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Volume 11
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Metals, Volume 11, Issue 12 (December 2021) – 183 articles

Cover Story (view full-size image): The spontaneous deposition of elemental sulfur (S8) caused the serious pitting corrosion of L360 pipelines steel during transport of “sour gas”. The hydrolysis of S8 at low temperature strengthened the environmental difference (especially pH value) under deposition of S8, besides crevice corrosion, resulting in the different corrosion behavior. Based on the combination of the localized corrosion intensity index (LCII) and morphology features, a novel model was proposed to better explain the pitting corrosion mechanisms of L360 pipelines steel under deposition of S8, especially the non-uniform deposition of S8. View this paper
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21 pages, 15302 KiB  
Article
Surface Characteristics and Corrosion Behavior of Carbon Steel Treated by Abrasive Blasting
by Aran Kim, Shigenobu Kainuma and Muye Yang
Metals 2021, 11(12), 2065; https://doi.org/10.3390/met11122065 - 20 Dec 2021
Cited by 20 | Viewed by 4800
Abstract
The effects of blasting with metallic steel grit and non-metallic alumina grit on steel surface characteristics were evaluated. These abrasives are generally used at construction sites and in vacuum blasting. Milled steel specimens were used to investigate the effect of the blasting conditions [...] Read more.
The effects of blasting with metallic steel grit and non-metallic alumina grit on steel surface characteristics were evaluated. These abrasives are generally used at construction sites and in vacuum blasting. Milled steel specimens were used to investigate the effect of the blasting conditions on surface properties. The effect of difference in surface properties on the adhesion strength and corrosion behavior were measured through adhesion tests, polarization curves, and electrochemical impedance spectroscopy. The limitations of blasting were evaluated using corroded steel specimens, as were the effects of corrosion products, salts, and abrasive material remaining on the blasted steel surface on the adhesion and corrosion resistance of paint. Steel grit more effectively increased the surface roughness than alumina grit; however, with both abrasive materials, the roughness increased with the blast projection angle. However, in the case of alumina grit, some abrasive material remained on the surface; thus, the actual roughness not including the residual abrasive material was more complex and greater than that of the sample blasted with steel grit. According to the adhesion strength test of painted and unpainted specimens, the adhesion force improved with increasing surface roughness and residual abrasive materials. Further, surface roughness was linearly correlated with the adhesion strength of unpainted specimens for both abrasive materials with blasting, and the adhesion strength force with alumina grit was approximately 1.4 times higher than that with steel grit, suggesting that increased roughness and residual abrasive material could benefit adhesion. According to the electrochemical test results, lower roughness and increased residual abrasive material owing to alumina grit on the steel surface enhanced the surface corrosion resistance, confirming the benefit of residual materials. Grinding left behind corrosion products and salts under the steel, resulting in the recurrence of rusting. However, the residue from blasting with alumina suppressed corrosion, thus improving the adhesion and corrosion resistance of the paint. Full article
(This article belongs to the Special Issue Influence of Surface Treatment on Corrosion Behavior of Steels)
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Figure 1

Figure 1
<p>Micrograph of abrasive materials: (<b>a</b>) Steel grit; (<b>b</b>) Aluminum grit.</p> Full article ">Figure 2
<p>Blasted surface and color scale according to roughness measured by a laser microscope: (<b>a</b>) MST, (<b>b</b>) MAL.</p> Full article ">Figure 3
<p>SEM-EDX analysis of the surface and cross-section of the specimens subjected to blasting: (<b>a</b>) MST, (<b>b</b>) MAL.</p> Full article ">Figure 4
<p>Fractal dimension (<span class="html-italic">D</span><sub>B</sub>) analysis of roughness from the SEM images of the cross-section of the blasted specimens: (<b>a</b>) MST, (<b>b</b>) MAL.</p> Full article ">Figure 5
<p>The results of adhesion strength.</p> Full article ">Figure 6
<p>Fracture surface of the unpainted steel after the adhesion test (MAL, Blasting angle 90°).</p> Full article ">Figure 7
<p>Relationship between surface roughness (<span class="html-italic">R</span><sub>a</sub>) and adhesion strength.</p> Full article ">Figure 8
<p>Fractal dimension (<span class="html-italic">D</span><sub>B</sub>) analysis shows the mechanical interlock at the interface between the glue and the blasted substrate: (<b>a</b>) MST, (<b>b</b>) MAL.</p> Full article ">Figure 9
<p>Fracture surface of the painted steel after the adhesion test (Blasting angle 90°): (<b>a</b>) MST, (<b>b</b>) MAL.</p> Full article ">Figure 10
<p>Potentiodynamic polarization curves of the blasted specimens in 3.5 wt% NaCl solution: (<b>a</b>) MST, (<b>b</b>) MAL.</p> Full article ">Figure 11
<p>Microscopic images of the surfaces of the blasted specimens before and after immersing in 3.5 wt% NaCl solution: (<b>a</b>) MST, (<b>b</b>) MAL.</p> Full article ">Figure 12
<p>Microscopic images of the cross-section of the blasted specimens after immersing in 3.5 wt% NaCl solution: (<b>a</b>) MST, (<b>b</b>) MAL.</p> Full article ">Figure 13
<p>EIS measurements of the blasted specimens after immersing in 3.5 wt% NaCl solution for 1 h: (<b>a</b>) MST, (<b>b</b>) MAL.</p> Full article ">Figure 14
<p>Equivalent circuits.</p> Full article ">Figure 15
<p>SEM-EDX images on the cross-section of the blasted specimens: (<b>a</b>) CST, (<b>b</b>) CAL.</p> Full article ">Figure 16
<p>Surface morphologies of the blasted surfaces after 0 and 24 h exposure: (<b>a</b>) CST, (<b>b</b>) CAL.</p> Full article ">Figure 17
<p>Fracture surface conditions after the adhesion tests on the unpainted surfaces after 0 and 24 h exposure: (<b>a</b>) CST, (<b>b</b>) CAL.</p> Full article ">Figure 18
<p>Fracture surface conditions after the adhesion tests on the painted surfaces after 0 and 24 h exposure: (<b>a</b>) CST, (<b>b</b>) CAL.</p> Full article ">Figure 19
<p>The results of adhesion strength.</p> Full article ">Figure 20
<p>EIS measurements of blasted specimens after immersing in 3.5 wt% NaCl solution for 24 h showed by Bode plots: (<b>a</b>) log |Z|-log <span class="html-italic">f</span>, (<b>b</b>) Phase angle-log <span class="html-italic">f</span>.</p> Full article ">Figure 21
<p>Surface morphology after immersing in a 3.5 wt% NaCl solution (CST 24 h): (<b>a</b>) CST Microscopic image, (<b>b</b>) Contour image.</p> Full article ">
19 pages, 6990 KiB  
Article
Comparative Multi-Modal, Multi-Scale Residual Stress Evaluation in SLM 3D-Printed Al-Si-Mg Alloy (RS-300) Parts
by Eugene S. Statnik, Fatih Uzun, Svetlana A. Lipovskikh, Yuliya V. Kan, Sviatoslav I. Eleonsky, Vladimir S. Pisarev, Pavel A. Somov, Alexey I. Salimon, Yuliya V. Malakhova, Aleksandr G. Seferyan, Dmitry K. Ryabov and Alexander M. Korsunsky
Metals 2021, 11(12), 2064; https://doi.org/10.3390/met11122064 - 20 Dec 2021
Cited by 13 | Viewed by 3883
Abstract
SLM additive manufacturing has demonstrated great potential for aerospace applications when structural elements of individual design and/or complex shape need to be promptly supplied. 3D-printable AlSi10Mg (RS-300) alloy is widely used for the fabrication of different structures in the aerospace industry. The importance [...] Read more.
SLM additive manufacturing has demonstrated great potential for aerospace applications when structural elements of individual design and/or complex shape need to be promptly supplied. 3D-printable AlSi10Mg (RS-300) alloy is widely used for the fabrication of different structures in the aerospace industry. The importance of the evaluation of residual stresses that arise as a result of the 3D-printing process’ complex thermal history is widely discussed in literature, but systematic assessment remains lacking for their magnitude, spatial distribution, and comparative analysis of different evaluation techniques. In this study, we report the results of a systematic study of residual stresses in 3D-printed double tower shaped samples using several approaches: the contour method, blind hole drilling laser speckle interferometry, X-ray diffraction, and Xe pFIB-DIC micro-ring-core milling analysis. We show that a high level of tensile and compressive residual stresses is inherited from SLM 3D-printing and retained for longer than 6 months. The stresses vary (from −80 to +180 MPa) over a significant proportion of the material yield stress (from −⅓ to ¾). All residual stress evaluation techniques considered returned comparable values of residual stresses, regardless of dramatically different dimensional scales, which ranged from millimeters for the contour method, laser speckle interferometry, and XRD down to small fractions of a mm (70 μm) for Xe pFIB-DIC ring-core drilling. The use of residual stress evaluation is discussed in the context of optimizing printing strategies to enhance mechanical performance and long-term durability. Full article
(This article belongs to the Special Issue Experimental, Modeling and Simulation of Residual Stress in Metallic and Composite Materials)
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Figure 1

Figure 1
<p>The appearance of the double tower shape sample in the as-printed state: (<b>a</b>) front view, (<b>b</b>) side view, (<b>c</b>) dimensions.</p> Full article ">Figure 2
<p>The appearance of the double tower sample (<b>a</b>) before and (<b>b</b>,<b>c</b>) after EDM cutting: (<b>b</b>) top view with the indicated cut plane (red line) and (<b>c</b>) sectional view.</p> Full article ">Figure 3
<p>The appearance of double tower shaped sample as 3D models: (<b>a</b>) before and (<b>b</b>) after EDM cutting in half with the indicated primary coordinate system according to which all used methods are aligned.</p> Full article ">Figure 4
<p>(<b>a</b>) general view of vertical outer face of the half-cut double tower perpendicular to the sectioning plane with the drill hole, and (<b>b</b>,<b>c</b>) interference fringe patterns obtained as a result of 1.9 mm diameter hole drilling: (<b>b</b>) the horizontal in-plane displacement component and (<b>c</b>) the vertical in-plane displacement component.</p> Full article ">Figure 5
<p>Appearance of the polished half-cut double tower sample: (<b>a</b>) mounted in conductive resin for FIB-SEM studies, (<b>b</b>) fixed in the goniometer-like sample holder for X-ray diffraction measurements.</p> Full article ">Figure 6
<p>Illustration of the distribution of the <span class="html-italic">zz</span>-component of residual stresses (<b>a</b>,<b>c</b>) and <span class="html-italic">z</span>-component of displacements (<b>b</b>,<b>d</b>) in the real (<b>a</b>,<b>b</b>) and continuously processed (<b>c</b>,<b>d</b>) geometry models, respectively.</p> Full article ">Figure 7
<p>Illustration of distribution of <span class="html-italic">xx</span>- (<b>a</b>,<b>c</b>) and <span class="html-italic">yy</span>- (<b>b</b>,<b>d</b>) components of residual stresses in the real (<b>a</b>,<b>b</b>) and continuous (<b>c</b>,<b>d</b>) geometry models, respectively.</p> Full article ">Figure 8
<p>Line plots of <span class="html-italic">xx</span> (s11), <span class="html-italic">yy</span> (s22), and <span class="html-italic">zz</span> (s33) components of residual stress along the (<b>a</b>) horizontal (<span class="html-italic">x</span>) and (<b>b</b>) vertical (<span class="html-italic">y</span>) lines illustrated by the dashed lines on the sectioned surface in Figure 12 below.</p> Full article ">Figure 9
<p>Locations of holes for the hole drilling method with the indicated axis.</p> Full article ">Figure 10
<p>Details of X-ray measurement of residual stresses at the cut plane: (<b>a</b>) measured <math display="inline"><semantics> <mrow> <mi>I</mi> <mo>~</mo> <mn>2</mn> <mi>θ</mi> </mrow> </semantics></math> plot for the specimen (highlighted red plane <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <mrow> <mn>420</mn> </mrow> <mo>)</mo> </mrow> </mrow> </semantics></math> was selected for the next calculations); (<b>b</b>) <math display="inline"><semantics> <mrow> <msub> <mi>ε</mi> <mrow> <mi>H</mi> <mi>K</mi> <mi>L</mi> </mrow> </msub> <mo>~</mo> <msup> <mrow> <mi>sin</mi> </mrow> <mn>2</mn> </msup> <mi>φ</mi> </mrow> </semantics></math> plot.</p> Full article ">Figure 11
<p>Details of Xe-pFIB ring-core drilling measurements of residual stresses at the cut plane: (<b>a</b>) combined EBSD map with the Euler’ colors and milled ring, and (<b>b</b>) measured relief strains with the fitting curve.</p> Full article ">Figure 12
<p>Mapping of residual stresses using Xe-pFIB ring-core drilling method.</p> Full article ">Figure 13
<p>Comparison of all used methods: (<b>a</b>) the locations of the measured points and the plotting line, superimposed on the contour map of the horizontal stress component reconstructed by the contour method; (<b>b</b>) residual stress line plot.</p> Full article ">Figure A1
<p>EBSD map (<b>a</b>) before and (<b>b</b>) after scratch elimination using inpainting procedure.</p> Full article ">
11 pages, 51149 KiB  
Article
Fatigue Behavior of Laser-Cut Sheet Metal Parts with Brazed-On Elements
by André Till Zeuner, Robert Kühne, Christiane Standke, David Köberlin, Thomas Wanski, Sebastian Schettler, Uwe Füssel and Martina Zimmermann
Metals 2021, 11(12), 2063; https://doi.org/10.3390/met11122063 - 20 Dec 2021
Cited by 2 | Viewed by 3564
Abstract
Laser cutting is used in the production of formed sheet metal components. However, the cyclic load capacity is reduced compared to other subtractive processes. Laser cutting results in a significant loss of fatigue strength; however, thermal joining has its own effect on the [...] Read more.
Laser cutting is used in the production of formed sheet metal components. However, the cyclic load capacity is reduced compared to other subtractive processes. Laser cutting results in a significant loss of fatigue strength; however, thermal joining has its own effect on the cyclic load capacity. Accordingly, brazing causes a significant reduction in the mechanical strength. However, the open question is what consequences a combination of both processes may have on the overall fatigue strength of sheet metals. Laser-cut samples of AISI 304 with and without a brazed-on element were investigated for their microstructure and mechanical properties. The brazing process was found to have an annealing effect on the microstructure. It was further observed that the fatigue behavior of brazed specimens is dominated by inhomogeneities at the surface of the filler metal fillet located in the geometric notch of the brazed joint. Fatigue strength decreased by almost 50% compared to as-cut specimens. As long as no shared diffusion zone is formed between the laser-cut and the brazed joint, the use of laser cutting for the production of such components appears to be reasonable and does not further contribute to the loss of cyclic strength. Full article
(This article belongs to the Special Issue Technology of Welding and Joining 2021)
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Figure 1

Figure 1
<p>Specimen geometry manufactured by laser cutting, with dimensions in mm.</p> Full article ">Figure 2
<p>Temperature–time curve of the conductive brazing process.</p> Full article ">Figure 3
<p>Geometry of the base specimen with brazed-on local reinforcement, with dimensions in mm.</p> Full article ">Figure 4
<p>Conductive brazing test stand developed at Technical University Dresden.</p> Full article ">Figure 5
<p>Microstructure of specimens investigated in the preliminary tests: (<b>a</b>) as-cut; (<b>b</b>) pre-deformed to a total strain of 28%; (<b>c</b>) pre-deformed to a total strain of 28% and heat-treated in accordance with a conductive brazing temperature regime.</p> Full article ">Figure 5 Cont.
<p>Microstructure of specimens investigated in the preliminary tests: (<b>a</b>) as-cut; (<b>b</b>) pre-deformed to a total strain of 28%; (<b>c</b>) pre-deformed to a total strain of 28% and heat-treated in accordance with a conductive brazing temperature regime.</p> Full article ">Figure 6
<p>Stress–strain curves for different processing conditions from preliminary tests prior to fatigue testing.</p> Full article ">Figure 7
<p>Results of the fatigue tests of as-cut and brazed specimens.</p> Full article ">Figure 8
<p>Fracture surface analysis of an as-cut specimen: (<b>a</b>) overview; (<b>b</b>) crack-initiating site at the burr of the laser-cut edge.</p> Full article ">Figure 9
<p>Fracture surface analysis of a brazed specimen: (<b>a</b>) overview; (<b>b</b>) crack-initiating defect.</p> Full article ">Figure 10
<p>Cross-section of the brazed joint of a failed specimen: (<b>a</b>) overview; (<b>b</b>) filler metal seam where the failure occurred; (<b>c</b>) filler metal seam featuring crack propagation into the base material.</p> Full article ">
14 pages, 5478 KiB  
Article
Hot Wear of Single Phase fcc Materials—Influence of Temperature, Alloy Composition and Stacking Fault Energy
by Aaron Berger, Maximilian Walter, Santiago Manuel Benito and Sebastian Weber
Metals 2021, 11(12), 2062; https://doi.org/10.3390/met11122062 - 20 Dec 2021
Cited by 1 | Viewed by 2557
Abstract
The severe sliding abrasion of single-phase metallic materials is a complex issue with a gaining importance in industrial applications. Different materials with different lattice structures react distinctly to stresses, as the material reaction to wear of counter and base body is mainly determined [...] Read more.
The severe sliding abrasion of single-phase metallic materials is a complex issue with a gaining importance in industrial applications. Different materials with different lattice structures react distinctly to stresses, as the material reaction to wear of counter and base body is mainly determined by the deformation behavior of the base body. For this reason, fcc materials in particular are investigated in this work because, as shown in previous studies, they exhibit better hot wear behavior than bcc materials. In particular, three austenitic steels are investigated, with pure Ni as well as Ni20Cr also being studied as benchmark materials. This allows correlations to be worked out between the hot wear of the material and their microstructural parameters. For this reason, wear tests are carried out, which are analyzed on the basis of the wear characteristics and scratch marks using Electron Backscatter Diffraction. X-ray experiments at elevated temperatures were also carried out to determine the microstructural parameters. It was found that the stacking fault energy, which influences the strain hardening potential, governs the hot wear behavior at elevated temperatures. These correlations can be underlined by analysis of the wear affected cross section, where the investigated materials have shown clear differences. Full article
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Figure 1

Figure 1
<p>Temperature dependent hardness (HV 0.5) of the investigated materials.</p> Full article ">Figure 2
<p>Friction coefficient of the investigated materials.</p> Full article ">Figure 3
<p>Maximal wear path of the investigated materials.</p> Full article ">Figure 4
<p>Scratch energy of the investigated materials.</p> Full article ">Figure 5
<p>Temperature dependent stacking fault energy of X5CrNi18-8 and X3CrNi25-20.</p> Full article ">Figure 6
<p>EBSD analysis of the subsurface microstructure of Ni and Ni20Cr. Both materials have been affected by sliding abrasion at 30, 400, and 700 °C. The cross section after 3000 wear cycles is displayed.</p> Full article ">Figure 7
<p>EBSD analysis of the subsurface microstructure of the investigated austenitic steels. All materials have been affected by sliding abrasion at 30, 400, and 700 °C. The cross section after 3000 wear cycles is shown.</p> Full article ">
9 pages, 6591 KiB  
Article
Surface Investigation of Ni81Fe19 Thin Film: Using ARXPS for Thickness Estimation of Oxidation Layers
by Zongsheng He, Ziyu Li, Xiaona Jiang, Chuanjian Wu, Yu Liu, Xinglian Song, Zhong Yu, Yifan Wang, Zhongwen Lan and Ke Sun
Metals 2021, 11(12), 2061; https://doi.org/10.3390/met11122061 - 20 Dec 2021
Cited by 6 | Viewed by 3030
Abstract
This work demonstrates the dependence between magnetic properties and the thickness of NiFe thin films. More importantly, a quantitative study of the surface composition of NiFe thin film exposed to atmospheric conditions has been carried out employing angle-resolved X-ray photoelectron spectroscopy (ARXPS). In [...] Read more.
This work demonstrates the dependence between magnetic properties and the thickness of NiFe thin films. More importantly, a quantitative study of the surface composition of NiFe thin film exposed to atmospheric conditions has been carried out employing angle-resolved X-ray photoelectron spectroscopy (ARXPS). In this study, we fabricated Ni81Fe19 (NiFe) thin films on Si (100) substrate using electron beam evaporation and investigated their surface morphologies, magnetic properties, and the thickness of the surface oxide layer. The coexistence of metallic and oxidized species on the surface are suggested by the depth profile of ARXPS spectra. The thickness of the oxidized species, including NiO, Ni(OH)2, Fe2O3, and Fe3O4, are also estimated based on the ARXPS results. This work provides an effective approach to clarify the surface composition, as well as the thickness of the oxide layer of the thin films. Full article
(This article belongs to the Special Issue Advances in Metal-Containing Magnetic Materials)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>AFM images of NiFe films with a thickness of 100 nm.</p> Full article ">Figure 2
<p>The magnetic hysteresis (<span class="html-italic">M-H</span>) loops of NiFe thin films with thickness from 90 to 120 nm: (<b>a</b>) the in-plane magnetic hysteresis loops, (<b>b</b>) the out-of-plane magnetic hysteresis loop.</p> Full article ">Figure 3
<p>The static magnetic properties of the Ni<sub>81</sub>Fe<sub>19</sub> thin films with thicknesses from 90 to 120 nm: (<b>a</b>) the saturation magnetization (4π<span class="html-italic">M</span><sub>s</sub>) and average grain size (<span class="html-italic">D</span>), (<b>b</b>) the in-plane (<span class="html-italic">H</span><sub>c//</sub>) and out-of-plane (<span class="html-italic">H</span><sub>c</sub><sub>⊥</sub>) coercivity.</p> Full article ">Figure 4
<p>Spectra of Ni 2p photoemission in the NiFe thin film with a thickness of 100 nm at different take-off angles (<span class="html-italic">α</span>).</p> Full article ">Figure 5
<p>ARXPS spectra of Ni 2p photoemission and the fitting results for NiFe thin film with a thickness of 100 nm at different take-off angles: (<b>a</b>) 20° and (<b>b</b>) 90°.</p> Full article ">Figure 6
<p>(<b>a</b>) ARXPS spectra of O 1s photoemission and the fitting results, (<b>b</b>) the results of ln(1 + <span class="html-italic">R</span>/<span class="html-italic">R</span><sub>∞</sub>) against 1/sin<span class="html-italic">α</span> for the nickel oxides in NiFe thin film.</p> Full article ">Figure 7
<p>Spectra at different take-off angle of Fe 2p in NiFe(100 nm) thin film.</p> Full article ">Figure 8
<p>ARXPS spectra of Fe 2p photoemission and the peak fitting results for NiFe thin film with a thickness of 100 nm at different take-off angles of: (<b>a</b>) 20° and (<b>b</b>) 90°.</p> Full article ">Figure 9
<p>(<b>a</b>) ARXPS spectra of O 1s photoemission and the fitting results, (<b>b</b>) the results of ln(1 + <span class="html-italic">R</span>/<span class="html-italic">R</span>∞) against 1/sinα for the iron oxides in NiFe thin film.</p> Full article ">Figure 10
<p>The schematic diagram of NiFe thin film: (<b>a</b>) unoxidized, (<b>b</b>) oxidized.</p> Full article ">
21 pages, 9835 KiB  
Article
Drilling Parameters Analysis on In-Situ Al/B4C/Mica Hybrid Composite and an Integrated Optimization Approach Using Fuzzy Model and Non-Dominated Sorting Genetic Algorithm
by Palanikumar Kayaroganam, Velavan Krishnan, Elango Natarajan, Senthilkumar Natarajan and Kanesan Muthusamy
Metals 2021, 11(12), 2060; https://doi.org/10.3390/met11122060 - 20 Dec 2021
Cited by 32 | Viewed by 3550
Abstract
In-situ hybrid metal matrix composites were prepared by reinforcing AA6061 aluminium alloy with 10 wt.% of boron carbide (B4C) and 0 wt.% to 6 wt.% of mica. Machinability of the hybrid aluminium metal matrix composite was assessed by conducting drilling with [...] Read more.
In-situ hybrid metal matrix composites were prepared by reinforcing AA6061 aluminium alloy with 10 wt.% of boron carbide (B4C) and 0 wt.% to 6 wt.% of mica. Machinability of the hybrid aluminium metal matrix composite was assessed by conducting drilling with varying input parameters. Surface texture of the hybrid composites and morphology of drill holes were examined through scanning electron microscope images. The influence of rotational speed, feed rate and % of mica reinforcement on thrust force and torque were studied and analysed. Statistical analysis and regression analysis were conducted to understand the significance of each input parameter. Reinforcement of mica is the key performance indicator in reducing the thrust force and torque in drilling of the selected material, irrespective of other parameter settings. Thrust force is minimum at mid-speed (2000 rpm) with the lowest feed rate (25 mm/min), but torque is minimum at highest speed (3000 rpm) with lowest feed rate (25 mm/min). Multi-objective optimization through a non-dominated sorting genetic algorithm has indicated that 1840 rpm of rotational speed, 25.3 mm/min of feed rate and 5.83% of mica reinforcement are the best parameters for obtaining the lowest thrust force of 339.68 N and torque of 68.98 N.m. Validation through experimental results confirms the predicted results with a negligible error (less than 0.1%). From the analysis and investigations, it is concluded that use of Al/10 wt.% B4C/5.83 wt.% mica composite is a good choice of material that comply with European Environmental Protection Directives: 2000/53/CE-ELV for the automotive sector. The energy and production cost of the components can be very much reduced if the found optimum drill parameters are adopted in the production. Full article
(This article belongs to the Special Issue Optimization and Analysis of Metal Cutting Processes)
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Figure 1

Figure 1
<p>Stir casting setup used.</p> Full article ">Figure 2
<p>Drill tool and drilled samples of different compositions.</p> Full article ">Figure 3
<p>Experimental setup measuring thrust force and torque during drilling.</p> Full article ">Figure 4
<p>SEM images of fabricated composites (1000×). (<b>A</b>) 90 wt.% Al/10 wt.% B<sub>4</sub>C/0 wt.% mica (<b>B</b>) 87 wt.% Al/10 wt.% B<sub>4</sub>C/3 wt.% mica. (<b>C</b>) 84 wt.% Al/10% B<sub>4</sub>C + 6% mica.</p> Full article ">Figure 5
<p>Functional elements of fuzzy interference system used in this research.</p> Full article ">Figure 6
<p>(<b>a</b>) Fuzzy inference system model (<b>b</b>) Triangular membership function with 3 subsets for inputs (<b>c</b>) Triangular membership function with 9 subsets for outputs.</p> Full article ">Figure 7
<p>Rule editor representing the developed Fuzzy inference system.</p> Full article ">Figure 8
<p>Effect of rotational speed and feed rate on thrust force.</p> Full article ">Figure 9
<p>Effect of reinforcement on thrust force and normal probability plot.</p> Full article ">Figure 10
<p>Response surface plots of thrust force (<b>a</b>) rotational speed, feed rate vs. thrust force (<b>b</b>) rotational speed, % reinforcement of mica vs. thrust force and (<b>c</b>) feed rate, % reinforcement of mica vs. thrust force.</p> Full article ">Figure 11
<p>Effect of rotational speed and feed rate on torque.</p> Full article ">Figure 12
<p>Effect of reinforcement and normal probability plot for torque.</p> Full article ">Figure 13
<p>Response surface plot for torque (<b>a</b>) rotational speed, feed rate vs. torque (<b>b</b>) rotational speed, % reinforcement of mica vs. torque and (<b>c</b>) feed rate, % reinforcement of mica vs. torque.</p> Full article ">Figure 14
<p>SEM images of drilled holes at A = 2000 rpm, B = 75 mm/min. (<b>a</b>) 0 wt.% of mica, (<b>b</b>) 3 wt.% of mica, (<b>c</b>) 6 wt.% of mica.</p> Full article ">Figure 15
<p>Comparison of experimental and predicted values of thrust force.</p> Full article ">Figure 16
<p>Comparison of experimental and predicted values of torque.</p> Full article ">Figure 17
<p>Attainment of combined objective function value during execution.</p> Full article ">Figure 18
<p>Pareto optimal fronts.</p> Full article ">
14 pages, 3419 KiB  
Article
Experimental Study on Factors Influencing the Strength Distribution of In Situ Cemented Tailings Backfill
by Xiaopeng Peng, Lijie Guo, Guangsheng Liu, Xiaocong Yang and Xinzheng Chen
Metals 2021, 11(12), 2059; https://doi.org/10.3390/met11122059 - 20 Dec 2021
Cited by 11 | Viewed by 2912
Abstract
Previous studies have found that the strength of in situ cemented tailings backfill usually presents an S-shaped distribution, which decreases first, then increases, and decreases thereafter along the direction of slurry flow. In this study, to explore the factors determining the distribution, a [...] Read more.
Previous studies have found that the strength of in situ cemented tailings backfill usually presents an S-shaped distribution, which decreases first, then increases, and decreases thereafter along the direction of slurry flow. In this study, to explore the factors determining the distribution, a similar model test of cemented tailings backfill was carried out. The distribution law of grain size composition and the cement content of backfill materials along the flow direction were experimentally studied, and the comprehensive factor influencing the strength distribution was analyzed. The results show that, firstly, near the feeding point, there are more coarse particles, whereas the content of fine particles is higher farther away. The measured maximum median particle size can be more than three times the minimum value. Secondly, the cement content increases gradually along the flow direction and reaches the peak at the end of the model, which can be more than twice the minimum value, indicating that the degree of segregation is significant. Thirdly, the strength distribution of cemented backfills is comprehensively determined by both the particle size distribution (PSD) and the cement content. The maximum value appears neither at the point with peak median particle size, nor at the point with the highest cement content. Lastly, there is a strong linear correlation between the strength of cemented backfills and the strength factor (SF), which is defined as the product of the uniformity coefficient and cement content of filling materials, indicating that the SF can be used to quantitatively reflect the comprehensive effects of PSD and cement content on the strength. As SF is a comprehensive quantitative index reflecting the distribution of strength, it will be further studied in later research to acquire more experimental results of the relationship between sample strength and SF, which will be meaningful for the quality evaluation of in situ cemented backfills, and the optimization of backfill system. Full article
(This article belongs to the Special Issue Green Low-Carbon Technology for Metalliferous Minerals)
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<p>Diagram of the flume model test system. (<b>a</b>) Picture of the flume model; (<b>b</b>) the components of the test system.</p> Full article ">Figure 2
<p>Particle size distribution (PSD) of the tailings used in the study.</p> Full article ">Figure 3
<p>Illustration of sampling scheme of model test from top view.</p> Full article ">Figure 4
<p>The final profile of backfill slurries after flowing in the test model.</p> Full article ">Figure 5
<p>The cemented tailings backfill (CTB) specimens drilled from different areas after the flume test. (<b>a</b>) The specimens obtained; (<b>b</b>) the uniaxial compressive strength (UCS) testing process; (<b>c</b>) the specimen after failure.</p> Full article ">Figure 6
<p>The strength distribution of backfill samples along the flowing direction.</p> Full article ">Figure 7
<p>Verification chart of the cement contents of sampling cores in different rows.</p> Full article ">Figure 8
<p>Comparison chart of the cement contents and UCS of samples in row 2 along the flowing direction.</p> Full article ">Figure 9
<p>Comparison chart of the median particle sizes and UCS of samples in row 2 along the flowing direction.</p> Full article ">Figure 10
<p>Comparison chart of the strength factor (SF) values and UCS of samples in row 2.</p> Full article ">
21 pages, 6299 KiB  
Article
MQL-Assisted Hard Turning of AISI D2 Steel with Corn Oil: Analysis of Surface Roughness, Tool Wear, and Manufacturing Costs
by Bogdan Arsene, Catalin Gheorghe, Flavius Aurelian Sarbu, Magdalena Barbu, Lucian-Ionel Cioca and Gavrila Calefariu
Metals 2021, 11(12), 2058; https://doi.org/10.3390/met11122058 - 19 Dec 2021
Cited by 24 | Viewed by 4237
Abstract
Precision hard turning (HT) gained more and more attention in the cutting industry in the last years due to continuous pressure of the global market for reducing costs, minimizing the environmental and health issues, and achieving a cleaner production. Therefore, dry cutting and [...] Read more.
Precision hard turning (HT) gained more and more attention in the cutting industry in the last years due to continuous pressure of the global market for reducing costs, minimizing the environmental and health issues, and achieving a cleaner production. Therefore, dry cutting and minimal quantity lubrication (MQL) became widely used in manufacturing to meet the environmental issues with respect to harmful cutting fluids (CFs). Vegetable oils, in MQL machining, are a promising solutions to petroleum-based CFs; however, the effects and performance on surface roughness and tool wear in HT with ceramic inserts remain unclear. To address this limitation, hardened AIDI D2 steel and pure corn oil, rich in saturated and monounsaturated fatty acids, cheap and widely available, have been used to conduct dry and MQL experiments at different cutting speed and feeds. Results show that corn oil is suitable as cutting lubricant in HT, creating a strong anti-wear and anti-friction lubricating film which improves the roughness with 10–15% and tool life with 15–20%, therefore reducing costs. Best surface roughness values (Ra = 0.151 μm, Rz = 0.887 μm, Rpk = 0.261 μm) were obtained at 180 m/min and 0.1 mm/rev. The analysis of variance shows that corn oil has statistical significance on roughness, validating the results. Full article
(This article belongs to the Special Issue Optimization and Analysis of Metal Cutting Processes)
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<p>Experiment conditions.</p> Full article ">Figure 2
<p>Main effect plots for Ra, Rz, and Rpk.</p> Full article ">Figure 3
<p>Mean Ra vs. feed and cutting speed in dry (<b>a</b>) and MQL (<b>b</b>).</p> Full article ">Figure 4
<p>Mean Rz vs. feed and cutting speed in dry (<b>a</b>) and MQL (<b>b</b>).</p> Full article ">Figure 5
<p>Mean Rpk vs. feed and cutting speed in dry (<b>a</b>) and MQL (<b>b</b>).</p> Full article ">Figure 6
<p>Tool wear progression and related Ra.</p> Full article ">Figure 7
<p>Flank wear in dry (<b>a</b>) and MQL (<b>b</b>) after 14 min cutting time.</p> Full article ">Figure 8
<p>Tool wear after 28 min cutting time in MQL.</p> Full article ">Figure 9
<p>Tool wear after 28 min cutting time in dry.</p> Full article ">
11 pages, 2353 KiB  
Article
The Effect of Tin Content on the Strength of a Carbon Fiber/Al-Sn-Matrix Composite Wire
by Sergei Galyshev, Valery Orlov, Bulat Atanov, Evgeniy Kolyvanov, Oleg Averichev and Tigran Akopdzhanyan
Metals 2021, 11(12), 2057; https://doi.org/10.3390/met11122057 - 19 Dec 2021
Cited by 10 | Viewed by 3421
Abstract
The effect of tin content in an Al-Sn alloy in the range from 0 to 100 at.% on its mechanical properties was studied. An increase in the tin content leads to a monotonic decrease in the microhardness and conditional yield stress of the [...] Read more.
The effect of tin content in an Al-Sn alloy in the range from 0 to 100 at.% on its mechanical properties was studied. An increase in the tin content leads to a monotonic decrease in the microhardness and conditional yield stress of the Al-Sn alloy from 305 to 63 MPa and from 32 to 5 MPa, respectively. In addition, Young’s modulus and the shear modulus of the Al-Sn alloy decreases from 65 to 52 GPa and from 24 to 20 GPa, respectively. The effect of tin content in the Al-Sn matrix alloy in the range from 0 to 50 at.% on the strength of a carbon fiber/aluminum-tin-matrix (CF/Al-Sn) composite wire subject to three-point bending was also investigated. Increasing tin content up to 50 at.% leads to a linear increase in the composite wire strength from 1450 to 2365 MPa, which is due to an increase in the effective fiber strength from 65 to 89 at.%. The addition of tin up to 50 at.% to the matrix alloy leads to the formation of weak boundaries between the matrix and the fiber. An increase in the composite wire strength is accompanied by an increase in the average length of the fibers pulled out at the fracture surface. A qualitative model of the relationship between the above parameters is proposed. Full article
(This article belongs to the Special Issue Aluminum and Magnesium Alloys and Composites: Forming, Preparation, and Processing)
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<p>Scheme for obtaining a CF/Al-Sn-wire.</p> Full article ">Figure 2
<p>Microstructure of an Al-Sn matrix alloy with tin content from 5 to 50 at.%: (<b>a</b>) 5% Sn, (<b>b</b>) 10% Sn, (<b>c</b>) 25% Sn, (<b>d</b>) 50% Sn.</p> Full article ">Figure 3
<p>Dependence of microhardness, conditional yield stress (<b>a</b>) and Young’s modulus and shear modulus (<b>b</b>) of the matrix Al-Sn alloy on the tin content.</p> Full article ">Figure 4
<p>Fracture surfaces of a CF/Al-Sn composite with tin content from 0 to 50 at.%: (<b>a</b>) 0% Sn, (<b>b</b>) 10% Sn, (<b>c</b>) 25% Sn, (<b>d</b>) 50% Sn.</p> Full article ">Figure 4 Cont.
<p>Fracture surfaces of a CF/Al-Sn composite with tin content from 0 to 50 at.%: (<b>a</b>) 0% Sn, (<b>b</b>) 10% Sn, (<b>c</b>) 25% Sn, (<b>d</b>) 50% Sn.</p> Full article ">Figure 5
<p>Curves of deformation of a CF/Al-Sn composite with tin content in the matrix of 0 and 50 at.%.</p> Full article ">Figure 6
<p>Dependence of the strength of a composite CF/Al-Sn wire subjected to three-point bending (<b>a</b>), effective fiber strength and length of the pulled-out part of the fiber at fracture (<b>b</b>) on the tin content.</p> Full article ">Figure 7
<p>Microstructure of a CF/Al-Sn composite with tin content from 0 to 50 at.%: (<b>a</b>) 0% Sn, (<b>b</b>) 10% Sn, (<b>c</b>) 25% Sn, (<b>d</b>) 50% Sn.</p> Full article ">Figure 7 Cont.
<p>Microstructure of a CF/Al-Sn composite with tin content from 0 to 50 at.%: (<b>a</b>) 0% Sn, (<b>b</b>) 10% Sn, (<b>c</b>) 25% Sn, (<b>d</b>) 50% Sn.</p> Full article ">Figure 8
<p>Schematic representation of the dependence of the composite strength <span class="html-italic">σ</span> on the shear strength of the interface between the matrix and the fiber τ.</p> Full article ">
14 pages, 101903 KiB  
Article
The Casting Rate Impact on the Microstructure in Al–Mg–Si Alloy with Silicon Excess and Small Zr, Sc Additives
by Evgenii Aryshenskii, Maksim Lapshov, Sergey Konovalov, Jurgen Hirsch, Vladimir Aryshenskii and Svetlana Sbitneva
Metals 2021, 11(12), 2056; https://doi.org/10.3390/met11122056 - 19 Dec 2021
Cited by 3 | Viewed by 2636
Abstract
The study investigates the effect of casting speed on the solidification microstructure of the aluminum alloy Al0.3Mg1Si with and without the additions of zirconium and scandium. Casting was carried out in steel, copper, and water-cooled chill molds with a [...] Read more.
The study investigates the effect of casting speed on the solidification microstructure of the aluminum alloy Al0.3Mg1Si with and without the additions of zirconium and scandium. Casting was carried out in steel, copper, and water-cooled chill molds with a crystallization rate of 20 °C/s, 10 °C/s, and 30 °C/s, respectively. For each casting mode, the grain structure was investigated by optical microscopy and the intermetallic particles were investigated by scanning and transmission microscopy; in addition, measurements of the microhardness and the electrical conductivity were carried out. An increase in the solidification rate promotes grain refinement in both alloys. At the same time, the ingot cooling rate differently affects the number of intermetallic particles. In an alloy without scandium–zirconium additives, an increase in the ingot cooling rate leads to a decrease in the number of dispersoids due to an increase in the solubility of the alloying elements in a supersaturated solid solution. With the addition of scandium and zirconium, the amount of dispersoids increases slightly. This is because increasing the solubility of the alloying elements in a supersaturated solid solution is leveled by a growth of the number of grain boundaries, promoting the formation of particles of the (AlSi)3ScZr type, including those of the L12 type. In addition, the increase in the crystallization rate increases the number of primary nonequilibrium intermetallic particles which have a eutectic nature. Full article
(This article belongs to the Special Issue Microstructure and Mechanical Properties of Aluminum Alloys)
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<p>Temperature depending on time for different chill molds.</p> Full article ">Figure 2
<p>Optical microstructure Al<sub>0.3</sub>Mg<sub>1</sub>Si—(<b>a</b>) steel chill mold; (<b>b</b>) copper chill mold; (<b>c</b>) copper chill mold of round section; Al<sub>0.3</sub>Mg<sub>1</sub>Si<sub>0.3</sub>Sc<sub>0.15</sub>Zr–(<b>d</b>) steel chill mold; (<b>e</b>) copper chill mold; (<b>f</b>) copper chill mold of round section.</p> Full article ">Figure 2 Cont.
<p>Optical microstructure Al<sub>0.3</sub>Mg<sub>1</sub>Si—(<b>a</b>) steel chill mold; (<b>b</b>) copper chill mold; (<b>c</b>) copper chill mold of round section; Al<sub>0.3</sub>Mg<sub>1</sub>Si<sub>0.3</sub>Sc<sub>0.15</sub>Zr–(<b>d</b>) steel chill mold; (<b>e</b>) copper chill mold; (<b>f</b>) copper chill mold of round section.</p> Full article ">Figure 3
<p>Grain size of alloys depending on the cooling rate.</p> Full article ">Figure 4
<p>Electrical conductivity of alloys depending on the cooling rate.</p> Full article ">Figure 5
<p>Microhardness of alloys depending on the cooling rate.</p> Full article ">Figure 6
<p>Micrographs of Al<sub>0.3</sub>Mg<sub>1</sub>Si alloy. Phases: (<b>a</b>) steel chill mold; (<b>b</b>) copper chill mold; (<b>c</b>) copper chill mold of round section. Intermetallic compounds: (<b>d</b>) steel chill mold; (<b>e</b>) copper chill mold; (<b>f</b>) copper chill mold of round section.</p> Full article ">Figure 7
<p>Micrograph of Al<sub>0.3</sub>Mg<sub>1</sub>Si<sub>0.3</sub>Sc<sub>0.15</sub>Zr alloy. Phases: (<b>a</b>) steel chill mold; (<b>b</b>) copper chill mold; (<b>c</b>) copper chill mold of round section. Intermetallic compounds: (<b>d</b>) steel chill mold; (<b>e</b>) copper chill mold; (<b>f</b>) copper chill mold of round section.</p> Full article ">Figure 8
<p>Number and size of dispersoids in alloys. Al<sub>0.3</sub>Mg<sub>1</sub>Si: (<b>a</b>) steel chill mold; (<b>b</b>) copper chill mold; (<b>c</b>) copper chill mold of round section. Al <sub>0.3</sub>Mg<sub>1</sub>Si<sub>0.3</sub>Sc<sub>0.15</sub>Zr: (<b>d</b>) steel chill mold; (<b>e</b>) copper chill mold; (<b>f</b>) copper chill mold of round section.</p> Full article ">Figure 9
<p>Al<sub>0.3</sub>Mg1Si<sub>0.15</sub>Zr<sub>0.3</sub>S steel chill mold (<b>a</b>); (<b>b</b>) Particles (AlSi)<sub>3</sub>ScZr and Al<sub>3</sub>Sc; (<b>c</b>) EDS profile line scan.</p> Full article ">Figure 10
<p>Al<sub>0.3</sub>Mg<sub>1</sub>Si<sub>0.15</sub>Zr<sub>0.3</sub>S copper chill mold (<b>a</b>); (<b>b</b>) Particles (AlSi)<sub>3</sub>ScZr; (<b>c</b>) EDS profile line scan.</p> Full article ">Figure 11
<p>Al<sub>0.3</sub>Mg<sub>1</sub>Si<sub>0.15</sub>Zr<sub>0.3</sub>S copper cone (<b>a</b>) Al<sub>3</sub>Sc particles, bright field; (<b>b</b>) Al<sub>3</sub>Sc particles, dark field in the superstructural reflection; (<b>c</b>) Al<sub>3</sub>Sc particles with zirconium, bright field; (<b>d</b>) EDS profile line scan.</p> Full article ">
17 pages, 5422 KiB  
Article
Effect of Multi-Step Austempering Treatment on the Microstructure and Mechanical Properties of a High Silicon Carbide-Free Bainitic Steel with Bimodal Bainite Distribution
by Mattia Franceschi, Alvise Miotti Bettanini, Luca Pezzato, Manuele Dabalà and Pascal J. Jacques
Metals 2021, 11(12), 2055; https://doi.org/10.3390/met11122055 - 19 Dec 2021
Cited by 20 | Viewed by 3025
Abstract
The effect of multi-step austempering treatments on the microstructure and mechanical properties of a novel medium carbon high silicon carbide-free bainitic steel was studied. Five different isothermal treatment processes were selected, including single-step isothermal treatments above martensite start temperature (at 350 °C and [...] Read more.
The effect of multi-step austempering treatments on the microstructure and mechanical properties of a novel medium carbon high silicon carbide-free bainitic steel was studied. Five different isothermal treatment processes were selected, including single-step isothermal treatments above martensite start temperature (at 350 °C and 370 °C, respectively), and three kinds of two-step routes (370 °C + 300 °C, 370 °C + 250 °C, and 350 °C + 250 °C). In comparison with single-step austempering treatment adopting a two-step process, a microstructure with a bimodal-size distribution of bainitic ferrite and without martensite was obtained. Bainitic transformation was studied using dilatometry both for single-step and two-step routes and the specimens were completely characterised by electron microscopy (SEM and TEM), X-ray diffraction (XRD) and standard tensile tests. The mechanical response of the samples subjected to two-step routes was superior to those treated at a single temperature. Full article
(This article belongs to the Special Issue Steel Heat Treatment)
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<p>Bainitic multi-step heat treatment cycles after austenitisation at 900 °C for 5 min.</p> Full article ">Figure 2
<p>(<b>A</b>) Thermal cycle for the dilatometric study of critical temperature Ac1, Ac3, Ms (black line), and dilatation curve (blue line); (<b>B</b>) relative change in length against temperature curve indicating austenitisation transformation and the temperature Ac1 and Ac3; the star (*) = indicates cementite precipitation during heating of martensite due to carbon rejection; (<b>C</b>) cooling branch of the thermal cycle indicating Ms temperature and martensitic transformation.</p> Full article ">Figure 3
<p>(<b>A</b>) Cooling branch of the dilatometric curve showing the change in length (∆l) against temperature for the single-step austempering treatment (<b>A</b>,<b>B</b>) Cooling branch of the dilatometric curve showing the change in length (∆l) against temperature for the single-step austempering treatment (<b>B</b>) dilatometric curve showing the change in length with time. The insert in (<b>B</b>) better shows the incubation time for both bainite transformations.</p> Full article ">Figure 4
<p>Cooling branch of the dilatometric curve showing the change in length (∆l) against temperature for the multi-step austempering treatments indicating the bainitic transformation during all the austempering steps.</p> Full article ">Figure 5
<p>(<b>A</b>,<b>C</b>,<b>E</b>) relative change in length vs. austempering time curve for the heat treatment MULTISTEP 1, MULTISTEP 2, MULTISTEP 3, respectively. (<b>B</b>,<b>D</b>,<b>F</b>) Derivative of the relative change in length against time treatment describing the kinetic of the bainitic transformation during MULTISTEP 1, MULTISTEP 2, MULTISTEP 3 respectively.</p> Full article ">Figure 6
<p>SEM micrograph of the sample treated with a single-step austempering until completion of the bainitic transformation at (<b>A</b>) 370 °C and (<b>B</b>) 350 °C. RAF: retained austenite film; BF: bainitic ferrite; RAB: retained austenite block; M: martensite.</p> Full article ">Figure 7
<p>(<b>A</b>–<b>C</b>) SEM micrograph of the sample treated according to route MULTISTEP 1, MULTISTEP2, MULTISTEP3, respectively. (<b>D</b>) High magnification SEM micrograph showing filmy of retained austenite and bainitic ferrite. B1: carbide-free bainite formed during the first isothermal holding step; B2: carbide-free bainite formed during the second isothermal holding step. RAF: retained austenite film; BF: bainitic ferrite; RAB: retained austenite block; M: martensite.</p> Full article ">Figure 8
<p>TEM micrographs on the samples treated according to MULTISTEP3 route (<b>A</b>) Bainitic sheave, (<b>B</b>) Retained austenite block between bainitic sheaves.</p> Full article ">Figure 9
<p>XRD patterns for the different heat treatment routes.</p> Full article ">Figure 10
<p>Engineering stress–engineering strain curves of the heat-treated steel samples (yield strength was calculated with the 0.2 criteria).</p> Full article ">Figure 11
<p>Fracture surfaces of the heat-treated steel samples after (<b>A</b>) MULTISTEP1, (<b>B</b>) MULTISTEP2, (<b>C</b>) MULTISTEP3.</p> Full article ">
38 pages, 14074 KiB  
Review
The Evolution of Intermetallic Compounds in High-Entropy Alloys: From the Secondary Phase to the Main Phase
by Junqi Liu, Xiaopeng Wang, Ajit Pal Singh, Hui Xu, Fantao Kong and Fei Yang
Metals 2021, 11(12), 2054; https://doi.org/10.3390/met11122054 - 18 Dec 2021
Cited by 34 | Viewed by 9551
Abstract
High-performance structural materials are critical to the development of transportation, energy, and aerospace. In recent years, newly developed high-entropy alloys with a single-phase solid-solution structure have attracted wide attention from researchers due to their excellent properties. However, this new material also has inevitable [...] Read more.
High-performance structural materials are critical to the development of transportation, energy, and aerospace. In recent years, newly developed high-entropy alloys with a single-phase solid-solution structure have attracted wide attention from researchers due to their excellent properties. However, this new material also has inevitable shortcomings, such as brittleness at ambient temperature and thermodynamic instability at high temperature. Efforts have been made to introduce a small number of intermetallic compounds into single-phase solid-solution high-entropy alloys as a secondary phase to their enhance properties. Various studies have suggested that the performance of high-entropy alloys can be improved by introducing more intermetallic compounds. At that point, researchers designed an intermetallic compound-strengthened high-entropy alloy, which introduced a massive intermetallic compound as a coherent strengthening phase to further strengthen the matrix of the high-entropy alloy. Inspired from this, Fantao obtained a new alloy—high-entropy intermetallics—by introducing different alloying elements to multi-principalize the material in a previous study. This new alloy treats the intermetallic compound as the main phase and has advantages of both structural and functional materials. It is expected to become a new generation of high-performance amphibious high-entropy materials across the field of structure and function. In this review, we first demonstrate the inevitability of intermetallic compounds in high-entropy alloys and explain the importance of intermetallic compounds in improving the properties of high-entropy alloys. Secondly, we introduce two new high-entropy alloys mainly from the aspects of composition design, structure, underlying mechanism, and performance. Lastly, the high-entropy materials containing intermetallic compound phases are summarized, which lays a theoretical foundation for the development of new advanced materials. Full article
(This article belongs to the Special Issue Design, Development and Characterization of Advanced Metallic Materials)
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<p>The identified crystal structures of high-entropy alloys adaptedted with permission from ref. [<a href="#B51-metals-11-02054" class="html-bibr">51</a>] Copyright 2021 Elsevier. (<b>A</b>) Body-centered cubic (BCC) structure. (<b>B</b>) Face-centered cubic (FCC) structure. (<b>C</b>) Hexagonal close-packed (HCP) structure.</p> Full article ">Figure 2
<p>The comparison of the microstructure between high-entropy alloys and nickel-based superalloys. (<b>A</b>) The morphology of γ’ phase in Al0.2CrFeCoNi2Cu0.2 HEA adapted with permission from ref. [<a href="#B54-metals-11-02054" class="html-bibr">54</a>]. Copyright 2021 Elsevier. (<b>B</b>) TEM images of the Al0.2CrFeCoNi2Cu0.2 HEA containing γ and γ′ phases adapted with permission from ref. [<a href="#B54-metals-11-02054" class="html-bibr">54</a>]. Copyright 2021 Elsevier. (<b>C</b>) The morphology of γ’ phase in nickel-based superalloy that contains 1.5 wt.% Ti adapted with permission from ref. [<a href="#B6-metals-11-02054" class="html-bibr">6</a>]. Copyright 2021 Elsevier. (<b>D</b>) The morphology of γ’ phase in nickel-based superalloy that contains 10 wt.% Mo [<a href="#B6-metals-11-02054" class="html-bibr">6</a>]. (<b>E</b>) The morphology of γ’ phase in nickel-based superalloy that contains 12 wt.% Mo adapted with permission from ref. [<a href="#B6-metals-11-02054" class="html-bibr">6</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 3
<p>Ashby plot of strength versus fracture toughness showing that CrCoNi-based, medium-entropy, and high-entropy alloys are among the most damage-tolerant materials on record adapted with permission from ref. [<a href="#B70-metals-11-02054" class="html-bibr">70</a>]. Copyright 2021 Springer Nature.</p> Full article ">Figure 4
<p>Scanning electron microscopy (SEM) image of CrMnFeCoNi HEA after heat treatment under different temperatures for 500 days adapted with the permission from ref. [<a href="#B42-metals-11-02054" class="html-bibr">42</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 5
<p>(<b>A</b>) Tensile properties of alloys alloy A (as-homogenized FeCoNiCr), alloy B as-homogenized (FeCoNiCr)94Ti2Al4, P1 (alloy B after first treatment), and P2 (alloy B after second treatment) at room temperature. (<b>B</b>) The map of ultimate tensile strength-ductility combinations of various advanced steels including P1 and P2 HEAs adapted with the permission from ref. [<a href="#B71-metals-11-02054" class="html-bibr">71</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 6
<p>The relationship between δ, ΔH <sub>mix</sub>, and the alloy phase composition adapted with permission from ref. [<a href="#B96-metals-11-02054" class="html-bibr">96</a>]. Copyright 2021 Wiley Online Library.</p> Full article ">Figure 7
<p>The relationship between high-entropy alloy phase composition and α<sub>2</sub> value adapted with permission from ref. [<a href="#B98-metals-11-02054" class="html-bibr">98</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 8
<p>(<b>A</b>) Phase diagram calculations of (CoCrNi)<sub>97-x</sub>Ti<sub>3</sub>Al<sub>x</sub> alloy (at.%) using Thermo-Calc software with TTNI8 database. (<b>B</b>) Phase diagram calculations of Al<sub>x</sub>Co<sub>1.5</sub>CrFeNi<sub>1.5</sub>Ti<sub>y</sub> alloy (at.%) using Thermo-Calc software with TTNI8 database. (<b>C</b>) The equilibrium phase mole fraction of the L1<sub>2</sub> phase as a function of temperature calculated using the thermal-calc method for the (FeCoNiCr)<sub>100-x-y</sub>Al<sub>x</sub>Ti<sub>y</sub> system (at.%). (<b>D</b>) The equilibrium phase mole fraction of the Heusler phase as a function of temperature calculated using the thermal-calc method for the (FeCoNiCr)<sub>100-x-y</sub>Al<sub>x</sub>Ti<sub>y</sub> system (at.%) adapted with permission from ref. [<a href="#B85-metals-11-02054" class="html-bibr">85</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 9
<p>Conceptual design and microstructural characterizations of the MCINPS alloys. (<b>A</b>) Schematic of the design concept of the MCINPS alloys. MCM, multicomponent matrix. (<b>B</b>) Scanning electron microscopy (SEM) image of the Al<sub>7</sub>Ti<sub>7</sub> alloy exhibiting the typical equiaxed grain structures. (<b>C</b>) SEM image of the Al<sub>7</sub>Ti<sub>7</sub> alloy revealing the uniform distribution of high-density L1<sub>2</sub> MCINP within the grain interior. (<b>D</b>) XRD patterns showing the phase compositions of the Al<sub>7</sub>Ti<sub>7</sub> alloy. a.u., arbitrary units. (<b>E</b>) TEM image of the Al<sub>7</sub>Ti<sub>7</sub> alloy showing the nanostructured morphology. The inset shows the corresponding SAED pattern. (<b>F</b>) Representative high-resolution TEM image confirming the interfacial coherency. adapted with permission from ref. [<a href="#B105-metals-11-02054" class="html-bibr">105</a>]. Copyright 2021 The American Association for the Advancement of Science.</p> Full article ">Figure 10
<p>Yield strength versus the product of strength and ductility of the MCINPS alloys compared with those of other high-performing materials showing that exceptional strength-ductility combination can be achieved in the MCINPS alloys at ambient temperature (Al<sub>7</sub>Ti<sub>7</sub> is the alloy of interest in this reference)adapted with permission from ref. [<a href="#B105-metals-11-02054" class="html-bibr">105</a>]. Copyright 2021 The American Association for the Advancement of Science.</p> Full article ">Figure 11
<p>(<b>A</b>) Mechanical properties of the L1<sub>2</sub>-strengthened (CoCrNi)<sub>94</sub>Al<sub>3</sub>Ti<sub>3</sub> alloy (at.%) compared to those of single-phase CoCrNi base alloy at ambient temperature. (<b>B</b>) Ultimate tensile strength at ambient temperature versus tensile elongation of the present alloy in comparison with commercial superalloys adapted with permission from ref. [<a href="#B85-metals-11-02054" class="html-bibr">85</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 12
<p>Microstructures of the Ni<sub>44.7</sub>Co<sub>23</sub>.<sub>7</sub>Fe<sub>8</sub>Cr10Al<sub>8.6</sub>Ti<sub>5</sub>-based (at.%) high-entropy superalloys aged at (<b>A</b>) 900 °C for 24 h and (<b>B</b>) 1000 °C for 24 h adapted with permission from ref. [<a href="#B114-metals-11-02054" class="html-bibr">114</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 13
<p>(<b>A</b>) Schematic illustration showing the complex precipitation pathway from the metastable L1<sub>2</sub> to the Heusler phase at the GBs. (<b>B</b>) Schematic illustration of the duplex-aging design for the GB stabilization adapted with permission from ref. [<a href="#B128-metals-11-02054" class="html-bibr">128</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 14
<p>(<b>A</b>) Elongation versus yield strength of the L1<sub>2</sub>-strengthened HEAs compared with the single-phase solid-solution HEAs at 293 K. (<b>B</b>) Elongation versus ultimate tensile strength of the L1<sub>2</sub>-strengthened HEAs compared with the single-phase solid-solution HEAs at 293 K. (<b>C</b>) Elongation versus yield strength of the L1<sub>2</sub>-strengthened HEAs compared with the single-phase solid-solution HEAs at 77 K. (<b>D</b>) Elongation versus ultimate tensile strength of the L1<sub>2</sub>-strengthened HEAs compared with the single-phase solid-solution HEAs at 77 K adapted with permission from ref. [<a href="#B128-metals-11-02054" class="html-bibr">128</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 15
<p>The strength contributions from different hardening mechanisms adapted with permission from ref. [<a href="#B71-metals-11-02054" class="html-bibr">71</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 16
<p>(<b>A</b>) Variations in precipitate shearing stress (σ<sub>sh</sub>) and Orowan dislocation looping stress (σ<sub>Or</sub>) as functions of the aging time. (<b>B</b>) Schematics of the Orowan dislocation looping and dislocation shearing mechanisms adapted with permission from ref. [<a href="#B131-metals-11-02054" class="html-bibr">131</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 17
<p>(<b>A</b>) Engineering stress–strain curves of the CoCrFeMnNi alloy at the six testing temperatures for the fine-grained (grain size 4.4 μm). (<b>B</b>) Engineering stress–strain curves of the CoCrFeMnNi alloy at the six testing temperatures for the coarse-grained (grain size 155 μm) adapted with permission from ref. [<a href="#B135-metals-11-02054" class="html-bibr">135</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 18
<p>Crystal structure model of L1<sub>2</sub> structure high-entropy intermetallics simulation adapted with permission from ref. [<a href="#B54-metals-11-02054" class="html-bibr">54</a>]. Copyright 2021 Elsevier. (<b>A</b>) I type high-entropy intermetallics. (<b>B</b>) II type high-entropy intermetallics. (<b>C</b>) III type high-entropy intermetallics.</p> Full article ">Figure 19
<p>Crystal structure model of B2 structure high-entropy intermetallics adapted from [<a href="#B55-metals-11-02054" class="html-bibr">55</a>]. (<b>A</b>) I type high-entropy intermetallics. (<b>B</b>) II type high-entropy intermetallics.</p> Full article ">Figure 20
<p>Three-dimensional compositional distributions and nanoscale interfacial cosegregation of the NDI-SMs. (<b>A</b>) Atom maps reconstructed using 3D-APT that show the distribution of each element. (<b>B</b>) Two-dimensional compositional contour maps revealing the multielement cosegregation behaviors of different elements within the disordered interfacial nanolayer (DINL). (<b>C</b>) One-dimensional compositional profile that quantitatively reveals the elemental distributions across the micrometer-scale ordered superlattice grain (OSG) and disordered interfacial nanolayer (DINL) adapted with permission from ref. [<a href="#B143-metals-11-02054" class="html-bibr">143</a>]. Copyright 2021 The American Association for the Advancement of Science.</p> Full article ">Figure 21
<p>Structure models of Ni<sub>44</sub>Co<sub>23</sub>Fe<sub>11</sub>Al<sub>8</sub>Ti<sub>14</sub> based on the results of MC simulation adapted with permission from ref. [<a href="#B141-metals-11-02054" class="html-bibr">141</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 22
<p>(<b>A</b>) Schematic illustrations of the HEIs with the B2 structure. (<b>B</b>) XRD patterns for seven HEIs specimens that exhibit primarily or completely single B2 phases after annealing at 1100 °C for 10 h adapted with permission from ref. [<a href="#B145-metals-11-02054" class="html-bibr">145</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 23
<p>(<b>A</b>) Schematic illustrations of the HEIs with the D0<sub>22</sub> structure. (<b>B</b>) XRD patterns of five HEIs specimens with primarily the D0<sub>22</sub> phase after annealing at 1300 °C for 10 h. The D0<sub>22</sub> phase is indexed, and the unindexed peaks with low intensity correspond to the secondary phases. The D0<sub>22</sub> phase is dominant in all five cases (albeit some minor secondary phases) adapted with permission from ref. [<a href="#B145-metals-11-02054" class="html-bibr">145</a>]. Copyright 2021 Elsevier.</p> Full article ">Figure 24
<p>Ashby plot showing strength–toughness relationships for engineering materials. Diagonal lines show the plastic-zone size, K<sub>c</sub><sup>2</sup>/πσ<sub>y</sub><sup>2</sup>, where K<sub>c</sub> is the fracture toughness and σ<sub>y</sub> is the yield strength adapted with permission from ref. [<a href="#B146-metals-11-02054" class="html-bibr">146</a>]. Copyright 2021 Springer Nature.</p> Full article ">Figure 25
<p>(<b>A</b>) SEM backscattered electron image of conventionally cast EHEA. (<b>B</b>) SEM backscattered electron image of directionally solidified EHEA with a hierarchical herringbone microstructure. The black arrows in (<b>B</b>) indicate the DS direction, and also the tensile loading direction in <a href="#metals-11-02054-f026" class="html-fig">Figure 26</a>A. (<b>C</b>) HAADF-STEM image and related SAED patterns of B2 and L1<sub>2</sub> phases. (<b>D</b>) SHE-XRD of B2 and L1<sub>2</sub> phases adapted with permission from ref. [<a href="#B147-metals-11-02054" class="html-bibr">147</a>]. Copyright 2021 The American Association for the Advancement of Science.</p> Full article ">Figure 26
<p>Tensile response at ambient temperature. (<b>A</b>) Engineering stress–strain curves of the directionally solidified EHEA compared with the conventionally cast EHEA. Inset shows the corresponding strain-hardening curves. MDIH and MBIH refer to multi-slip dislocation-induced hardening and microband-induced hardening, respectively. (<b>B</b>) Yield strength versus uniform strain of the directionally solidified EHEAs compared with those of previously reported as-cast eutectic and near-eutectic HEAs. (N-) EHEAs refer to eutectic and near-eutectic HEAs. The conventional (N-) EHEAs include directly cast and arc-melting eutectic and near-eutectic HEAs adapted with permission from ref. [<a href="#B147-metals-11-02054" class="html-bibr">147</a>]. Copyright 2021 The American Association for the Advancement of Science.</p> Full article ">Figure 27
<p>(<b>A</b>) Bright-field TEM image showing the polycrystalline morphology. (Inset) A corresponding selected-area electron diffraction pattern collected from the grain interior, which shows the L1<sub>2</sub>-type ordered structure. (<b>B</b>) Atomic-resolution HAADF-STEM image and corresponding EDX maps taken from the inner L1<sub>2</sub>-type OSG, revealing the sublattice occupations. (<b>C</b>) High-resolution HAADF-STEM image revealing the ultrathin disordered layer at the grain boundaries with a nanoscale thickness. The images on the right show the corresponding fast Fourier transform (FFT) patterns. (<b>D</b>) EDX maps showing the compositional distribution of the DINL. (<b>E</b>) Superlattice architecture with nanoscale disordered interfaces. © Schematic illustration highlighting the nanoscale interfacially disordered structure adapted with permission from ref. [<a href="#B143-metals-11-02054" class="html-bibr">143</a>]. Copyright 2021 The American Association for the Advancement of Science.</p> Full article ">Figure 28
<p>Mechanical properties and thermal stability of the NDI-SMs. (<b>A</b>) Tensile stress–strain curve of the NDI-SM tested at 20 °C in air. The stress–strain curve of the high-strength Ni<sub>3</sub>Al-type (Ni<sub>3</sub>Al-2.5 at.% B) alloy is also plotted for a direct comparison. (Inset) Tensile fractography showing the ductile dimpled structures. (<b>B</b>) Yield strength (σ<sub>y</sub>) versus uniform elongation (ɛ<sub>u</sub>) of the present NDI-SM compared with various conventional bulk-ordered alloys adapted with permission from ref. [<a href="#B143-metals-11-02054" class="html-bibr">143</a>]. Copyright 2021 The American Association for the Advancement of Science.</p> Full article ">Figure 29
<p>Electrocatalytic performance of the HEI for HER in 1.0 m KOH solutions. (<b>A</b>) Comparison of overpotentials at 10 mA cm<sup>−2</sup> versus Tafel slopes for ≈60 different catalysts in 1.0 m KOH, showing that the L1<sub>2</sub> HEI represents a low-cost alternative to compete with noble metal catalysts (inset is the comparison of raw materials costs). (<b>B</b>) Relationship between dealloying time-dependent structural evolution and catalytic performance variation, showing the significant improvement for hydrogen evolution reactions (HER) performance as the dealloying process isolated the L1<sub>2</sub> phase adapted with permission from ref. [<a href="#B163-metals-11-02054" class="html-bibr">163</a>]. Copyright 2021 Wiley Online Library.</p> Full article ">Figure 30
<p>Spatial morphology, structural, and multicomponent nature of the D15h HEI. (<b>A</b>) SEM image of the D15h HEI with a dendritic-like morphology. (<b>B</b>) Aberration-corrected HAADF-STEM image viewed along the [001] zone axis showing the L1<sub>2</sub>-type A<sub>3</sub>B structure. The inset shows the corresponding SAED pattern. (<b>C</b>) High-magnification HAADF-STEM image accompanied by atomic-resolution elemental maps clearly showing the ordered crystallographic structure and site occupancy of the L1<sub>2</sub>-type structure (AlTi atoms on the vertices and FeCoNi atoms on the face center in an L1<sub>2</sub> unit cell). (<b>D</b>) DFT modeling of the atomic ordering and site occupancy of the L1<sub>2</sub> HEI, emphasizing the site-isolated structure (A: AlTi, B: FeCoNi) adapted with permission from ref. [<a href="#B163-metals-11-02054" class="html-bibr">163</a>]. Copyright 2021 Wiley Online Library.</p> Full article ">
16 pages, 5552 KiB  
Article
Enhancement of Uniform Elongation by Temperature Change during Tensile Deformation in a 0.2C TRIP Steel
by Noriyuki Tsuchida and Stefanus Harjo
Metals 2021, 11(12), 2053; https://doi.org/10.3390/met11122053 - 18 Dec 2021
Cited by 2 | Viewed by 2992
Abstract
It is important to control the deformation-induced martensitic transformation (DIMT) up to the latter part of the deformation to improve the uniform elongation (U.El) through the TRIP effect. In the present study, tensile tests with decreasing deformation temperatures were conducted to achieve continuous [...] Read more.
It is important to control the deformation-induced martensitic transformation (DIMT) up to the latter part of the deformation to improve the uniform elongation (U.El) through the TRIP effect. In the present study, tensile tests with decreasing deformation temperatures were conducted to achieve continuous DIMT up to the latter part of the deformation. As a result, the U.El was improved by approximately 1.5 times compared with that in the tensile test conducted at 296 K. The enhancement of the U.El in the temperature change test was discussed with the use of neutron diffraction experiments. In the continuous DIMT behavior, a maximum transformation rate of about 0.4 was obtained at a true strain (ε) of 0.2, which was larger than that in the tensile test at 296 K. The tensile deformation behavior of ferrite (α), austenite (γ), and deformation-induced martensite (α′) phases were investigated from the viewpoint of the fraction weighted phase stress. The tensile test with a decreasing deformation temperature caused the increase of the fraction weighted phase stress of α and that of α′, which was affected by the DIMT behavior, resulting in the increase in the work hardening, and also controlled the ductility of α and α′, resulting in the enhancement of the U.El. Especially, the α phase contributed to maintaining high strength instead of α′ at a larger ε. Therefore, not only the DIMT behavior but also the deformation behavior of γ, α, and α′ are important in order to improve U.El due to the TRIP effect. Full article
(This article belongs to the Special Issue Advances in High-Strength Low-Alloy Steels)
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<p>Schematic illustrations about test A (<b>a</b>) and B (<b>b</b>).</p> Full article ">Figure 2
<p>EBSD phase mapping image (<b>a</b>) and orientation color map (<b>b</b>) in the 0.2C TRIP steel.</p> Full article ">Figure 3
<p>Nominal stress–nominal strain curves of the 0.2C TRIP steel obtained by tensile tests at various deformation temperatures.</p> Full article ">Figure 4
<p>0.2% proof strength, tensile strength, and uniform elongation as a function of temperature in the 0.2C TRIP steel.</p> Full article ">Figure 5
<p>True stress (<span class="html-italic">σ</span>) and work-hardening rate (<span class="html-italic">dσ</span>/<span class="html-italic">dε</span>) as a function of true strain at various deformation temperatures.</p> Full article ">Figure 6
<p>Nominal stress–nominal strain curves obtained by test A. Here, the nominal stress–nominal strain curve at 296 K is also shown as a dashed line.</p> Full article ">Figure 7
<p>Tensile strength and uniform elongation as a function of deformation temperature obtained by test A and the tensile tests. Here, the temperature of test A means the reloading temperature after tensile deformation at 373 K.</p> Full article ">Figure 8
<p>Volume fraction of deformation-induced martensite as a function of true strain obtained by the tensile tests at various deformation temperatures (<b>a</b>) and test A (<b>b</b>).</p> Full article ">Figure 9
<p>Nominal stress–nominal strain curves obtained by test B and the tensile test at 296 K.</p> Full article ">Figure 10
<p>Deformation temperature (<b>a</b>) and true stress (<span class="html-italic">σ</span>) or work-hardening rate (<span class="html-italic">dσ</span>/<span class="html-italic">dε</span>) (<b>b</b>) as functions of true strain obtained by test B and the tensile test at 296 K.</p> Full article ">Figure 11
<p>Volume fraction of deformation-induced martensite (<b>a</b>) and transformation rate by using Equation (4) (<b>b</b>) as a function of true strain obtained by test B and the tensile test at 296 K [<a href="#B7-metals-11-02053" class="html-bibr">7</a>].</p> Full article ">Figure 12
<p>Residual phase strains of austenite (<span class="html-italic">γ</span>), ferrite (<span class="html-italic">α</span>), and deformation-induced martensite (<span class="html-italic">α′</span>) phases as functions of true strain obtained by test B and the tensile test at 296 K [<a href="#B7-metals-11-02053" class="html-bibr">7</a>].</p> Full article ">Figure 13
<p>Phase strains of austenite (<span class="html-italic">γ</span>), ferrite (<span class="html-italic">α</span>), and deformation-induced martensite (<span class="html-italic">α′</span>) phases as functions of true strain obtained by <span class="html-italic">in situ</span> neutron diffraction experiments during tensile deformation at 296 K (<b>a</b>) [<a href="#B7-metals-11-02053" class="html-bibr">7</a>], 245 K (<b>b</b>), 188 K (<b>c</b>), and 128 K (<b>d</b>).</p> Full article ">Figure 14
<p>Volume fraction of deformation-induced martensite as a function of true strain obtained by <span class="html-italic">in situ</span> neutron diffraction experiments during tensile deformation at 296 K [<a href="#B7-metals-11-02053" class="html-bibr">7</a>], 245 K, 188 K, and 128 K. Here, the dashed lines are calculated by using Equation (4).</p> Full article ">Figure 15
<p>Calculated fraction-weighted phase stresses for ferrite phase (<b>a</b>) and deformation-induced martensite one (<b>b</b>) as a function of true strain in test B, the tensile tests at 296 K and 128 K.</p> Full article ">
15 pages, 5199 KiB  
Article
Effect of Ta and W Additions on Microstructure and Mechanical Properties of Tilt-Cast Ti-45Al-5Nb-2C Alloy
by Juraj Lapin and Kateryna Kamyshnykova
Metals 2021, 11(12), 2052; https://doi.org/10.3390/met11122052 - 18 Dec 2021
Cited by 6 | Viewed by 2979
Abstract
The effect of Ta and W additions on microstructure and mechanical properties of tilt-cast Ti-45Al-5Nb-2C (at.%) alloy was investigated. Three alloys with nominal composition Ti-45Al-5Nb-2C-2X (in at.%), where X is Ta or W, were prepared by vacuum induction melting in graphite crucibles followed [...] Read more.
The effect of Ta and W additions on microstructure and mechanical properties of tilt-cast Ti-45Al-5Nb-2C (at.%) alloy was investigated. Three alloys with nominal composition Ti-45Al-5Nb-2C-2X (in at.%), where X is Ta or W, were prepared by vacuum induction melting in graphite crucibles followed by tilt casting into graphite moulds. The microstructure of the tilt-cast alloys consists of the α2(Ti3Al) + γ(TiAl) lamellar grains, single γ phase, (Ti,Nb,X)2AlC particles with a small amount of (Ti,Nb,X)C, and β/B2 phase identified only in W containing alloy. The EDS analysis shows that Ta segregates into the carbide particles and reduces dissolution of Nb in both (Ti,Nb,Ta)C and (Ti,Nb,Ta)2AlC phases. The alloying with W reduces Nb content in both carbide phases and leads to stabilisation of β/B2 phase in the lamellar α2 + γ regions. The alloying with Ta and W does not affect the volume fraction of the carbide particles but influences their size and morphology. While the alloying with Ta and W has no significant effect on Vickers hardness and the indentation elastic modulus of the studied alloys, the addition of Ta affects the nanohardness and elastic modulus of the (Ti,Nb,Ta)2AlC phase. The addition of W significantly increases the Vickers microhardness of the lamellar α2 + γ regions. Full article
(This article belongs to the Special Issue TiAl-Based Alloys and Their Applications)
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<p>BSE micrographs showing the microstructure of tilt-cast alloys: (<b>a</b>,<b>b</b>) Nb, (<b>c</b>,<b>d</b>) NbTa, and (<b>e</b>,<b>f</b>) NbW.</p> Full article ">Figure 2
<p>The typical XRD patterns of tilt-cast Nb, NbTa, and NbW alloys.</p> Full article ">Figure 3
<p>SEM micrograph and corresponding EDS map analysis of the tilt-cast Nb alloy.</p> Full article ">Figure 4
<p>SEM micrograph and corresponding EDS map analysis of the tilt-cast NbTa alloy.</p> Full article ">Figure 5
<p>SEM micrograph and corresponding EDS map analysis of the tilt-cast NbW alloy.</p> Full article ">Figure 6
<p>Measured grain size and interlamellar spacing in Nb, NbTa, and NbW alloys after tilt casting: (<b>a</b>) log-normal distribution curves of grain size, (<b>b</b>) log-normal distribution curves of α<sub>2</sub>-α<sub>2</sub> interlamellar spacing. The types of the distribution curves are marked in the figure.</p> Full article ">Figure 7
<p>Vickers hardness and elastic modulus of tilt-cast Nb, NbTa, and NbW alloys: (<b>a</b>) Vicker hardness HV30, (<b>b</b>) indentation elastic modulus. The studied alloys are marked in the figures.</p> Full article ">Figure 8
<p>Nanohardness and elastic modulus of Ti<sub>2</sub>AlC particles in the tilt-cast Nb, NbTa, and NbW alloys: (<b>a</b>) nanohardness, (<b>b</b>) indentation elastic modulus. The studied alloys are marked in the figures.</p> Full article ">Figure 9
<p>Vickers microhardness of lamellar α<sub>2</sub> + γ regions in the tilt-cast Nb, NbTa, and NbW alloys: (<b>a</b>) Vickers microhardness HV0.05, (<b>b</b>) dependence of Vickers microhardness HV0.05 on α<sub>2</sub>-α<sub>2</sub> interlamellar spacing. The studied alloys are marked in the figures.</p> Full article ">
13 pages, 3466 KiB  
Article
Formation of Complex Inclusions in Gear Steels for Modification of Manganese Sulphide
by Haseeb Ahmad, Baojun Zhao, Sha Lyu, Zongze Huang, Yingtie Xu, Sixin Zhao and Xiaodong Ma
Metals 2021, 11(12), 2051; https://doi.org/10.3390/met11122051 - 18 Dec 2021
Cited by 8 | Viewed by 4421
Abstract
Suitable MnS inclusions in gear steel can significantly improve the steel machinability and reduce the manufacturing costs. Two gear steel samples with different sulphur contents were prepared via aluminium deoxidation followed by calcium treatment. The shape, size, composition and percentage distribution of the [...] Read more.
Suitable MnS inclusions in gear steel can significantly improve the steel machinability and reduce the manufacturing costs. Two gear steel samples with different sulphur contents were prepared via aluminium deoxidation followed by calcium treatment. The shape, size, composition and percentage distribution of the inclusions present in the steel samples were analyzed using an electron probe micro-analysis (EPMA) technique. The average diameter of MnS precipitated on an oxide inclusion is less than 5 µm. It was found that the steel with high sulphur content contains a greater number of elongated MnS precipitates than low sulphur steel. Moreover, there are more oxide inclusions such as calcium-aluminates and spinels with a small amount of solid solution of (Ca,Mn)S in low content sulphur steel after calcium treatment, which indicates the modification of solid alumina inclusions into liquid aluminates. The typical inclusions generated in high sulphur steel are sulphide encapsulating oxide inclusions and some core oxides were observed as spinel. The formation mechanisms of complex inclusions with different sulphur and calcium contents are discussed. The results are in good agreement with thermodynamic calculations. Full article
(This article belongs to the Special Issue Fundamentals of Advanced Pyrometallurgy)
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<p>Typical morphology and composition of inclusions in both samples. (<b>a</b>) GS1 sample, (<b>b</b>) GS2 sample.</p> Full article ">Figure 2
<p>The size of different types of inclusions in the gear steels. (<b>a</b>) GS1 sample, (<b>b</b>) GS2 sample.</p> Full article ">Figure 3
<p>Pure MnS size comparison between the two samples.</p> Full article ">Figure 4
<p>The composition distribution of Al<sub>2</sub>O<sub>3</sub>-MgO-CaO inclusions on the Al<sub>2</sub>O<sub>3</sub>-MgO-CaO phase diagram. (<b>a</b>) GS1 sample, (<b>b</b>) GS2 sample.</p> Full article ">Figure 5
<p>Calculated liquidus projection of CaO-Al<sub>2</sub>O<sub>3</sub>-CaS system in low CaS region.</p> Full article ">Figure 6
<p>Equilibrium precipitation of inclusions and evolution during steel solidification. (<b>a</b>) GS1 sample, (<b>b</b>) GS2 sample.</p> Full article ">Figure 7
<p>Mass percentage of CaS and MnS in the solid solution of (Ca,Mn)S. (<b>a</b>) GS1 sample, (<b>b</b>) GS2 Sample.</p> Full article ">Figure 8
<p>The schematic diagram of inclusions formation at the different stages of Al deoxidation, Ca treatment and solidification.</p> Full article ">Figure 9
<p>Duplex inclusions formed by (Ca,Mn)S precipitated on and collided with Al<sub>2</sub>O<sub>3</sub>-MgO inclusions.</p> Full article ">Figure 10
<p>Stability diagram of inclusions formed in the Fe-C-Si-Mn-Cr-S-O-Al-Ca system at 1600 °C. (<b>a</b>) GS1 sample, (<b>b</b>) GS2 sample.</p> Full article ">
21 pages, 33585 KiB  
Article
Identification of a Spatio-Temporal Temperature Model for Laser Metal Deposition
by Matthias Kahl, Sebastian Schramm, Max Neumann and Andreas Kroll
Metals 2021, 11(12), 2050; https://doi.org/10.3390/met11122050 - 18 Dec 2021
Cited by 7 | Viewed by 3311
Abstract
Laser-based additive manufacturing enables the production of complex geometries via layer-wise cladding. Laser metal deposition (LMD) uses a scanning laser source to fuse in situ deposited metal powder layer by layer. However, due to the excessive number of influential factors in the physical [...] Read more.
Laser-based additive manufacturing enables the production of complex geometries via layer-wise cladding. Laser metal deposition (LMD) uses a scanning laser source to fuse in situ deposited metal powder layer by layer. However, due to the excessive number of influential factors in the physical transformation of the metal powder and the highly dynamic temperature fields caused by the melt pool dynamics and phase transitions, the quality and repeatability of parts built by this process is still challenging. In order to analyze and/or predict the spatially varying and time dependent thermal behavior in LMD, extensive work has been done to develop predictive models usually by using finite element method (FEM). From a control-oriented perspective, simulations based on these models are computationally too expensive and are thus not suitable for real-time control applications. In this contribution, a spatio-temporal input–output model based on the heat equation is proposed. In contrast to other works, the parameters of the model are directly estimated from measurements of the LMD process acquired with an infrared (IR) camera during processing specimens using AISI 316 L stainless steel. In order to deal with noisy data, system identification techniques are used taking different disturbing noise into account. By doing so, spatio-temporal models are developed, enabling the prediction of the thermal behavior by means of the radiance measured by the IR camera in the range of the considered processing parameters. Furthermore, in the considered modeling framework, the computational effort for thermal prediction is reduced compared to FEM, thus enabling the use in real-time control applications. Full article
(This article belongs to the Special Issue Additive Manufacturing Processes in Metals)
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<p>Schematic of the experimental setup.</p> Full article ">Figure 2
<p>Camera setup during the LMD process. The protective enclosure is only opened for the photography of the setup.</p> Full article ">Figure 3
<p>Produced specimen.</p> Full article ">Figure 4
<p>Temperature evolution in five distinct points of the first weld bead.</p> Full article ">Figure 5
<p>Resultant false color image of the specimen for time instance at the end of the build process.</p> Full article ">Figure 6
<p>False color image with predefined grid before (<b>a</b>) and after (<b>b</b>) spatial downsampling at time instance <math display="inline"><semantics> <mrow> <mi>k</mi> <mo>=</mo> <mn>4000</mn> </mrow> </semantics></math>.</p> Full article ">Figure 7
<p>Thermal evolution in one discrete point <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <msub> <mi>s</mi> <mn>1</mn> </msub> <mo>,</mo> <msub> <mi>s</mi> <mn>2</mn> </msub> <mo>)</mo> </mrow> <mo>=</mo> <mrow> <mo>(</mo> <mn>7</mn> <mo>,</mo> <mn>2</mn> <mo>)</mo> </mrow> </mrow> </semantics></math> after spatial downsampling. Left axis (blue): Measured temperature (solid line) and corrected temperature values (dashed line). Right axis (red): Corresponding laser power. (<b>a</b>) First 3 s of experiment; (<b>b</b>) Hole duration of experiment.</p> Full article ">Figure 8
<p>Illustration of required input and output lags (filled circles) for calculation of <math display="inline"><semantics> <mrow> <mi>y</mi> <mo>(</mo> <mi>k</mi> <mo>,</mo> <mi>s</mi> <mo>)</mo> </mrow> </semantics></math> based on Equation (<a href="#FD8-metals-11-02050" class="html-disp-formula">8</a>) on two-dimensional lattice. (<b>a</b>) Index set <math display="inline"><semantics> <msub> <mi mathvariant="script">M</mi> <mi>y</mi> </msub> </semantics></math>; (<b>b</b>) index set <math display="inline"><semantics> <msub> <mi mathvariant="script">M</mi> <mi>u</mi> </msub> </semantics></math>.</p> Full article ">Figure 9
<p>Illustration of spatial discretization for one weld bead.</p> Full article ">Figure 10
<p>Illustration of spatial discretization for one layer: (<b>a</b>) Schematic of specimen with grid, (<b>b</b>) corresponding false color image with grid.</p> Full article ">Figure 11
<p>Illustration of spatial discretization of training data for single track case: (<b>a</b>) Schematic of specimen with grid (white circles represent fictional boundary grid points), (<b>b</b>) corresponding false color image with grid.</p> Full article ">Figure 12
<p><span class="html-italic">y</span> versus <math display="inline"><semantics> <mover accent="true"> <mi>y</mi> <mo>^</mo> </mover> </semantics></math> plots for single track case. (<b>a</b>) ARX model predictions on training data; (<b>b</b>) ARX model predictions on validation data; (<b>c</b>) OE model predictions on training data; (<b>d</b>) OE model predictions on validation data.</p> Full article ">Figure 13
<p>Time series for single track case at spatial instance <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <msub> <mi>s</mi> <mn>1</mn> </msub> <mo>,</mo> <msub> <mi>s</mi> <mn>2</mn> </msub> <mo>)</mo> </mrow> <mo>=</mo> <mrow> <mo>(</mo> <mn>5</mn> <mo>,</mo> <mn>2</mn> <mo>)</mo> </mrow> </mrow> </semantics></math> for both, training and validation data. Left axis (blue): Observed temperature <span class="html-italic">y</span> (solid line), predicted temperature <math display="inline"><semantics> <mover accent="true"> <mi>y</mi> <mo>^</mo> </mover> </semantics></math> (dashed line), and absolute value of residuals <math display="inline"><semantics> <mi>ϵ</mi> </semantics></math> (dotted line). Right axis (red): Corresponding laser power. (<b>a</b>) ARX model, training data; (<b>b</b>) ARX model, validation data; (<b>c</b>) OE model, training data; (<b>d</b>) OE model, validation data.</p> Full article ">Figure 13 Cont.
<p>Time series for single track case at spatial instance <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <msub> <mi>s</mi> <mn>1</mn> </msub> <mo>,</mo> <msub> <mi>s</mi> <mn>2</mn> </msub> <mo>)</mo> </mrow> <mo>=</mo> <mrow> <mo>(</mo> <mn>5</mn> <mo>,</mo> <mn>2</mn> <mo>)</mo> </mrow> </mrow> </semantics></math> for both, training and validation data. Left axis (blue): Observed temperature <span class="html-italic">y</span> (solid line), predicted temperature <math display="inline"><semantics> <mover accent="true"> <mi>y</mi> <mo>^</mo> </mover> </semantics></math> (dashed line), and absolute value of residuals <math display="inline"><semantics> <mi>ϵ</mi> </semantics></math> (dotted line). Right axis (red): Corresponding laser power. (<b>a</b>) ARX model, training data; (<b>b</b>) ARX model, validation data; (<b>c</b>) OE model, training data; (<b>d</b>) OE model, validation data.</p> Full article ">Figure 14
<p><span class="html-italic">y</span> versus <math display="inline"><semantics> <mover accent="true"> <mi>y</mi> <mo>^</mo> </mover> </semantics></math> plots for multiple tracks experiment. (<b>a</b>) ARX model predictions on training data; (<b>b</b>) ARX model predictions on validation data; (<b>c</b>) OE model predictions on training data; (<b>d</b>) OE model predictions on validation data.</p> Full article ">Figure 15
<p>Time series for multiple tracks case at spatial instances <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <msub> <mi>s</mi> <mn>1</mn> </msub> <mo>,</mo> <msub> <mi>s</mi> <mn>2</mn> </msub> <mo>)</mo> </mrow> <mo>=</mo> <mrow> <mo>{</mo> <mrow> <mo>(</mo> <mn>3</mn> <mo>,</mo> <mn>3</mn> <mo>)</mo> </mrow> <mo>,</mo> <mrow> <mo>(</mo> <mn>7</mn> <mo>,</mo> <mn>7</mn> <mo>)</mo> </mrow> <mo>}</mo> </mrow> </mrow> </semantics></math> on training and <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <msub> <mi>s</mi> <mn>1</mn> </msub> <mo>,</mo> <msub> <mi>s</mi> <mn>2</mn> </msub> <mo>)</mo> </mrow> <mo>=</mo> <mrow> <mo>(</mo> <mn>12</mn> <mo>,</mo> <mn>12</mn> <mo>)</mo> </mrow> </mrow> </semantics></math> on validation data. Left axis (blue): Observed temperature <span class="html-italic">y</span> (solid line), predicted temperature <math display="inline"><semantics> <mover accent="true"> <mi>y</mi> <mo>^</mo> </mover> </semantics></math> (dashed line), and absolute value of residual <math display="inline"><semantics> <mi>ϵ</mi> </semantics></math> (dotted line). Right axis (red): Corresponding laser power. (<b>a</b>) Time series at spatial instance (3,3) on training data; (<b>b</b>) Time series at spatial instance (7,7) on training data; (<b>c</b>) Time series at spatial instance (12,12) on validation data.</p> Full article ">Figure 15 Cont.
<p>Time series for multiple tracks case at spatial instances <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <msub> <mi>s</mi> <mn>1</mn> </msub> <mo>,</mo> <msub> <mi>s</mi> <mn>2</mn> </msub> <mo>)</mo> </mrow> <mo>=</mo> <mrow> <mo>{</mo> <mrow> <mo>(</mo> <mn>3</mn> <mo>,</mo> <mn>3</mn> <mo>)</mo> </mrow> <mo>,</mo> <mrow> <mo>(</mo> <mn>7</mn> <mo>,</mo> <mn>7</mn> <mo>)</mo> </mrow> <mo>}</mo> </mrow> </mrow> </semantics></math> on training and <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <msub> <mi>s</mi> <mn>1</mn> </msub> <mo>,</mo> <msub> <mi>s</mi> <mn>2</mn> </msub> <mo>)</mo> </mrow> <mo>=</mo> <mrow> <mo>(</mo> <mn>12</mn> <mo>,</mo> <mn>12</mn> <mo>)</mo> </mrow> </mrow> </semantics></math> on validation data. Left axis (blue): Observed temperature <span class="html-italic">y</span> (solid line), predicted temperature <math display="inline"><semantics> <mover accent="true"> <mi>y</mi> <mo>^</mo> </mover> </semantics></math> (dashed line), and absolute value of residual <math display="inline"><semantics> <mi>ϵ</mi> </semantics></math> (dotted line). Right axis (red): Corresponding laser power. (<b>a</b>) Time series at spatial instance (3,3) on training data; (<b>b</b>) Time series at spatial instance (7,7) on training data; (<b>c</b>) Time series at spatial instance (12,12) on validation data.</p> Full article ">
17 pages, 8169 KiB  
Article
Corrosion Nature in [CoN/AlN]n Multilayers Obtained from Laser Ablation
by Julio Caicedo, Neufer Bonilla and Willian Aperador
Metals 2021, 11(12), 2049; https://doi.org/10.3390/met11122049 - 17 Dec 2021
Cited by 2 | Viewed by 2423
Abstract
The aim of this work is the improvement of the electrochemical behavior of industrial steel using [CoN/AlN]n multilayered system via reactive Pulsed Laser Deposition (PLD) technique with a Nd: YAG laser (λ = 1064 nm) on Silicon (100) and AISI 302 steel substrates. [...] Read more.
The aim of this work is the improvement of the electrochemical behavior of industrial steel using [CoN/AlN]n multilayered system via reactive Pulsed Laser Deposition (PLD) technique with a Nd: YAG laser (λ = 1064 nm) on Silicon (100) and AISI 302 steel substrates. In this work was varied systematically the bilayer period (Λ) and the coatings were characterized by X-ray diffraction (XRD), Fourier transform infrared spectroscopy (FTIR), scanning electron microscopy (SEM), and the chemical composition was determined by means of X-ray photoelectron spectroscopy (XPS). The maximum corrosion resistance for the coating with (Λ) equal to 34.7 nm, corresponding to n = 30 bilayered. The polarization resistance and corrosion rate were around 7.62 × 105 kOhm × cm2 and 7.25 × 10−5 mm/year, these values were 6.3 × 105 and 78.6 times better than those showed by the uncoated 302 stainless steel substrate (1.2 kOhm × cm2 and 0.0057 mm/year), respectively. The improvement of the electrochemical behavior of the steel 302 coated with this [CoN/AlN]n can be attributed to the presence of several interfaces that act as obstacles for the inward and outward diffusions of Cl− ions, generating an increment in the corrosion resistance. The electrochemical results found in the [CoN/AlN]n open a possibility of future applications in mechanical devices that require high demands in service conditions. Full article
(This article belongs to the Special Issue Corrosion and Surface Modification of Metallic Materials)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>Image associate to PLD (pulsed laser deposition) device and the experimental processes used for the deposition of coatings by means of PLD.</p> Full article ">Figure 2
<p>Diffractograms of the AlN and CoN single layer and [AlN/CoN]n multilayer coatings with different number of bilayers (n = 10, 20, 30). The dotted lines indicate the position of the peaks obtained from international index files (JCPDF).</p> Full article ">Figure 3
<p>XPS depth spectra of the coatings: (<b>a</b>) AlN single layer and (<b>b</b>) CoN single layer.</p> Full article ">Figure 4
<p>High resolution XPS spectra for AlN and CoN layers: (<b>a</b>) N1s signal in the (AlN) layer, (<b>b</b>) N1s signal in the (CoN) layer, (<b>c</b>) Al2p signal in the (AlN) layer and (<b>d)</b> Co2p signal in the (CoN) layer.</p> Full article ">Figure 5
<p>Elemental percentages according to the coating stoichiometry (<b>a</b>) aluminum nitride layer (Al<sub>74</sub>N<sub>26</sub>), (<b>b</b>) cobalt nitride layer (Co<sub>69</sub>N<sub>31</sub>).</p> Full article ">Figure 6
<p>FTIR spectra obtained from the single layer coatings: (<b>a</b>) Aluminum nitride (AlN) and (<b>b</b>) Cobalt nitride (CoN).</p> Full article ">Figure 7
<p>Unit cell crystallographic simulation obtained using CaRIne Crystallography 3.1 software: (<b>a</b>) AlN materials, (<b>b</b>) CoN materials and (<b>c</b>) [AlN/CoN]n mulyilayers.</p> Full article ">Figure 8
<p>Diffraction patterns obtained for all coatings: (<b>a</b>) AlN diffraction pattern (experimental), (<b>b</b>) AlN diffraction pattern (simulated), (<b>c</b>) CoN diffraction pattern (experimental) and (<b>d</b>) CoN diffraction (simulated).</p> Full article ">Figure 8 Cont.
<p>Diffraction patterns obtained for all coatings: (<b>a</b>) AlN diffraction pattern (experimental), (<b>b</b>) AlN diffraction pattern (simulated), (<b>c</b>) CoN diffraction pattern (experimental) and (<b>d</b>) CoN diffraction (simulated).</p> Full article ">Figure 9
<p>SEM micrographs of the cross section of the [AlN/CoN]n multilayer coatings: (<b>a</b>) [AlN/CoN]<sub>10</sub> Λ = 104.3 nm and (<b>b</b>) [AlN/CoN]<sub>30</sub> (Λ = 34.7 nm).</p> Full article ">Figure 10
<p>Electrochemical impedance diagrams for AISI 302 steel coated with AlN and CoN single layers and multilayer systems [AlN/CoN]n.</p> Full article ">Figure 11
<p>Equivalent circuit used to fit the experimental data in the multilayer systems [AlN/CoN]n.</p> Full article ">Figure 12
<p>Tafel curves for AISI 302 steel coated with AlN and CoN single layers and [AlN/CoN]n multilayer systems.</p> Full article ">Figure 13
<p>Bilayers number effect on the [AlN/CoN]n multilayers system: (<b>a</b>) polarization resistance and (<b>b</b>) corrosion rate.</p> Full article ">Figure 14
<p>Electrochemical properties: (<b>a</b>) Porosity factor of coatings and (<b>b</b>) coating efficiency percentage.</p> Full article ">Figure 15
<p>SEM micrographs of the corrosive process for uncoated and coated steel substrates: (<b>a</b>) uncoated stainless steel, (<b>b</b>) cobalt nitride (single layer), (<b>c</b>) aluminum nitride (single layer) and (<b>d</b>) multilayer [AlN/CoN]<sub>30</sub>.</p> Full article ">Figure 15 Cont.
<p>SEM micrographs of the corrosive process for uncoated and coated steel substrates: (<b>a</b>) uncoated stainless steel, (<b>b</b>) cobalt nitride (single layer), (<b>c</b>) aluminum nitride (single layer) and (<b>d</b>) multilayer [AlN/CoN]<sub>30</sub>.</p> Full article ">Figure 16
<p>Correlation between corrosion rate and coating efficiency for the [AlN/CoN]<sub>n</sub> multilayer coatings deposited as a function of the increasing bilayer number (n).</p> Full article ">
9 pages, 7927 KiB  
Article
Inconel 713C Coating by Cold Spray for Surface Enhancement of Inconel 718
by Kaiqiang Wu, Sin Wei Chee, Wen Sun, Adrian Wei-Yee Tan, Sung Chyn Tan, Erjia Liu and Wei Zhou
Metals 2021, 11(12), 2048; https://doi.org/10.3390/met11122048 - 17 Dec 2021
Cited by 16 | Viewed by 3483
Abstract
Inconel 713C is a nickel-based superalloy usually considered as a material of poor weldability due to its susceptibility to hot cracking in the heat-affected zones. Cold spray, a solid-state deposition technology that does not involve melting, can be proposed as a methodology to [...] Read more.
Inconel 713C is a nickel-based superalloy usually considered as a material of poor weldability due to its susceptibility to hot cracking in the heat-affected zones. Cold spray, a solid-state deposition technology that does not involve melting, can be proposed as a methodology to deposit Inconel 713C for surface enhancement of other target components. In this study, Inconel 713C coating was deposited on Inconel 718 substrate with a high-pressure cold spray system. The coating was characterized in terms of microstructure, hardness, and wear properties. The cold-sprayed Inconel 713C coating has a low porosity level and refined grain structures. Microhardness of the Inconel 713C coating was much higher than the Inconel 718 substrate. The sliding wear tests showed that the wear resistance of the cold-sprayed Inconel 713C coating is three times higher than the Inconel 718 substrate, making the coating a suitable protective layer. The main wear mechanisms of the coating include oxidation, tribo-film formation, and adhesive wear. Full article
(This article belongs to the Special Issue Advances in Welding, Joining and Surface Coating Technology)
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<p>(<b>a</b>) SEM image of Inconel 713C feedstock powder. (<b>b</b>) BSE-SEM image of the cross-section of a single Inconel 713C particle.</p> Full article ">Figure 2
<p>(<b>a</b>) OM image and (<b>b</b>) BSE-SEM images of the Inconel 713C coating on Inconel 718 substrate.</p> Full article ">Figure 3
<p>Microstructure analysis of the Inconel 713C coatings: (<b>a</b>) BSE-SEM image. EBSD images: (<b>b</b>) band contrast image showing grains, (<b>c</b>) IPF image showing grain orientations, and (<b>d</b>) Kernel average misorientation (KAM) map.</p> Full article ">Figure 4
<p>Microhardness of Inconel 713C coating and Inconel 718 substrate at different locations.</p> Full article ">Figure 5
<p>Specific wear rates of cold spray Inconel 713C coating and Inconel 718 substrate.</p> Full article ">Figure 6
<p>SEM micrographs of wear tracks of (<b>a</b>) cold spray Inconel 713C coating and (<b>b</b>) Inconel 718 substrate.</p> Full article ">Figure 7
<p>SEM micrographs showing adhesive wear and abrasive grooves on the wear tracks of (<b>a</b>) cold spray Inconel 713C coating and (<b>b</b>) Inconel 718 substrate.</p> Full article ">
13 pages, 3023 KiB  
Article
Fracture Toughness Characteristics of High-Manganese Austenitic Steel Plate for Application in a Liquefied Natural Gas Carrier
by Gyubaek An, Jeongung Park, Hongkyu Park and Ilwook Han
Metals 2021, 11(12), 2047; https://doi.org/10.3390/met11122047 - 17 Dec 2021
Cited by 8 | Viewed by 3381
Abstract
High-manganese austenitic steel was developed to improve the fracture toughness and safety of steel under cryogenic temperatures, and its austenite structure was formed by increasing the Mn content. The developed high-manganese austenitic steel was alloyed with austenite-stabilizing elements (e.g., C, Mn, and Ni) [...] Read more.
High-manganese austenitic steel was developed to improve the fracture toughness and safety of steel under cryogenic temperatures, and its austenite structure was formed by increasing the Mn content. The developed high-manganese austenitic steel was alloyed with austenite-stabilizing elements (e.g., C, Mn, and Ni) to increase cryogenic toughness. It was demonstrated that 30 mm thickness high-manganese austenitic steel, as well as joints welded with this steel, had a sufficiently higher fracture toughness than the required toughness values evaluated under the postulated stress conditions. High-manganese austenitic steel can be applied to large offshore and onshore LNG storage and fuel tanks located in areas experiencing cryogenic conditions. Generally, fracture toughness decreases at lower temperatures; therefore, cryogenic steel requires high fracture toughness to prevent unstable fractures. Brittle fracture initiation and arrest tests were performed using 30 mm thickness high-manganese austenitic steel and SAW joints. The ductile fracture resistance of the weld joints (weld metal, fusion line, fusion line + 2 mm) was investigated using the R-curve because a crack in the weld joint tends to deviate into the weld metal in the case of undermatched joints. The developed high-manganese austenitic steel showed little possibility of brittle fracture and a remarkably unstable ductile fracture toughness. Full article
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<p>Macrosection and microstructure of 30 mm thick steel plate. (<b>a</b>) Schematic of the groove shape. (<b>b</b>) Macrosection of SAW joint. (<b>c</b>) Microstructure of SAW joint.</p> Full article ">Figure 2
<p>Test specimens to evaluate the suitability of high-manganese austenitic steel. (<b>a</b>) CTOD specimen geometry of the weld joint. (<b>b</b>) Duplex ESSO test specimen geometry using the base metal. (<b>c</b>) Duplex ESSO test specimen geometry using the base metal. (<b>d</b>) Wide-plate test specimen geometry of the weld joint using the base metal.</p> Full article ">Figure 3
<p>Tensile test results of base metal at the head and tail of the steel plate.</p> Full article ">Figure 4
<p>Charpy impact test at −196 °C using base metal and strain-aged conditions.</p> Full article ">Figure 5
<p>Tensile test results using the SA weld joint.</p> Full article ">Figure 6
<p>Charpy impact test at −196 °C with weld joints.</p> Full article ">Figure 7
<p>CTOD test results at −165 °C using the base metal and weld joints. (In P-V curve, the yellow line shows the maximum load and displacement; the blue line shows the P-V curve.)</p> Full article ">Figure 8
<p>Duplex ESSO test results for the base metal and weld joints at −196 °C.</p> Full article ">Figure 9
<p>Example of wide-plate test R-curve recorded at room temperature.</p> Full article ">Figure 10
<p>Compared results using the three-point bending test and the wide-plate test on fracture resistance.</p> Full article ">Figure 11
<p>Fracture resistance of FL and FL + 2 in three-point bending test at RT and −165 °C.</p> Full article ">
21 pages, 9171 KiB  
Article
Numerical Investigation of Segregation Evolution during the Vacuum Arc Remelting Process of Ni-Based Superalloy Ingots
by Jiajun Cui, Baokuan Li, Zhongqiu Liu, Fengsheng Qi, Beijiang Zhang and Ji Zhang
Metals 2021, 11(12), 2046; https://doi.org/10.3390/met11122046 - 17 Dec 2021
Cited by 16 | Viewed by 3867
Abstract
Segregation defects greatly affect the service performance and working life of castings during the vacuum arc remelting (VAR) process. However, the corresponding research on the prediction of segregation defects is still not comprehensive. Through considering the influence of water-cooled crucible on the electromagnetic [...] Read more.
Segregation defects greatly affect the service performance and working life of castings during the vacuum arc remelting (VAR) process. However, the corresponding research on the prediction of segregation defects is still not comprehensive. Through considering the influence of water-cooled crucible on the electromagnetic field inside an ingot, a full-scale model for the comprehensive prediction of freckles and macrosegregation defects during the VAR process is developed in this paper. The macroscopic solute transport phenomenon and the segregation behavior of Ni-5.8 wt% Al-15.2 wt% Ta alloy are predicted. The results indicate that the freckles are mainly concentrated in the lower region of the ingot. With the growth of the ingot, the solute enrichment channels gradually develop into solute enrichment regions, and the channel segregation evolves into macrosegregation. The Lorentz force mainly affects the flow pattern at the top of the molten pool, while the complex flow of multiple vortices is dominated by thermosolutal buoyancy. The maximum and minimum relative segregation ratio inside the ingot can reach 290% and −90%, respectively, and the positive segregation region accounts for about 79% of the total volume. This paper provides a new perspective for understanding the segregation behavior inside the ingot by studying the segregation evolution during the VAR process. Full article
(This article belongs to the Special Issue Numerical Modeling of Metallurgical Processes: Continuous Casting and Electroslag Remelting)
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<p>Schematic diagram of the VAR process.</p> Full article ">Figure 2
<p>Coupled computational relationships for multiple physical fields during the VAR process.</p> Full article ">Figure 3
<p>Schematic diagram of mesh used in simulation: (<b>a</b>) local longitudinal section; (<b>b</b>) top view.</p> Full article ">Figure 4
<p>Schematic view of the boundary conditions: (<b>a</b>) simplified boundary condition; (<b>b</b>) actual boundary condition in this paper.</p> Full article ">Figure 5
<p>Contour of current density distribution changing with the arc radius when the current of the side arc accounts for 50% from this paper (left half) and the literature [<a href="#B3-metals-11-02046" class="html-bibr">3</a>] (right half): the ratio of arc radius to ingot radius is (<b>a</b>) 0.5; (<b>b</b>) 0.6; (<b>c</b>) 0.7.</p> Full article ">Figure 6
<p>Comparison of <span class="html-italic">C</span><sub>relative</sub> in the cavity: (<b>a</b>) at 400 s from the literature; (<b>b</b>) this paper; (<b>c</b>) at 0.025 m from the bottom of the cavity after solidification.</p> Full article ">Figure 7
<p>The distribution of the electromagnetic field inside the ingot: (<b>a</b>) current density using actual boundary; (<b>b</b>) magnetic induction intensity using actual boundary; (<b>c</b>) current density using simplified boundary; (<b>d</b>) magnetic induction intensity using simplified boundary.</p> Full article ">Figure 8
<p>The distribution of the Lorentz force during the VAR process.</p> Full article ">Figure 9
<p>Results at solidification time <span class="html-italic">t</span> = 1500 s: (<b>a</b>) temperature distribution; (<b>b</b>) velocity field in the channels and liquid fraction at X = 0; (<b>c</b>) streamlines and solute distribution; (<b>d</b>) thermosolutal buoyancy.</p> Full article ">Figure 10
<p>Results at time t = 3500 s: (<b>a</b>) streamlines and solute distribution; (<b>b</b>) temperature distribution.</p> Full article ">Figure 11
<p>Results after solidification: (<b>a</b>) solute distribution and segregation behavior; (<b>b</b>) temperature distribution.</p> Full article ">Figure 12
<p>Solute distribution of horizontal sections inside the ingot after solidification: (<b>a</b>) <span class="html-italic">Z</span> = 0.1 m; (<b>b</b>) <span class="html-italic">Z</span> = 0.25 m; (<b>c</b>) <span class="html-italic">Z</span> = 0.35 m; (<b>d</b>) <span class="html-italic">Z</span> = 0.45 m.</p> Full article ">Figure 13
<p>Variation of volume-averaged segregation extent during the VAR process.</p> Full article ">Figure 14
<p>Variation of relative segregation ratio along the centerline and midradius-line after solidification.</p> Full article ">Figure 15
<p>Corresponding volume of each range of relative segregation ratio after solidification.</p> Full article ">
17 pages, 11122 KiB  
Article
Effect of Sigma Phase in Wire Arc Additive Manufacturing of Superduplex Stainless Steel
by Odd M. Akselsen, Ruben Bjørge, Håkon Wiik Ånes, Xiaobo Ren and Bård Nyhus
Metals 2021, 11(12), 2045; https://doi.org/10.3390/met11122045 - 17 Dec 2021
Cited by 8 | Viewed by 3399
Abstract
In the present study, the thermal program in wire and arc additive manufacturing has been varied in terms of heat input and interpass temperature. Three walls were completed with subsequent Charpy V impact toughness and crack-tip opening displacement fracture toughness, together with a [...] Read more.
In the present study, the thermal program in wire and arc additive manufacturing has been varied in terms of heat input and interpass temperature. Three walls were completed with subsequent Charpy V impact toughness and crack-tip opening displacement fracture toughness, together with a detailed microstructure characterization using light microscopy and scanning and transmission electron microscopy. The results clearly demonstrate that the formation of sigma phase may deteriorate the toughness of superduplex components. Such formation may take place under prolonged cooling time, which may occur when subsequent passes are deposited with too high interpass temperatures. This transformation behavior may limit the productivity in additive manufacturing of such steels and care must be taken in selection of proper combination of arc energy and interpass temperature. Full article
(This article belongs to the Special Issue Wire Arc Additive Manufacturing of Metallic Components)
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<p>Microstructure of the building plate (austenite is white, ferrite is brownish).</p> Full article ">Figure 2
<p>(<b>a</b>) Macrographs of walls (cross sections). (<b>b</b>) Schematic illustration of the walls with the Charpy and crack tip opening displacement (CTOD) specimens indicated.</p> Full article ">Figure 3
<p>Hardness in different regions. Error bars indicate standard deviation.</p> Full article ">Figure 4
<p>Charpy V results; tested at −46 °C. Data for −20 °C taken from Lervåg et al. [<a href="#B19-metals-11-02045" class="html-bibr">19</a>]. (<b>a</b>) Effect of mean heat input, and (<b>b</b>) effect of mean interpass temperature. The results from the present study are compared with previous data for superduplex steels, indicating that there is a certain temperature effect; the previous results are from testing at −20 °C while the current walls are tested at −46 °C. The previously published data by Lervåg et al. [<a href="#B19-metals-11-02045" class="html-bibr">19</a>] illustrate that the toughness of the walls is nearly independent of the gross arc energy due to limited microstructure variations; the Charpy values are all close to 100 J. However, it is reasonable to assume that the major reason for the toughness fall is due to σ formation, especially for the 0.63 kJ/mm heat input combined with high interpass temperature.</p> Full article ">Figure 5
<p>CTOD fracture toughness; tested at −46 °C. Data for −20 °C taken from Lervåg et al. [<a href="#B19-metals-11-02045" class="html-bibr">19</a>]. (<b>a</b>) Effect of mean heat input, and (<b>b</b>) effect of mean interpass temperature (left).</p> Full article ">Figure 6
<p>Epitaxial solidification in Wall 1.</p> Full article ">Figure 7
<p>Microstructure of primary (non-reheated) region, next last layer of Wall 1.</p> Full article ">Figure 8
<p>Reheated layer microstructure of Wall 1.</p> Full article ">Figure 9
<p>Ferrite content in different wall regions. Error bars indicate standard deviation.</p> Full article ">Figure 10
<p>HAZ microstructure.</p> Full article ">Figure 11
<p>Sigma phase content as a function of mean interpass temperature. The black curve is an exponential fit.</p> Full article ">Figure 12
<p>EBSD phase maps from Wall 1; ferrite is green, austenite is red.</p> Full article ">Figure 13
<p>EBSD phase maps from Wall 1 showing the presence of secondary austenite. Ferrite is green, austenite is red. (<b>b</b>) is close-up of square in (<b>a</b>). Arrows point to γ<sub>2</sub>.</p> Full article ">Figure 14
<p>Sigma phase in Wall 3; (<b>a</b>) Backscattered electron (BSE) image showing the microstructure (Sigma is white), (<b>b</b>) EDX maps of Fe, Cr, Ni, and Mo.</p> Full article ">Figure 14 Cont.
<p>Sigma phase in Wall 3; (<b>a</b>) Backscattered electron (BSE) image showing the microstructure (Sigma is white), (<b>b</b>) EDX maps of Fe, Cr, Ni, and Mo.</p> Full article ">Figure 15
<p>EBSD phase map showing an area from Wall 3 with significant amounts of sigma phase. The center column shows four experimental patterns taken from the positions indicated in <a href="#metals-11-02045-t003" class="html-table">Table 3</a>. sample (<a href="#metals-11-02045-f016" class="html-fig">Figure 16</a>) also revealed the presence of the chi phase, which has a similar composition to the sigma phase. However, the fact that this phase was not found in the SEM investigation suggests that this phase is present in much lower quantities than the sigma phase.</p> Full article ">Figure 16
<p>(<b>a</b>) HAADF STEM image and corresponding EDX maps from an area of Wall 3 with several particles. The Mo-rich particles were identified as the bcc chi phase. The oxide particle might be of the bcc Mn<sub>3</sub>Cr<sub>2</sub>(SiO<sub>4</sub>)<sub>3</sub> phase (Ottonello, 1996 [<a href="#B25-metals-11-02045" class="html-bibr">25</a>]). (<b>b</b>) Diffraction pattern from the [011] zone axis of the chi particle pointed to by a black arrow in (<b>a</b>). The Miller indices of the circled reflections are indicated. (<b>c</b>) Diffraction pattern from the same chi particle near the [-111] zone axis, tilted 35° from (<b>b</b>).</p> Full article ">Figure 17
<p>SEM image of Cr nitrides in HAZ of building plate.</p> Full article ">
19 pages, 11990 KiB  
Article
Influence of High Pressure Sliding and Rotary Swaging on Creep Behavior of P92 Steel at 500 °C
by Petr Kral, Jiri Dvorak, Vaclav Sklenicka, Zenji Horita, Yoichi Takizawa, Yongpeng Tang, Lenka Kunčická, Marie Kvapilova and Marie Ohankova
Metals 2021, 11(12), 2044; https://doi.org/10.3390/met11122044 - 16 Dec 2021
Cited by 6 | Viewed by 2158
Abstract
High-pressure sliding (HPS) and rotary swaging (RS) at room temperature were used to form severely deformed microstructures in martensitic creep-resistant P92 steel. The deformed microstructures contained markedly different ratios of low- and high-angle grain boundaries (LAGBs/HAGBs). The application of the RS method, with [...] Read more.
High-pressure sliding (HPS) and rotary swaging (RS) at room temperature were used to form severely deformed microstructures in martensitic creep-resistant P92 steel. The deformed microstructures contained markedly different ratios of low- and high-angle grain boundaries (LAGBs/HAGBs). The application of the RS method, with an imposed equivalent strain of 1.4, led to the formation of a heterogeneous microstructure with a high number of LAGBs, while the HPS method, with an imposed equivalent strain of 7.8, led to the formation of a relatively homogeneous ultrafine-grained microstructure with a significant predominance of HAGBs. Microstructure analyses after creep testing showed that the microstructure of RS- and HPS-processed P92 steel is quite stable, but a slight coarsening of subgrains and grains during creep testing can be observed. Constant load tensile creep tests at 500 °C and initial stresses ranging from 300 to 900 MPa revealed that the specimens processed by HPS exhibited higher creep strength (slower minimum creep rate) and ductility compared to the coarse-grained and RS-processed P92 steel. However, the HPS-processed P92 steel also exhibited lower values of stress exponent n than the other investigated states of P92 steel. For this reason, the differences in minimum creep rates determined for different states decrease with decreasing values of applied stress, and at applied stresses lower than 500 MPa, the creep resistance of the RS-processed state is higher than the creep resistance of the HPS-processed state. Full article
(This article belongs to the Special Issue Severe Plastic Deformation Techniques of Metal Alloys)
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<p>Microstructure of P92 steel before creep testing (<b>a</b>) CG state, (<b>b</b>) RS-processed state, (<b>c</b>) HPS-processed state and (<b>d</b>) density of GND.</p> Full article ">Figure 2
<p>Dependences of strain rate vs. strain (<b>a</b>) CG state, (<b>b</b>) RS-processed state, (<b>c</b>) HPS-processed state and (<b>d</b>) influence of annealing on creep behavior.</p> Full article ">Figure 3
<p>Stress dependences of (<b>a</b>) minimum creep rates and (<b>b</b>) time to fracture for CG, RS and HPS-processed states tested at 500 °C.</p> Full article ">Figure 4
<p>Microstructure of CG state tested at 500 °C and 400 MPa (<b>a</b>) grip part, (<b>b</b>) near the fracture, (<b>c</b>) misorientation distributions measured in different locations and (<b>d</b>) pole figures for {110} planes in the grip part and near the fracture, X direction is parallel to the stress axis.</p> Full article ">Figure 5
<p>Density of GND in CG state tested at 500 °C and 400 MPa determined in various locations of the tensile specimen.</p> Full article ">Figure 6
<p>Microstructure of RS-processed P92 steel tested at 500 °C and 600 MPa (<b>a</b>) grip part, (<b>b</b>) near the fracture, (<b>c</b>) misorientation distributions measured in different locations and (<b>d</b>) pole figures for {110} planes in the grip part and the near fracture, X direction is parallel to the stress axis.</p> Full article ">Figure 6 Cont.
<p>Microstructure of RS-processed P92 steel tested at 500 °C and 600 MPa (<b>a</b>) grip part, (<b>b</b>) near the fracture, (<b>c</b>) misorientation distributions measured in different locations and (<b>d</b>) pole figures for {110} planes in the grip part and the near fracture, X direction is parallel to the stress axis.</p> Full article ">Figure 7
<p>Density of GND in RS-processed P92 steel tested at 500 °C determined in various locations of the tensile specimen.</p> Full article ">Figure 8
<p>Substructure and dislocation structure of RS-processed P92 steel tested at 700 MPa (<b>a</b>) grip part–substructure inside of long band, zone axis near [101], (<b>b</b>) grip part–dislocation inside subgrain, zone axis near [111], (<b>c</b>) gauge length–dislocation inside subgrain, zone axis near [111].</p> Full article ">Figure 9
<p>Microstructure of the HPS-processed P92 steel tested at 500 °C and 700 MPa (<b>a</b>) grip part, (<b>b</b>) near the fracture, (<b>c</b>) misorientation distributions measured in different locations and (<b>d</b>) pole figures for {110} planes in the grip part and near the fracture, X direction is parallel to the stress axis.</p> Full article ">Figure 10
<p>Density of GND in HPS-processed P92 steel tested at 500 °C determined in various locations of the tensile specimen.</p> Full article ">Figure 11
<p>Microstructure after creep testing at 300 MPa and creep strain about 0.04 (<b>a</b>) gauge length, (<b>b</b>) dislocations inside larger grain, zone axis near [111] and (<b>c</b>) dislocations in interior of fine grain, zone axis near [111].</p> Full article ">Figure 12
<p>Microstructure in the gauge length of the HPS-processed specimen tested at 500 °C and 300 Mpa. (<b>a</b>) Formation of the Laves phase, (<b>b</b>) distribution of Cr and (<b>c</b>) distribution of W in the microstructure.</p> Full article ">Figure 13
<p>Microstructure of HPS-processed P92 steel annealed at 650 °C for 500 h and tested at 500 °C and 400 MPa (<b>a</b>) grip part, (<b>b</b>) near the fracture, (<b>c</b>) misorientation distributions measured in different locations and (<b>d</b>) pole figures for {110} planes in the grip part and near the fracture, X direction is parallel to the stress axis.</p> Full article ">Figure 14
<p>Density of GND in HPS-processed P92 steel annealed at 650 °C for 500 h and tested at 400 MPa determined in various locations of the tensile specimen.</p> Full article ">Figure 15
<p>Microstructure characteristics measured in grip part (empty symbols) and gauge length (full symbols) of CG and SPD-processed P92 steel after creep testing at different stresses. (<b>a</b>) Subgrain, the crosses indicate the subgrain size measured by TEM in gauge length (<b>b</b>) grain size. The subgrain and grain sizes are compared with the expected stationary subgrain size with different value of <span class="html-italic">G</span> [<a href="#B30-metals-11-02044" class="html-bibr">30</a>,<a href="#B31-metals-11-02044" class="html-bibr">31</a>].</p> Full article ">
11 pages, 7312 KiB  
Article
Thermal Stability and Mechanical Properties of Al-Zn and Al-Bi-Zn Alloys Deformed by ECAP
by Hailong Jia, Yinan Piao, Kaining Zhu, Chaoran Yin, Wenqiang Zhou, Feng Li and Min Zha
Metals 2021, 11(12), 2043; https://doi.org/10.3390/met11122043 - 16 Dec 2021
Cited by 1 | Viewed by 2580
Abstract
It is well known that ultrafine grained and nanocrystalline materials show enhanced strength, while they are susceptible to thermally induced grain coarsening. The present work aims to enhance the thermal stability of ultrafine Al grains produced by equal channel angular pressing (ECAP) via [...] Read more.
It is well known that ultrafine grained and nanocrystalline materials show enhanced strength, while they are susceptible to thermally induced grain coarsening. The present work aims to enhance the thermal stability of ultrafine Al grains produced by equal channel angular pressing (ECAP) via dynamically precipitation. Detailed characterization by electron backscatter diffraction (EBSD) and transmission electron microscopy (TEM) has been carried out to reveal the microstructural evolution during both ECAP and post-ECAP annealing. After five passes of ECAP, both Al-8Zn and Al-6Bi-8Zn alloys show an ultrafine grain structure together with dynamic precipitated nanoscale Zn particles along grain boundaries. Upon annealing at 200 °C, ultrafine grains in the Al-8Zn and Al-6Bi-8Zn alloys show a remarkable thermal stability compared to the Al-8Bi alloy, which is mainly due to the presence of nanoscale Zn precipitates along grain boundaries. The present work reveals that nanoscale Zn particles have a positive effect on preserving the ultrafine grains during annealing, which is useful for the design of UFG Al alloys with improved thermal stability. Full article
(This article belongs to the Special Issue Composition Design, Structure, and Plastic Deformation of Al and Mg Alloys)
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<p>(<b>a</b>) Sketch of the ECAP die, showing the shear direction (SD), shear plane normal (SPN), and (<b>b</b>) EBSD, hardness, and tensile test measurement planes (ED-ND) with reference to the ECAP die geometry.</p> Full article ">Figure 2
<p>Microstructures of the as-cast Al-8Zn (<b>a</b>,<b>b</b>) and Al-6Bi-8Zn (<b>c</b>,<b>d</b>) alloys. (<b>a</b>,<b>c</b>) EBSD-IPF maps, (<b>b</b>,<b>d</b>) SEM-BSE images.</p> Full article ">Figure 3
<p>Microstructures of the 5P Al-8Zn and Al-6Bi-8Zn alloys. (<b>a</b>,<b>b</b>) EBSD, (<b>c</b>,<b>d</b>) TEM, and (<b>e</b>,<b>f</b>) STEM. In boundary maps, green, red, and blue lines depict boundaries with misorientations in the range of 2° ≤ θ &lt; 5°, 5° ≤ θ &lt; 15° and 15° ≤ θ &lt; 180°, respectively.</p> Full article ">Figure 4
<p>EBSD IPF maps and boundary maps of the annealed (at 200 °C) Al-8Zn and Al-6Bi-8Zn alloys. (<b>a</b>,<b>b</b>) 2 h, (<b>c</b>,<b>d</b>) 8 h, and (<b>e</b>,<b>f</b>) 96 h. In boundary maps, green, red, and blue lines depict boundaries with misorientations in the rage of 2° ≤ θ &lt; 5°, 5° ≤ θ &lt; 15°, and 15° ≤ θ &lt; 180°, respectively.</p> Full article ">Figure 5
<p>The evolution of average grain size for the as-deformed and annealed (at 200 °C) Al-8Zn and Al-6Bi-8Zn alloys.</p> Full article ">Figure 6
<p>{1 1 0} and {1 1 1} pole figures of ECAP-processed Al-8Zn and Al-6Bi-8Zn alloys. (<b>a</b>) and (<b>b</b>) 1P, (<b>c</b>) and (<b>d</b>) 3P, (<b>e</b>) and (<b>f</b>) 5P, and (<b>g</b>) is a schematic of the ECAP die and its relevant coordinate system. The dotted lines in the pole figures indicate the ideal shear plane (SD and SPN are abbreviations for shear plane and shear plane normal).</p> Full article ">Figure 7
<p>{1 1 0} and {1 1 1} pole figures of ECAP-processed (<b>a</b>–<b>c</b>) Al-8Zn alloy and (<b>d</b>–<b>f</b>) Al-6Bi-8Zn alloy. The dotted lines indicate the ideal shear plane, and SPN is the shear plane normal.</p> Full article ">Figure 8
<p>Evolution of Vickers hardness of the 5P Al-8Zn and Al-6Bi-8Zn alloys as a function of annealing time at 200 °C.</p> Full article ">Figure 9
<p>Tensile engineering stress–strain curves of the Al-8Zn and Al-6Bi-8Zn alloys.</p> Full article ">Figure 10
<p>Fracture morphologies. (<b>a</b>) SE images of the 5P Al-8Zn sample, (<b>b</b>) BSE image of the 5P Al-6Bi-8Zn sample.</p> Full article ">
17 pages, 11732 KiB  
Article
Basic Tool Design Guidelines for Friction Stir Welding of Aluminum Alloys
by Elizabeth Hoyos and María Camila Serna
Metals 2021, 11(12), 2042; https://doi.org/10.3390/met11122042 - 16 Dec 2021
Cited by 18 | Viewed by 6209
Abstract
Friction Stir Welding (FSW) is a solid-state welding process that has multiple advantages over fusion welding. The design of tools for the FSW process is a factor of interest, considering its fundamental role in obtaining sound welds. There are some commercially available alternatives [...] Read more.
Friction Stir Welding (FSW) is a solid-state welding process that has multiple advantages over fusion welding. The design of tools for the FSW process is a factor of interest, considering its fundamental role in obtaining sound welds. There are some commercially available alternatives for FSW tools, but unlike conventional fusion welding consumables, their use is limited to very specific conditions. In this work, equations to act as guidelines in the design process for FSW tools are proposed for the 2XXX, 5XXX, 6XXX, and 7XXX aluminum series and any given thickness to determine: pin length, pin diameter, and shoulder diameter. Over 80 sources and 200 tests were used and detailed to generate these expressions. As a verification approach, successful welds by authors outside the scope of the original review and the tools used were evaluated under this development and used as case studies or verification for the guidelines. Variations between designs made using the guidelines and those reported by other researchers remain under 21%. Full article
(This article belongs to the Special Issue Advances in Friction Stir Welding and Processing)
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<p>Types of shoulders.</p> Full article ">Figure 2
<p>Types of pins.</p> Full article ">Figure 3
<p>Pin diameter vs. thickness for series: (<b>a</b>) 2XXX; (<b>b</b>) 5XXX; (<b>c</b>) 6XXX, and (<b>d</b>) 7XXX.</p> Full article ">Figure 4
<p>Summary of trend lines for pin diameter vs. thickness.</p> Full article ">Figure 5
<p>Pin length vs. thickness for series: (<b>a</b>) 2XXX; (<b>b</b>) 5XXX; (<b>c</b>) 6XXX; and (<b>d</b>) 7XXX.</p> Full article ">Figure 6
<p>Summary of trend lines for pin length vs. thickness.</p> Full article ">Figure 7
<p>Shoulder diameter vs. thickness for series: (<b>a</b>) 2XXX; (<b>b</b>) 5XXX; (<b>c</b>) 6XXX; and (<b>d</b>) 7XXX.</p> Full article ">Figure 7 Cont.
<p>Shoulder diameter vs. thickness for series: (<b>a</b>) 2XXX; (<b>b</b>) 5XXX; (<b>c</b>) 6XXX; and (<b>d</b>) 7XXX.</p> Full article ">Figure 8
<p>Summary of trend lines for shoulder diameter vs. thickness.</p> Full article ">Figure 9
<p>Tool design: (<b>a</b>) pin; (<b>b</b>) shoulder.</p> Full article ">Figure 10
<p>X-ray of an AA7075-T6 aluminum FSW weld [<a href="#B117-metals-11-02042" class="html-bibr">117</a>].</p> Full article ">Figure 11
<p>Test plate dimensions (all units in mm).</p> Full article ">Figure 12
<p>Tool 1 design: (<b>a</b>) shoulder and (<b>b</b>) pin.</p> Full article ">Figure 13
<p>Tool 1 trial—X-ray of an AA 6061-T6 aluminum FSW weld.</p> Full article ">Figure 14
<p>Tool 2 trial—X-ray of an AA 6061-T6 aluminum FSW weld.</p> Full article ">Figure 15
<p>Ultrasound results with indication for Tool 2 trial (EPOCH 4 ultrasound system).</p> Full article ">Figure 16
<p>Cavity location for Tool 2 trial, according to ultrasound results (EPOCH 4 ultrasound system).</p> Full article ">
8 pages, 2334 KiB  
Article
Simulation for Cu Atom Diffusion Leading to Fluctuations in Solder Properties and Cu6Sn5 Growth during Multiple Reflows
by Min Shang, Chong Dong, Haoran Ma, Yunpeng Wang and Haitao Ma
Metals 2021, 11(12), 2041; https://doi.org/10.3390/met11122041 - 16 Dec 2021
Cited by 1 | Viewed by 2917
Abstract
The multiple reflows process is widely used in 3D packaging in the field of electronic packaging. The growth behavior of interfacial intermetallic compound (IMC) is more important to the reliability of solder joints. In this paper, experimental measurement combined with simulation calculation were [...] Read more.
The multiple reflows process is widely used in 3D packaging in the field of electronic packaging. The growth behavior of interfacial intermetallic compound (IMC) is more important to the reliability of solder joints. In this paper, experimental measurement combined with simulation calculation were preformed to investigate the evolution of Cu concentration in solders during multiple reflows, as well as its effects on the growth behavior of IMC and solder properties. The concentration of Cu in solder fluctuated, increasing with the increase of reflow times, which led to the fluctuation in the growth rate of the IMC. Furthermore, the Vickers hardness and melting point of the solder fluctuated during the multiple reflow processes due to the fluctuation in the Cu concentration. The data generated during this study could help to develop machine learning tools in relation to the study of interfacial microstructure evolution during multiple reflows. Full article
(This article belongs to the Special Issue Data-Driven Approaches in Modeling of Intermetallics)
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<p>Schematics of the experimental solder bump (<b>a</b>) and the temperature profile of the nine-time reflow process (<b>b</b>).</p> Full article ">Figure 2
<p>Top-view and cross-section morphologies of Cu<sub>6</sub>Sn<sub>5</sub> during the nine-time reflow process.</p> Full article ">Figure 3
<p>EPMA results for the Cu mass fraction in solder (<b>a</b>), plots of the measured and calculated Cu mass fractions (<b>b</b>) and simulations for the Cu distribution in solder balls (<b>c</b>) during the nine-time reflow process.</p> Full article ">Figure 4
<p>Sketch for the growth behavior of Cu<sub>6</sub>Sn<sub>5</sub> during multiple reflows. (<b>a</b>) Isothermal stage, (<b>b</b>) Cooling stage, (<b>c</b>) Next isothermal stage.</p> Full article ">Figure 5
<p>Evolutions of IMC thickness and growth rate (<b>a</b>), Vickers hardness of the solder bumps (<b>b</b>), DSC results (<b>c</b>) and melting points of the reflowed solders (<b>d</b>) during the nine-time reflow process.</p> Full article ">
9 pages, 2690 KiB  
Article
High-Temperature Corrosion Performance of FeAl-Based Alloys Containing Carbon in Molten Salt
by Munish Kumar, Ravi Kant, Suresh Chand, Ujjwal Prakash, Shankar Sehgal, Kuldeep Kumar Saxena, Joao Paulo Davim and Chander Prakash
Metals 2021, 11(12), 2040; https://doi.org/10.3390/met11122040 - 16 Dec 2021
Cited by 3 | Viewed by 2876
Abstract
Corrosion behavior of FeAl-based alloys containing carbon produced through arc melting in argon atmosphere has been studied at 500 °C to 700 °C. The samples were tested in the aggressive environment of molten salts (80%V2O5/20%Na2SO4). [...] Read more.
Corrosion behavior of FeAl-based alloys containing carbon produced through arc melting in argon atmosphere has been studied at 500 °C to 700 °C. The samples were tested in the aggressive environment of molten salts (80%V2O5/20%Na2SO4). The corrosion behavior was observed by weight change method and the layer products formed were examined by using X-ray diffraction (XRD), scanning electron microscopy (SEM) and energy-dispersive spectroscopy (EDS). The different phase components were observed in the surface layer after the test in Fe-22Al alloy. A protective Al2O3 layer was confirmed for Fe-22Al alloy containing carbon only. However, an additional TiO layer was also observed in Fe-22Al alloy containing carbon with Ti addition. The microstructural and XRD examinations revealed that this additional TiO layer protects better against penetration of corrosive media. The corrosion resistance behavior of FeAl-based alloys were addressed on the basis of microstructural evidence. Full article
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<p>SEM images of FeAl-based alloys (<b>a</b>) C01T0 with fine Fe<sub>3</sub>AlC<sub>0.5</sub>, (<b>b</b>) C1T0 with graphite and coarse Fe<sub>3</sub>AlC<sub>0.5</sub>, (<b>c</b>) C01T1 with Fe<sub>3</sub>AlC<sub>0.5</sub> and TiC and (<b>d</b>) C1T5 with Fe<sub>3</sub>AlC<sub>0.5</sub> and TiC.</p> Full article ">Figure 2
<p>EDS analysis of (<b>a</b>) needle like Fe<sub>3</sub>AlC<sub>0.5</sub> carbide in C01T0, (<b>b</b>) TiC, (<b>c</b>) matrix, and (<b>d</b>) needle like Fe<sub>3</sub>AlC<sub>0.5</sub> carbide in C1T5.</p> Full article ">Figure 2 Cont.
<p>EDS analysis of (<b>a</b>) needle like Fe<sub>3</sub>AlC<sub>0.5</sub> carbide in C01T0, (<b>b</b>) TiC, (<b>c</b>) matrix, and (<b>d</b>) needle like Fe<sub>3</sub>AlC<sub>0.5</sub> carbide in C1T5.</p> Full article ">Figure 3
<p>Variation of the weight gain of FeAl-based alloys with time at exposure of temperature (<b>a</b>) 500 °C, (<b>b</b>) 600 °C, and (<b>c</b>) 700 °C (lines joining the data points are for visual aid only).</p> Full article ">Figure 4
<p>Comparative study showing the weight gain of (<b>a</b>) C01T0 and (<b>b</b>) C1T5 at different temperatures (Lines joining the data points are for visual aid only).</p> Full article ">Figure 5
<p>X-ray diffraction (XRD) patterns of oxidized samples.</p> Full article ">Figure 6
<p>SEM Micrographs of (<b>a</b>) C01T0 and (<b>b</b>) C1T5 after oxidation at 700 °C /70 h and EDS elements analysis of cross-sections show the presence of Fe, O and Al in C01T0 and Fe, O, Ti and Al in C1T5.</p> Full article ">
18 pages, 45774 KiB  
Article
A Novel Approach to Inhibit Intergranular Corrosion in Ferritic Stainless Steel Welds Using High-Speed Laser Cladding
by Niklas Sommer, Lukas Grimm, Christian Wolf and Stefan Böhm
Metals 2021, 11(12), 2039; https://doi.org/10.3390/met11122039 - 15 Dec 2021
Cited by 5 | Viewed by 3444
Abstract
Ferritic stainless steels are prone to localized corrosion phenomena such as pitting corrosion or intergranular corrosion, in particular when jointed by fusion welding processes. State-of-the-art techniques to avoid intergranular corrosion mainly consist of alternating alloy concepts or post-weld heat-treatments—all of which are associated [...] Read more.
Ferritic stainless steels are prone to localized corrosion phenomena such as pitting corrosion or intergranular corrosion, in particular when jointed by fusion welding processes. State-of-the-art techniques to avoid intergranular corrosion mainly consist of alternating alloy concepts or post-weld heat-treatments—all of which are associated with increased production costs. Hence, the present investigation seeks to introduce a novel approach for the inhibition of intergranular corrosion in ferritic stainless steel welds through the use of high-speed laser cladding. Here, vulnerable sites prone to intergranular corrosion along the weld seam area are coated with a chemically resistant alloy, whereby an overlap is achieved. Optical and electron microscopy as well as computer tomography and tensile tests reveal that the detrimental effects of intergranular corrosion in both stabilized and unstabilized ferritic stainless steel are substantially reduced. In addition to that, the effects of varying overlap widths on the identified corrosion phenomena are studied. Moreover, the resulting dilution and precipiation phenomena at the clad–sheet interface are thoroughly characterized by electron backscatter diffraction and energy dispersive X-ray spectroscopy, whereby interrelationships to corrosion resistance can be drawn. As a result of this investigation, the number of techniques for the inhibition of intergranular corrosion is enlarged, and substantial cost-saving potentials in the manufacturing industry are unlocked. Full article
(This article belongs to the Special Issue Influence of Surface Treatment on Corrosion Behavior of Steels)
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<p>Schematic detailing the principle of a novel approach to inhibit intergranular corrosion of ferritic stainless steel welds by high-speed laser cladding.</p> Full article ">Figure 2
<p>Particle size distribution of the powder batch used in the present investigation. Nominal distribution 20–53 <math display="inline"><semantics> <mo mathvariant="normal">μ</mo> </semantics></math><math display="inline"><semantics> <mi mathvariant="normal">m</mi> </semantics></math>, partially adopted with permission from ref. [<a href="#B37-metals-11-02039" class="html-bibr">37</a>]. Copyright 2021 Niklas Sommer.</p> Full article ">Figure 3
<p>Initial weld seam macrostructure of (<b>a</b>) AISI 430 and (<b>b</b>) AISI 430Ti. Laser power 1201<math display="inline"><semantics> <mi mathvariant="normal">W</mi> </semantics></math>, welding speed 90 mm·s<sup>−1</sup>.</p> Full article ">Figure 4
<p>Schematic depicting (<b>a</b>) sample extraction and (<b>b</b>) specimen geometry employed for tensile-testing, modified with permission from ref. [<a href="#B31-metals-11-02039" class="html-bibr">31</a>]. Copyright 2020 Niklas Sommer.</p> Full article ">Figure 5
<p>Etched micrographs depicting the weld macrostructure of laser-cladded (<b>a</b>,<b>b</b>) AISI 430 and (<b>c</b>,<b>d</b>) AISI 430Ti. <b>Left</b>: overlap 1.6, <b>right</b>: overlap 2.4. Colored rectangles mark ROI for further investigation.</p> Full article ">Figure 6
<p>Etched micrographs of ROI along the clad–weld interface as marked in <a href="#metals-11-02039-f005" class="html-fig">Figure 5</a>. (<b>a</b>,<b>b</b>): AISI 430, overlap 2.4. (<b>c</b>,<b>d</b>): AISI 430Ti, overlap 2.4.</p> Full article ">Figure 7
<p>(<b>a</b>) Combined image-quality mapping (IQM) and inverse pole figure mapping (IPFM) of the clad–weld metal transition zone of AISI 430 with an overlap of 2.4. (<b>b</b>) Combined IQM- and phase mapping (PM), (<b>c</b>) IQM with singular display of orthohombic/hexagonal chromium-carbide phases. Orientations in IPFM plotted perpendicular to traverse direction. (<b>d</b>) EDS-map for molybdenum and (<b>e</b>) EDS-map for nickel.</p> Full article ">Figure 8
<p>(<b>a</b>) Combined IQM and IPFM of the clad–base metal and weld metal transition zone of AISI 430Ti, laser-cladded with an overlap of 2.4. (<b>b</b>) Combined IQM and PM. Orientations in IPFM plotted perpendicular to traverse direction. (<b>c</b>) EDS-map for molybdenum and (<b>d</b>) EDS-map for nickel.</p> Full article ">Figure 9
<p>Evolution of (<b>a</b>) yield strength and (<b>b</b>) elongation at break in dependency on the corrosion testing duration of laser-cladded AISI 430. Mean values of 12 samples each, error bars represent standard error of the mean.</p> Full article ">Figure 10
<p>Laser-cladded weld macrostructure of AISI 430 following corrosion testing for (<b>a</b>) 2<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math>, (<b>b</b>) 4<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math>, (<b>c</b>) 6<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math> (<b>d</b>) 8<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math>, (<b>e</b>) 10<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math> and (<b>f</b>) 20<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math>. Overlap 2.4. Colored rectangle represents ROI for further investigation.</p> Full article ">Figure 11
<p>Etched micrographs of a laser-cladded AISI 430 weld seam following corrosion testing. Exposure duration 10<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math>, overlap 2.4. ROI analog to markings in previous figure.</p> Full article ">Figure 12
<p>Top: Reconstructed <math display="inline"><semantics> <mi>μ</mi> </semantics></math>-CT-models of laser-cladded AISI 430 following corrosion testing for 2<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math> and 10<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math>. Bottom: Inverted illustration depicting excavations along the weld seam and clad interface. Overlap 2.4.</p> Full article ">Figure 13
<p>SEM fractography images of laser-cladded AISI 430 with corresponding EDS maps of the fracture surfaces: (<b>a</b>) sample 1, (<b>b</b>) sample 6, (<b>c</b>) stress–strain curves of the respective samples.</p> Full article ">Figure 14
<p>Evolution of (<b>a</b>) yield strength and (<b>b</b>) elongation at break in dependency on corrosion testing duration of laser-cladded AISI 430Ti. Mean values of 12 samples each, error bars represent standard error of the mean.</p> Full article ">Figure 15
<p>Etched micrographs of a laser-cladded AISI 430Ti weld seam following corrosion testing. Exposure duration 20<math display="inline"><semantics> <mi mathvariant="normal">h</mi> </semantics></math>, overlap 2.4. Colored rectangles represent ROI for further inspection.</p> Full article ">
14 pages, 6695 KiB  
Article
Effect of Post-Deposition Solution Treatment and Ageing on Improving Interfacial Adhesion Strength of Cold Sprayed Ti6Al4V Coatings
by Dibakor Boruah and Xiang Zhang
Metals 2021, 11(12), 2038; https://doi.org/10.3390/met11122038 - 15 Dec 2021
Cited by 7 | Viewed by 2937 | Correction
Abstract
This study aims at investigating the effect of post-deposition solution treatment and ageing (STA) on improving the interfacial adhesion strength in cold spray (CS) Ti6Al4V coatings deposited on Ti6Al4V substrates, measured by the adhesive-free collar-pin pull-off (CPP) test. Solution treatment was performed at [...] Read more.
This study aims at investigating the effect of post-deposition solution treatment and ageing (STA) on improving the interfacial adhesion strength in cold spray (CS) Ti6Al4V coatings deposited on Ti6Al4V substrates, measured by the adhesive-free collar-pin pull-off (CPP) test. Solution treatment was performed at 940 °C for 1 h and ageing was carried out at 480 °C for 8 h. Investigations were carried out for specimens with three different pre-treatments of the substrate surface, namely grit-blasted, as-machined (faced on lathe machine), and ground. Additionally, the effect of post-deposition STA was studied in terms of phase analysis, microstructure, and porosity level. It was observed that STA led to complete interfacial mixing resulting in significantly improved adhesion strength (by more than 520%) with the maximum measured value of greater than 766 MPa for ground substrates, reaching 81% of the ultimate tensile strength of mill annealed Ti6Al4V. Full article
(This article belongs to the Special Issue Modern Cold Spray Technique)
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Figure 1

Figure 1
<p>Characteristics of the gas-atomized Ti6Al4V powder: (<b>a</b>) size distribution, (<b>b</b>) morphology, and (<b>c</b>) microstructure [<a href="#B1-metals-11-02038" class="html-bibr">1</a>,<a href="#B20-metals-11-02038" class="html-bibr">20</a>].</p> Full article ">Figure 2
<p>3D surface profile measured by focus variation technique for different surface preparation conditions: (<b>a</b>) grit-blasted, (<b>b</b>) as-machined, and (<b>c</b>) ground (adopted from [<a href="#B1-metals-11-02038" class="html-bibr">1</a>]).</p> Full article ">Figure 3
<p>The adhesive-free collar-pin pull-off (CPP) test method: (<b>a</b>) specimen preparation process, (<b>b</b>) experimental set-up [<a href="#B1-metals-11-02038" class="html-bibr">1</a>].</p> Full article ">Figure 4
<p>Effect of STA on the interfacial adhesion strength for different substrate surface preparations: (<b>a</b>) stress vs. displacement curve, showing one example for each condition; (<b>b</b>) interfacial adhesion strength (average of three tests). (Note: AD-As-deposited, STA-Solution Treated and Aged, IF-Interface Failure, CF-Cohesion Failure).</p> Full article ">Figure 5
<p>Tested collar-pin pull-off (CPP) specimens: (<b>a</b>) failure mechanisms showing interface failure and cohesion failure, (<b>b</b>) an enlarged view of the ‘pin’ with cohesion failure, and further magnified images of its sidewall surface showing no evidence of metallurgical bonding between the ‘collar’ and ‘pin’ sidewall surface after STA treatment.</p> Full article ">Figure 6
<p>XRD patterns of Ti6Al4V in the sample cross-section covering various materials: mill annealed substrate, feedstock powder, cold spray deposits (as-deposited, solution treated, and aged).</p> Full article ">Figure 7
<p>Cross-section microstructure images of the CS Ti6Al4V deposit/substrate interface, in as-deposited (AD) and solution treatment and aged (STA) conditions, with three different surface preparation conditions: (<b>a</b>–<b>c</b>) grit-blasted, (<b>d</b>–<b>f</b>) as-machined, (<b>g</b>–<b>i</b>) ground (of which (<b>a</b>,<b>b</b>,<b>d</b>,<b>e</b>,<b>g</b>,<b>h</b>) are optical micrographs), and (<b>c</b>,<b>f</b>,<b>i</b>) are SEM images. (<b>a</b>,<b>d</b>,<b>g</b>) are adopted from [<a href="#B1-metals-11-02038" class="html-bibr">1</a>].</p> Full article ">Figure 8
<p>Adhesion strength in terms of % of mill annealed Ti6Al4V’ s ultimate tensile strength (UTS), before and after thermal treatments, and comparison with the literature. (Note: IF-Interface Failure, and CF-Cohesion Failure).</p> Full article ">
21 pages, 6530 KiB  
Article
Effects of Residual Stresses on the Fatigue Lifetimes of Self-Piercing Riveted Joints of AZ31 Mg Alloy and Al5052 Al Alloy Sheets
by Young-In Lee and Ho-Kyung Kim
Metals 2021, 11(12), 2037; https://doi.org/10.3390/met11122037 - 15 Dec 2021
Cited by 12 | Viewed by 2609
Abstract
During the self-piercing riveting (SPR) process, residual stress develops due to the high plastic deformation of the sheet materials. In this study, the effect of the residual stress on the fatigue lifetime of SPR joints with dissimilar magnesium AZ31 alloy and aluminum Al5052 [...] Read more.
During the self-piercing riveting (SPR) process, residual stress develops due to the high plastic deformation of the sheet materials. In this study, the effect of the residual stress on the fatigue lifetime of SPR joints with dissimilar magnesium AZ31 alloy and aluminum Al5052 alloy sheets was evaluated. The residual stress distribution was derived through a simulation of the SPR process by the FEA (finite element analysis). The measured values by the X-ray diffraction technique confirmed that the validity of the simulation has a maximum error of 17.2% with the experimental results. The fatigue strength of the SPR joint was evaluated at various loading angles using tensile-shear and cross-shaped specimens. It was found that the compressive residual stresses of the joint reduce the stress amplitude by 13% at 106 cycles lifetime, resulting in extension of its lifetime to approximately 3.4 million cycles from 106 cycles lifetime. Finally, it was confirmed that the fatigue life of SPR joints was appropriately predicted within a factor of three using the relationship between the fatigue life and the equivalent stress intensity factor. The fatigue resistance of the magnesium AZ31 alloy on the upper sheet was found to govern fatigue lifetimes of SPR joints of dissimilar magnesium AZ31 alloy sheets. Full article
(This article belongs to the Special Issue Research and Development of Lightweight Metal Automotive Components)
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Figure 1
<p>Illustration of the self-piercing riveting process.</p> Full article ">Figure 2
<p>Geometries and dimensions of the (<b>a</b>) tensile-shear and (<b>b</b>) cross-shaped specimens (mm).</p> Full article ">Figure 3
<p>Loading angle fixture for the cross-shaped SPR joint specimen.</p> Full article ">Figure 4
<p>Half-section view of the FEA model for SPR joining.</p> Full article ">Figure 5
<p>Stress-strain curves of AZ31 and Al5052 alloys with the Johnson–cook deformation model.</p> Full article ">Figure 6
<p>Observed the cross-section of the SPR joint by SEM and measured points of the residual stress.</p> Full article ">Figure 7
<p>Overview of FE models and a detail SPR joint for the cross-shaped specimen.</p> Full article ">Figure 8
<p>The punch force-displacement curve and stress distribution at the stages (<b>A</b>,<b>B</b>) during the SPR process.</p> Full article ">Figure 8 Cont.
<p>The punch force-displacement curve and stress distribution at the stages (<b>A</b>,<b>B</b>) during the SPR process.</p> Full article ">Figure 9
<p>Residual stress (<b>a</b>) <span class="html-italic">σ<sub>r</sub></span> and (<b>b</b>) <span class="html-italic">σ<sub>θ</sub></span> distributions for the SPR joint after spring-back.</p> Full article ">Figure 10
<p>Comparison of the static load versus the displacement curves of SPR specimens at different loading angles.</p> Full article ">Figure 11
<p>Comparison of the load amplitude against the number of cycle plots for SPR joints at different loading angles.</p> Full article ">Figure 12
<p>Experimental fatigue lifetimes of the SPR joint specimens under various loading conditions as a function of (<b>a</b>) the von-Mises stress, (<b>b</b>) the maximum principal stress, and (<b>c</b>) the SWT fatigue parameter.</p> Full article ">Figure 13
<p>Experimental fatigue lifetimes of the SPR joint specimens under various loading conditions as a function of the equivalent stress intensity factor amplitude.</p> Full article ">Figure 14
<p>Comparison of fatigue lifetimes of SPR joint profile models according to a joining analysis with and without residual stress as a function of the maximum principal stress.</p> Full article ">Figure 15
<p>Effects of compressive residual stress on the fatigue lifetime of an SPR joint.</p> Full article ">Figure 16
<p>Comparison of the fatigue lifetimes of the T.M-B.A and T.M-B.S [<a href="#B9-metals-11-02037" class="html-bibr">9</a>] SPR joints under three loading conditions.</p> Full article ">Figure 17
<p>Fatigue lifetime predictions of the T.M-B.A and T.M-B.S [<a href="#B9-metals-11-02037" class="html-bibr">9</a>] SPR joints under three loading conditions adopting the relationship of T.M-B.A.</p> Full article ">Figure 18
<p>Fatigue lifetime predictions of the T.M-B.A and T.M-B.S SPR joints under three loading conditions adopting the relationship of T.M-B.S [<a href="#B9-metals-11-02037" class="html-bibr">9</a>].</p> Full article ">
10 pages, 2440 KiB  
Article
The Role of Silica in the Chlorination–Volatilization of Cobalt Oxide by Using Calcium Chloride
by Peiwei Han, Zhengchen Li, Xiang Liu, Jingmin Yan and Shufeng Ye
Metals 2021, 11(12), 2036; https://doi.org/10.3390/met11122036 - 15 Dec 2021
Cited by 1 | Viewed by 2209
Abstract
The role of silica in the chlorination–volatilization of cobalt oxide, using calcium chloride, is investigated in this paper. It is found that the Co volatilization percentage of the CoO–Fe2O3–CaCl2 system is not larger than 12.1%. Silica plays an [...] Read more.
The role of silica in the chlorination–volatilization of cobalt oxide, using calcium chloride, is investigated in this paper. It is found that the Co volatilization percentage of the CoO–Fe2O3–CaCl2 system is not larger than 12.1%. Silica plays an important role in the chlorination–volatilization of cobalt oxide by using calcium chloride. In the CoO–SiO2–Fe2O3–CaCl2 system, the Co volatilization percentage is initially positively related to the molar ratio of SiO2 to CaCl2, and remains almost constant when the molar ratio of SiO2 to CaCl2 rises from zero to eight. The critical molar ratios of SiO2 to CaCl2 are 1 and 2 when the molar ratios of CaCl2 to CoO are 8.3 and 16.6, respectively. The Co volatilization percentage remains almost constant with the increase in CaO concentration, and decreases when Al2O3 and MgO are added. Ca2SiO3Cl2 is determined after roasting at 1073 K and 1173 K, and disappears at temperatures in excess of 1273 K in the calcines from the CoO–SiO2–CaCl2 system. CaSiO3 always exists in the calcines at temperatures in excess of 973 K. Full article
(This article belongs to the Special Issue Fundamentals of Advanced Pyrometallurgy)
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<p>Effect of CaCl<sub>2</sub> dosage on the Co volatilization percentage.</p> Full article ">Figure 2
<p>The standard Gibbs free energy of reactions: (<b>a</b>) solid or liquid CaCl<sub>2</sub>; (<b>b</b>)gaseous CaCl<sub>2</sub>.</p> Full article ">Figure 3
<p>Effect of the molar ratio of SiO<sub>2</sub> to CaCl<sub>2</sub> on the Co volatilization percentage.</p> Full article ">Figure 4
<p>Effect of (<b>a</b>) Al<sub>2</sub>O<sub>3</sub>, (<b>b</b>) MgO and (<b>c</b>) CaO on Co volatilization percentage.</p> Full article ">Figure 5
<p>XRD results of calcines from the CoO–Fe<sub>2</sub>O<sub>3</sub>–CaCl<sub>2</sub> system.</p> Full article ">Figure 6
<p>The equilibrated chlorine partial pressure of the following reaction: 2CaCl<sub>2</sub> + O<sub>2</sub> + 2CO<sub>2</sub> = 2CaCO<sub>3</sub> + 2Cl<sub>2</sub>.</p> Full article ">Figure 7
<p>XRD results of calcines from the CoO–SiO<sub>2</sub>–CaCl<sub>2</sub> system (<b>a</b>) at 973 K; (<b>b</b>) at 1073 K and 1173 K; (<b>c</b>) between 1273 K and 1473 K.</p> Full article ">Figure 8
<p>XRD results of calcines from the CoO–SiO<sub>2</sub>–Fe<sub>2</sub>O<sub>3</sub>–(Al<sub>2</sub>O<sub>3</sub>/MgO)–CaCl<sub>2</sub> system at 1273 K.</p> Full article ">
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