoulder fillet radius was clarified. It was observed that an increase in the shoulder fillet radius decreased the welding defects and flash formation; however, it increased the minimum thickness of the weld except when the shoulder fillet radius was 0.5 mm. The grain size of the stir zone increased with increasing shoulder fillet radius from 0.5 mm to 6 mm. The ultimate tensile strength (UTS) of the stir zone decreased with increasing shoulder fillet radius from 0.5 mm to 1 mm, increased from 1 mm to 3 mm, and remained constant from 3 mm to 6 mm. The results indicate that a shoulder fillet radius larger than 3 mm is effective in decreasing flash formation and maintaining a constant weld thickness. Full article
(This article belongs to the Section Welding and Joining)
►▼ Show Figures

Figure 1

Figure 1
<p>Schematic illustration of BT-FSW.</p> Full article ">Figure 2
<p>Appearance of the FSW setup.</p> Full article ">Figure 3
<p>Schematic illustration of a bobbin tool.</p> Full article ">Figure 4
<p>Schematic illustrations of the welding path and sampling location and size for (<b>a</b>) a within-weld tensile test and cross-sectional observation, (<b>b</b>) whole-weld tensile test, (<b>c</b>) a within-weld tensile test specimen, and (<b>d</b>) a whole-weld tensile test specimen.</p> Full article ">Figure 5
<p>Process windows of shoulder fillet radii <span class="html-italic">R</span><sub>sf</sub> of (<b>a</b>) 0.5 mm, (<b>b</b>) 3 mm, and (<b>c</b>) 6 mm (the red dotted line delineates the weldable area without welding defects, and the blue dotted delineates the sound joint area without welding defects or flash formation.</p> Full article ">Figure 6
<p>Appearances of the upper surfaces of the BT-FSW welds.</p> Full article ">Figure 6 Cont.
<p>Appearances of the upper surfaces of the BT-FSW welds.</p> Full article ">Figure 7
<p>Appearances of weld surface and cross-sectional images of welds with different shoulder fillet radii <span class="html-italic">R</span><sub>sf</sub> (<span class="html-italic">N</span> = 2000 rpm, <span class="html-italic">V</span> = 2000 mm/min, <span class="html-italic">F</span><sub>s</sub> = 8 mm, <span class="html-italic">d</span><sub>s</sub> = 1.8 mm).</p> Full article ">Figure 8
<p>Relation between minimum thickness, average height of flash, and the shoulder fillet radii <span class="html-italic">R</span><sub>sf</sub> (<span class="html-italic">N</span> = 2000 rpm, <span class="html-italic">V</span>= 2000 mm/min, <span class="html-italic">F</span><sub>s</sub> = 8 mm, <span class="html-italic">d</span><sub>s</sub> = 1.8 mm).</p> Full article ">Figure 9
<p>Cross-sectional microstructure of flash formed in the RS edge of the stir zones welded with the shoulder fillet radius <span class="html-italic">R</span><sub>sf</sub> of (<b>a</b>) 0.5 mm and (<b>b</b>) 2 mm. (<span class="html-italic">N</span> = 2000 rpm, <span class="html-italic">V</span> = 2000 mm/min, <span class="html-italic">F</span><sub>s</sub> = 8 mm, <span class="html-italic">d</span><sub>s</sub> = 1.8 mm).</p> Full article ">Figure 10
<p>Optical microscope images of (<b>a</b>) base metal and stir zones with different shoulder fillet radii <span class="html-italic">R</span><sub>sf</sub> of (<b>b</b>) 0.5 mm, (<b>c</b>) 1 mm, (<b>d</b>) 2 mm, (<b>e</b>) 3 mm, and (<b>f</b>) 6 mm, respectively. (<span class="html-italic">N</span> = 2000 rpm, <span class="html-italic">V</span> = 2000 mm/min, <span class="html-italic">F</span><sub>s</sub> = 8 mm, tool gap = 1.8 mm).</p> Full article ">Figure 11
<p>Relation between the grain size of the stir zone and the shoulder fillet radius <span class="html-italic">R</span><sub>sf</sub> (<span class="html-italic">N</span> = 2000 rpm, <span class="html-italic">V</span> = 2000 mm/min, <span class="html-italic">F</span><sub>s</sub> = 8 mm, <span class="html-italic">d</span><sub>s</sub> = 1.8 mm).</p> Full article ">Figure 12
<p>Relation between shoulder fillet radii <span class="html-italic">R</span><sub>sf</sub> and tensile properties of whole-weld specimens (<span class="html-italic">N</span> = 2000 rpm, <span class="html-italic">V</span> = 2000 mm/min, <span class="html-italic">F</span><sub>s</sub> = 8 mm, <span class="html-italic">d</span><sub>s</sub> = 1.8 mm).</p> Full article ">Figure 13
<p>Appearance of whole-joint tensile specimens with different shoulder fillet radii <span class="html-italic">R</span><sub>sf</sub>.</p> Full article ">Figure 14
<p>Relation between shoulder fillet radii <span class="html-italic">R</span><sub>sf</sub> and tensile properties of within-weld specimens (<span class="html-italic">N</span> = 2000 rpm, <span class="html-italic">V</span> = 2000 mm/min, <span class="html-italic">F</span><sub>s</sub> = 8 mm, <span class="html-italic">d</span><sub>s</sub> = 1.8 mm).</p> Full article ">
10 pages, 12679 KiB  
Article
Laser Beam Welding of CubeSat 1U Structure Parts Obtained by Powder Bed Fusion
by Rafael Humberto Mota de Siqueira, Diego Javier Celentano and Milton Sergio Fernandes de Lima
Metals 2022, 12(11), 1992; https://doi.org/10.3390/met12111992 - 21 Nov 2022
Viewed by 1496
Abstract
This study contributes to a possible methodology for manufacturing CubeSats using additive manufacturing and laser beam welding. Titanium connectors were constructed by selective laser melting and electron beam melting and characterized from a topological point of view. The connectors can be joined to [...] Read more.
This study contributes to a possible methodology for manufacturing CubeSats using additive manufacturing and laser beam welding. Titanium connectors were constructed by selective laser melting and electron beam melting and characterized from a topological point of view. The connectors can be joined to titanium tubes for the construction of CubeSats via laser spot welding. The fiber laser welds exhibited full penetration using pulses with 400 J of energy. The welds showed titanium acicular martensite grains with recesses and pores. The average hardness of the cast zone was 350 HV, which is close to the hardness of the connectors (400 HV) and more rigid than that of the tubes (100 HV). Spot welding has proven to be useful in resisting forces above 2000 N, which is sufficient for CubeSat frame space applications. Full article
(This article belongs to the Special Issue Laser Welding and Welding Joint Quality Assessment - State of Art)
►▼ Show Figures

Figure 1

Figure 1
<p>Computer drawing of the connector (<b>a</b>) and its dimensions (<b>b</b>).</p> Full article ">Figure 2
<p>Shapes and dimensions of the connectors: Isometric view of the connector pieces: (<b>a</b>) SLM and (<b>b</b>) EBM. Scale at the right. Actual pieces: (<b>c</b>) SLM and (<b>d</b>) EBM connectors.</p> Full article ">Figure 2 Cont.
<p>Shapes and dimensions of the connectors: Isometric view of the connector pieces: (<b>a</b>) SLM and (<b>b</b>) EBM. Scale at the right. Actual pieces: (<b>c</b>) SLM and (<b>d</b>) EBM connectors.</p> Full article ">Figure 3
<p>Example of the actual CubeSat structure ready for the welds. Each edge measures 100 mm.</p> Full article ">Figure 4
<p>Laser spot welds for superposed connector to tube geometry: (<b>a</b>) SLM and (<b>b</b>) EBM.</p> Full article ">Figure 5
<p>Laser spot welds cross-section macrographs: (<b>a</b>) SLM and (<b>b</b>) EBM. The letters refer to fusion zone (FZ), connector (C), and tube (T). The arrows indicate, approximately, the laser beam optical axis. Micrographs of FZ: (<b>c</b>) SLM and (<b>d</b>) EBM.</p> Full article ">Figure 5 Cont.
<p>Laser spot welds cross-section macrographs: (<b>a</b>) SLM and (<b>b</b>) EBM. The letters refer to fusion zone (FZ), connector (C), and tube (T). The arrows indicate, approximately, the laser beam optical axis. Micrographs of FZ: (<b>c</b>) SLM and (<b>d</b>) EBM.</p> Full article ">Figure 6
<p>Vickers hardness values as a function of the weld center distance for SLM and EBM cases. Each region is noted at bottom.</p> Full article ">Figure 7
<p>Tensile force (F) versus displacement (d) for two representative cases of SLM and EBM and for the unwelded base Ti tube.</p> Full article ">Figure 8
<p>Fractured surfaces after tensile test for two representative cases: (<b>a</b>) SLM and (<b>b</b>) EBM.</p> Full article ">
11 pages, 3366 KiB  
Article
Analysis of Texture and Anisotropic Elastic Properties of Additively Manufactured Ni-Base Alloys
by Thomas Obermayer, Christian Krempaszky and Ewald Werner
Metals 2022, 12(11), 1991; https://doi.org/10.3390/met12111991 - 21 Nov 2022
Cited by 5 | Viewed by 1664
Abstract
Additive manufacturing of metallic materials generates strong crystallographic textures, leading to anisotropic elastic properties on the macroscopic scale. The impact of the processing parameters on the resulting texture requires suitable techniques for the prediction and the experimental determination of elastic properties to exploit [...] Read more.
Additive manufacturing of metallic materials generates strong crystallographic textures, leading to anisotropic elastic properties on the macroscopic scale. The impact of the processing parameters on the resulting texture requires suitable techniques for the prediction and the experimental determination of elastic properties to exploit the anisotropy in the design process. Within this study mechanical as well as microstructure based approaches are applied on a batch of specimens manufactured from IN718 by selective laser melting to assess the elastic behavior on macroscropic scale. Tensile loading experiments and the impulse excitation technique are applied for the determination of elastic properties without additional constitutive data. Furthermore, the elastic behavior is estimated from single-crystal elastic properties and texture data measured by electron backscatter diffraction and high energy X-ray diffraction. The results of the applied approaches are discussed and compared, allowing also to assess the homogeneity of the elastic properties within the batch of specimens. Full article
(This article belongs to the Special Issue Advanced Techniques for Metallurgical Characterization)
►▼ Show Figures

Figure 1

Figure 1
<p>(<b>a</b>) Batch of cylindrical specimens showing the locations of specimens for analysis by EBSD (specimen A–C) and by HE-XRD (specimen D–F). (<b>b</b>) Machined specimens for tensile loading experiments and IET.</p> Full article ">Figure 2
<p>Stripes, scan vectors and scanning sequence of FO1 scanning strategy.</p> Full article ">Figure 3
<p>(<b>a</b>) Stress and strain plotted over time in a measurement cycle. (<b>b</b>) Evaluation of the Young’s modulus from a single measurement cycle.</p> Full article ">Figure 4
<p>Sketch of the experimental setup for determination of <math display="inline"><semantics> <msub> <mi>E</mi> <mn>1</mn> </msub> </semantics></math>, with respect to the specimen basis (<math display="inline"><semantics> <msub> <mi mathvariant="bold">e</mi> <mn>1</mn> </msub> </semantics></math>, <math display="inline"><semantics> <msub> <mi mathvariant="bold">e</mi> <mn>3</mn> </msub> </semantics></math>) from the fundamental flexural eigenfrequency <math display="inline"><semantics> <msub> <mi>f</mi> <mi mathvariant="normal">f</mi> </msub> </semantics></math>.</p> Full article ">Figure 5
<p>Distance measure <math display="inline"><semantics> <msup> <mi>d</mi> <mi mathvariant="normal">E</mi> </msup> </semantics></math>, shown in (<b>a</b>), and distance measure <math display="inline"><semantics> <msup> <mi>d</mi> <mi mathvariant="double-struck">C</mi> </msup> </semantics></math>, shown in (<b>b</b>), between elasticity tensors determined from different measurement areas and the corresponding elasticity tensors determined from measurement areas of <math display="inline"><semantics> <mrow> <mn>4</mn> <mspace width="0.166667em"/> <msup> <mi>mm</mi> <mn>2</mn> </msup> </mrow> </semantics></math> in case of specimen A, B and C.</p> Full article ">Figure 6
<p>(<b>a</b>) Texture measurement by the Debey-Scherrer method. (<b>b</b>) Diffraction image.</p> Full article ">Figure 7
<p>Comparison of Young’s moduli in direction of the cylinder axes obtained by different approaches. Tensile loading (TL), Impulse Excitation Technique (IET), Electron Backscatter Diffraction (EBSD), High Energy XRD (HE-XRD).</p> Full article ">Figure 8
<p>(<b>a</b>) Pole figures of specimens A–C, determined by EBSD from measurement areas of <math display="inline"><semantics> <mrow> <mn>4</mn> <mspace width="0.166667em"/> <msup> <mi>mm</mi> <mn>2</mn> </msup> </mrow> </semantics></math>. (<b>b</b>) Pole figures of specimens D–F, determined by HE-XRD.</p> Full article ">
17 pages, 9889 KiB  
Article
Macrosegregation Evolution in Eutectic Al-Si Alloy under the Influence of a Rotational Magnetic Field
by Kassab Al-Omari, András Roósz, Arnold Rónaföldi, Mária Svéda and Zsolt Veres
Metals 2022, 12(11), 1990; https://doi.org/10.3390/met12111990 - 21 Nov 2022
Cited by 2 | Viewed by 1680
Abstract
Using magnetic stirring during solidification provides a good opportunity to control the microstructure of alloys, thus controlling their physical properties. However, magnetic stirring is often accompanied by a change in local concentrations, and new structures form which could harm the physical properties. This [...] Read more.
Using magnetic stirring during solidification provides a good opportunity to control the microstructure of alloys, thus controlling their physical properties. However, magnetic stirring is often accompanied by a change in local concentrations, and new structures form which could harm the physical properties. This research paper investigated the effect of forced melt flow by a rotating magnetic field (RMF) on the macrostructure of an Al-Si eutectic alloy. To serve this purpose, Al-12.6 wt% Si alloy samples were solidified in a vertical Bridgman-type furnace equipped with a rotating magnetic inductor to induce the flow in the melt. The diameter and length of the sample are 8 mm and 120 mm, respectively. The solidification parameters are a temperature gradient (G) of 6 K/m, and the solid/liquid front velocity (v) of 0.1 mm/s. These samples were divided into parts during the solidification process, where some of these parts are solidified under the effect of RMF stirring while others are solidified without stirring. The structure obtained after solidification showed a distinct impact of stirring by RMF; new phases have been solidified which were not originally present in the structure before stirring. Besides the eutectic structure, the new phases are the primary aluminum and the primary silicon. The Si concentration and the volume fraction of each phase were measured using Energy-Dispersive Spectroscope (EDS)and new image processing techniques. The experimental results reveal that applying the RMF during the solidification has a distinct effect on the macrostructure of Al-Si eutectic alloys. Indeed, the RMF provokes macro-segregation, reduces the amount of eutectic structure, and changes the sample’s Si concentration distribution. Full article
(This article belongs to the Special Issue Solidification Process of Alloys under Magnetic Field)
►▼ Show Figures

Figure 1

Figure 1
<p>(<b>a</b>) The primary and secondary melt flow in the unidirectionally solidified sample. (<b>b</b>) simulated spiral flow in the pure liquid phase, the flow paths colored by the velocity magnitude (m/s). (<b>c</b>) “Christmas Tree Like” (CTL) macrostructure.</p> Full article ">Figure 2
<p>Sketch of solidification facility; 1: Sample, 2: alumina capsule, 3: quartz tube, 4: copper cooling core, 5: furnace with four heating zones, 6: step motor, 7: RMF inductor, 8: water cooling, 9: basement [<a href="#B5-metals-12-01990" class="html-bibr">5</a>].</p> Full article ">Figure 3
<p>The solidified macro-structure for the samples in four steps (M = 50×). (<b>A</b>) Sample A; (<b>B</b>) Sample B; (<b>C</b>) Sample C; (<b>D</b>) Sample D.</p> Full article ">Figure 4
<p>Micrograph of the solidified samples during step 1 (M = 50×).</p> Full article ">Figure 5
<p>Macrostructures of solidified samples during step 2 of the solidification experiments (M = 50×); (<b>A</b>) step 2 of sample A, when 20 mT magnetic induction was applied. (<b>B</b>) step 2 of sample B, when 60 mT magnetic induction was applied. (<b>C</b>) step 2 of sample C, when 100 mT magnetic induction was applied. (<b>D</b>) step 2 of sample D, when 150 mT magnetic induction was applied.</p> Full article ">Figure 6
<p>Macrostructures of solidified samples during step four of the solidification experiments (M = 50×); (<b>A</b>) step 4 of sample A, when 40 mT magnetic induction was applied. (<b>B</b>) step 4 of sample B, when 80 mT magnetic induction was applied. (<b>C</b>) step 4 of sample C, when 120 mT magnetic induction was applied.</p> Full article ">Figure 7
<p>The volume fraction of the eutectic structure and primary phases (primary aluminum and silicon) during the solidification experiments steps with changing magnetic induction B.</p> Full article ">Figure 8
<p>Si-concentration distribution along the diameter in C sample at steps 1 and 2.</p> Full article ">Figure 9
<p>Si-concentration distribution along the diameter in C sample at steps 3 and 4.</p> Full article ">Figure 10
<p>Solidification mechanism with and without the presence of RMF, (<b>a</b>) Eutectic composition alloy without RMF, (<b>b</b>) Eutectic alloys with RMF.</p> Full article ">Figure 11
<p>The micrograph photo of sample C shows (<b>I</b>) a pure irregular Al-Si eutectic structure in step 1 and (<b>II</b>) a solidified structure in step 2 when 100 mT magnetic field is applied, including (<b>II.a</b>) primary aluminum with inter-dendritic eutectic structure at the sample edges. (<b>II.b</b>) a eutectic structure solidified as CTL segregation pattern with sidearm freckles. (<b>II.c</b>) Primary silicon arc-like colony.</p> Full article ">
12 pages, 10933 KiB  
Article
Atomistic Investigation of Titanium Carbide Ti8C5 under Impact Loading
by Kang Xia, Haifei Zhan, Jianli Shao, Jiaqiu Wang, Zhuoqun Zheng, Xinjie Zhang and Zhiyong Li
Metals 2022, 12(11), 1989; https://doi.org/10.3390/met12111989 - 20 Nov 2022
Cited by 1 | Viewed by 1786
Abstract
Titanium carbides attract attention from both academic and industry fields because of their intriguing mechanical properties and proven potential as appealing candidates in the variety of fields such as nanomechanics, nanoelectronics, energy storage and oil/water separation devices. A recent study revealed that the [...] Read more.
Titanium carbides attract attention from both academic and industry fields because of their intriguing mechanical properties and proven potential as appealing candidates in the variety of fields such as nanomechanics, nanoelectronics, energy storage and oil/water separation devices. A recent study revealed that the presence of Ti8C5 not only improves the impact strength of composites as coatings, but also possesses significant strengthening performance as an interlayer material in composites by forming strong bonding between different matrices, which sheds light on the design of impact protection composite materials. To further investigate the impact resistance and strengthening mechanism of Ti8C5, a pilot Molecular Dynamics (MD) study utilizing comb3 potential is carried out on a Ti8C5 nanosheet by subjecting it to hypervelocity impacts. The deformation behaviour of Ti8C5 and the related impact resist mechanisms are assessed in this research. At a low impact velocity ~0.5 km/s, the main resonance frequency of Ti8C5 is 11.9 GHz and its low Q factor (111.9) indicates a decent energy damping capability, which would eliminate the received energy in an interfacial reflection process and weaken the shock waves for Ti8C5 strengthened composites. As the impact velocity increases above the threshold of 1.8 km/s, Ti8C5 demonstrates brittle behaviour, which is signified by its insignificant out-of-plane deformation prior to crack initiation. When tracking atomic Von Mises stress distribution, the elastic wave propagation velocity of Ti8C5 is calculated to be 5.34 and 5.90 km/s for X and Y directions, respectively. These figures are inferior compared with graphene and copper, which indicate slower energy delocalization rates and thus less energy dissipation via deformation is expected prior to bond break. However, because of its relatively small mass density comparing with copper, Ti8C5 presents superior specific penetration. This study provides a fundamental understanding of the deformation and penetration mechanisms of titanium carbide nanosheets under impact, which is crucial in order to facilitate emerging impact protection applications for titanium carbide-related composites. Full article
(This article belongs to the Special Issue Deformation and Fracture of Condensed Materials in Extreme Conditions)
►▼ Show Figures

Figure 1

Figure 1
<p>Impact simulation setup for the Ti<sub>8</sub>C<sub>5</sub> model, (010) with the plane of the Ti<sub>8</sub>C<sub>5</sub> facing the projectile. (<b>a</b>) Top view of the schematic setup. Insert shows the unit cell thickness of the (001) plane Ti<sub>8</sub>C<sub>5</sub>, and atoms labelled in silver and bule are carbon and titanium, respectively; (<b>b</b>) Side view of the schematic setup; the thickness of the model is equal to the thickness of the Ti<sub>8</sub>C<sub>5</sub> unit cell; (<b>c</b>) Energy change in projectile as a function of time for impact velocity of 2 km/s. ∆<span class="html-italic">E<sub>ball,ke</sub></span> and ∆<span class="html-italic">E<sub>ball,pe</sub></span> represent the kinetic and potential energy change of the projectile, respectively. The green balls represent the diamond projectile at a simulation time of 0, 2 and 3.9 ps. Significant deformation is not observed in the projectile during the impact process.</p> Full article ">Figure 2
<p>Von Mises stress distribution of nanosheets under impact velocity of 3km/s. (<b>a</b>–<b>c</b>) Atomic configuration of Ti<sub>8</sub>C<sub>5</sub>: (<b>a</b>) Von Mises stress distribution pattern at a simulation time of 1.3 ps; the insert represents stress accumulation and formation of cracks in the impact region; (<b>b</b>) Stress redistribution/propagation at a simulation time of 2.3 ps; (<b>c</b>) Final atomic configuration at 4.0 ps. (<b>d</b>–<b>f</b>) Atomic configuration of copper: (<b>d</b>) Von Mises stress distribution pattern at a simulation time of 2.0 ps; (<b>e</b>) Stress propagation phase at 3.0 ps; (<b>f</b>) final atomic configuration at 6.0 ps.</p> Full article ">Figure 3
<p>Impact deformation and discrete debris of nanosheets under various impact velocities. Amount of debris after impact for (<b>a</b>,<b>b</b>) Ti<sub>8</sub>C<sub>5</sub>, (<b>c</b>,<b>d</b>) Copper and (<b>e</b>,<b>f</b>) Graphene are presented, respectively.</p> Full article ">Figure 4
<p>Z-direction oscillation of adjacent carbon and titanium atoms located at the highlighted position in <a href="#metals-12-01989-f001" class="html-fig">Figure 1</a>a of the Ti<sub>8</sub>C<sub>5</sub> nanosheet under impact velocities of (<b>a</b>) 3 km/s and (<b>b</b>) 6 km/s; z-direction oscillation of 2 carbon atoms in graphene nanosheets (similar position as the Ti<sub>8</sub>C<sub>5</sub>) under impact velocities of (<b>c</b>) 3 km/s and (<b>d</b>) 6 km/s, respectively; z-direction oscillation of 2 copper atoms (similar position as the Ti<sub>8</sub>C<sub>5</sub>) under impact velocities of (<b>e</b>) 3 km/s and (<b>f</b>) 6 km/s, respectively.</p> Full article ">Figure 5
<p>Oscillation of external energy of nanosheet after the separation of projectile and nanosheet. (<b>a</b>) External energy variation of Ti<sub>8</sub>C<sub>5</sub> over time period of 600 ps and (<b>b</b>) The corresponding frequency spectrum of Ti<sub>8</sub>C<sub>5</sub> over time period of 600 ps.</p> Full article ">Figure 6
<p>Von Mises atomic stress distribution of Ti<sub>8</sub>C<sub>5</sub> prior to bond break, when it subject to impact velocity: (<b>a</b>) 3 km/s; (<b>b</b>) 4 km/s; (<b>c</b>) 5 km/s; (<b>d</b>) 6 km/s. (<b>e</b>) Cumulative density function (CDF) of the Von Mises atomic stress distribution before the crack initiation for different nanosheets under diverse impact velocities.</p> Full article ">Figure 7
<p>Stress distribution diagram for (<b>a</b>) copper and (<b>b</b>) graphene subject to an impact velocity of 3 km/s; (<b>c</b>) Cumulative density function (CDF) of the Von Mises atomic stress distribution before the crack initiation for different nanosheets under an impact velocity of 3km/s (red, blue and brown lines represent copper, graphene and Ti<sub>8</sub>C<sub>5</sub>, respectively).</p> Full article ">Figure 8
<p>Performance of nanosheets’ impact with projectiles of various velocity amplitudes. (<b>a</b>) Penetration energy and (<b>b</b>) Specific penetration energy as a function of impact velocity for Ti<sub>8</sub>C<sub>5</sub>, graphene and Copper nanosheets.</p> Full article ">Figure 9
<p>Energy breakdown for various nanosheets at the moment of perforation. Figures (<b>a</b>–<b>c</b>) represent the kinetic energy gain (<math display="inline"><semantics> <mrow> <mo>Δ</mo> <mi>K</mi> <msub> <mi>e</mi> <mrow> <mi>N</mi> <mi>S</mi> </mrow> </msub> </mrow> </semantics></math>) and the potential energy gain (<math display="inline"><semantics> <mrow> <mo>Δ</mo> <mi>P</mi> <msub> <mi>e</mi> <mrow> <mi>N</mi> <mi>S</mi> </mrow> </msub> </mrow> </semantics></math>) for Ti<sub>8</sub>C<sub>5</sub>, copper and the graphene nanosheet, respectively.</p> Full article ">
18 pages, 3151 KiB  
Review
Research Progress of Magnetic Field Regulated Mechanical Property of Solid Metal Materials
by Yujun Hu, Hongjin Zhao, Xuede Yu, Junwei Li, Bing Zhang and Taotao Li
Metals 2022, 12(11), 1988; https://doi.org/10.3390/met12111988 - 20 Nov 2022
Cited by 9 | Viewed by 1870
Abstract
During the material preparation process, the magnetic field can act with high intensity energy on the material without contact and affect its microstructure and properties. This non-contact processing method, which can change the microstructure and properties of material without affecting the shape and [...] Read more.
During the material preparation process, the magnetic field can act with high intensity energy on the material without contact and affect its microstructure and properties. This non-contact processing method, which can change the microstructure and properties of material without affecting the shape and size of products, has become an important technical means to develop new materials and optimize the properties of materials. It has been widely used in scientific research and industrial production. In recent years, the magnetic field assisted processing of difficult-to-deform materials or improving the performance of complex and precision parts has been rapidly and widely concerned by scholars at home and abroad. This paper reviews the research progress of magnetic field regulating the microstructure, and properties of solid metal materials. The effects of magnetic field-assisted heat treatment, magnetic field assisted stretching, and magnetic field independent treatment on the microstructure and properties of solid metal materials are introduced. The mechanism of the magnetic field effect on the properties of metal materials is summarized, and future research on the magnetic field effect on solid metal has been prospected. Full article
(This article belongs to the Special Issue Sustainable Manufacturing of Metallic Materials and Structures: Design, Processing and Characterization)
►▼ Show Figures

Figure 1

Figure 1
<p>The properties and microstructure of Ti–6Al–4V by heat treatment under the magnetic field. Reprinted with permission from Ref. [<a href="#B11-metals-12-01988" class="html-bibr">11</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 2
<p>The microhardness of beryllium bronze alloys by aging under magnetic field: (<b>a</b>) Cu–0.5 wt % Be; (<b>b</b>) Cu–1.0 wt % Be; (<b>c</b>) Cu–1.6 wt % Be; (<b>d</b>) Cu–1.9 wt % Be-0.33 wt % Ni. Reprinted with permission from Refs. [<a href="#B16-metals-12-01988" class="html-bibr">16</a>,<a href="#B19-metals-12-01988" class="html-bibr">19</a>]. Copyright 2020 Springer Nature.</p> Full article ">Figure 3
<p>The microhardness of aluminum alloys by aging under magnetic field: (<b>a</b>) Al-based alloys with Zn, Mg, and Cu dopants; (<b>b</b>) Al-based alloys with Cu, Si, Mn, and Mg dopants; (<b>c</b>) Al-based alloys with Si, Cu, and Fe dopants; (<b>d</b>) Al-based alloys with Cu, Mg, and Mn dopants. Data from Ref. [<a href="#B20-metals-12-01988" class="html-bibr">20</a>].</p> Full article ">Figure 4
<p>The tensile properties of 2024 aluminum alloy under magnetic field. Data from Ref. [<a href="#B23-metals-12-01988" class="html-bibr">23</a>].</p> Full article ">Figure 5
<p>The tensile properties of 7075 aluminum alloy under magnetic field. Data from Ref. [<a href="#B25-metals-12-01988" class="html-bibr">25</a>].</p> Full article ">Figure 6
<p>The tensile properties of 7055 aluminum alloy under magnetic field. Data from Ref. [<a href="#B26-metals-12-01988" class="html-bibr">26</a>].</p> Full article ">Figure 7
<p>The tensile properties of Ti-6.2Al-4.1V-2.3Zr titanium alloy under magnetic field, Data from Ref. [<a href="#B28-metals-12-01988" class="html-bibr">28</a>].</p> Full article ">Figure 8
<p>The dynamic recrystallization process of alloy under magnetic field.</p> Full article ">Figure 9
<p>TEM images of AZ31 magnesium alloy’s precipitated phase under pulsed magnetic field treatment. (<b>a</b>) 0 T; (<b>b</b>) B = 1 T; (<b>c</b>) B = 3 T; (<b>d</b>) B = 5 T. Reprinted with permission from Ref [<a href="#B52-metals-12-01988" class="html-bibr">52</a>].</p> Full article ">Figure 10
<p>TEM images of aluminum alloy precipitates are treated by alternating magnetic field. (<b>a</b>) untreated, (<b>b</b>) treated. Reprinted with permission from Ref. [<a href="#B33-metals-12-01988" class="html-bibr">33</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 11
<p>TEM images of NAB alloy precipitates are treated by alternating magnetic field (<b>a</b>) untreated, (<b>b</b>) treated. Reprinted with permission from Ref. [<a href="#B33-metals-12-01988" class="html-bibr">33</a>]. Copyright 2021 Elsevier.</p> Full article ">
17 pages, 8876 KiB  
Article
Structure and Properties of Al-Co-Cr-Fe-Ni High-Entropy Alloy Subjected to Electron–Ion Plasma Treatment
by Yurii Ivanov, Victor Gromov, Sergey Konovalov, Vladimir Shugurov, Mikhail Efimov, Anton Teresov, Elizaveta Petrikova, Irina Panchenko and Yulia Shliarova
Metals 2022, 12(11), 1987; https://doi.org/10.3390/met12111987 - 20 Nov 2022
Cited by 1 | Viewed by 1972
Abstract
High-entropy alloys (HEAs) are a new class of materials consisting of at least five elements in equiatomic or near-equiatomic ratio. HEAs are subjected to various types of surface treatment to improve their properties. One of the most promising methods of surface hardening is [...] Read more.
High-entropy alloys (HEAs) are a new class of materials consisting of at least five elements in equiatomic or near-equiatomic ratio. HEAs are subjected to various types of surface treatment to improve their properties. One of the most promising methods of surface hardening is electron beam processing. This study aims to examine the structure, elemental, and phase composition of the AlCrFeCoNi HEA surface layer after the deposition of a (B + Cr) film and irradiation with a pulsed electron beam. HEA samples of non-equiatomic composition (33.4 Al; 8.3 Cr; 17.1 Fe; 5.4 Co; 35.7 Ni, at. %), fabricated by wire-arc additive manufacturing (WAAM), were used as study objects. Modification of the HEA surface layer was carried out by a complex method combining deposition of (B + Cr) film samples on the surface and irradiation with a pulsed electron beam in an argon medium. The mode of modification was identified. It makes it possible to increase microhardness (almost two times) and wear resistance (more than five times), reduce the friction coefficient of the HEA surface layer by 1.3 times due to the decrease in the average grain size, formation of particles of borides and oxyborides of complex elemental composition, the introduction of boron atoms into the crystal lattice of HEA. Full article
(This article belongs to the Special Issue Additive Technologies, Advanced Joining Technology and Study of Weld Joints)
►▼ Show Figures

Figure 1

Figure 1
<p>Structure of the additive manufactured HEA.</p> Full article ">Figure 2
<p>Electron microscopic image of the HEA structure (<b>a</b>); (<b>b</b>)—energy spectra of the HEA section shown in (<b>a</b>).</p> Full article ">Figure 3
<p>Electron microscopic image of the HEA structure (<b>a</b>); (<b>b</b>–<b>f</b>)—images of a sample section (<b>a</b>), obtained in the characteristic X-ray radiation of Cr (<b>b</b>), Fe (<b>c</b>), Ni (<b>d</b>), Al (<b>e</b>), Co (<b>f</b>) atoms.</p> Full article ">Figure 4
<p>Results of X-ray microspectral analysis of HEA sample section (<b>a</b>), obtained by the “along the line” method (blue line: Co; yellow line: Cr; green line: Fe; red line: Al; purple line: Ni); (<b>b</b>,<b>c</b>)—distribution of the intensities of characteristic X-ray emission of Al (<b>b</b>) and Cr (<b>c</b>) atoms along the line indicated in (<b>a</b>).</p> Full article ">Figure 5
<p>Fragment of the HEA X-ray diffraction pattern before modification (arrows indicate indices).</p> Full article ">Figure 6
<p>Electron microscopic image of the surface structure of the “film (Cr + B)/(HEA) substrate” system irradiated with a pulsed electron beam at electron beam energy density (J/cm<sup>2</sup>) of 20 (<b>a</b>), 30 (<b>b</b>), 40 (<b>c</b>).</p> Full article ">Figure 7
<p>Electron microscopic image of the surface structure of the “film (Cr + B)/(HEA) substrate” system irradiated with a pulsed electron beam (20 J/cm<sup>2</sup>) (<b>a</b>); (<b>b</b>,<b>c</b>)—images of the sample area (<b>a</b>) obtained in the characteristic X-ray emission of Cr (<b>b</b>) and B (<b>c</b>) atoms.</p> Full article ">Figure 8
<p>Electron microscopic image of the surface structure of the “film (Cr + B)/(HEA) substrate” system irradiated with a pulsed electron beam (30 J/cm<sup>2</sup>) (<b>a</b>); (<b>b</b>–<b>d</b>)—images of the sample area (<b>a</b>), obtained in the characteristic X-ray emission of Cr (<b>b</b>), Al (<b>c</b>) and O (<b>d</b>) atoms.</p> Full article ">Figure 9
<p>Electron microscopic image of the surface structure of the “film (Cr + B)/(HEA) substrate” system irradiated with a pulsed electron beam at electron beam energy density (J/cm<sup>2</sup>) of 30 (<b>a</b>) and 40 (<b>b</b>).</p> Full article ">Figure 10
<p>Fragment of the X-ray diffraction pattern of HEA surface layer modified by irradiating the “film/substrate” system with a pulsed electron beam at E<sub>S</sub> = 30 J/cm<sup>2</sup> (arrows indicate indices) (<b>a</b>); (<b>b</b>)—dependence of the HEA lattice parameter on the energy density of the electron beam.</p> Full article ">Figure 11
<p>Electron microscopic image (STEM analysis method) of a foil section.</p> Full article ">Figure 12
<p>Structure of the surface layer of the “film (B + Cr)/(HEA) substrate” system irradiated with a pulsed electron beam (E<sub>S</sub> = 20 J/cm<sup>2</sup>); (<b>a</b>) crystalline structure of particles; (<b>b</b>) amorphous structure of the layer.</p> Full article ">Figure 13
<p>Image of the surface layer structure of the “film/substrate” system irradiated with a pulsed electron beam (E<sub>S</sub> = 20 J/cm<sup>2</sup>), obtained in the characteristic X-ray radiation of chromium, iron, cobalt, nickel atoms. (<b>a</b>)—Cr; (<b>b</b>)—Fe; (<b>c</b>)—Co; (<b>d</b>)—Ni.</p> Full article ">Figure 14
<p>Electron microscopic image of the structure of HEA surface layer formed as a result of irradiation of the “film/substrate” system with a pulsed electron beam (E<sub>S</sub> = 20 J/cm<sup>2</sup>); (<b>a</b>)—bright-field image; (<b>b</b>)—microelectron diffraction pattern obtained from this section of the foil; (<b>c</b>)—dark-field image obtained in [111]CrB + [131]Cr<sub>3</sub>NiB<sub>6</sub> reflections. In (<b>b</b>), the arrow indicates the reflection in which the dark-field image was obtained (<b>c</b>).</p> Full article ">Figure 15
<p>Electron microscopic image of the structure of HEA surface layer formed as a result of irradiation of the “film/substrate” system with a pulsed electron beam; (<b>a</b>)—foil section bounded by the selector diaphragm of the electron microscope; (<b>b</b>)—microelectron diffraction pattern obtained from area (<b>a</b>). The arrows in (<b>b</b>) indicate the reflections in which the dark-field images shown in <a href="#metals-12-01987-f016" class="html-fig">Figure 16</a> (reflections 1 and 2) and in <a href="#metals-12-01987-f017" class="html-fig">Figure 17</a> (reflections 3 and 4).</p> Full article ">Figure 16
<p>Electron microscopic image of the structure of HEA surface layer formed as a result of irradiation of the “film/substrate” system with a pulsed electron beam; (<b>a</b>)—dark-field image obtained in [211]Cr + [172]Cr2Ni3B6 reflections; (<b>b</b>)—dark-field image obtained in [081]Cr3NiB6 reflection. The reflections in which these dark-field images were obtained are shown in <a href="#metals-12-01987-f015" class="html-fig">Figure 15</a>b (reflections No. 1 and No. 2, respectively).</p> Full article ">Figure 17
<p>Electron microscopic image of the structure of HEA surface layer formed as a result of irradiation of the “film/substrate” system with a pulsed electron beam; (<b>a</b>)—dark-field obtained in [110]CrB + [131] Cr<sub>2</sub>Ni<sub>3</sub>B<sub>6</sub> reflection; (<b>b</b>)—dark-field image obtained in [002]HEA + [211]FeB reflections. The reflections in which these dark-field images were obtained are shown in <a href="#metals-12-01987-f015" class="html-fig">Figure 15</a>b (reflections No. 3 and No. 4, respectively).</p> Full article ">Figure 18
<p>Microhardness dependence of the surface layer of “film/substrate” system on the energy density of the electron beam. The HEA microhardness in the initial state is 4.7 GPa.</p> Full article ">Figure 19
<p>Dependence of the wear parameter (<b>a</b>) and friction coefficient (<b>b</b>) of the surface layer of “film/substrate” system on the energy density of the electron beam. HEA wear parameter in the initial state is 14 × 10<sup>−5</sup> mm<sup>3</sup>/N × m, the friction coefficient is 0.65.</p> Full article ">
13 pages, 3871 KiB  
Article
Experimental Investigation and Numerical Simulation of the Fluidity of A356 Aluminum Alloy
by Hyeon-Sik Bang, Hyeok-In Kwon, Sung-Bean Chung, Dae-Up Kim and Min-Su Kim
Metals 2022, 12(11), 1986; https://doi.org/10.3390/met12111986 - 20 Nov 2022
Cited by 2 | Viewed by 1899
Abstract
The fluidity of A356 aluminum alloy was experimentally determined at the melt temperatures and vacuum degrees by a series of suction fluidity tests. In order to achieve different cooling rates during the test, quartz tubes, as well as stainless steel tubes, were employed [...] Read more.
The fluidity of A356 aluminum alloy was experimentally determined at the melt temperatures and vacuum degrees by a series of suction fluidity tests. In order to achieve different cooling rates during the test, quartz tubes, as well as stainless steel tubes, were employed as the fluidity channels. As the melt temperature increased from 650 to 730 °C, fluidity lengths either linearly increased from 26 to 36 cm or parabolically increased from 13 to 29 cm when quartz tubes or stainless steel tubes were employed, respectively. As the vacuum degree of the fluidity test increased from 0.005 to 0.03 MPa, fluidity increased from 25 to 43 cm in quartz tubes while the smaller increase in fluidity from 20 to 31 cm was observed in stainless steel tubes. Shorter fluidity lengths in stainless steel tubes than those in quartz tubes under the same fluidity measurement condition were due to faster solidification speed confirmed by microstructural analysis. In order to predict the fluidity of the A356 alloy obtained from the suction fluidity tests, a mathematical model was developed based on heat and mass transfer equations coupled with thermodynamic calculations by ChemApp software. The simulation results show good agreement with the fluidity length obtained in the present study. From a series of model calculations, the effects of casting parameters on the fluidity of the A356 melt were discussed. Full article
(This article belongs to the Special Issue Computational Simulation and Numerical Modeling of Metal Refining and Casting Process)
►▼ Show Figures

Figure 1

Figure 1
<p>Schematic diagram of the experimental apparatus of the present suction fluidity test.</p> Full article ">Figure 2
<p>The experimental procedure of the present suction fluidity test.</p> Full article ">Figure 3
<p>Schematic diagram of the calculation process of the present fluidity model.</p> Full article ">Figure 4
<p>The effect of (<b>a</b>) melt temperature and (<b>b</b>) superheat on the fluidity of A356 and A383 alloy.</p> Full article ">Figure 5
<p>The effect of (<b>a</b>) vacuum degree and (<b>b</b>) both vacuum degree and melt temperature on the fluidity of the A356 alloy.</p> Full article ">Figure 6
<p>The effect of tube materials on the fluidity in the A356 alloy with respect to (<b>a</b>) melt temperature and (<b>b</b>) vacuum degree.</p> Full article ">Figure 7
<p>Microstructure of the fluidity samples of the A356 alloy obtained at 650 °C of the melt temperature and 0.01 MPa of the vacuum degree in a quartz tube (<b>left</b>) and a stainless steel tube (<b>right</b>) after air cooling.</p> Full article ">Figure 8
<p>Calculation results of the present fluidity model with different heat transfer coefficients: (<b>a</b>) fluidity length and (<b>b</b>) temperature change at the suction front of the fluidity sample.</p> Full article ">Figure 9
<p>Comparison of the measured fluidity lengths with the calculation result of the present fluidity model at (<b>a</b>) different melt temperatures, (<b>b</b>) different vacuum degrees.</p> Full article ">Figure 10
<p>Prediction of the fluidity of the A356 melts at arbitrary superheat and initial velocity conditions with different heat transfer coefficients: (<b>a</b>) h = 3.8 kW/m<sup>2</sup>·K, (<b>b</b>) h = 5 kW/m<sup>2</sup>·K.</p> Full article ">
11 pages, 3834 KiB  
Article
Structural Phase Transformation of Rail Steel in Compression
by Krestina Aksenova, Victor Gromov, Yurii Ivanov, Rongshan Qin and Ekaterina Vashchuk
Metals 2022, 12(11), 1985; https://doi.org/10.3390/met12111985 - 20 Nov 2022
Cited by 3 | Viewed by 1704
Abstract
The analysis of structure and defective substructure of rail steel in uniaxial compression to a degree of 50% is carried out. It is revealed that cold hardening has a multi-stage character and is accompanied by fragmentations of pearlite grains which is in field [...] Read more.
The analysis of structure and defective substructure of rail steel in uniaxial compression to a degree of 50% is carried out. It is revealed that cold hardening has a multi-stage character and is accompanied by fragmentations of pearlite grains which is in field as the degree of deformation increases and reaches ≈ 0.4 volume of the foil studied at ε = 50%. The fragments being formed in ferrite plates are separated by low-angle boundaries. The average size of the fragmented ferrite decreases from 240 nm at ε = 15% to 200 nm at ε = 50%. Concurrently with the ferrite fragmentation, fragments of cementite are also observed. It is found that the sizes of the cementite fragments are in a range of 15 to 20 nm and depend weakly on the degree of sample deformation. The cementite fragmentation is caused by deformation-induced carbon dissolution and dislocation-induced fracture. The carbon atoms diffuse from cementite crystal to dislocations, which move through an interplanar space to form particles of tertiary cementite at nanoscale (2–4 nm). It is found that the increase in the degree of deformation is accompanied by a decrease in the scalar and an excess dislocation density. A physical interpretation of the observations has been given. Full article
(This article belongs to the Special Issue Deformation of Metals and Alloys: Theory, Simulations and Experiments)
►▼ Show Figures

Figure 1

Figure 1
<p>A stress-strain curve of rail steel in the uniaxial compression (<b>a</b>) and its treatment in coordinates <span class="html-italic">σ = f</span> (<span class="html-italic">ε</span> <sup>0.2</sup>) (<b>b</b>). The dotted lines in (<b>a</b>) indicate the elastic limit <math display="inline"><semantics> <mrow> <msub> <mi>σ</mi> <mn>0</mn> </msub> </mrow> </semantics></math> of the material. Arrows indicate the positions of samples used for studying the structural-phase state of steel on the deformation curve.</p> Full article ">Figure 2
<p>A TEM image of the steel at ε = 50%. (<b>a</b>)–fragmented structure; in (<b>b</b>) arrows indicate the low-angle boundaries being present in the ferrite plates of a pearlite colony.</p> Full article ">Figure 3
<p>A TEM image of the steel deformed structure (ε = 50%); (<b>a</b>)—bright field; (<b>b</b>)—microelectron diffraction pattern obtained from a foil portion whose image is given on (<b>a</b>); (<b>c</b>,<b>d</b>)—dark fields obtained in cementite reflections [121] Fe<sub>3</sub>C (<b>c</b>) and [211] Fe<sub>3</sub>C (<b>d</b>). On (<b>b</b>) arrows indicate reflections of dark fields, 1—(<b>c</b>), 2—(<b>d</b>).</p> Full article ">Figure 4
<p>A TEM image of the steel deformed structure (ε = 50%); (<b>a</b>)—bright field; (<b>b</b>)—dark field obtained in the reflection [012]Fe<sub>3</sub>C+ [110]α-Fe; <b>c</b>—microelectron diffraction pattern. On (<b>c</b>) arrows indicate the reflection of a dark field formation.</p> Full article ">Figure 5
<p>A TEM image of the steel dislocation substructure in the initial state (<b>a</b>) and after the compressive deformation at ε = 50% (<b>b</b>).</p> Full article ">Figure 6
<p>A TEM image of the steel structure (<b>a</b>); (<b>b</b>)—the microelectron diffraction pattern obtained from the given foil portion. The arrows indicate the reflection determined by azimuthal components of the whole disorientation angle of the steel structure.</p> Full article ">Figure 7
<p>A TEM image of the dislocation substructure being formed near cementite particles (indicated by arrows) at ε = 50%.</p> Full article ">Figure 8
<p>A structure of the deformed rail steel (ε = 50%). Arrows indicate the bend extinction contours. (<b>a</b>) 1 μm; (<b>b</b>) 200 nm.</p> Full article ">
11 pages, 3730 KiB  
Article
Selective Separation Recovery of Copper and Arsenic from the Leaching Solution of Copper Soot
by Zhizhao Yang, Yongbing Zhang, Hepeng Zhou, Xianping Luo, Xuekun Tang and Zishuai Liu
Metals 2022, 12(11), 1983; https://doi.org/10.3390/met12111983 - 20 Nov 2022
Cited by 3 | Viewed by 1707
Abstract
Through the main chemical reaction of metal ions and S2−, a new type of sulfide precipitant was first prepared and used to realize the selective separation recovery of copper and arsenic from the leaching solution of copper soot. It is proven [...] Read more.
Through the main chemical reaction of metal ions and S2−, a new type of sulfide precipitant was first prepared and used to realize the selective separation recovery of copper and arsenic from the leaching solution of copper soot. It is proven by experimental results that the prepared sulfide precipitant could realize the efficient separation recovery of copper and arsenic. Indeed, the copper sulfide slag with Cu grade of about 47% and arsenic trisulfide slag with As operation recovery of about 98% could be obtained. The results of chemical reaction energy calculation analysis and SEM images analysis illustrate that the selective separation recovery of copper and arsenic mainly depended on the chemical reactions of sulfide precipitation. The ions of S2− and HS− produced by the prepared sulfide precipitant could react with Cu2+ and arsenic components to form CuS and As2S, respectively, in the copper and arsenic recovery procedure. In addition, the smaller solubility of CuS and the lower rate of S2− engendered by the sulfide precipitant were key to achieving the efficient separation and recovery of copper and arsenic. Full article
(This article belongs to the Topic Green Low-Carbon Technology for Metalliferous Minerals)
►▼ Show Figures

Figure 1

Figure 1
<p>SEM images of the prepared sulfide precipitant. (<b>a</b>) 10,000×; (<b>b</b>) 200,000×.</p> Full article ">Figure 2
<p>Schematic diagram of copper precipitation recovery experiment setup: 1—water bath thermostat, 2—the prepared of sulfide precipitant, 3—thermometer, 4—magnetic rotor and 5—three mouth flask.</p> Full article ">Figure 3
<p>Effect of different molar ratios of sulfide precipitant to copper-on-copper selective recovery.</p> Full article ">Figure 4
<p>Effect of different pH values on copper selective recovery.</p> Full article ">Figure 5
<p>Effect of different reaction times on copper selective recovery.</p> Full article ">Figure 6
<p>Effect of different molar ratios of sulfide precipitant to arsenic-on-arsenic selective recovery.</p> Full article ">Figure 7
<p>Effect of different pH value on arsenic selective recovery.</p> Full article ">Figure 8
<p>Effect of different reaction times on arsenic selective recovery.</p> Full article ">Figure 9
<p>SEM images of copper sulfide slag. (<b>a</b>) 50,000×; (<b>b</b>) 200,000×.</p> Full article ">Figure 10
<p>SEM images of arsenic trisulfide slag. (<b>a</b>) 5000×; (<b>b</b>) 20,000×.</p> Full article ">
18 pages, 6550 KiB  
Article
Application of Optimizing Slab Corner Shapes to Reduce Edge Seam Defect on Heavy Plates
by Minglin Wang, Hui Zhang, Wenbo Zhao, Heping Liu and Xuebing Wang
Metals 2022, 12(11), 1984; https://doi.org/10.3390/met12111984 - 19 Nov 2022
Viewed by 1995
Abstract
The edge seam defect is a common defect in hot rolling heavy plates. It can be improved by optimizing the corner shapes of slabs. Based on a numerical analysis of the effects of the slab corner shape on the temperature distribution after the [...] Read more.
The edge seam defect is a common defect in hot rolling heavy plates. It can be improved by optimizing the corner shapes of slabs. Based on a numerical analysis of the effects of the slab corner shape on the temperature distribution after the slab’s exit from the heating furnace, three rolling methods are proposed for controlling the two-chamfered slab corner shape. The stress and deformation of the corner of the slab during the two-chamfered rolling process are investigated using a numerical simulation. The results show that a two-chamfered shape slab has the smallest temperature drop during the cooling process, and the slab corner can maintain higher temperature and uniformity, which is beneficial for controlling the deformation during the rolling process. Among the three kinds of two-chamfered rolling methods, frontal rolling using a two-roller has the smallest rolling force and rolling resistance to the casting machine, followed by horizontal rolling and then vertical rolling, which has the largest. The favorable slab corner in a two-chamfered shape can be obtained by frontal rolling using a two-roller. Industrial trials confirm that an edge seam defect rate of less than 5% in heavy plates can be achieved under the condition of a large broadside ratio. Full article
(This article belongs to the Special Issue Computational Methods in Metallic Materials Manufacturing Processes)
►▼ Show Figures

Figure 1

Figure 1
<p>Morphology of edge seam defects on heavy plates. (<b>a</b>) Long edge seam crack; (<b>b</b>) intermittent edge seam crack; (<b>c</b>) reticulated edge seam crack.</p> Full article ">Figure 2
<p>Formation and inversion of edge creases.</p> Full article ">Figure 3
<p>Schematic diagram of horizontal rolling at the chamfered slab corner.</p> Full article ">Figure 4
<p>Schematic diagram of vertical rolling at the chamfered slab corner.</p> Full article ">Figure 5
<p>Schematic diagram of frontal rolling using a two-roller at the chamfered slab corner.</p> Full article ">Figure 6
<p>The geometrical dimensions of four corner-shape slabs: (<b>a</b>) RS; (<b>b</b>) OCS; (<b>c</b>) STCS; (<b>d</b>) LTCS.</p> Full article ">Figure 7
<p>The model of different rolling methods. (<b>a</b>) Horizontal rolling; (<b>b</b>) vertical rolling; (<b>c</b>) frontal rolling using a two-roller.</p> Full article ">Figure 8
<p>Enlarged view and the selected nodes along the slab corner before deformation.</p> Full article ">Figure 9
<p>The physical property data of Q345B. (<b>a</b>) Specific heat; (<b>b</b>) thermal conductivity; (<b>c</b>) Young’s modulus; (<b>d</b>) thermal expansion coefficient.</p> Full article ">Figure 10
<p>Temperature distribution of slab corners of different shapes. (<b>a</b>) Right-angle slab; (<b>b</b>) one-chamfered slab; (<b>c</b>) small two-chamfered slab; (<b>d</b>) large two-chamfered slab.</p> Full article ">Figure 11
<p>Temperature distribution at different heights from the narrow face’s center of various shape slabs. (<b>a</b>) Temperature distribution; (<b>b</b>) point designation diagram.</p> Full article ">Figure 12
<p>Deformation of the slab corners under horizontal rolling with a chamfered length of 20 mm.</p> Full article ">Figure 13
<p>Deformation morphology of the slab corners under vertical rolling with a chamfered length of 20 mm.</p> Full article ">Figure 14
<p>Deformation morphology of the slab corners under frontal rolling using a two-roller with a chamfered length of 20 mm.</p> Full article ">Figure 15
<p>Stress on a slab corner under two-chamfered rolling process. (<b>a</b>) Horizontal rolling; (<b>b</b>) vertical rolling; (<b>c</b>) frontal rolling using a two-roller.</p> Full article ">Figure 16
<p>Industrial applications of two-chamfered equipment. (<b>a</b>) Horizontal rolling; (<b>b</b>) frontal rolling with a two-roller.</p> Full article ">Figure 17
<p>A two-chamfered shape slab. (<b>a</b>) Hot slab; (<b>b</b>) cold slab.</p> Full article ">Figure 18
<p>Typical morphology of an edge seam defect on heavy plates produced by the two-chamfered slab under frontal rolling with a two-roller.</p> Full article ">Figure 19
<p>Small cracks on the slope face of a two-chamfered slab under horizontal rolling.</p> Full article ">
18 pages, 87388 KiB  
Article
Microstructure Evolution Behavior of Spray-Deposited 7055 Aluminum Alloy during Hot Deformation
by Di Feng, Rui Xu, Jichen Li, Wenjie Huang, Jingtao Wang, Ying Liu, Linxiang Zhao, Chengbo Li and Hao Zhang
Metals 2022, 12(11), 1982; https://doi.org/10.3390/met12111982 - 19 Nov 2022
Cited by 7 | Viewed by 1765
Abstract
The evolution behaviors of the second phase, substructure and grain of the spray-deposited 7055 aluminum alloy during hot compression at 300~470 °C were studied by scanning electron microscopy (SEM), electron backscatter diffraction (EBSD) and transmission electron microscopy (TEM). Results show that the AlZnMgCu [...] Read more.
The evolution behaviors of the second phase, substructure and grain of the spray-deposited 7055 aluminum alloy during hot compression at 300~470 °C were studied by scanning electron microscopy (SEM), electron backscatter diffraction (EBSD) and transmission electron microscopy (TEM). Results show that the AlZnMgCu phase resulting from the deposition process dissolves gradually with the increase in deformation temperature, but the Al7Cu2Fe phase remains unchanged. The plastic instability of the spray-deposited 7055 aluminum alloy occurs at 470 °C with a 1~5 s−1 strain rate range. Partial dynamic recrystallization (PDRX) adjacent to the original high angle grain boundaries (HAGBs) not only occurs at 300~400 °C with the low strain rates ranging from 0.001 to 0.1 s−1 but also at 450 °C with a high strain rate of 5 s−1. Continuous dynamic recrystallization (CDRX) appears at 450 °C with a low strain rate of 0.001 s−1. The primary nucleation mechanism of PDRX includes the rotation of the subgrain adjacent to the original HAGBs and the subgrain boundary migration. The homogeneous misorientation increase in subgrains is the crucial nucleation mechanism of CDRX. At 300~400 °C, the residual coarse particle stimulated (PSN) nucleation can also be observed. Full article
(This article belongs to the Special Issue Aluminum Alloys and Aluminum-Based Matrix Composites)
►▼ Show Figures

Figure 1

Figure 1
<p>True stress-true strain curves under different strain rate and temperature conditions (Data are corrected based on deformation temperature rise effects for 5 s<sup>−1</sup> compressed samples) (<b>a</b>) 0.001 s<sup>−1</sup>; (<b>b</b>) 0.01 s<sup>−1</sup>; (<b>c</b>) 0.1 s<sup>−1</sup>; (<b>d</b>) 1 s<sup>−1</sup>; (<b>e</b>) 5 s<sup>−1</sup>.</p> Full article ">Figure 2
<p>SEM images of 7055 aluminum alloy under spray-deposited state (<b>a</b>), 300 °C/5 s<sup>−1</sup> (<b>b</b>), 300 °C/0.001 s<sup>−1</sup> (<b>c</b>), 400 °C/0.1 s<sup>−1</sup> (<b>d</b>), 450 °C/5 s<sup>−1</sup> (<b>e</b>) and 450 °C/0.001 s<sup>−1</sup> (<b>f</b>).</p> Full article ">Figure 3
<p>Element plane scanning images of the second phase of spray-deposited 7055 aluminum alloy under hot compression conditions of 450 °C/0.001 s<sup>−1</sup> (<b>a</b>), 450 °C/5 s<sup>−1</sup> (<b>b</b>), 400 °C/0.1 s<sup>−1</sup> (<b>c</b>), 300 °C/0.001 s<sup>−1</sup> (<b>d</b>) and 300 °C/5 s<sup>−1</sup> (<b>e</b>).</p> Full article ">Figure 3 Cont.
<p>Element plane scanning images of the second phase of spray-deposited 7055 aluminum alloy under hot compression conditions of 450 °C/0.001 s<sup>−1</sup> (<b>a</b>), 450 °C/5 s<sup>−1</sup> (<b>b</b>), 400 °C/0.1 s<sup>−1</sup> (<b>c</b>), 300 °C/0.001 s<sup>−1</sup> (<b>d</b>) and 300 °C/5 s<sup>−1</sup> (<b>e</b>).</p> Full article ">Figure 4
<p>The XRD patterns of the spray-deposited specimen and the one deformed under 450 °C/0.001 s<sup>−1</sup>, respectively.</p> Full article ">Figure 5
<p>EBSD images of grain morphology of spray-deposited 7055 aluminum alloy (<b>a</b>) and under hot compression conditions of 300 °C/5 s<sup>−1</sup> (<b>b</b>), 300 °C/0.001 s<sup>−1</sup> (<b>c</b>), 400 °C/0.1 s<sup>−1</sup> (<b>d</b>), 450 °C/5 s<sup>−1</sup> (<b>e</b>) and 450 °C/0.001 s<sup>−1</sup> (<b>f</b>) (Black line represents the high angle grain boundaries (HAGBs) which are higher than 15°, and the red line represents the low angle grain boundaries (LAGBs) ranging from 2° to 15°).</p> Full article ">Figure 6
<p>TEM images of the sub-microstructure of spray-deposited 7055 aluminum alloy under hot compression conditions of 300 °C/5 s<sup>−1</sup> (<b>a</b>), 300 °C/0.001 s<sup>−1</sup> (<b>b</b>), 400 °C/0.1 s<sup>−1</sup> (<b>c</b>), 450 °C/5 s<sup>−1</sup> (<b>d</b>) and 450 °C/0.001 s<sup>−1</sup> (<b>e</b>).</p> Full article ">Figure 7
<p>TEM images of interaction between dislocation and second phase of spray-deposited 7055 aluminum alloy under hot compression conditions of 300 °C/5 s<sup>−1</sup> (<b>a</b>), 300 °C/0.001 s<sup>−1</sup> (<b>b</b>), 400 °C/0.1 s<sup>−1</sup> (<b>c</b>), 450 °C/5 s<sup>−1</sup> (<b>d</b>) and 450 °C/0.001 s<sup>−1</sup> (<b>e</b>,<b>f</b>); P<sub>i</sub> (i = 1~6) represents the particles that remained after hot compressed.</p> Full article ">Figure 8
<p>TEM images of DRV microstructure of spray-deposited 7055 aluminum alloy under hot compression conditions of 300 °C/5 s<sup>−1</sup> (<b>a</b>), 300 °C/0.001 s<sup>−1</sup> (<b>b</b>), 400 °C/0.1 s<sup>−1</sup> (<b>c</b>), 450 °C/5 s<sup>−1</sup> (<b>d</b>) and 450 °C/0.001 s<sup>−1</sup> (<b>e</b>).</p> Full article ">Figure 9
<p>TEM images of DRX microstructure of spray-deposited 7055 aluminum alloy under hot compression conditions of 300 °C/0.001 s<sup>−1</sup> (<b>a</b>), 400 °C/0.1 s<sup>−1</sup> (<b>b</b>), 450 °C/5 s<sup>−1</sup> (<b>c</b>) and 450 °C/0.001 s<sup>−1</sup> (<b>d</b>).</p> Full article ">Figure 10
<p>IPF images of the deformed microstructure of spray deposited 7055 aluminum alloy under hot compression conditions of 300 °C/0.001 s<sup>−1</sup> (<b>a</b>), 400 °C/0.1 s<sup>−1</sup> (<b>b</b>), 450 °C/5 s<sup>−1</sup> (<b>c</b>) and 450 °C/0.001 s<sup>−1</sup> (<b>d</b>). (<b>e</b>) Representation of the color code used to identify the crystallographic orientation on a standard stereographic (Arrows 1, 2, 3 and 4 in each image are the misorientation cumulative routes).</p> Full article ">Figure 11
<p>Misorientation cumulative images between deformed subgrains of spray-deposited 7055 aluminum alloy under hot compression conditions of 300 °C/0.001 s<sup>−1</sup> (<b>a</b>), 400 °C/0.1 s<sup>−1</sup> (<b>b</b>), 450 °C/5 s<sup>−1</sup> (<b>c</b>) and 450 °C/0.001 s<sup>−1</sup> (<b>d</b>). (Curves 1, 2, 3 and 4 in each image are the misorientation cumulative routes in <a href="#metals-12-01982-f010" class="html-fig">Figure 10</a>).</p> Full article ">Figure 12
<p>Diagrammatic sketch of the microstructure evolution of spray-deposited 7055 aluminum alloy. (<b>a</b>) Dynamic recovery and a few second phases resolve. (<b>b</b>) Dynamic recrystallization based on the “subgrain rotation”, “subgrain boundary migration” and “PSN” nucleated mechanism. (<b>c</b>) Dynamic recrystallization based on “homogeneous misorientation increase of subgrain”, and a few phases exist.</p> Full article ">
16 pages, 6898 KiB  
Article
Impact Abrasive Wear of Cr/W-DLC/DLC Multilayer Films at Various Temperatures
by Wen Zhong, Haoyu Wang, Lei Ma and Changhua Zhang
Metals 2022, 12(11), 1981; https://doi.org/10.3390/met12111981 - 19 Nov 2022
Cited by 5 | Viewed by 1622
Abstract
Diamond-like carbon (DLC) films are widely used in key parts of nuclear reactors as a protective coating. A study on the abrasive wear property of Cr/W-DLC/DLC multilayer films was performed at various temperatures. Results show that the mechanism of impact wear under no [...] Read more.
Diamond-like carbon (DLC) films are widely used in key parts of nuclear reactors as a protective coating. A study on the abrasive wear property of Cr/W-DLC/DLC multilayer films was performed at various temperatures. Results show that the mechanism of impact wear under no sand condition is mainly plastic deformation. The multilayer film still has excellent impact wear resistance and favorable adhesion with 308L stainless steel substrate at elevated temperatures under no sand conditions. Sand particles destroy the surface of the multilayer film due to the effect of cutting and ploughing, leading to a nine-fold increase in the wear area. The impact wear mechanism changes into abrasive wear with sand addition. Oxidation wear exists on 308L stainless steel substrate material due to the removal of the multilayer film at high temperatures. More energy is absorbed for plastic deformation and material removal under sand conditions, resulting in lower rebound velocity and peak contact force than under no sand conditions. The temperature leads to the softening of the substrate; thus, the specimens become more prone to plastic deformation and material removal. Full article
(This article belongs to the Special Issue Tribological Properties and Surface Modification of Metallic Materials)
►▼ Show Figures

Figure 1

Figure 1
<p>Schematic of high-temperature impact abrasive wear test equipment: (<b>a</b>) Equipment diagram; (<b>b</b>) physical drawing of equipment; (<b>c</b>) surface topography of sand; (<b>d</b>) energy conversion process.</p> Full article ">Figure 1 Cont.
<p>Schematic of high-temperature impact abrasive wear test equipment: (<b>a</b>) Equipment diagram; (<b>b</b>) physical drawing of equipment; (<b>c</b>) surface topography of sand; (<b>d</b>) energy conversion process.</p> Full article ">Figure 2
<p>SEM micrographs and EDS analysis results of the surface and cross-section of Cr/W-DLC/DLC multilayers: (<b>a</b>) Surface topography; (<b>b</b>) cross-sectional topography; (<b>c</b>) EDS mapping of the cross-section.</p> Full article ">Figure 3
<p>Raman spectra (<b>a</b>) and XRD pattern (<b>b</b>) of the Cr/W-DLC/DLC multilayer films.</p> Full article ">Figure 4
<p>The impact velocity and force–time diagrams for different experimental conditions (impact cycle is 10<sup>4</sup>): (<b>a</b>) No sand, velocity−time, (<b>b</b>) No sand, contact force−time, (<b>c</b>) sand, velocity−time, (<b>d</b>) sand, contact force−time.</p> Full article ">Figure 5
<p>Surface topography of the surface topography of Cr/W-DLC/DLC multilayer films under no sand condition at (<b>a</b>) RT, (<b>b</b>) 100 °C, (<b>c</b>) 300 °C, and (<b>d</b>) 500 °C.</p> Full article ">Figure 6
<p>Surface topography of Cr/W-DLC/DLC multilayer films under the sand condition at (<b>a</b>) RT, (<b>b</b>) 100 °C, (<b>c</b>) 300 °C, and (<b>d</b>) 500 °C.</p> Full article ">Figure 7
<p>Wear scar profiles of Cr/W-DLC/DLC multilayer films under (<b>a</b>) no-sand condition and (<b>b</b>) sand condition at various temperatures.</p> Full article ">Figure 8
<p>Wear scar topography and elemental distribution of Cr/W-DLC/DLC multilayer films at RT under no sand condition: (<b>a</b>–<b>c</b>) wear scar topography; (<b>d</b>) elemental distribution.</p> Full article ">Figure 9
<p>Wear scar topography and elemental distribution of Cr/W-DLC/DLC multilayer films at 500 °C under no-sand condition: (<b>a</b>–<b>c</b>) wear scar topography; (<b>d</b>) elemental distribution.</p> Full article ">Figure 10
<p>Wear scar topography and elemental distribution of Cr/W-DLC/DLC multilayer films at RT under the sand condition: (<b>a</b>–<b>c</b>) wear scar topography; (<b>d</b>) elemental distribution.</p> Full article ">Figure 11
<p>Wear scar topography and elemental distribution of Cr/W-DLC/DLC multilayer films at 500 °C under the sand condition: (<b>a</b>–<b>c</b>) wear scar topography; (<b>d</b>) elemental distribution.</p> Full article ">Figure 12
<p>The impact mechanism of Cr/W-DLC/DLC multilayer under (<b>a</b>) no-sand condition and (<b>b</b>) sand condition.</p> Full article ">Figure 13
<p>Cross-sectional topography of Cr/W-DLC/DLC multilayer films at 500 °C under (<b>a</b>) no-sand condition and (<b>b</b>) sand condition.</p> Full article ">
18 pages, 13341 KiB  
Article
Atmospheric Corrosion Evolution of Carbon Steel AISI 1020 along a Longitude Transect in the Atacama Desert
by Luis Cáceres, Alvaro Soliz and Felipe Galleguillos
Metals 2022, 12(11), 1980; https://doi.org/10.3390/met12111980 - 19 Nov 2022
Cited by 2 | Viewed by 1906
Abstract
Carbon steel AISI 1020 was exposed to environmental conditions along a transect of the Atacama Desert to gather experimental evidence to identify the local atmospheric mechanism that triggers corrosion through a buildup of water layer formation on the metal surface in addition to [...] Read more.
Carbon steel AISI 1020 was exposed to environmental conditions along a transect of the Atacama Desert to gather experimental evidence to identify the local atmospheric mechanism that triggers corrosion through a buildup of water layer formation on the metal surface in addition to corrosion evolution. Coupons initially left in selected sites were periodically collected to determine weight loss and surface attributes by scanning electron microscopy and X-ray diffraction. In addition, meteorological conditions were measured in addition to a fog water collector in one site. During the study period, the predominant conditions were the absence of rain, clear skies, and large daily oscillations in temperature and relative humidity. The evidence indicates a water film formation on a metal surface either from a vertical water flux as fog water droplets and/or by the dew water harvesting mechanism. The uptakes of oxygen and chlorides during the corrosion process were highest in the coastal site P0 and gradually decreased with the increasing distance from the coast. This is attributed to both humidity and saline marine fog intrusion from the coast. The oxide layer evolved to form a compact layer with main constituents of lepidocrocite, goethite, and lesser amounts of akageneite. The corrosion depth can be modelled by a simple power function d=AtB with B < 1, indicating a deceleration process. Full article
►▼ Show Figures

Figure 1

Figure 1
<p>Map of experimental sites.</p> Full article ">Figure 2
<p>Fog water collector positioned nearby the P3 site. A 1 × 1 m<sup>2</sup> polypropylene Raschel mesh was mounted in a metal frame structure. It was equipped with a <span class="html-italic">T</span>/<span class="html-italic">RH</span> Hobo logger sensor (bottom right side) and a Hobo Onset logging rain gauge model RG2-M (bottom left side).</p> Full article ">Figure 3
<p>One-week temperature (<b>a</b>) and relative humidity (<b>b</b>) measurements from ground-level sensors.</p> Full article ">Figure 4
<p>Wind directions and speeds at the coastal station (P0) measured between 22 January 2022 and 24 January 2022.</p> Full article ">Figure 5
<p>Daily average temperature (<b>a</b>) and relative humidity (<b>b</b>) values from 24 October 2021 to 18 June 2022, measurements by ground-level sensors.</p> Full article ">Figure 6
<p>Temperature (<b>a</b>) and relative humidity (<b>b</b>) vertical gradients in 3 different sites expressed as differences between measured values from sensors positioned at the ground and 5 m above the ground, respectively (refer to <a href="#metals-12-01980-t001" class="html-table">Table 1</a>).</p> Full article ">Figure 7
<p>Precipitated fog water events in a fog water collector at site P3 (refer to <a href="#metals-12-01980-t001" class="html-table">Table 1</a>).</p> Full article ">Figure 8
<p>Time of wetness values determined from the measured relative humidity values at all sensors.</p> Full article ">Figure 9
<p>Corrosion rates measured at 32, 74, and 241 days of atmospheric exposure of coupons at the P0, P1, P2, P3, and P4 sites.</p> Full article ">Figure 10
<p>SEM photograph of a polished coupon.</p> Full article ">Figure 11
<p>SEM photographs of carbon steel samples in position P0 at 32 days (<b>a</b>), 74 days (<b>b</b>), 120 days (<b>c</b>), and 241 days (<b>d</b>) of exposure.</p> Full article ">Figure 12
<p>SEM photographs of carbon steel samples in position P1 at 32 days (<b>a</b>), 74 days (<b>b</b>), 120 days (<b>c</b>), and 241 days (<b>d</b>) of exposure.</p> Full article ">Figure 13
<p>SEM photographs of carbon steel samples in position P2 at 32 days (<b>a</b>), 74 days (<b>b</b>), 120 days (<b>c</b>), and 241 days (<b>d</b>) of exposure.</p> Full article ">Figure 13 Cont.
<p>SEM photographs of carbon steel samples in position P2 at 32 days (<b>a</b>), 74 days (<b>b</b>), 120 days (<b>c</b>), and 241 days (<b>d</b>) of exposure.</p> Full article ">Figure 14
<p>SEM photographs of carbon steel samples in position P3 at 32 days (<b>a</b>), 74 days (<b>b</b>), 120 days (<b>c</b>), and 241 days (<b>d</b>) of exposure.</p> Full article ">Figure 15
<p>SEM photographs of carbon steel samples in position P4 at 32 days (<b>a</b>), 74 days (<b>b</b>), 120 days (<b>c</b>), and 241 days (<b>d</b>) of exposure.</p> Full article ">Figure 16
<p>XRD patterns of carbon steel coupons after 241-day exposure in all sites.</p> Full article ">
15 pages, 6747 KiB  
Article
Structural Optimization, Fabrication, and Corrosion Behaviors of Biodegradable Mg-Nd-Zn-Zr Alloy Hemostatic Clip
by Lin Mao, Xin Zheng, Yongji Tian, Yiling Shi, Xiaochen Zhang and Chengli Song
Metals 2022, 12(11), 1979; https://doi.org/10.3390/met12111979 - 19 Nov 2022
Cited by 3 | Viewed by 1665
Abstract
In this study, the process of ligating blood vessels via biodegradable Mg alloy hemostatic clips with toothless, transverse teeth, and embedded teeth was simulated through finite element analysis (FEA). The results showed that the transverse tooth clip caused the minimum stress (0.81489 MPa) [...] Read more.
In this study, the process of ligating blood vessels via biodegradable Mg alloy hemostatic clips with toothless, transverse teeth, and embedded teeth was simulated through finite element analysis (FEA). The results showed that the transverse tooth clip caused the minimum stress (0.81489 MPa) to blood vessels. Furthermore, the effects of clips with transverse teeth of different parameters, including lower tooth length, tooth height, and tooth pitch, on clamped blood vessels were studied. The numerical simulation results showed that the three optimal parameters for clips with transverse teeth were 0.2, 0.1, and 0.1 mm, respectively. Then, the optimally designed clip based on the Mg–Nd–Zn–Zr alloy was manufactured and evaluated using immersion tests. Results from the corrosion behavior study showed that closed clips (0.118 ± 0.041 mg·cm−2·day−1) corroded slightly faster than open clips (0.094 ± 0.041 mg·cm−2·day−1). Moreover, micromorphological observations showed that no cracks appeared on the closed clips, indicating that the Mg alloy had excellent performance and avoided stress corrosion cracking (SCC). Thus, the new type of Mg alloy clip kept good blood vessel closure during FEA and exhibited no corrosion cracking during the degradation process, making it a promising candidate for applications with biodegradable hemostatic clips. Full article
(This article belongs to the Section Biobased and Biodegradable Metals)
►▼ Show Figures

Figure 1

Figure 1
<p>Structural dimensions of the designed hemostatic clips: (<b>a</b>) NS; (<b>b</b>) TS; (<b>c</b>) FS; (<b>d</b>) Schematic diagram of tooth profile parameters (mm).</p> Full article ">Figure 2
<p>Finite element analysis results of hemostatic clips and blood vessels: (<b>a</b>–<b>c</b>) Deformation cloud images of the clips (NS, TS, and FS); (<b>d</b>) Equivalent plastic strain distribution cloud map of TS clip after occluding the vessel; (<b>e</b>) Strain distribution of the vessel clamped by TS clip; (<b>f</b>) Strain distribution of the vessel clamped by TS clip (cross-sectional view).</p> Full article ">Figure 3
<p>Numerical results of the finite element analysis of the hemostatic clips with different structural designs: (<b>a</b>) Maximum equivalent plastic strains of the Mg alloy clips and maximum equivalent stress of blood vessels; (<b>b</b>) Variation curves of the maximum equivalent plastic strain of the Mg-based clip with the lower tooth length and maximum equivalent stress of blood vessels; (<b>c</b>) Curves of the plastic strain of the clip with tooth height and maximum stress of blood vessels; (<b>d</b>) Curves of the maximum strain of the clip with tooth pitch and maximum stress of blood vessels.</p> Full article ">Figure 4
<p>Mg alloy clips and the phase composition: (<b>a</b>) Image of hemostatic clips manufactured from Mg-Nd-Zn-Zr alloy; (<b>b</b>) XRD pattern of Mg-Nd-Zn-Zr alloy.</p> Full article ">Figure 5
<p>Microstructure of the as-extruded Mg-Nd-Zn-Zr alloy clip: (<b>a</b>) Regions for optical observation, (<b>b</b>) Typical microstructure of the Mg-Nd-Zn-Zr alloy; (<b>c</b>) Microstructure of spectrum 1 after the clip is closed; (<b>d</b>) Microstructure of spectrum 2 after the clip is closed.</p> Full article ">Figure 6
<p>Immersion test of Mg-Nd-Zn-Zr alloy clips in artificial plasma for 14 days: (<b>a</b>) Volume of <math display="inline"><semantics> <mrow> <msub> <mi mathvariant="normal">H</mi> <mn>2</mn> </msub> </mrow> </semantics></math> produced by the Mg alloy clips, (<b>b</b>) pH value variations of artificial plasma during immersion of the Mg alloy clips.</p> Full article ">Figure 7
<p>Corrosion behavior of Mg-Nd-Zn-Zr alloy clips in artificial plasma: (<b>a</b>) Macroscopic corrosion morphology of the closed clip; (<b>b</b>) Micromorphology of the top region of the closed clip after 14 days of immersion; (<b>c</b>) Amplified morphology of corrosion products; (<b>d</b>) Surface morphology of the top region of the closed clip after removing the corrosion products, (<b>e</b>) Surface morphology of the tail end of the closed clip after cleaning the corrosion products; (<b>f</b>) Surface morphology of the open clip after cleaning the corrosion products; (<b>g</b>) Corrosion rates of the closed and open clips achieved by weight loss calculation after degradation for 14 days.</p> Full article ">Figure 8
<p>SEM-EDS results of the corrosion products on the surface of Mg-Nd-Zn-Zr alloy clips.</p> Full article ">Figure 9
<p>XPS spectra for the corrosion products on Mg alloy hemostatic clips: (<b>a</b>) XPS surface survey scan; (<b>b</b>) Mg1s; (<b>c</b>) Ca2p; (<b>d</b>) P2p.</p> Full article ">Figure 10
<p>Schematic diagram of biocorrosion of Mg-Nd-Zn-Zr alloy clips: (<b>a</b>) Anodic reaction and cathodic reaction; (<b>b</b>) Formation of Mg hydroxide; (<b>c</b>) Chloride ions destroy the protective layer and formation of apatite; (<b>d</b>) Mg alloy is in the state of dynamic degradation.</p> Full article ">
15 pages, 9698 KiB  
Article
Effects of Simulated PWHT on the Microstructure and Mechanical Properties of 2.25Cr1Mo0.25V Steel for a Hydrogenation Reactor
by Yanmei Li, Yonghao Cui, Jimou Zhang, Minghui Song and Chen Xu
Metals 2022, 12(11), 1978; https://doi.org/10.3390/met12111978 - 19 Nov 2022
Cited by 1 | Viewed by 3056
Abstract
The effect of post-welding heat treatment (PWHT) on the microstructure and mechanical properties of large-thickness 2.25Cr1Mo0.25V steel was investigated through simulated post-welding heat treatment (SPWHT). The results showed that an increase in the SPWHT time decreased the toughness, hardness, and strength of the [...] Read more.
The effect of post-welding heat treatment (PWHT) on the microstructure and mechanical properties of large-thickness 2.25Cr1Mo0.25V steel was investigated through simulated post-welding heat treatment (SPWHT). The results showed that an increase in the SPWHT time decreased the toughness, hardness, and strength of the steel. After Min.SPWHT, the high-temperature tensile strength decreased more significantly, and the damage of Min.SPWHT to the high-temperature tensile strength reached approximately 80% of the Max.SPWHT. The microstructure of the tested steel before and after SPWHT consisted of granular bainite and lath bainite. After SPWHT, intergranular carbides were precipitated as coarsened carbides, carbide clusters, and chains of carbides; alloy element segregation occurred, and the segregation of Mo was the most serious, followed by Cr, and V. The precipitation behavior of the carbides and the increase in the effective grain size caused by the widening of the bainite–ferrite lath worked together and resulted in the decline of the impact toughness; the reduction in the solid solution and precipitation strengthening effects were the main factors in the strength reduction of the tested steel. In the high-temperature tensile tests, defects first appeared around the coarse carbides and carbide clusters. Controlling the size of the intergranular large-size carbides and the degree of cluster precipitation in the NT state structure may be a means of obtaining higher strength of the base metal subjected to PWHT. Full article
(This article belongs to the Special Issue Advanced Technology in Microalloyed Steels)
►▼ Show Figures

Figure 1

Figure 1
<p>The sampling position and analysis method of the sample.</p> Full article ">Figure 2
<p>(<b>a</b>) Hardness, (<b>b</b>) −30 °C impact energy, (<b>c</b>) yield strength, and (<b>d</b>) tensile strength of 2.25Cr1Mo0.25V steel before and after SPWHT.</p> Full article ">Figure 3
<p>(<b>a</b>) Metallographic picture, (<b>b</b>) low-magnification picture, and (<b>c</b>) high-magnification picture of the microstructure of the 2.25Cr1Mo0.25V steel in the NT state (as-receive state).</p> Full article ">Figure 4
<p>(<b>a</b>) Metallographic picture, (<b>b</b>) low-magnification picture, and (<b>c</b>) high-magnification picture of the microstructure of the 2.25Cr1Mo0.25V steel in Min.SPWHT; (<b>d</b>) Metallographic picture, (<b>e</b>) low-magnification picture, and (<b>f</b>) high-magnification picture of the microstructure of the 2.25Cr1Mo0.25V steel in Max.SPWHT.</p> Full article ">Figure 5
<p>TEM images of the microstructure in the precipitates of the 2.25Cr1Mo0.25V steel in the (<b>a</b>–<b>c</b>) NT state and (<b>d</b>–<b>f</b>) Max.SPWHT state.</p> Full article ">Figure 6
<p>Average size change of carbide in the 2.25Cr1Mo0.25V steel after Max.SPWHT.</p> Full article ">Figure 7
<p>Distribution of C, Cr, Mo, and V in the 2.25Cr1Mo0.25V steel in the (<b>a1</b>–<b>a5</b>) NT, (<b>b1</b>–<b>b5</b>) Min.SPWHT, and (<b>c1</b>–<b>c5</b>) Max.SPWHT states by EPMA.</p> Full article ">Figure 8
<p>EDS spectra of spectrum 1 and 8 marked in <a href="#metals-12-01978-f009" class="html-fig">Figure 9</a>.</p> Full article ">Figure 9
<p>The position of the measured point in the (<b>a</b>) NT specimen, (<b>b</b>) Min.SPWHT specimen, and (<b>c</b>) Max.SPWHT specimen in the EDS test.</p> Full article ">Figure 10
<p>(<b>a</b>) The picture of the 550 °C tensile specimen subjected to Max.SPWHT; (<b>b</b>) the reduction rate of the area of the specimen after the high-temperature tension test.</p> Full article ">Figure 11
<p>The micro-morphology of the fracture of the tensile specimens of the 2.25Cr1Mo0.25V steel in the (<b>a</b>–<b>c</b>) NT, (<b>d</b>–<b>f</b>) Min.SPWHT, and (<b>g</b>–<b>i</b>) Max.SPWHT states at 450–550 °C.</p> Full article ">Figure 12
<p>SEM images of the necking area of the (<b>a</b>) 450 °C, (<b>b</b>) 500 °C, and (<b>c</b>) 550 °C tensile samples. All samples were subjected to Max.SPWHT.</p> Full article ">Figure 13
<p>TEM images showing (<b>a</b>) many carbides forming on the dislocation lines, (<b>b</b>) the large-sized carbides at the grain boundary in the necking area of the 450 °C tensile samples, (<b>c</b>,<b>d</b>) carbides in the necking area of the 500 °C tensile samples, and (<b>e</b>,<b>f</b>) carbides in necking area of the 550 °C tensile samples. All samples were subjected to Max.SPWHT.</p> Full article ">
14 pages, 3744 KiB  
Article
Evaluation Method and Application of Cold Rolled Strip Flatness Quality Based on Multi-Objective Decision-Making
by Qiuna Wang, Jingdong Li, Xiaochen Wang, Quan Yang and Zedong Wu
Metals 2022, 12(11), 1977; https://doi.org/10.3390/met12111977 - 19 Nov 2022
Cited by 4 | Viewed by 1662
Abstract
Flatness is a vital quality index that determines the dimensional accuracy of the cold-rolled strip. This paper designs a local shape wave extraction algorithm and a fuzzy classification algorithm for overall flatness defect classification based on cosine distance. By introducing the small displacement [...] Read more.
Flatness is a vital quality index that determines the dimensional accuracy of the cold-rolled strip. This paper designs a local shape wave extraction algorithm and a fuzzy classification algorithm for overall flatness defect classification based on cosine distance. By introducing the small displacement buckling theory of thin plates, the plate stress buckling model of overall and local shape waves is studied, and the critical buckling elongation difference of the overall shape and the local shape under the given conditions are obtained. Finally, using the multi-objective decision-making evaluation method, a comprehensive evaluation model of the flatness quality is established. The model is applied to the actual cold rolling production. The on-site flatness data are used to verify the flatness quality determination model both locally and overall. The results show that the model can accurately identify the local and overall flatness defects of cold-rolled strips, realizes the accurate identification and evaluation of the cold-rolled flatness quality, and provides strong support for the optimization of rolling process parameters and the improvement of the quality of thin strip products. Full article
(This article belongs to the Section Computation and Simulation on Metals)
►▼ Show Figures

Figure 1

Figure 1
<p>Common symmetrical flatness defects. (<b>a</b>) Center waves; (<b>b</b>) double-sided waves.</p> Full article ">Figure 2
<p>The diagram to determine the number of valid channels for the flatness detection roll.</p> Full article ">Figure 3
<p>Diagram of the flatness data reading results.</p> Full article ">Figure 4
<p>Local wave shapes on site.</p> Full article ">Figure 5
<p>Principle of the external local wave shape defect recognition and extraction.</p> Full article ">Figure 6
<p>Principle of the internal local wave shape defect recognition and extraction.</p> Full article ">Figure 7
<p>The flow chart of flatness pattern recognition calculation. (<span class="html-italic">PP<sub>k</sub></span> is is the result of calculating two mutually inverted flatness defects according to Equation (19), <span class="html-italic">k</span> = 1,2,3,4).</p> Full article ">Figure 8
<p>The flow chart of flatness eigenvalues.</p> Full article ">Figure 9
<p>The interface of the flatness quality evaluation system.</p> Full article ">
18 pages, 1731 KiB  
Article
Optimization of Process Parameters for Powder Bed Fusion Additive Manufacturing Using a Linear Programming Method: A Conceptual Framework
by Alexander Khaimovich, Andrey Balyakin, Maxim Oleynik, Artem Meshkov and Vitaly Smelov
Metals 2022, 12(11), 1976; https://doi.org/10.3390/met12111976 - 19 Nov 2022
Cited by 6 | Viewed by 2654
Abstract
At present, the selection of optimal technological parameters for laser powder bed fusion (LPBF) is determined by the requirements of the fusion process. The main parameters that are commonly varied include laser power (P), scanning speed (v), hatch spacing [...] Read more.
At present, the selection of optimal technological parameters for laser powder bed fusion (LPBF) is determined by the requirements of the fusion process. The main parameters that are commonly varied include laser power (P), scanning speed (v), hatch spacing (h), and layer thickness (t). The productivity of the LPBF process (the increment in the fused volume of the material) is equal to the product of the last three parameters, and the mechanical properties are largely determined by the volumetric fusion energy density, which is equal to the ratio of laser power to productivity. While ensuring maximum process productivity, it is possible to obtain acceptable quality characteristics—mechanical properties, surface roughness, etc.—for a certain range of LPBF technological parameters. In these cases, several quality characteristics act as constraints on the optimization process, and productivity and the key quality characteristics become components of the objective function. Therefore, this article proposes a formalized representation of the optimization problem for the LPBF process, including the derivation of the objective function with the constraint matrix, and provides a solution to the problem using the linear programming (LP) method. The advantages of the proposed method include the guaranteed convergence of the solution with an unlimited number of constraints; the disadvantage concerns the adequacy of the solution, which is limited by a relatively narrow range of parameter changes. The proposed method was tested in determining the optimal LPBF parameters for an HN58MBYu powder LP model that included 13 constraints and an objective function with two target parameters. The obtained results made it possible to increase the productivity by 15% relative to the basic technological parameters. Full article
(This article belongs to the Section Additive Manufacturing)
►▼ Show Figures

Figure 1

Figure 1
<p>Samples from heat-resistant HN58MBYu alloy (<b>a</b>); Selective Laser Melting System<sup>®</sup> 280HL (<b>b</b>).</p> Full article ">Figure 2
<p>Span diagram for the mechanical properties of LPBF specimens grouped by their thickness.</p> Full article ">Figure 3
<p>The observed and predicted values of the relative elongation.</p> Full article ">Figure 4
<p>The observed and predicted yield strength values.</p> Full article ">Figure 5
<p>Determination of the minimum depth of the melting pool <span class="html-italic">с</span> was calculated as the distance from the layer surface to the heat-affected zone: (<b>a</b>) LPBF single track (<span class="html-italic">P</span> = 200 W, <span class="html-italic">v</span> = 500 mm/s), (<b>b</b>) LPBF single track (<span class="html-italic">P</span> = 200 W, <span class="html-italic">v</span> = 750 mm/s).</p> Full article ">
10 pages, 3541 KiB  
Article
Effects of Vacancy Defects on Electrical and Optical Properties of ZnO/WSe2 Heterostructure: First-Principles Study
by Xi Yong, Ao Wang, Lichuan Deng, Xiaolong Zhou and Jintao Li
Metals 2022, 12(11), 1975; https://doi.org/10.3390/met12111975 - 18 Nov 2022
Cited by 1 | Viewed by 1568
Abstract
In this work, based on the first principles calculation of density functional theory (DFT), we studied the band structure changes of monolayer ZnO and ZnO/WSe2 before and after vacancy generation, and systematically studied the vacancy formation energy, band structure, density of states, [...] Read more.
In this work, based on the first principles calculation of density functional theory (DFT), we studied the band structure changes of monolayer ZnO and ZnO/WSe2 before and after vacancy generation, and systematically studied the vacancy formation energy, band structure, density of states, electronic density difference and optical properties of ZnO/WSe2 heterostructure before and after vacancy generation. The results show that the band structures of ZnO, WSe2, and ZnO/WSe2 heterostructure are changed after the formation of Zn, O, W, and Se vacancies. The bandgap of the ZnO/WSe2 heterostructure can be effectively controlled, the transition from direct to indirect bandgap semiconductor will occur, and the heterostructure will show metallic properties. The optical properties of heterostructure have also changed significantly, and the absorption capacity of heterostructure to infrared light has been greatly increased with red shift and blue shift respectively. The generation of vacancy changes the electrical and optical properties of ZnO/WSe2 heterostructure, which provides a feasible strategy for adjusting the photoelectric properties of two-dimensional optoelectronic nano devices and has good potential and broad application prospects. Full article
(This article belongs to the Special Issue Preparation, Properties, Computational Simulations of Precious and Rare Metal Materials and Their Compounds)
►▼ Show Figures

Figure 1

Figure 1
<p>(<b>a</b>) Schematic diagram of possible vacancies in monolayer ZnO (<b>b</b>) Schematic diagram of possible vacancies in monolayer WSe<sub>2</sub> (<b>c</b>) Schematic diagram of possible vacancies in heterojunction ZnO/WSe<sub>2</sub>. V<sub>Zn</sub>, V<sub>O</sub>, V<sub>W,</sub> and V<sub>Se</sub> represent the possible vacancy positions of Zn, O, W and Se.</p> Full article ">Figure 2
<p>The band structure of ZnO, ZnO−V<sub>O</sub> and ZnO−V<sub>Zn</sub>. Fermi levels are represented by dashed lines, and the conduction band minimum (CBM) and valence band maximum (VBM) are marked by black stars.</p> Full article ">Figure 3
<p>The band structure of WSe<sub>2</sub>, WSe<sub>2</sub>−V<sub>W,</sub> and WSe<sub>2</sub>−V<sub>Se</sub>. Fermi levels are represented by dashed lines, and the CBM and VBM are marked by black stars.</p> Full article ">Figure 4
<p>The band structure of ZnO/WSe<sub>2</sub> heterostructure, ZnO/WSe<sub>2</sub>−V<sub>Zn</sub>, ZnO/WSe<sub>2</sub>−V<sub>O</sub>, ZnO/WSe<sub>2</sub>−V<sub>W</sub>, and ZnO/WSe<sub>2</sub>−V<sub>Se</sub>. Fermi levels are represented by dashed lines, and the CBM and VBM are marked by black stars.</p> Full article ">Figure 5
<p>The partial density of states (PDOS) of ZnO/WSe<sub>2</sub> heterostructure, ZnO/WSe<sub>2</sub>-V<sub>Zn</sub>, ZnO/WSe<sub>2</sub>-V<sub>O</sub>, ZnO/WSe<sub>2</sub>-V<sub>W</sub>, and ZnO/WSe<sub>2</sub>-V<sub>Se</sub>. (<b>a</b>) ZnO/WSe<sub>2</sub> heterostructure; (<b>b</b>) ZnO/WSe<sub>2</sub>-V<sub>Zn</sub>; (<b>c</b>) ZnO/WSe<sub>2</sub>-V<sub>O</sub>; (<b>d</b>) ZnO/WSe<sub>2</sub>-V<sub>W</sub>; (<b>e</b>) ZnO/WSe<sub>2</sub>-V<sub>Se</sub>.</p> Full article ">Figure 6
<p>The electron differential density of after ZnO/WSe<sub>2</sub> heterostructure, ZnO/WSe<sub>2</sub>−V<sub>Zn</sub>, ZnO/WSe<sub>2</sub>−V<sub>O</sub>, ZnO/WSe<sub>2</sub>−V<sub>W</sub>, and ZnO/WSe<sub>2</sub>−V<sub>Se</sub>. The red represents charge accumulation and the blue represents charge depletion. (<b>a</b>) ZnO/WSe<sub>2</sub> heterostructure; (<b>b</b>) ZnO/WSe<sub>2</sub>−V<sub>Zn</sub> heterostructure; (<b>c</b>) ZnO/WSe<sub>2</sub>−V<sub>O</sub> heterostructure; (<b>d</b>) ZnO/WSe<sub>2</sub>−V<sub>W</sub> heterostructure; (<b>e</b>) ZnO/WSe<sub>2</sub>−V<sub>Se</sub> heterostructure.</p> Full article ">Figure 6 Cont.
<p>The electron differential density of after ZnO/WSe<sub>2</sub> heterostructure, ZnO/WSe<sub>2</sub>−V<sub>Zn</sub>, ZnO/WSe<sub>2</sub>−V<sub>O</sub>, ZnO/WSe<sub>2</sub>−V<sub>W</sub>, and ZnO/WSe<sub>2</sub>−V<sub>Se</sub>. The red represents charge accumulation and the blue represents charge depletion. (<b>a</b>) ZnO/WSe<sub>2</sub> heterostructure; (<b>b</b>) ZnO/WSe<sub>2</sub>−V<sub>Zn</sub> heterostructure; (<b>c</b>) ZnO/WSe<sub>2</sub>−V<sub>O</sub> heterostructure; (<b>d</b>) ZnO/WSe<sub>2</sub>−V<sub>W</sub> heterostructure; (<b>e</b>) ZnO/WSe<sub>2</sub>−V<sub>Se</sub> heterostructure.</p> Full article ">Figure 7
<p>(<b>a</b>) The imaginary part of the dielectric function and (<b>b</b>) the light absorption coefficient of ZnO/WSe<sub>2</sub> heterostructure, ZnO/WSe<sub>2</sub>−V<sub>Zn</sub>, ZnO/WSe<sub>2</sub>−V<sub>O</sub>, ZnO/WSe<sub>2</sub>−V<sub>W</sub>, and ZnO/WSe<sub>2</sub>−V<sub>Se</sub>.</p> Full article ">
12 pages, 3536 KiB  
Article
Tin Removal from Tin-Bearing Iron Concentrate with a Roasting in an Atmosphere of SO2 and CO
by Lei Li, Zhipeng Xu and Shiding Wang
Metals 2022, 12(11), 1974; https://doi.org/10.3390/met12111974 - 18 Nov 2022
Cited by 1 | Viewed by 1260
Abstract
The tin could be volatilized and removed effectively from the tin-bearing iron concentrate while roasted in an atmosphere of SO2 and CO. The reduction of SO2 by CO occurred in preference to the SnO2 and Fe3O4, [...] Read more.
The tin could be volatilized and removed effectively from the tin-bearing iron concentrate while roasted in an atmosphere of SO2 and CO. The reduction of SO2 by CO occurred in preference to the SnO2 and Fe3O4, and the generated S2 could sulfurize the SnO2 to an evaporable SnS, which resulted in the tin volatilization. However, the Fe3O4 could be sulfurized simultaneously, and a phase of iron sulfide was formed, retaining in the roasted iron concentrate. It decreased the quality of the iron concentrate. In addition, the formation of Sn-Fe alloy was accelerated as the roasting temperature exceeded 1100 °C, which decreased the Sn removal ratio. An appropriate SO2 partial pressure and roasting temperature should be controlled. Under the condition of the roasting temperature of 1050 °C, SO2 partial pressure of 0.003, CO partial pressure of 0.85, and residence time of 60 min, the tin content in the roasted iron concentrate was decreased to 0.032 wt.% and the sulfur residual content was only 0.062 wt.%, which meets the standard of iron concentrate for BF ironmaking. Full article
(This article belongs to the Special Issue Metal Recovery and Separation from Wastes)
►▼ Show Figures

Figure 1

Figure 1
<p>XRD pattern (<b>a</b>) and EPMA analysis (<b>b</b>) of the Sn-bearing iron concentrate.</p> Full article ">Figure 2
<p>Schematic illustration of the experimental apparatus. (1-Mass flow meter; 2-Gas mixer; 3-Pressure gauge; 4-Filter; 5-Resistive heater; 6-Corundum reactor; 7-Temperature controller).</p> Full article ">Figure 3
<p>(<b>a</b>) Equilibrium phase composition of 2 mol CO+ 1 mol Fe<sub>3</sub>O<sub>4</sub> + 1 mol SnO<sub>2</sub> roasted with different amounts of SO<sub>2</sub> at 1100 °C; (<b>b</b>) Gibbs free energy changes for reaction (1)–(5) at 500–1100 °C; (<b>c</b>,<b>d</b>) Equilibrium phase composition of 1.5 mol SO<sub>2</sub>+ 1 mol Fe<sub>3</sub>O<sub>4</sub> + 1 mol SnO<sub>2</sub> roasted with different amounts of CO at 1100 °C.</p> Full article ">Figure 4
<p>(<b>a</b>) The changes of CO partial pressure (P<sub>CO</sub>) with the increase of SO<sub>2</sub> partial pressure (P<sub>SO2</sub>); (<b>b</b>) The predominance area diagram of Fe−Sn−S−O at 1000 °C ; (<b>c</b>) Effects of SO<sub>2</sub> partial pressure on the Sn volatilization ratio from the tin-bearing iron concentrate; (<b>d</b>) The S content in the roasted iron concentrate under different SO<sub>2</sub> partial pressure.</p> Full article ">Figure 5
<p>(<b>a</b>) XRD patterns of the roasted residue at the SO<sub>2</sub> partial pressure of 0.001 and 0.003 respectively; (<b>b</b>) SEM-EDS result of the roasted residue at the SO<sub>2</sub> partial pressure of 0.003.</p> Full article ">Figure 6
<p>Effects of roasting temperature on the Sn volatilization ratio (<b>a</b>) from the tin-bearing iron concentrate and S content in the roasted iron concentrate (<b>b</b>).</p> Full article ">Figure 7
<p>XRD results of the roasted residues at 1050 °C (<b>a</b>) and 1100 °C (<b>b</b>), respectively, for 40 min with the P<sub>SO2</sub> of 0.003; SEM-EDS results of the roasted residues at 1050 °C (<b>c</b>) and 1100 °C (<b>d</b>), respectively, for 40 min with the P<sub>SO2</sub> of 0.003.</p> Full article ">Figure 8
<p>Effect of residence time on Sn volatilization ratio from the tin-bearing iron concentrate (<b>a</b>) and S content in the roasted residue (<b>b</b>).</p> Full article ">Figure 9
<p>(<b>a</b>) Relationship between [1 − 2X/3 − (1 − X) <sup>2/3</sup>] and different residence time; (<b>b</b>) Relationship between [1 − (1 − X)<sup>1/3</sup>] and different residence time; (<b>c</b>) Relationship between [1 + (1 − X)<sup>1/3</sup> 2 (1 − X)<sup>2/3</sup>] and t/ [1 − (1 − X)<sup>1/3</sup>].</p> Full article ">
13 pages, 15034 KiB  
Article
Effect of Solution Heat Treatment on the Porosity Growth of Nickel-Based P/M Superalloys
by Hengyong Bu, Lu Chen and Yonghua Duan
Metals 2022, 12(11), 1973; https://doi.org/10.3390/met12111973 - 18 Nov 2022
Cited by 5 | Viewed by 1765
Abstract
Thermal-induced porosity (TIP) is one of the major defects in powder metallurgy (P/M) superalloys, and it seriously affects the performance of P/M superalloys. The effects of solution heat treatment on the growth of the TIP of the nickel-based P/M superalloy FGH97 were investigated. [...] Read more.
Thermal-induced porosity (TIP) is one of the major defects in powder metallurgy (P/M) superalloys, and it seriously affects the performance of P/M superalloys. The effects of solution heat treatment on the growth of the TIP of the nickel-based P/M superalloy FGH97 were investigated. A series of solution heat treatment tests were carried out at holding temperatures ranging from 1150 to 1200 °C, with holding times ranging from 0.5 to 8 h. The results showed that the holding time, temperature, and the initial volume of porosity are the primary factors influencing porosity growth, and the volume fraction of TIPs increases by increasing the temperature or extending the holding time. The porosity growth models were constructed based on the porosity statistics combined with a nonlinear fitting method. To evaluate the accuracy of the proposed models, the correlation coefficient (R) and average absolute relative error (AARE) were calculated between the predicted and experimental values. The unbiased AARE values were 2.06% and 3.99% for the average value of TIP and the worst value of TIP, respectively, which imply that the proposed porosity growth models have greater accuracy and can be used to illustrate TIP behavior in solution heat treatment. Full article
(This article belongs to the Special Issue Heat Resistant Steels and Alloys)
►▼ Show Figures

Figure 1

Figure 1
<p>Schematic of solution heat treatment tests for the P/M FGH97 superalloy.</p> Full article ">Figure 2
<p>OM images of the as-HIPed P/M FGH97 superalloys, (<b>a</b>) Ni1#, (<b>b</b>) Ni2#, and (<b>c</b>) Ni3#.</p> Full article ">Figure 2 Cont.
<p>OM images of the as-HIPed P/M FGH97 superalloys, (<b>a</b>) Ni1#, (<b>b</b>) Ni2#, and (<b>c</b>) Ni3#.</p> Full article ">Figure 3
<p>SEM morphology of the as-HIPed Ni3# specimen.</p> Full article ">Figure 4
<p>Initial TIP values of the as-HIPed P/M FGH97 superalloys.</p> Full article ">Figure 5
<p>OM images of the Ni1# superalloy after solution heat treatment, (<b>a</b>) 1150 °C × 0.5 h, (<b>b</b>) 1150 °C × 2 h, (<b>c</b>) 1150 °C × 4 h, (<b>d</b>) 1150 °C × 8 h, (<b>e</b>) 1175 °C × 0.5 h, (<b>f</b>) 1175 °C × 2 h, (<b>g</b>) 1175 °C × 4 h, (<b>h</b>) 1175 °C × 8 h, (<b>i</b>) 1200 °C × 0.5 h, (<b>j</b>) 1200 °C × 2 h, (<b>k</b>) 1200 °C × 4 h, and (<b>l</b>) 1200 °C × 8 h.</p> Full article ">Figure 5 Cont.
<p>OM images of the Ni1# superalloy after solution heat treatment, (<b>a</b>) 1150 °C × 0.5 h, (<b>b</b>) 1150 °C × 2 h, (<b>c</b>) 1150 °C × 4 h, (<b>d</b>) 1150 °C × 8 h, (<b>e</b>) 1175 °C × 0.5 h, (<b>f</b>) 1175 °C × 2 h, (<b>g</b>) 1175 °C × 4 h, (<b>h</b>) 1175 °C × 8 h, (<b>i</b>) 1200 °C × 0.5 h, (<b>j</b>) 1200 °C × 2 h, (<b>k</b>) 1200 °C × 4 h, and (<b>l</b>) 1200 °C × 8 h.</p> Full article ">Figure 6
<p>OM images of the Ni2# superalloy after solution heat treatment, (<b>a</b>) 1150 °C × 0.5 h, (<b>b</b>) 1150 °C × 2 h, (<b>c</b>) 1150 °C × 4 h, (<b>d</b>) 1175 °C × 0.5 h, (<b>e</b>) 1175 °C × 2 h, (<b>f</b>) 1175 °C × 4 h, (<b>g</b>) 1200 °C × 0.5 h, (<b>h</b>) 1200 °C × 2 h, and (<b>i</b>) 1200 °C × 4 h.</p> Full article ">Figure 7
<p>OM images of the Ni3# superalloy after solution heat treatment, (<b>a</b>) 1150 °C × 0.5 h, (<b>b</b>) 1150 °C × 2 h, (<b>c</b>) 1150 °C × 4 h, (<b>d</b>) 1150 °C × 8 h, (<b>e</b>) 1175 °C × 0.5 h, (<b>f</b>) 1175 °C × 2 h, (<b>g</b>) 1175 °C × 4 h, (<b>h</b>) 1175 °C × 8 h, (<b>i</b>) 1200 °C × 0.5 h, (<b>j</b>) 1200 °C × 2 h, (<b>k</b>) 1200 °C × 4 h, and (<b>l</b>) 1200 °C × 8 h.</p> Full article ">Figure 8
<p>The shape of the Ni1# specimen before and after 1200 °C × 8 h solution heat treatment.</p> Full article ">Figure 9
<p>Microstructure of the FGH97 superalloy with different volumes of porosity after etching; (<b>a</b>) Ni1#, (<b>b</b>) Ni2#, and (<b>c</b>) Ni3#.</p> Full article ">Figure 10
<p>Evolution of the TIP as a function of holding temperature and holding time; (<b>a</b>) the average value of the TIP and (<b>b</b>) the worst value of the TIP.</p> Full article ">Figure 11
<p>Correlation between the experimental and predicted values of porosity growth models; (<b>a</b>) the average value of the TIP and (<b>b</b>) the worst value of the TIP.</p> Full article ">
15 pages, 6059 KiB  
Article
Microstructure Refinement by a Combination of Heat Treatment and Thixoforming Process Followed by Severe Plastic Deformation and Their Effects on Al-Si Alloy Hardness
by Mohamed Abdelgawad Gebril, Mohd Zaidi Omar, Intan Fadhlina Mohamed, Norinsan Kamil Othman and Osama M. Irfan
Metals 2022, 12(11), 1972; https://doi.org/10.3390/met12111972 - 18 Nov 2022
Cited by 4 | Viewed by 1382
Abstract
This study fabricated a thixoformed Al-7% Si alloy using the cooling slope technique and subjected it to heat treatment before processing with severe plastic deformation to determine the effect of this combination method on the microstructure refinement and hardness of Al-Si alloys (300 [...] Read more.
This study fabricated a thixoformed Al-7% Si alloy using the cooling slope technique and subjected it to heat treatment before processing with severe plastic deformation to determine the effect of this combination method on the microstructure refinement and hardness of Al-Si alloys (300 Series). Each as-cast and thixoformed Al-Si alloy sample was subjected to equal-channel angular pressing (ECAP) and high-pressure torsion (HPT) individually at room temperature before and after heat treatment. ECAP was conducted in a mould with a 120° channel angle via route A, and HPT was applied with 0.75 and 5 turns. The heat-treated thixoformed Al-Si alloy subjected to the HPT process had an ultra-fine grain microstructure and showed a fine and homogeneous redistribution of the eutectic phase in the Al matrix. For the as-cast alloy, the hardness of the heat-treated thixoformed Al-7% Si alloy increased from 63 HV to 124 and 215 Hv after two ECAP passes and five turns of HPT due to the reduced and redistributed eutectic phase in the Al matrix. Subjecting the Al-7% Si alloy to a combination of semisolid and heat treatment processes before subjecting it to severe plastic deformation resulted in microstructural refinement and improved the hardness of the Al-Si alloy. The results indicate that HPT is a more effective method than ECAP for increasing the hardness of the thixoformed Al-Si alloy due to microstructure refinement. Full article
(This article belongs to the Special Issue Advances in Casting, Thermomechanical and Heat Treatments of Aluminum Alloy)
►▼ Show Figures

Figure 1

Figure 1
<p>Schematic of (<b>a</b>) cooling slope casting, (<b>b</b>) ECAP die, and (<b>c</b>) HPT facility.</p> Full article ">Figure 2
<p>Schematic representation of the HPT disc and the location for measuring the disc hardness and the TEM area.</p> Full article ">Figure 3
<p>Optical micrographs of the pre-T6 (<b>a</b>) as-cast and (<b>b</b>) thixoformed samples. SEM micrograph of the thixoformed (<b>c</b>,<b>d</b>) post-T6 Si particles of thixoformed samples.</p> Full article ">Figure 4
<p>Optical micrograph of ECAPed Al-Si alloy samples after two passes. The (<b>a</b>) as-cast, (<b>b</b>) heat-treated as-cast, (<b>c</b>) heat-treated thixoformed, and (<b>d</b>) SEM enlargement of heat-treated thixoformed sample after two passes and (<b>e</b>) cracks on the thixoformed sample surface.</p> Full article ">Figure 5
<p>Optical micrographs at the centre and the edge of (<b>a</b>,<b>b</b>) as-cast and (<b>c</b>,<b>d</b>) thixoformed Al-Si alloy after 0.75 turns.</p> Full article ">Figure 6
<p>Optical micrograph of the centre and the edge of heat-treated (<b>a</b>,<b>b</b>) as-cast and (<b>c</b>,<b>d</b>) thixoformed Al-Si alloy after 0.75 turns.</p> Full article ">Figure 7
<p>Morphological map of the Si particles (green) in the Al matrix (red) before heat treatment at the centre and the edge of the as-cast (<b>a</b>,<b>b</b>) and thixoformed (<b>c</b>,<b>d</b>) after five turns.</p> Full article ">Figure 8
<p>Morphological map of the Si particles (green) in the Al matrix (red) after heat treatment at the centre and the edge of the as-cast (<b>a</b>,<b>b</b>) and thixoformed (<b>c</b>,<b>d</b>) after five turns.</p> Full article ">Figure 9
<p>Mapping of HPTed heat-treated (<b>a</b>) as-cast (<b>b</b>) Al, (<b>c</b>) Si, (<b>d</b>) Mg, (<b>e</b>) Fe, (<b>f</b>) Ti, and (<b>g</b>) Cu in the edge position of Al-Si alloy after five turns.</p> Full article ">Figure 10
<p>Mapping of HPTed heat-treated (<b>a</b>) and thixoformed (<b>b</b>) Al, (<b>c</b>) Si, (<b>d</b>) Mg, (<b>e</b>) Fe, (<b>f</b>) Ti, and (<b>g</b>) Cu in the edge position of Al-Si alloy respectively after five turns.</p> Full article ">Figure 10 Cont.
<p>Mapping of HPTed heat-treated (<b>a</b>) and thixoformed (<b>b</b>) Al, (<b>c</b>) Si, (<b>d</b>) Mg, (<b>e</b>) Fe, (<b>f</b>) Ti, and (<b>g</b>) Cu in the edge position of Al-Si alloy respectively after five turns.</p> Full article ">Figure 11
<p>TEM image of subgrains in HPTed heat-treated thixoformed sample after five turns.</p> Full article ">Figure 12
<p>TEM image of a heat-treated thixoformed sample after five turns of HPT.</p> Full article ">Figure 13
<p>Microhardness of the Al-Si alloy (<b>a</b>) after ECAP, (<b>b</b>) 0.75 turns and (<b>c</b>) 5 turns of HPT process.</p> Full article ">
11 pages, 4611 KiB  
Article
Comparison of Submillimeter Spot Ablation of Copper and Nickel by Multipulse Picosecond and Femtosecond Laser
by Mingyu Li, Jifei Ye, Lan Li, Bangdeng Du, Ying Wang, Heyan Gao and Chenghao Yu
Metals 2022, 12(11), 1971; https://doi.org/10.3390/met12111971 - 18 Nov 2022
Viewed by 1473
Abstract
The current transmission and reflection laser ablation micropropulsion modes have the problem of a complex working medium supply system in engineering. Therefore, we propose large-spot laser ablation with a one-dimensional supply mode. In order to verify this ablation mode, a multipulse ablation experiment [...] Read more.
The current transmission and reflection laser ablation micropropulsion modes have the problem of a complex working medium supply system in engineering. Therefore, we propose large-spot laser ablation with a one-dimensional supply mode. In order to verify this ablation mode, a multipulse ablation experiment of submillimeter-scale light spots was carried out on the surface of pretreated copper and nickel under the atmosphere using an ultrafast laser with a pulse width of 290 fs and 10 ps. The results show that femtosecond laser multipulse ablation (FLMA) leads to the grain refinement of copper, the crater quality of the two metals under FLMA is better, and picosecond laser multipulse ablation (PLMA) causes the crater of nickel to form a dense remelting bulge that affects laser absorption; both metals have obvious heat-affected zones after FLMA and PLMA, the heat-affected zones of nickel are 5–10% larger than those of copper, and the ablation depth of copper is deeper. Under the same conditions, the ablation mass of copper is smaller than that of nickel, and the specific impulse performance of laser ablation micropropulsion is better. Full article
(This article belongs to the Special Issue Laser Materials Processing Technology)
►▼ Show Figures

Figure 1

Figure 1
<p>Schematic diagram of the laser ablation experimental setup.</p> Full article ">Figure 2
<p>Surface ablation morphologies obtained by Laser confocal microscopy for Ni and Cu specimens under different pulse energies and number of pulses for: (<b>a</b>) PLMA; (<b>b</b>) FLMA.</p> Full article ">Figure 3
<p>3D profiles obtained by laser confocal microscopy for Ni and Cu specimens under different pulse energies and number of pulses for: (<b>a</b>) PLMA; (<b>b</b>) FLMA.</p> Full article ">Figure 4
<p>Surface ablation morphologies obtained by SEM for Ni and Cu specimens under different pulse widths and same number of pulses (1000) for: (<b>a</b>) 100 μJ; (<b>b</b>) 150 μJ; (<b>c</b>) 200 μJ.</p> Full article ">Figure 5
<p>Surface ablation morphology of copper after laser irradiation with energy of 200 μJ, pulse width of 290 fs and pulse number of 1000 ((<b>a</b>) is SEM image, (<b>b</b>) is EDS analysis of oxygen element, (<b>c</b>) is optical microscopy photo).</p> Full article ">Figure 6
<p>Regional variations in surface morphology of laser ablation of copper and nickel with different energies and pulse widths for 1000 pulse numbers.</p> Full article ">Figure 7
<p>Height profile of copper and nickel after laser irradiation with different pulse widths of energy 200 <math display="inline"><semantics> <mrow> <mi mathvariant="sans-serif">μ</mi> <mi mathvariant="normal">J</mi> </mrow> </semantics></math> and pulse number of 1000. ((<b>a</b>,<b>b</b>,<b>c</b>,<b>d</b>) are the surface ablative morphologies of copper and nickel craters; (<b>a1</b>,<b>b1</b>,<b>c1</b>,<b>d1</b>) are 2D ablative height profiles of copper and nickel craters; (<b>a2</b>,<b>b2</b>,<b>c2</b>,<b>d2</b>) are 3D ablative height profiles of copper and nickel craters).</p> Full article ">Figure 8
<p>Variation trends of copper and nickel ablation depth under the action of ps and fs lasers.</p> Full article ">Figure 9
<p>Variation trends in ablation mass of copper and nickel.</p> Full article ">
11 pages, 1835 KiB  
Article
Microstructure, Mechanical Properties, and Fish-Scaling Resistance of a Ti-Nb Microalloyed Hot-Rolled Enamel Steel
by Yi Zhang, Bo Yu, Jian Zhang, Yu Du, Xiaonan Wang, Hongyan Wu, Xiuhua Gao and Linxiu Du
Metals 2022, 12(11), 1970; https://doi.org/10.3390/met12111970 - 18 Nov 2022
Cited by 2 | Viewed by 1542
Abstract
Currently, the fish-scaling resistance of most hot-rolled enamel steels is improved by adding Ti to form fine TiC carbides as hydrogen traps. Given that the hydrogen capture capacity of NbC is higher than that of TiC, the manufacture of hot-rolled enamel steels via [...] Read more.
Currently, the fish-scaling resistance of most hot-rolled enamel steels is improved by adding Ti to form fine TiC carbides as hydrogen traps. Given that the hydrogen capture capacity of NbC is higher than that of TiC, the manufacture of hot-rolled enamel steels via Ti-Nb microalloying has a promising future. In the present study, a Ti-Nb microalloyed hot-rolled enamel steel was developed, and its microstructure, mechanical properties, and fish-scaling resistance were studied by optical microscopy, transmission electron microscopy, tensile test, and hydrogen permeation test. The results show that the microstructure of hot-rolled experimental steel is composed of ferrite and fine carbides, with a large number of fine precipitates uniformly distributed in the ferrite grains. After the first and second enamel firings, the average sizes of ferrite grain and precipitates gradually increase, the yield strength decreases from 711 ± 9 MPa to 471 ± 17 MPa and 409 ± 8 MPa, the tensile strength decreases from 761 ± 7 MPa to 524 ± 15 MPa and 490 ± 12 MPa, and the elongation increases from 21.0 ± 2.8% to 27.8 ± 1.8% and 33.9 ± 1.1%. The hydrogen permeation value (TH value) decreases from 35.9 min/mm2 to 6.8 min/mm2 and 3.9 min/mm2 after the first and second enamel firings. That is, the fish-scaling resistance of hot-rolled experimental steel is significantly reduced after enamel firing, which is caused by the coarsening of precipitates, resulting in a significant reduction in the density of irreversible hydrogen traps (from 1.21 × 1025 cm−3 to 6.50 × 1023 cm−3 and 4.27 × 1023 cm−3). A large amount of semi-coherent precipitates is the key to obtaining the good fish-scaling resistance of hot-rolled enamel steel. Full article
(This article belongs to the Special Issue Advanced Technology in Microalloyed Steels)
►▼ Show Figures

Figure 1

Figure 1
<p>OM microstructures of different samples: (<b>a</b>) HR sample; (<b>b</b>) EF1 sample; (<b>c</b>) EF2 sample.</p> Full article ">Figure 1 Cont.
<p>OM microstructures of different samples: (<b>a</b>) HR sample; (<b>b</b>) EF1 sample; (<b>c</b>) EF2 sample.</p> Full article ">Figure 2
<p>TEM images of precipitates in (<b>a</b>) HR sample; (<b>b</b>) EF1 sample; (<b>c</b>) EF2 sample.</p> Full article ">Figure 3
<p>Size distribution of the precipitates in (<b>a</b>) HR sample; (<b>b</b>) EF1 sample; (<b>c</b>) EF2 sample.</p> Full article ">Figure 4
<p>Current–time curves of hydrogen permeation test: (<b>a</b>) HR sample; (<b>b</b>) EF1 sample; (<b>c</b>) EF2 sample.</p> Full article ">Figure 5
<p>Charge–time curves of hydrogen permeation test: (<b>a</b>) HR sample; (<b>b</b>) EF1 sample; (<b>c</b>) EF2 sample.</p> Full article ">Figure 6
<p>Relationship between <span class="html-italic">TH</span>, <span class="html-italic">D</span><sub>eff</sub>, and <span class="html-italic">N</span><sub>T</sub>, <span class="html-italic">N</span><sub>ir</sub>: (<b>a</b>) <span class="html-italic">TH</span>–<span class="html-italic">N</span><sub>T</sub> curve; (<b>b</b>) <span class="html-italic">TH</span>–<span class="html-italic">N</span><sub>ir</sub> curve; (<b>c</b>) <span class="html-italic">D</span><sub>eff</sub>–<span class="html-italic">N</span><sub>T</sub> curve; (<b>d</b>) <span class="html-italic">D</span><sub>eff</sub>–<span class="html-italic">N</span><sub>ir</sub> curve.</p> Full article ">
11 pages, 2273 KiB  
Article
Kinetics of Bainite Transformation in Multiphase High Carbon Low-Silicon Steel with and without Pre-Existing Martensite
by Zeinab Babasafari, Alexey V. Pan, Farshid Pahlevani, Suk Chun Moon, Madeleine Du Toit and Rian Dippenaar
Metals 2022, 12(11), 1969; https://doi.org/10.3390/met12111969 - 18 Nov 2022
Cited by 2 | Viewed by 1910
Abstract
In the present study, the isothermal decomposition of austenite to bainite in 1.0 wt% carbon, 0.21% silicon steel during the partitioning step of a quenching and partitioning (Q&P) heat treatment has been investigated in a dilatometer in the temperature range of 200 to [...] Read more.
In the present study, the isothermal decomposition of austenite to bainite in 1.0 wt% carbon, 0.21% silicon steel during the partitioning step of a quenching and partitioning (Q&P) heat treatment has been investigated in a dilatometer in the temperature range of 200 to 350 °C and compared to conventional austempering heat treatment. The bainite transformation was shortened by about 75% in the presence of pre-existing martensite (QP). The kinetics of bainite transformation is described by the well-known Avrami equation. The calculated parameter ‘n’ in the Avrami equation shows that bainite forms in the absence of pre-existing martensite (TT) at a constant nucleate rate, while in the presence of pre-existing martensite, nucleation is interface controlled. The overall bainite transformation activation energy, calculated by the Avrami equation, ranges from 64 to 110 kJ/mol. The outcomes of this investigation provide guidelines for the development of multiphase microstructures, including pre-existing martensite and bainite in high-carbon low-silicon steel, within an industrially acceptable time scale and mechanical performance. Full article
(This article belongs to the Topic Energy-Saving and Emission Reduction in Metallurgy)
►▼ Show Figures

Figure 1

Figure 1
<p>Schematic view of the different heat treatment schedules (temperature vs. time) used in this work.</p> Full article ">Figure 2
<p>Dilatation during bainitic formation in one-step (TT) and two-step (QP) heat treatments: (<b>a</b>) 200 °C; (<b>b</b>) 250 °C; (<b>c</b>) 300 °C; (<b>d</b>) 350 °C.</p> Full article ">Figure 3
<p>Bainite transformation ratio: (<b>a</b>) quenching partitioning; (<b>b</b>) isothermal.</p> Full article ">Figure 4
<p>Hardness measurement at different tempering temperatures.</p> Full article ">Figure 5
<p>Linear fitting of log (−ln (1 − x(t)) and log(t): (<b>a</b>) QP (<b>b</b>) TT.</p> Full article ">Figure 6
<p>Arrhenius plot of ln(t<sub>i</sub>) versus 1/T used to determine the activation energy: (<b>a</b>) QP; (<b>b</b>) TT.</p> Full article ">
22 pages, 13796 KiB  
Article
Effects of Build Orientations on Microstructure Evolution, Porosity Formation, and Mechanical Performance of Selective Laser Melted 17-4 PH Stainless Steel
by Mohammad Azlan Aripin, Zainuddin Sajuri, Nashrah Hani Jamadon, Amir Hossein Baghdadi, Junaidi Syarif, Intan Fadhlina Mohamed and Ahmad Muhammad Aziz
Metals 2022, 12(11), 1968; https://doi.org/10.3390/met12111968 - 17 Nov 2022
Cited by 10 | Viewed by 2619
Abstract
In this study, the effect of phase, microstructure, and porosity in Selective Laser Melting (SLM) on hardness, tensile, and fracture behavior of 17-4 PH was investigated. The increasing interest in SLM in producing complex parts has encouraged the industry to produce performance parts, [...] Read more.
In this study, the effect of phase, microstructure, and porosity in Selective Laser Melting (SLM) on hardness, tensile, and fracture behavior of 17-4 PH was investigated. The increasing interest in SLM in producing complex parts has encouraged the industry to produce performance parts, such as martensitic 17-4 PH stainless steel. However, the microstructure and mechanical behavior of SLM 17-4PH is not fully understood by researchers. Understanding the microstructure profile is complex because it is driven by thermal history and porosity. Both elements vary, based on the build directions, further hindering researchers from fully understanding the mechanical properties. To fabricate specimens in three different building orientations (0°, 45°, and 90°), 17-4 powder was used. Two phases, namely, austenite and martensite, with 90° build direction, retained more austenite, due to the reheating process on a smaller base area. The optical microstructure revealed several elements that were distinct for SLM processing, including circular, columnar lath, wave melt pool, and porosity. Columnar lath was found to grow continuously across different melt pools. Hardness was found to be higher for 0° than for 90°, due to higher martensite content. Tensile strength was highest for 0°, at 958 MPa, higher than at 45° and 90° at 743 and 614 MPa, respectively. Porosity analysis validated that 90° had all three types of porosities and, specifically, the crescent type, which held un-melted powders. All types of porosities were found in fractography analysis. Full article
(This article belongs to the Special Issue Additive Manufacturing of Metallic Materials: Characterization, Properties, and Modeling)
►▼ Show Figures

Figure 1

Figure 1
<p>Possible phases of 17-4 PH shown in red dot according to Schaffler Diagram (reproduced with permission from Elsevier, 2022) [<a href="#B15-metals-12-01968" class="html-bibr">15</a>,<a href="#B16-metals-12-01968" class="html-bibr">16</a>].</p> Full article ">Figure 2
<p>(<b>a</b>) EOSINT machine, (<b>b</b>) schematic diagram of processing parameters, (<b>c</b>) schematic diagram of layers and laser paths, and (<b>d</b>) build orientation of the specimen.</p> Full article ">Figure 3
<p>Powder morphology reveal by SEM observation at (<b>a</b>) 250 and (<b>b</b>) 1000 magnification scale.</p> Full article ">Figure 4
<p>Particle size distribution of the 17-4 PH utilized in this research D<sub>10</sub> = 33 μm, D<sub>50</sub> = 43 μm and D<sub>90</sub> = 52 μm.</p> Full article ">Figure 5
<p>The fraction of liquid and solid phase vs. temperature for 17-4 PH.</p> Full article ">Figure 6
<p>Volume fraction of phases recorded at (<b>a</b>) 0°, (<b>b</b>) 45° and (<b>c</b>) 90°.</p> Full article ">Figure 7
<p>Microstructure profile for 0° build direction.</p> Full article ">Figure 8
<p>Microstructure profile for 90° build direction.</p> Full article ">Figure 9
<p>Microstructure profile for 45° build direction.</p> Full article ">Figure 10
<p>Columnar lath generation acress multiple meltpools.</p> Full article ">Figure 11
<p>Identification of (<b>a</b>,<b>b</b>) austenite phase in build direction and transverse direction and (<b>c</b>,<b>d</b>) martensite phase in build and transverse direction.</p> Full article ">Figure 11 Cont.
<p>Identification of (<b>a</b>,<b>b</b>) austenite phase in build direction and transverse direction and (<b>c</b>,<b>d</b>) martensite phase in build and transverse direction.</p> Full article ">Figure 12
<p>Hardness value for 0° and 90° build direction.</p> Full article ">Figure 13
<p>The types of porosity observed from three different areas of 90° build direction; (<b>a</b>) crescent shape, (<b>b</b>) unmelted powder, (<b>c</b>) keyhole and (<b>d</b>) spherical porosity.</p> Full article ">Figure 14
<p>The types of porosity observed from three different areas of 45° build direction; (<b>a</b>) crescent shape, (<b>b</b>) unmelted powder, (<b>c</b>) keyhole and (<b>d</b>) spherical porosity.</p> Full article ">Figure 15
<p>The type of porosity observed from three different areas of 0° build direction; (<b>a</b>) keyhole and (<b>b</b>) spherical porosity.</p> Full article ">Figure 16
<p>Schematic diagram of 14-7 microstructure produced by SLM method.</p> Full article ">Figure 17
<p>Stress strain curve for all build direction.</p> Full article ">Figure 18
<p>The direction of the melt pool boundary for each build direction respective to the load direction.</p> Full article ">Figure 19
<p>Crack coalescence and propagation for 90°.</p> Full article ">Figure 20
<p>Fracture surface overall under SEM observation including flat and cleavage surface.</p> Full article ">Figure 21
<p>SEM of a fractured surface under 200× and 5000× magnification for (<b>a</b>) cleavage surface and (<b>b</b>) flat surface.</p> Full article ">
10 pages, 2389 KiB  
Article
Effect of Secondary Phases on Multi-Step Phase Transitions and Magnetocaloric Properties in MnFe-Based Alloys
by A-Young Lee, Min-Ha Lee, Song-Yi Kim, JunHee Han, Ki-Hoon Kang and Jong-Woo Kim
Metals 2022, 12(11), 1967; https://doi.org/10.3390/met12111967 - 17 Nov 2022
Viewed by 1559
Abstract
This study investigated the effect of the secondary phases on multi-step phase transitions and the magnetocaloric properties depending on the Ge content in the MnFeCoPSiGe alloys. Two-step phase transitions were observed by the variations of the Fe2P-type hexagonal structure (first-order) and [...] Read more.
This study investigated the effect of the secondary phases on multi-step phase transitions and the magnetocaloric properties depending on the Ge content in the MnFeCoPSiGe alloys. Two-step phase transitions were observed by the variations of the Fe2P-type hexagonal structure (first-order) and secondary phases (second-order). The Curie temperature alters with non-linear behavior consistent with change of the lattice parameters. In addition, the magnetic entropy change decreased with the increase of the Ge content and, subsequently, fractions of the secondary phases. However, the morphological variation of microstructure, distributed as a circular-type shape of the Fe2P-type hexagonal structure in the Ge-rich matrix, increased the magnetic entropy change. Therefore, the addition of Ge enables the control of the Curie temperature to be applicable for high temperature operating devices. The control of the secondary phases and morphology of the microstructure are crucial to improve the phase transition and magnetic entropy change. Full article
(This article belongs to the Special Issue Phase Transition and Magnetic Effect of Magnetic Alloy)
►▼ Show Figures

Figure 1

Figure 1
<p>FE-SEM images and corresponding mapping images of the MnFe<sub>0.99</sub>Co<sub>0.01</sub>P<sub>0.4</sub>Si<sub>0.6-x</sub>Ge<sub>x</sub> alloys, (<b>a</b>) Ge<sub>0.0</sub>, (<b>b</b>) Ge<sub>0.1</sub>, (<b>c</b>) Ge<sub>0.3</sub>, (<b>d</b>) Ge<sub>0.5</sub>, (<b>e</b>) Ge<sub>0.6</sub>. Arrows indicate the rich phase regions. The results are summarized in <a href="#metals-12-01967-t001" class="html-table">Table 1</a>.</p> Full article ">Figure 2
<p>The XRD patterns of the MnFe<sub>0.99</sub>Co<sub>0.01</sub>P<sub>0.4</sub>Si<sub>0.6-x</sub>Ge<sub>x</sub> (x = 0, 0.1, 0.3, 0.5, 0.6) alloys.</p> Full article ">Figure 3
<p>(<b>a</b>) Lattice parameters and unit cell volume, (<b>b</b>) phase fraction of the MnFe<sub>0.99</sub>Co<sub>0.01</sub>P<sub>0.4</sub>Si<sub>0.6-x</sub>Ge<sub>x</sub> (x = 0, 0.1, 0.3, 0.5, 0.6) alloys.</p> Full article ">Figure 4
<p>The magnetocaloric properties of the MnFe<sub>0.99</sub>Co<sub>0.01</sub>P<sub>0.4</sub>Si<sub>0.6-x</sub>Ge<sub>x</sub> alloys; M-T curve under applied magnetic field (0.01 T) and S-T curve at magnetic field <span class="html-italic">ΔH</span> = 2 T of (<b>a</b>,<b>b</b>) Ge<sub>0.0</sub>, (<b>c</b>,<b>d</b>) Ge<sub>0.1</sub> and Ge<sub>0.3</sub>, (<b>e</b>,<b>f</b>) Ge<sub>0.5</sub> and Ge<sub>0.6</sub> alloys.</p> Full article ">
Show export options expand_more Show export options expand_less
Select all
Export citation of selected articles as:
 
Displaying articles 1-222
Previous Issue
Next Issue
Back to TopTop