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Metals
Volume 10
Issue 2
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Metals, Volume 10, Issue 2 (February 2020) – 141 articles

Cover Story (view full-size image): The unidirectional cellular structure (UniPore,) with long and uniform unidirectional cells, was investigated using an explosive compaction technique, and we fabricated a composite UniPore structure composed of copper and stainless steel pipes. UniPore was investigated for applying a heat exchanger in combination with highly heat-conductive copper. Corrosion resistant stainless steel was the fabricated composite. View this paper.
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16 pages, 7256 KiB  
Article
FAST-forge of Titanium Alloy Swarf: A Solid-State Closed-Loop Recycling Approach for Aerospace Machining Waste
by Nicholas S. Weston and Martin Jackson
Metals 2020, 10(2), 296; https://doi.org/10.3390/met10020296 - 24 Feb 2020
Cited by 26 | Viewed by 8350
Abstract
Titanium alloys have excellent properties, but components are very expensive due to the high levels of processing required, such as vacuum melting, multi-stage forging, and machining. As a result, forged titanium alloy components are largely exclusive to the aerospace industry, where a high [...] Read more.
Titanium alloys have excellent properties, but components are very expensive due to the high levels of processing required, such as vacuum melting, multi-stage forging, and machining. As a result, forged titanium alloy components are largely exclusive to the aerospace industry, where a high strength-to-weight ratio, corrosion resistance, and excellent fatigue resistance are essential. However, a typical buy-to-fly ratio for such components is approximately 9:1, as much of the forged billet is machined to swarf. The quantity of waste titanium alloy swarf generated is increasing as aircraft orders, and the titanium components contained within them, are increasing. In this paper, waste swarf material has been recycled using the two-step solid-state FAST-forge process, which utilizes field assisted sintering technology (FAST) followed by hot forging. Cleaned Ti-6Al-4V swarf was fully consolidated using the FAST process at sub-transus and super-transus temperatures, followed by hot forging at sub-transus temperatures at different strain rates. It was demonstrated that swarf-derived Ti-6Al-4V FAST billets have equivalent hot forging flow behaviour and resultant microstructures when directly compared to equivalently processed conventional expensive hydride–dehydride powder, and previously reported Kroll-derived melt-wrought material. This demonstrates that titanium swarf is a good quality feedstock for downstream processing. Additionally, FAST-forge is a viable processing route for the closed-loop recycling of machining waste for next-generation components in vehicles and non-aerospace applications, which is game changing for the economics of titanium alloy components. Full article
(This article belongs to the Special Issue Towards the Development of Affordable Titanium Alloy Components)
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Figure 1
<p>Schematic representing the solid-state closed-loop recycling approach, incorporating the FAST-<span class="html-italic">forge</span> process (where FAST stands for field assisted sintering technology), for sustainable production of low-cost titanium alloy components for multiple sectors.</p> Full article ">Figure 2
<p>Photograph of the Ti-6Al-4V machining swarf feedstock (also referred to as scrap, turnings, or chips) procured for this study. The swarf has gone through a commercial grading and cleaning operation to remove contamination after the machining process.</p> Full article ">Figure 3
<p>(<b>a</b>) Photograph of the 60 mm diameter disc of Ti-6Al-4V swarf consolidated via FAST at 1200 °C and 20 MPa for 30 min; (<b>b</b>) Schematic showing the dimensions of the FAST disc, where it was sectioned, and the location and orientation the test specimens were extracted from; (<b>c</b>) Schematic showing the dimensions of the cylindrical axisymmetric hot compression test specimens.</p> Full article ">Figure 4
<p>Cross-polarized light micrographs of the Ti-6Al-4V swarf feedstock. Two distinct microstructural varieties of swarf were identified: one in a mill-annealed condition with equiaxed α grains (<b>a</b>) low magnification; (<b>b</b>) high magnification; and one in a β annealed condition with transformed β grains containing an α colony structure (<b>c</b>) low magnification; (<b>d</b>) high magnification.</p> Full article ">Figure 5
<p>Cross-polarized light micrographs showing the microstructural evolution of the Ti-6Al-4V swarf specimens with increasing FAST dwell temperature and/or dwell time, when using a fixed dwell pressure and heating rate: (<b>a</b>) 950 °C for 0 min; (<b>b</b>) 1200 °C for 0 min; (<b>c</b>) 950 °C for 5 min; (<b>d</b>) 1200 °C for 5 min; (<b>e</b>) 950 °C for 30 min; (<b>f</b>) 1200 °C for 30 min.</p> Full article ">Figure 6
<p>Cross-polarized light micrographs showing the microstructure of a 60 mm diameter disc of Ti-6Al-4V swarf after FAST processing at 1200 °C for 30 min: (<b>a</b>) low magnification; (<b>b</b>) high magnification. For comparison, cross-polarized light micrographs showing the microstructure of a 100 mm diameter disc of Ti-6Al-4V HDH powder after FAST processing at 1200 °C for 30 min: (<b>c</b>) low magnification; (<b>d</b>) high magnification.</p> Full article ">Figure 7
<p>Photographs of (<b>a</b>) side view and (<b>b</b>) top view of the Ti-6Al-4V swarf specimens consolidated via FAST at 1200 °C for 30 min after a subsequent hot compression at 950 °C (top row), 900 °C (middle row), and 850 °C (bottom row), and a strain rate of 0.01 s<sup>−1</sup> (left column), 0.1 s<sup>−1</sup> (middle column), and 1 s<sup>−1</sup> (right column). The deformed specimens have been abrasively sectioned in half to allow metallographic preparation.</p> Full article ">Figure 8
<p>Cross-polarized light micrographs of higher-strain central regions in Ti-6Al-4V swarf specimens consolidated via FAST at 1200 °C for 30 min after a subsequent hot compression at 850, 900, or 950 °C, and a strain rate of 0.01, or 0.1, or 1 s<sup>−1</sup> (as labelled).</p> Full article ">Figure 9
<p>Cross-polarized light micrographs of lower strain dead zone regions in Ti-6Al-4V swarf specimens consolidated via FAST at 1200 °C for 30 min after a subsequent hot compression at 850, 900, or 950 °C, and a strain rate of 0.01, or 0.1, or 1 s<sup>−1</sup> (as labelled).</p> Full article ">Figure 10
<p>Flow curves showing true stress vs. true strain behaviour for Ti-6Al-4V swarf consolidated via FAST under axisymmetric compression at 850, 900, and 950 °C: (<b>a</b>) deformed with a strain rate of 1 s<sup>−1</sup>; (<b>b</b>) deformed with a strain rate of 0.1 s<sup>−1</sup>; (<b>c</b>) deformed with a strain rate of 0.01 s<sup>−1</sup>. Flow curves for Ti-6Al-4V HDH powder consolidated via FAST are also shown for comparison.</p> Full article ">Figure 11
<p>Cross-polarized light micrographs of the microstructural evolution with increasing strain from the top/edge to the centre of a specimen after hot compression at 950 °C and 0.01 s<sup>−1</sup>; test specimens were produced from Ti-6Al-4V swarf consolidated via FAST (<b>top</b>) and Ti-6Al-4V HDH powder consolidated via FAST (<b>bottom</b>).</p> Full article ">
6 pages, 211 KiB  
Editorial
Refining and Casting of Steel
by Karel Gryc and Jan Falkus
Metals 2020, 10(2), 295; https://doi.org/10.3390/met10020295 - 24 Feb 2020
Cited by 1 | Viewed by 2500
Abstract
Steel was the most requested material all over the world during the past fast technically evolving centuries [...] Full article
(This article belongs to the Special Issue Refining and Casting of Steel)
11 pages, 4401 KiB  
Article
The Interdependence of the Degree of Precipitation and Dislocation Density during the Thermomechanical Treatment of Microalloyed Niobium Steel
by Stoja Rešković, Ljerka Slokar Benić and Martina Lovrenić-Jugović
Metals 2020, 10(2), 294; https://doi.org/10.3390/met10020294 - 24 Feb 2020
Cited by 5 | Viewed by 2931
Abstract
In this paper, thermomechanical processing of niobium microalloyed steel was performed with the purpose of determining the interaction between niobium precipitates and dislocations, as well as determining the influence of the temperature of final deformation on the degree of precipitation and dislocation density. [...] Read more.
In this paper, thermomechanical processing of niobium microalloyed steel was performed with the purpose of determining the interaction between niobium precipitates and dislocations, as well as determining the influence of the temperature of final deformation on the degree of precipitation and dislocation density. Two variants of thermomechanical processing with different final rolling temperatures were carried out. Samples were studied using electrochemical isolation with an atomic absorption spectrometer, transmission electron microscopy, X-ray diffraction analysis, and universal tensile testing with a thermographic camera. The results show that the increase in the density of dislocations before the onset of intense precipitation is insignificant because the recrystallization process takes place simultaneously. It increases with the onset of strain-induced precipitation. In this paper, it is shown that niobium precipitates determine the density of dislocations. The appearance of Lüders bands was noticed as a consequence of the interaction between niobium precipitates and dislocations during the subsequent cold deformation. In both variants of the industrial process performed on the cold deformed strip, Lüders bands appeared. Full article
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<p>(<b>a</b>) Schematic of the final deformation (FD) and (<b>b</b>) the change in the degree of deformation.</p> Full article ">Figure 2
<p>Degree of precipitation of niobium in each stage of thermomechanical processing. (PD: pre-deformation, FD-S: final deformation—start, FD-E: final deformation—end, C: cooling).</p> Full article ">Figure 3
<p>TEM micrograph of niobium precipitates in pre-deformation, reproduced from [<a href="#B10-metals-10-00294" class="html-bibr">10</a>], with permission from Taylor&amp;Francis, 2013.</p> Full article ">Figure 4
<p>Change in precipitation degree (<b>a</b>) through passes and (<b>b</b>) with the temperature of thermomechanical processing.</p> Full article ">Figure 5
<p>TEM micrograph showing the onset of strain-induced precipitation at 963 °C, low temperature (LT) variant.</p> Full article ">Figure 6
<p>Change in the dislocation density by stages of final thermomechanical processing.</p> Full article ">Figure 7
<p>Relation between the degree of precipitation, the dislocation density, and temperature for (<b>a</b>) the LT variant and (<b>b</b>) the HT variant.</p> Full article ">Figure 8
<p>TEM micrograph of the interaction between precipitates and dislocations [<a href="#B20-metals-10-00294" class="html-bibr">20</a>].</p> Full article ">Figure 9
<p>TEM micrograph of dislocations accumulating on the rows of precipitations.</p> Full article ">Figure 10
<p>Microstructure of the hot rolled strips of Nb microalloyed steel. (<b>a</b>) LT variant. (<b>b</b>) HT variant.</p> Full article ">Figure 11
<p>The stress–strain curve and distribution of temperature change at the beginning of the plastic material flow in the field of the appearance of Lüders bands.</p> Full article ">
13 pages, 3509 KiB  
Article
Selective Copper Recovery by Acid Leaching from Printed Circuit Board Waste Sludge
by Ha Bich Trinh, Seunghyun Kim and Jaeryeong Lee
Metals 2020, 10(2), 293; https://doi.org/10.3390/met10020293 - 23 Feb 2020
Cited by 12 | Viewed by 6291
Abstract
The most challenging issue associated with recycling the sludge generated from printed circuit boards (PCBs) is the separation of copper (Cu) from iron (Fe), using multi-stage leaching, or adding oxidizing and precipitating agents. Herein we investigated simple acid leaching to effectively extract copper [...] Read more.
The most challenging issue associated with recycling the sludge generated from printed circuit boards (PCBs) is the separation of copper (Cu) from iron (Fe), using multi-stage leaching, or adding oxidizing and precipitating agents. Herein we investigated simple acid leaching to effectively extract copper and limit iron dissolution. Selective copper leaching was achieved with all the acids studied, including HCl, HNO3, and H2SO4. The lower concentration of acid solutions resulted in a larger difference in leachabilities between Cu and Fe. Among three leachates, the H2SO4 solution performed effectively on the selective leaching of Cu and Fe. Adjusting the pulp density to 4% and the H2SO4 concentration at ~0.2 M, accomplished ~95% Cu leaching and reduced the Fe extraction to less than 5%. Kinetic studies revealed that Cu leaching followed the ash diffusion-controlled mechanism. Aactivation energy (Ea) of 9.8 kJ/mol was determined for the first 10 min of leaching. Further, leaching up to 60 min corresponded to a mixed control model, increasing the Ea to 20.9 kJ/mol. The change in the control model with regard to the two leaching stages can be attributed to the Cu hydroxide and metal phases present in the original sample. A simple, economically attractive H2SO4 acid leaching process was demonstrated, recovering Cu efficiently and selectively from PCBs waste sludge under moderate conditions. Full article
(This article belongs to the Special Issue Separation and Leaching for Metals Recovery)
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<p>X-ray diffraction of the sludge sample before and after heating at 500 °C.</p> Full article ">Figure 2
<p>DTA-TGA curves of copper sludge sample.</p> Full article ">Figure 3
<p>Morphology and SEM-EDS of the copper sludge sample.</p> Full article ">Figure 4
<p>Cu and Fe leachability with variation of (<b>a</b>) HCl concentration, (<b>b</b>) HNO<sub>3</sub> concentration, (<b>c</b>) H<sub>2</sub>SO<sub>4</sub> concentration.</p> Full article ">Figure 5
<p>Effect of H<sub>2</sub>SO<sub>4</sub> concentration on Cu and Fe leachability.</p> Full article ">Figure 6
<p>Effect of temperature and time on Cu leachability.</p> Full article ">Figure 7
<p>Plot of <math display="inline"><semantics> <mrow> <mrow> <mo>[</mo> <mn>1</mn> <mo>−</mo> <mn>3</mn> </mrow> <msup> <mrow> <mrow> <mo>(</mo> <mrow> <mrow> <mn>1</mn> <mo>−</mo> <mi>x</mi> </mrow> </mrow> <mo>)</mo> </mrow> </mrow> <mrow> <mfrac bevelled="true"> <mn>2</mn> <mn>3</mn> </mfrac> </mrow> </msup> <mrow> <mo>+</mo> <mn>2</mn> <mo>(</mo> <mn>1</mn> <mo>−</mo> <mi>x</mi> <mo>)</mo> <mo>]</mo> </mrow> </mrow> </semantics></math> versus leaching time at different temperature.</p> Full article ">Figure 8
<p>Plot of <math display="inline"><semantics> <mrow> <mo stretchy="false">[</mo> <msup> <mrow> <mrow> <mn>1</mn> <mo>−</mo> <mn>3</mn> <mo>(</mo> <mn>1</mn> <mo>−</mo> <mi>x</mi> <mo>)</mo> </mrow> </mrow> <mrow> <mfrac bevelled="true"> <mn>2</mn> <mn>3</mn> </mfrac> </mrow> </msup> <mrow> <mo>+</mo> <mtext> </mtext> <mn>2</mn> <mo>(</mo> <mn>1</mn> <mo>−</mo> <mi>x</mi> <mo>)</mo> <mo>]</mo> </mrow> </mrow> </semantics></math> versus leaching time (2–10 min) at different temperatures.</p> Full article ">Figure 9
<p>Plot of <math display="inline"><semantics> <mrow> <mo stretchy="false">[</mo> <msup> <mrow> <mrow> <mn>1</mn> <mo>−</mo> <mn>3</mn> <mo>(</mo> <mn>1</mn> <mo>−</mo> <mi>x</mi> <mo>)</mo> </mrow> </mrow> <mrow> <mfrac bevelled="true"> <mn>2</mn> <mn>3</mn> </mfrac> </mrow> </msup> <mrow> <mo>+</mo> <mtext> </mtext> <mn>2</mn> <mo>(</mo> <mn>1</mn> <mo>−</mo> <mi>x</mi> <mo>)</mo> <mo>]</mo> </mrow> </mrow> </semantics></math> versus leaching time (10–60 min) at different temperatures.</p> Full article ">Figure 10
<p>Plot of lnK versus 1/T to obtain the apparent activation energy.</p> Full article ">
12 pages, 21640 KiB  
Article
Effect of SLM Processing Parameters on Microstructures and Mechanical Properties of Al0.5CoCrFeNi High Entropy Alloys
by Kun Sun, Weixiang Peng, Longlong Yang and Liang Fang
Metals 2020, 10(2), 292; https://doi.org/10.3390/met10020292 - 23 Feb 2020
Cited by 59 | Viewed by 5785
Abstract
Selective laser melting (SLM) to fabricate Al0.5CoCrFeNi high entropy alloys with pre-mixed powders was studied in this paper. The influences of process parameters including laser power, scanning speed, and hatch spacing on the relative density of high-entropy alloys (HEAs) were investigated. [...] Read more.
Selective laser melting (SLM) to fabricate Al0.5CoCrFeNi high entropy alloys with pre-mixed powders was studied in this paper. The influences of process parameters including laser power, scanning speed, and hatch spacing on the relative density of high-entropy alloys (HEAs) were investigated. A relative density of 99.92% can be achieved by optimizing the SLM process parameters with laser power 320 W, scanning speed 800 mm/s, and hatch spacing of 60 μm, respectively. Moreover, the microstructure of the HEAs was also studied using scanning electron microscopy (SEM) and x-ray diffraction (XRD). It was found that the microstructure of the HEAs was only composed of face-centered cubic and body-centered cubic phases, without complex intermetallic compounds. The mechanical properties of the HEAs were also characterized. At ambient temperature, the alloys had a high yield strength of about 609 MPa, tensile strength about 878 MPa, and hardness about 270 HV. Through a comparison with the corresponding alloys fabricated by vacuum induction melting, it can be concluded that the high entropy alloys fabricated by SLM had fine microstructures and improved mechanical properties. Full article
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Figure 1
<p>SEM micrograph of the powders: (<b>a</b>) Al, (<b>b</b>) Co, (<b>c</b>) Cr, (<b>d</b>) Fe, (<b>e</b>) Ni, and (<b>f</b>) mixed.</p> Full article ">Figure 2
<p>The schematic illustration of the XJRP SLM molding machine.</p> Full article ">Figure 3
<p>Schematic illustration of the SLM processing parameters.</p> Full article ">Figure 4
<p>Schematic illustration of the laser scanning pattern: “orthogonal scanning strategy”.</p> Full article ">Figure 5
<p>Photograph of the SLM-processed sample and tensile specimen.</p> Full article ">Figure 6
<p>The SEM images of Al<sub>0.5</sub>CoCrFeNi HEAs at different laser power: (<b>a</b>) 160 W, (<b>b</b>) 200 W, (<b>c</b>) 240 W, (<b>d</b>) 280 W, (<b>e</b>) 320 W at a scan speed of 800 mm/s and hatch spacing of 60 μm, and (<b>f</b>) image of the selected area illustrated by a square in (<b>e</b>) and the elemental distribution maps of the area.</p> Full article ">Figure 7
<p>Schematic illustration of the balling effect during SLM processing.</p> Full article ">Figure 8
<p>The influence of the process parameters on the relative density of the SLM specimens: (<b>a</b>) laser power, (<b>b</b>) scanning speed, and (<b>c</b>) hatch spacing.</p> Full article ">Figure 9
<p>The relationship between energy density and relative density of the SLM specimens.</p> Full article ">Figure 10
<p>The micrographs of the specimens fabricated by SLM with the optimized parameters: the vertical cross (VC) sections (<b>a</b>), (<b>c</b>), and (<b>e</b>) and the horizontal cross (HC) sections (<b>b</b>), (<b>d</b>), and (<b>f</b>).</p> Full article ">Figure 11
<p>XRD patterns of the cross sections of the Al<sub>0.5</sub>CoCrFeNi HEAs.</p> Full article ">Figure 12
<p>The uniaxial tensile stress–strain curve of the SLM Al<sub>0.5</sub>CoCrFeNi HEAs.</p> Full article ">Figure 13
<p>The SEM images of the fracture surface for the SLM specimen at room temperature: (<b>a</b>) low magnifications, (<b>b</b>) high magnifications.</p> Full article ">
15 pages, 4127 KiB  
Article
The Elusive Thomson Effect in Thermoelectric Devices. Experimental Investigation from 363 K to 213 K on Various Peltier Modules
by Valter Giaretto and Elena Campagnoli
Metals 2020, 10(2), 291; https://doi.org/10.3390/met10020291 - 23 Feb 2020
Cited by 5 | Viewed by 3295
Abstract
At steady state, in the governing equation of one-stage thermoelectric cooler, the heat resulting from Fourier conduction is balanced by heat generation due to the Joule and Thomson effects inside semiconductors. Since the heat flux observed at the junction of a semiconductor, r [...] Read more.
At steady state, in the governing equation of one-stage thermoelectric cooler, the heat resulting from Fourier conduction is balanced by heat generation due to the Joule and Thomson effects inside semiconductors. Since the heat flux observed at the junction of a semiconductor, r pair includes the Peltier effect and the Fourier heat flux caused by both the aforementioned contributions, the Thomson effect is easily masked by the Joule heat, which makes it elusive. With the aim of highlighting the contribution of the Thomson effect, measurements were carried out in the temperature range from 363 K to 213 K on different Peltier modules. The temperature dependence of the Seebeck and Thomson coefficients was evaluated as well as the electrical resistivity, and thermal conductivity of the Peltier modules examined. The results obtained show that the temperature dependence of the thermoelectric properties can reduce the cooling capacity of the Peltier module compared to what is declared in the technical datasheets of the commercial devices. The analyses allow us to conclude that an increase in the Thomson effect could have a positive effect on the performance of the Peltier only if it were possible to reduce the Joule contribution simultaneously. Full article
(This article belongs to the Special Issue Thermoelectric Compounds: Processing, Properties and Applications)
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Figure 1
<p>Scheme of the experimental arrangement. DX 4090 is the z-meter, TTC is the Thermostatic Test Chamber, SAC is the Still Air Chamber, and DUT is the Device Under Test. Solid and double red and black lines are the feeding cables, while the dashed ones represent the signal wires.</p> Full article ">Figure 2
<p>(<b>a</b>) Assembled <span class="html-italic">p</span>-<span class="html-italic">n</span> series resistance <span class="html-italic">R</span> given by Equation (1) versus the test reference temperature for each Peltier module; (<b>b</b>) Equivalent electrical resistivity of the assembled thermoelectric materials versus test reference temperature for each Peltier module. The dashed line represents a possible trend in temperature of the average electrical resistivity for the <span class="html-italic">p</span> and <span class="html-italic">n</span> doped Bi<sub>2</sub>Te<sub>3</sub> bulk materials, as found in the literature [<a href="#B26-metals-10-00291" class="html-bibr">26</a>].</p> Full article ">Figure 3
<p>Measured dimensionless figure of merit for the investigated Peltier modules versus the test reference temperature.</p> Full article ">Figure 4
<p>Examples of the transient temperature difference measured on the external surface of the ceramic plates versus time for the two cases with maximum and minimum reference temperature. Solid lines represent the trend calculated with a lumped model that uses the characteristic time <math display="inline"><semantics> <mrow> <msub> <mi>t</mi> <mn>0</mn> </msub> </mrow> </semantics></math> provided by the DX 4090; dashed lines represent the asymptotic values <math display="inline"><semantics> <mrow> <mo>Δ</mo> <msub> <mi>T</mi> <mrow> <mi>P</mi> <mo>,</mo> <mo>∞</mo> </mrow> </msub> </mrow> </semantics></math>. (<b>a</b>) module a.2; (<b>b</b>) module c.</p> Full article ">Figure 5
<p>Seebeck voltage of a single pair (total Seebeck voltage divided by the number of pairs <span class="html-italic">N</span><sub>P</sub>) versus the estimated temperature difference <math display="inline"><semantics> <mrow> <mo>Δ</mo> <msub> <mi>T</mi> <mi>S</mi> </msub> </mrow> </semantics></math> of the Seebeck junction.</p> Full article ">Figure 6
<p>(<b>a</b>) Equivalent single junction Seebeck coefficients versus test reference temperature for modules a.2, b and c. The fitting equation and regression coefficients are shown in the diagram box. The solid lines represent the fit trends, and the dashed line is the theoretical trend mentioned by Goldsmid [<a href="#B26-metals-10-00291" class="html-bibr">26</a>]; (<b>b</b>) Thomson coefficient trends for module c (<span class="html-italic">i</span>) and modules a.2 and b (<span class="html-italic">ii</span>). Parameters <span class="html-italic">m</span> and <span class="html-italic">n</span> are those shown in (<b>a</b>), and <math display="inline"><semantics> <mrow> <msub> <mi>τ</mi> <mn>0</mn> </msub> </mrow> </semantics></math> values are the Thomson coefficients at <math display="inline"><semantics> <mrow> <msub> <mi>T</mi> <mn>0</mn> </msub> </mrow> </semantics></math>.</p> Full article ">Figure 7
<p>(<b>a</b>) Thermal conductance versus the reference temperature of the test for modules a.2, b, and c.; (<b>b</b>) Equivalent thermal conductivity of the assembled thermoelectric materials versus the reference temperature of the test for modules a.2, b, and c. The dashed line represents a possible trend in temperature of the average thermal conductivity for <span class="html-italic">p</span> and <span class="html-italic">n</span> doped Bi<sub>2</sub>Te<sub>3</sub> bulk materials, as found in the literature [<a href="#B27-metals-10-00291" class="html-bibr">27</a>].</p> Full article ">Figure 8
<p>Cold side: ratio between Thomson heat and Joule effect versus the relative cooling capacity, solid line case (i), dashed line case (ii). (<b>a</b>) <math display="inline"><semantics> <mrow> <msub> <mi>T</mi> <mi>H</mi> </msub> <mo>=</mo> <mn>298.15</mn> <mo> </mo> <mi>K</mi> </mrow> </semantics></math>; (<b>b</b>) <math display="inline"><semantics> <mrow> <msub> <mi>T</mi> <mi>H</mi> </msub> <mo>=</mo> <mn>323.15</mn> <mo> </mo> <mi>K</mi> </mrow> </semantics></math>.</p> Full article ">Figure 9
<p><math display="inline"><semantics> <mrow> <msub> <mi>T</mi> <mi>H</mi> </msub> <mo>=</mo> <mn>298.15</mn> <mo> </mo> <mi>K</mi> </mrow> </semantics></math>. Relative cooling capacity <math display="inline"><semantics> <mrow> <msub> <mi>Q</mi> <mi>C</mi> </msub> <mo>/</mo> <msub> <mi>Q</mi> <mrow> <mi>M</mi> <mi>a</mi> <mi>x</mi> </mrow> </msub> </mrow> </semantics></math> against the relative temperature difference <math display="inline"><semantics> <mrow> <mi mathvariant="sans-serif">Δ</mi> <mi>T</mi> <mo>/</mo> <mi mathvariant="sans-serif">Δ</mi> <msub> <mi>T</mi> <mrow> <mi>M</mi> <mi>a</mi> <mi>x</mi> </mrow> </msub> </mrow> </semantics></math> for different relative feeding currents <math display="inline"><semantics> <mrow> <mi>I</mi> <mo>/</mo> <msub> <mi>I</mi> <mrow> <mi>M</mi> <mi>A</mi> <mi>X</mi> </mrow> </msub> </mrow> </semantics></math>. The thick solid lines represent the performed experiment; the dashed lines refer to thermoelectric properties independent of temperature, the thin solid lines indicate the scenario in which the Seebeck coefficient depends on the temperature, while both the electrical resistivity and the thermal conductivity are temperature independent: (<b>a</b>) case (<span class="html-italic">i</span>); (<b>b</b>) case (<span class="html-italic">ii</span>).</p> Full article ">Figure 10
<p><math display="inline"><semantics> <mrow> <msub> <mi>T</mi> <mi>H</mi> </msub> <mo>=</mo> <mn>323.15</mn> <mo> </mo> <mi>K</mi> </mrow> </semantics></math>. Relative cooling capacity <math display="inline"><semantics> <mrow> <msub> <mi>Q</mi> <mi>C</mi> </msub> <mo>/</mo> <msub> <mi>Q</mi> <mrow> <mi>M</mi> <mi>a</mi> <mi>x</mi> </mrow> </msub> </mrow> </semantics></math> against the relative temperature difference <math display="inline"><semantics> <mrow> <mi mathvariant="sans-serif">Δ</mi> <mi>T</mi> <mo>/</mo> <mi mathvariant="sans-serif">Δ</mi> <msub> <mi>T</mi> <mrow> <mi>M</mi> <mi>a</mi> <mi>x</mi> </mrow> </msub> </mrow> </semantics></math> for different relative feeding currents <math display="inline"><semantics> <mrow> <mi>I</mi> <mo>/</mo> <msub> <mi>I</mi> <mrow> <mi>M</mi> <mi>A</mi> <mi>X</mi> </mrow> </msub> </mrow> </semantics></math>. The thick solid lines represent the performed experiment, the dashed lines refer to thermoelectric properties independent of temperature, the thin solid lines indicate the scenario in which the Seebeck coefficient depends on the temperature, while both the electrical resistivity and the thermal conductivity are temperature independent: (<b>a</b>) case (<span class="html-italic">i</span>); (<b>b</b>) case (<span class="html-italic">ii</span>).</p> Full article ">Figure 11
<p><math display="inline"><semantics> <mrow> <msub> <mi>T</mi> <mi>H</mi> </msub> <mo>=</mo> <mn>298.15</mn> <mo> </mo> <mi>K</mi> </mrow> </semantics></math>. Coefficient of performance against the relative temperature difference <math display="inline"><semantics> <mrow> <mi mathvariant="sans-serif">Δ</mi> <mi>T</mi> <mo>/</mo> <mi mathvariant="sans-serif">Δ</mi> <msub> <mi>T</mi> <mrow> <mi>M</mi> <mi>a</mi> <mi>x</mi> </mrow> </msub> </mrow> </semantics></math> at different relative feeding current <math display="inline"><semantics> <mrow> <mi>I</mi> <mo>/</mo> <msub> <mi>I</mi> <mrow> <mi>M</mi> <mi>A</mi> <mi>X</mi> </mrow> </msub> </mrow> </semantics></math>, thick solid lines represent the performed experiment, dashed lines refer to temperature-independent thermoelectric properties, thin solid lines denote the scenario in which the Seebeck coefficient is temperature-dependent and both electrical resistivity and thermal conductivity are temperature independent: (<b>a</b>) case (<span class="html-italic">i</span>); (<b>b</b>) case (<span class="html-italic">ii</span>).</p> Full article ">Figure 12
<p><math display="inline"><semantics> <mrow> <msub> <mi>T</mi> <mi>H</mi> </msub> <mo>=</mo> <mn>323.15</mn> <mo> </mo> <mi>K</mi> </mrow> </semantics></math>. Coefficient of performance against the relative temperature difference <math display="inline"><semantics> <mrow> <mi mathvariant="sans-serif">Δ</mi> <mi>T</mi> <mo>/</mo> <mi mathvariant="sans-serif">Δ</mi> <msub> <mi>T</mi> <mrow> <mi>M</mi> <mi>a</mi> <mi>x</mi> </mrow> </msub> </mrow> </semantics></math> at different relative feeding current <math display="inline"><semantics> <mrow> <mi>I</mi> <mo>/</mo> <msub> <mi>I</mi> <mrow> <mi>M</mi> <mi>A</mi> <mi>X</mi> </mrow> </msub> </mrow> </semantics></math>, thick solid lines represent the performed experiment, dashed lines refer to temperature-independent thermoelectric properties, thin solid lines denote the scenario in which the Seebeck coefficient is temperature-dependent and both electrical resistivity and thermal conductivity are temperature independent: (<b>a</b>) case (<span class="html-italic">i</span>); (<b>b</b>) case (<span class="html-italic">ii</span>).</p> Full article ">
19 pages, 20553 KiB  
Article
Surface Quality and Free Energy Evaluation of s275 Steel by Shot Blasting, Abrasive Water Jet Texturing and Laser Surface Texturing
by Fermin Bañon, Alejandro Sambruno, Moises Batista, Bartolome Simonet and Jorge Salguero
Metals 2020, 10(2), 290; https://doi.org/10.3390/met10020290 - 22 Feb 2020
Cited by 18 | Viewed by 3409
Abstract
Surface modification by different technologies prior to joining operations or improving tribological properties is a point of great interest. Improving surface activation by increasing the roughness of the metal is a relationship that is becoming more defined. In turn, an increase in surface [...] Read more.
Surface modification by different technologies prior to joining operations or improving tribological properties is a point of great interest. Improving surface activation by increasing the roughness of the metal is a relationship that is becoming more defined. In turn, an increase in surface wettability by evaluating contact angles indicates surface activation by obtaining a high surface free energy. Technologies such as shot blasting and laser surface texturing (LST) have generated several scientific studies where they have identified the influence of parameters on the formation of rough surfaces with defined patterns. However, the application of abrasive water jet texturing (AWJT) has been little studied as an alternative. This article compares these technologies in the texturing of a carbon steel s275 in order to identify the relationship between surface quality and surface activation. It has been determined that AWJT produces the highest Rt values close to 64 µm with a cross feed of 0.45 mm and a traverse speed of 5000 mm/min. Furthermore, LST obtains the best values of free surface energy by combining a power of 20 W with a frequency of 20 kHz and a sweeping speed of 10 mm/s. Finally, contour diagrams have been obtained which relate these variables to the texturing parameters. Full article
(This article belongs to the Special Issue Surface Modification Technology in Metals)
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Figure 1
<p>Measurement process and evaluation of contact angle on textured surfaces: (<b>a</b>) equipment used for contact angle measurement; (<b>b</b>) Drop deposition procedure on textured surfaces; (<b>c</b>) Contact angle measurement by image processing software.</p> Full article ">Figure 2
<p>Metallography at 10x of the surfaces obtained for: (<b>a</b>) Shot blasting (Corundum particles); (<b>b</b>) Non-textured; (<b>c</b>) AWJT; (<b>d</b>) Laser Power 5 W; (<b>e</b>) Laser Power 20 W.</p> Full article ">Figure 3
<p>Remains of abrasive particles adhered to the surface after texturing.</p> Full article ">Figure 4
<p>SEM images of abrasive waterjet textured surfaces at 60× for: (<b>a</b>) 0.1 mm and 5000 mm/min; (<b>b</b>) 0.1 mm and 7000 mm/min; (<b>c</b>) 0.1 mm and 9000 mm/min; (<b>d</b>) 0.3 mm and 5000 mm/min; (<b>e</b>) 0.3 mm and 7000 mm/min; (<b>f</b>) 0.3 mm and 9000 mm/min; (<b>g</b>) 0.45 mm and 5000 mm/min; (<b>h</b>) 0.45 mm and 7000 mm/min; (<b>i</b>) 0.45 mm and 9000 mm/min.</p> Full article ">Figure 5
<p>SEM imaging of laser-textured surfaces at 500× for: (<b>a</b>) 5 W and 250 mm/s; (<b>b</b>) 20 W and 10 mm/s.</p> Full article ">Figure 6
<p>Rt values obtained in blasting tests.</p> Full article ">Figure 7
<p>Rt values obtained in abrasive water jet textured tests.</p> Full article ">Figure 8
<p>Rt values obtained in laser textured tests.</p> Full article ">Figure 9
<p>Roughness profiles obtained for: (<b>a</b>) 5 W and 10 mm/s; (<b>b</b>) 5 W and 250 mm/s.</p> Full article ">Figure 10
<p>Surface free energy values for water jet texturing tests.</p> Full article ">Figure 11
<p>Increase of the contact angle for a separation of 0.3 mm and: (<b>a</b>) 5000 mm/min; (<b>b</b>) 7000 mm/min; (<b>c</b>) 9000 mm/min.</p> Full article ">Figure 12
<p>Contact angle reduction for a traverse speed of 7000 mm/min and: (<b>a</b>) 0.1 mm; (<b>b</b>) 0.3 mm; (<b>c</b>) 0.45 mm.</p> Full article ">Figure 13
<p>Surface free energy values for laser texturing tests.</p> Full article ">Figure 14
<p>Surface differences obtained by laser texturing at 20×: (<b>a</b>) 5 W and 250 mm/s; (<b>b</b>) 20 W and 10 mm/s.</p> Full article ">Figure 15
<p>Contact angles obtained in laser texturing for: (<b>a</b>) 5 W and 250 mm/s; (<b>b</b>) 20 W and 10 mm/s.</p> Full article ">Figure 16
<p>Contour diagrams of the surface free energy values obtained for: (<b>a</b>) Abrasive water jet texturing; (<b>b</b>) Laser texturing.</p> Full article ">Figure 17
<p>Contour diagrams of the surface quality obtained for: (<b>a</b>) Abrasive water jet texturing; (<b>b</b>) Laser surface texturing.</p> Full article ">
15 pages, 13191 KiB  
Article
Influence of Elastomer Layers in the Quality of Aluminum Parts on Finishing Operations
by Antonio Rubio-Mateos, Asuncion Rivero, Eneko Ukar and Aitzol Lamikiz
Metals 2020, 10(2), 289; https://doi.org/10.3390/met10020289 - 22 Feb 2020
Cited by 7 | Viewed by 2872
Abstract
In finishing processes, the quality of aluminum parts is mostly influenced by static and dynamic phenomena. Different solutions have been studied toward a stable milling process attainment. However, the improvements obtained with the tuning of process parameters are limited by the system stiffness [...] Read more.
In finishing processes, the quality of aluminum parts is mostly influenced by static and dynamic phenomena. Different solutions have been studied toward a stable milling process attainment. However, the improvements obtained with the tuning of process parameters are limited by the system stiffness and external dampers devices interfere with the machining process. To deal with this challenge, this work analyzes the suitability of elastomer layers as passive damping elements directly located under the part to be machined. Thus, exploiting the sealing properties of nitrile butadiene rubber (NBR), a suitable flexible vacuum fixture is developed, enabling a proper implementation in the manufacturing process. Two different compounds are characterized under axial compression and under finishing operations. The compression tests present the effect of the feed rate and the strain accumulative effect in the fixture compressive behavior. Despite the higher strain variability of the softer rubber, different milling process parameters, such as the tool feed rate, can lead to a similar compressive behavior of the fixture regardless the elastomer hardness. On the other hand, the characterization of these flexible fixtures is completed over AA2024 floor milling of rigid parts and compared with the use of a rigid part clamping. These results show that, as the cutting speed and the feed rate increases, due to the strain evolution of the rubber, the part quality obtained tend to equalize between the flexible and the rigid clamping of the workpiece. Due to the versatility of the NBR for clamping different part geometries without new fixture redesigns, this leads to a competitive advantage of these flexible solutions against the classic rigid vacuum fixtures. Finally, a model to predict the grooving forces with a bull-nose end mill regardless of the stiffness of the part support is proposed and validated for the working range. Full article
(This article belongs to the Special Issue Metal Machining—Recent Advances, Applications and Challenges)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>Adapted rubber: (<b>a</b>) vacuum channels distribution. (<b>b</b>) Rubber layer implementation as a vacuum fixture on milling tests.</p> Full article ">Figure 2
<p>Compression test procedure: (<b>a</b>) set-up scheme. (<b>b</b>) Load application zone.</p> Full article ">Figure 3
<p>Milling test set-up with and without a rubber layer.</p> Full article ">Figure 4
<p>Accumulative stress and feed rate increase effects on both rubber materials: (<b>a</b>) rubber A and (<b>b</b>) rubber B.</p> Full article ">Figure 5
<p>Thickness error evolution based on the <math display="inline"><semantics> <mrow> <msub> <mi>v</mi> <mi>c</mi> </msub> </mrow> </semantics></math> variation.</p> Full article ">Figure 6
<p>Roughness evolution: (<b>a</b>) based on the cutting speed and (<b>b</b>) based on the feed per tooth.</p> Full article ">Figure 7
<p>Stability Lobe Diagrams (SLD) variation: (<b>a</b>) complete and (<b>b</b>) zoomed on the studied zone.</p> Full article ">Figure 8
<p>Fast Fourier Transform of the <math display="inline"><semantics> <mrow> <msub> <mi>F</mi> <mi>z</mi> </msub> </mrow> </semantics></math> signal for the test at <math display="inline"><semantics> <mrow> <msub> <mi>f</mi> <mi>z</mi> </msub> </mrow> </semantics></math> = 0.1 mm/tooth and <math display="inline"><semantics> <mrow> <msub> <mi>a</mi> <mi>p</mi> </msub> </mrow> </semantics></math> = 0.8 mm, under different spindle speeds: (<b>a</b>) No rubber - 2000 rpm, (<b>b</b>) No rubber - 4000 rpm, (<b>c</b>) No rubber - 6000 rpm, (<b>d</b>) Rubber A - 2000 rpm, (<b>e</b>) Rubber A - 4000 rpm, (<b>f</b>) Rubber A - 6000 rpm, (<b>h</b>) Rubber B - 2000 rpm, (<b>i</b>) Rubber B - 4000 rpm and (<b>j</b>) Rubber B - 6000 rpm.</p> Full article ">Figure 9
<p>Axial mean loads predicted (lines) and experimental data against the material removal rate (MRR) as a function of the cutting speed (<math display="inline"><semantics> <mrow> <msub> <mi>v</mi> <mi>c</mi> </msub> </mrow> </semantics></math>).</p> Full article ">
23 pages, 4008 KiB  
Review
The Challenge of Digitalization in the Steel Sector
by Teresa Annunziata Branca, Barbara Fornai, Valentina Colla, Maria Maddalena Murri, Eliana Streppa and Antonius Johannes Schröder
Metals 2020, 10(2), 288; https://doi.org/10.3390/met10020288 - 21 Feb 2020
Cited by 108 | Viewed by 25588
Abstract
Digitalization represents a paramount process started some decades ago, but which received a strong acceleration by Industry 4.0 and now directly impacts all the process and manufacturing sectors. It is expected to allow the European industry to increase its production efficiency and its [...] Read more.
Digitalization represents a paramount process started some decades ago, but which received a strong acceleration by Industry 4.0 and now directly impacts all the process and manufacturing sectors. It is expected to allow the European industry to increase its production efficiency and its sustainability. In particular, in the energy-intensive industries, such as the steel industry, digitalization concerns the application of the related technologies to the production processes, focusing on two main often overlapping directions: Advanced tools for the optimization of the production chain and specific technologies for low-carbon and sustainable production. Furthermore, the rapid evolution of the technologies in the steel sector require the continuous update of the skills of the industrial workforce. The present review paper, resulting from a recent study developed inside a Blueprint European project, introduces the context of digitalization and some important definitions in both the European industry and the European iron and steel sector. The current technological transformation is depicted, and the main developments funded by European Research Programs are analyzed. Moreover, the impact of digitalization on the steel industry workforce are considered together with the foreseen economic developments. Full article
(This article belongs to the Special Issue Challenges and Prospects of Steelmaking Towards the Year 2050)
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<p>Research Fund for Coal and Steel (RFCS) Projects and the developed enabling technologies.</p> Full article ">Figure 2
<p>Number of RFCS projects by enabling technologies.</p> Full article ">Figure 3
<p>Number of projects related to low-carbon technologies.</p> Full article ">
22 pages, 6941 KiB  
Article
Multi-Objective Optimization of Intermediate Roll Profile for a 6-High Cold Rolling Mill
by Xin Jin, Chang-sheng Li, Yu Wang, Xiao-gang Li, Tian Gu and Yong-guang Xiang
Metals 2020, 10(2), 287; https://doi.org/10.3390/met10020287 - 21 Feb 2020
Cited by 14 | Viewed by 4563
Abstract
The multi-objective optimization of the SmartCrown intermediate roll profile for a cold rolling mill was proposed in this paper in order to improve the strip flatness quality. A coupling model of roll profile and strip flatness was established, and the roll gap profile, [...] Read more.
The multi-objective optimization of the SmartCrown intermediate roll profile for a cold rolling mill was proposed in this paper in order to improve the strip flatness quality. A coupling model of roll profile and strip flatness was established, and the roll gap profile, roll gap crown adjustment range, rolls contact pressure, and strip flatness under different intermediate roll profile parameters were calculated based on the coupling model. The results showed that the roll gap crown adjustment range and rolls contact pressure difference increased with increasing roll profile parameters. The roll profile parameters were multi-optimized based on the non-dominated sorting genetic algorithm II (NSGA-II). The minimum rolls contact pressure difference and maximum roll gap crown adjustment range were taken as the objective function of multi-objective optimization. The optimal roll profile parameters were applied to a six-high five stand tandem cold rolling mills, which improved the flatness quality of the DP780 steel strip. Full article
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<p>The (<b>a</b>) roll spalling and (<b>b</b>) flatness defects.</p> Full article ">Figure 1 Cont.
<p>The (<b>a</b>) roll spalling and (<b>b</b>) flatness defects.</p> Full article ">Figure 2
<p>The diagram of loading, deflection, shifting and discrete of rolls.</p> Full article ">Figure 3
<p>The diagram of rolling process.</p> Full article ">Figure 4
<p>The deflection and flatten of rolls.</p> Full article ">Figure 5
<p>The flow diagram of contact stress between rolls.</p> Full article ">Figure 6
<p>Strip section profile (<b>a</b>) before and (<b>b</b>) after rolling.</p> Full article ">Figure 7
<p>The verification result of (<b>a</b>) rolling force and (<b>b</b>) strip flatness.</p> Full article ">Figure 8
<p>The effect of <span class="html-italic">A</span> on (<b>a</b>) profile and (<b>b</b>) crown adjustment range of intermediate roll gap.</p> Full article ">Figure 9
<p>The effect of <span class="html-italic">B</span> on (<b>a</b>) profile and (<b>b</b>) crown adjustment range of intermediate roll gap.</p> Full article ">Figure 10
<p>The effect of <span class="html-italic">C</span> on (<b>a</b>) profile and (<b>b</b>) crown adjustment range of intermediate roll gap.</p> Full article ">Figure 11
<p>The relationship between <span class="html-italic">Fai</span> and sine function.</p> Full article ">Figure 12
<p>The effect of <span class="html-italic">Fai</span> on (<b>a</b>) profile and (<b>b</b>) crown adjustment range of intermediate roll gap.</p> Full article ">Figure 13
<p>The contact pressure between work roll and intermediate roll under different (<b>a</b>) <span class="html-italic">A</span>, (<b>b</b>) <span class="html-italic">B</span>, (<b>c</b>) <span class="html-italic">C,</span> and (<b>d</b>) <span class="html-italic">Fai</span>.</p> Full article ">Figure 14
<p>The contact pressure between intermediate roll and backup roll under different (<b>a</b>) <span class="html-italic">A</span>, (<b>b</b>) <span class="html-italic">B</span>, (<b>c</b>) <span class="html-italic">C,</span> and (<b>d</b>) <span class="html-italic">Fai</span>.</p> Full article ">Figure 15
<p>The strip flatness under different (<b>a</b>) <span class="html-italic">A</span>, (<b>b</b>) <span class="html-italic">B</span>, (<b>c</b>) <span class="html-italic">C</span>, and (<b>d</b>) <span class="html-italic">Fai</span>.</p> Full article ">Figure 16
<p>The (<b>a</b>) evolution of objective function and the (<b>b</b>) Pareto-optimal solution.</p> Full article ">Figure 17
<p>The (<b>a</b>) roll gap crown adjustment range and (<b>b</b>) contact pressure between work roll and intermediate roll (QWI) of optimal roll profile.</p> Full article ">Figure 18
<p>The measured flatness of a certain moment. (<b>a</b>) Roll profile 1 and (<b>b</b>) Roll profile 2.</p> Full article ">Figure 19
<p>The flatness measured value of whole coil. (<b>a</b>) Roll profile 1 and (<b>b</b>) Roll profile 2.</p> Full article ">
19 pages, 10332 KiB  
Article
Metallurgical Characterization of the Interfaces in Steel Plates Clad with Austenitic Steel or High Ni Alloys by Hot Rolling
by Fabio Giudice, Severino Missori, Francesco Murdolo and Andrea Sili
Metals 2020, 10(2), 286; https://doi.org/10.3390/met10020286 - 21 Feb 2020
Cited by 7 | Viewed by 4561
Abstract
An integrated experimental-theoretical approach to the metallurgical characterization of the interfaces in steel plates clad by hot rolling is proposed. Three different couplings of materials have been studied: ASTM A 515 Gr.60 low carbon steel clad with austenitic stainless steel AISI 304L; extra [...] Read more.
An integrated experimental-theoretical approach to the metallurgical characterization of the interfaces in steel plates clad by hot rolling is proposed. Three different couplings of materials have been studied: ASTM A 515 Gr.60 low carbon steel clad with austenitic stainless steel AISI 304L; extra low carbon steel ASTM A283 clad with high Ni content Alloy 59; and, low carbon steel AISI 1010 clad with Cu-Ni Monel 400. Experimental investigations, which are addressed to analyse the microstructural changes near the interfaces and identify the present phases, have been carried out through scanning electron microscopy (SEM) observations, microanalytical measurements by energy dispersive spectroscopy (EDS), and Vickers microhardness tests. In all of the cases examined, the zones that are affected by detrimental microstructural changes results in being considerably less thick than the overall cladding layer. Simulations that are based on theoretical diffusion modelling have been integrated to the experimental characterization by introducing a cladding parameter that acts on the diffusion bonding efficiency, in order to evaluate the effects of process temperature and time variations on diffusion bonding efficiency and stability. In particular, this analytical investigation has shown how the shorter is the duration of the diffusion transient and the higher the temperature, the lower results the sensitivity of the diffusion processes to temperature fluctuations. Full article
(This article belongs to the Special Issue Clad Metals: Fabrication, Properties and Applications)
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<p>ASTM A515 Gr.60/AISI 304 L interface.</p> Full article ">Figure 2
<p>Austenitic steel side: (<b>a</b>) zone of carbide precipitation, (<b>b</b>) detail of the austenitic deformed grains close to the cladding line.</p> Full article ">Figure 3
<p>The results of Vickers microhardness test along a transversal line to the interface ASTM A515 Gr.60/AISI 304 L.</p> Full article ">Figure 4
<p>Results of energy dispersive spectroscopy (EDS) measurements along a line orthogonal to the interface ASTM A515 Gr.60/AISI 304 L. Concentration profiles according to the distance from the initial interface: (<b>a</b>) Ni, (<b>b</b>) Cr.</p> Full article ">Figure 5
<p>Labelled scanning electron microscopy (SEM) micrograph of the austenitic/ferritic interface.</p> Full article ">Figure 6
<p>Results of punctual EDS measurements at the austenitic/ferritic interface: (<b>a</b>) SEM micrograph with indication of four points; and, (<b>b</b>) WRC-92 diagram utilized for phase identification.</p> Full article ">Figure 7
<p>Sensitized region at the austenitic side near the cladding line (ASTM Test A262-Practice E).</p> Full article ">Figure 8
<p>ASTM A283/Alloy 59 interface.</p> Full article ">Figure 9
<p>Microstructure of the Alloy 59.</p> Full article ">Figure 10
<p>The results of Vickers microhardness test along a transversal line to the interface ASTM A283/Alloy59.</p> Full article ">Figure 11
<p>Results of EDS measurements along a line orthogonal to the interface ASTM A2830/Alloy 59. Concentration profiles vs. the distance from the initial interface: (a) Ni; (<b>b</b>) Cr; (<b>c</b>) Mo; and, (<b>d</b>) Fe.</p> Full article ">Figure 12
<p>AISI 1010/Monel 400 interface.</p> Full article ">Figure 13
<p>Results of Vickers microhardness test along a line orthogonal to the interface AISI 1010/Monel 400.</p> Full article ">Figure 14
<p>Results of EDS measurements along a line orthogonal to the interface AISI 1010/Monel 400. Concentration profiles: (<b>a</b>) Ni; (<b>b</b>) Cr; (<b>c</b>) Fe.</p> Full article ">Figure 15
<p>Experimental and calculated concentration profiles along a line orthogonal to the interface ASTM A515 Gr.60/AISI 304 L.</p> Full article ">Figure 16
<p>Trends of parameter <span class="html-italic">D∙t</span> vs <span class="html-italic">T</span> for different values of diffusion transient duration.</p> Full article ">Figure 17
<p>Concentration profiles of nickel along a line orthogonal to the interface ASTM A515 Gr.60/AISI 304 L calculated for different fluctuations of <span class="html-italic">T</span><sub>eq.</sub> (<span class="html-italic">t</span> = 3600 s).</p> Full article ">Figure 18
<p>Concentration profiles of carbon along a line orthogonal to the interface ASTM A515 Gr.60/AISI 304L calculated for ±100 K fluctuations of <span class="html-italic">T</span><sub>eq.</sub>: comparison between <span class="html-italic">t</span> = 3600 s and <span class="html-italic">t</span> = 1800 s.</p> Full article ">Figure 19
<p>Concentration profiles of nickel along a line orthogonal to the interface AISI 1010/Monel 400 Cu-Ni alloy, calculated for different fluctuations of <span class="html-italic">T</span><sub>eq.</sub> (<span class="html-italic">t</span> = 3600 s).</p> Full article ">
13 pages, 3799 KiB  
Article
Effects of Microstructure, Mechanical and Physical Properties on Machinability of Graphite Cast Irons
by Jiangzhuo Ren, Fengzhang Ren, Fengjun Li, Linkai Cui, Yi Xiong and Alex A. Volinsky
Metals 2020, 10(2), 285; https://doi.org/10.3390/met10020285 - 21 Feb 2020
Cited by 8 | Viewed by 4535
Abstract
Flake (FGI) and spheroidal (SGI) graphite cast irons are often used to produce workpieces, which often need to be machined. Machinability differences under various machining methods are the basis for choosing machining equipment and technology. In this work, FGI and SGI were used [...] Read more.
Flake (FGI) and spheroidal (SGI) graphite cast irons are often used to produce workpieces, which often need to be machined. Machinability differences under various machining methods are the basis for choosing machining equipment and technology. In this work, FGI and SGI were used to produce tractor front brackets, and the machinability of both materials under turning and drilling processes was compared. The machinability (turning and drilling ability) has been evaluated in terms of machining load, chips shape, surface roughness, and tool temperature. The influence of materials microstructure and thermal conductivity on the machinability was analyzed. In the turning process, the cutting force and its standard deviation of the FGI were larger than the SGI due to the higher volume fraction of pearlite. The surface roughness was similar in both materials. In the drilling process, the even action of the friction and cutting force on the bit turned into similar drilling loads for both materials. Higher friction and lower thermal conductivity caused a higher bit temperature in SGI drilling compared to FGI. The chip breaking was worse in SGI drilling, where the longer chips scratched the internal surface of the holes, resulting in the higher surface roughness. Full article
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<p>The front bracket casting and its tested parts.</p> Full article ">Figure 2
<p>Schematic diagram of samples cutting sites and hardness measurement locations.</p> Full article ">Figure 3
<p>Cutting tool and octagonal ring dynamometer installation.</p> Full article ">Figure 4
<p>Calibration of measuring force system.</p> Full article ">Figure 5
<p>Measurement scene.</p> Full article ">Figure 6
<p>The microstructure of two materials: (<b>a</b>) HT250 and (<b>b</b>) QT450.</p> Full article ">Figure 7
<p>The turning chips appearance at 2 mm cutting depth: (<b>a</b>) HT250 and (<b>b</b>) QT450.</p> Full article ">Figure 8
<p>The drilling chips appearance for the Φ 14.5 mm bit diameter with 800 rpm rotating speed and 0.5 mm/rev feed rate: (<b>a</b>) HT250 and (<b>b</b>) QT450.</p> Full article ">Figure 9
<p>Three-dimensional shape of the holes internal surface of the HT250 borehole: (<b>a</b>) Φ 14.5 mm bit and (<b>b</b>) Φ 19 mm bit; QT450 borehole: (<b>c</b>) Φ 14.5 mm bit and (<b>d</b>) Φ 19 mm bit (1.6 mm measured length).</p> Full article ">
12 pages, 3497 KiB  
Article
PbO-SiO2 Based Glass Coating of PbI2 Doped PbTe
by Yatir Sadia, Idan Koron and Yaniv Gelbstein
Metals 2020, 10(2), 284; https://doi.org/10.3390/met10020284 - 21 Feb 2020
Cited by 1 | Viewed by 3334
Abstract
Thermoelectrics is one promising way of increasing the efficiency of machines and devices by reusing some of the waste heat produced. One obstacle for commercialization is the need to coat the materials to prevent sublimation and oxidation of the thermoelectric materials. Such coatings [...] Read more.
Thermoelectrics is one promising way of increasing the efficiency of machines and devices by reusing some of the waste heat produced. One obstacle for commercialization is the need to coat the materials to prevent sublimation and oxidation of the thermoelectric materials. Such coatings were designed for PbI2 doped PbTe using a (SiO2)0.68(PbO)0.3(B2O3)0.01(Na2O)0.01 based glass designed for operation at 500 °C. In this research various conditions of the coating process were examined. The effect of the atmosphere on the bonding and densification of the coating was studied using argon, vacuum and air. From the three air shows, the best bonding characteristics were from a better flow of glass and increased bonding between the oxidized PbTe layer and glass. This also created a PbO rich glass in the interface between the glass and the PbTe sample. The effect of 0, 3, and 6 wt. % NaCl additive to the solution was tested and showed that NaCl achieves better coverage due to high green body density, reaction of NaCl with the glass and removal of remaining CO2 from the glass in the form of decomposing Na2CO3. In addition, when testing the time and temperature, it was shown that the temperature of 520 °C was the minimum needed for high densification of the glass, but a duration shorter than 30 min did not allow for bonding of the glass to the substrate despite adequate densification. Finely, to obtain a well bonded coating with full coverage over the sample, the glass was coated with 6% NaCl in air at 520 °C for 30 min. Full article
(This article belongs to the Special Issue Thermoelectric Compounds: Processing, Properties and Applications)
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<p>Some of the challenges in turning thermoelectric devices into generators. In red the challenges of sublimation and oxidation are marked, with the hot-side of the element showing sublimation. A suggested coating is marked in a dashed red line.</p> Full article ">Figure 2
<p>(<b>a</b>) XRD patterns of the coated samples, following 520 °C/30 min. thermal treatment under argon (black), vacuum (red) and air (blue) atmospheres. (<b>b</b>) Dilatometry measuremnts of the glass powder in air (black) and argon (green) with a pre-sintered pellete (red) as reference, showing densification at lower temperatures in air and that densification is expected at a little above 500 °C.</p> Full article ">Figure 3
<p>(<b>a</b>) SEM micrographs of glass coated PbTe, following 520 °C/30 min. thermal treatment under vacuum (<b>a</b>,<b>c</b>) and air (<b>b</b>,<b>d</b>) atmospheres. (<b>a</b>,<b>b</b>) are secondary electrons images while (<b>c</b>,<b>d</b>) are back scattered electrons images.</p> Full article ">Figure 4
<p>(<b>a</b>) XRD patterns of the coated samplesupon addition of 3 wt. % (black) and 6 wt. % NaCl (red). (<b>b</b>) Dilatometry measuremnts following different factors variation; black and red are glass in air and pre-sintered pellete redrawn from <a href="#metals-10-00284-f002" class="html-fig">Figure 2</a>b for a comperison; Dried glass suspention (green); dried glass suspension with 6% NaCl (yellow); glass powder with NaCl powder and no water (purple) and NaCl powder only (pink).</p> Full article ">Figure 5
<p>Scanning electron microscopy (SEM) micrographs of glass coated PbTe, following 3 wt. % (<b>a</b>,<b>c</b>) and 6 wt. % (<b>b</b>,<b>d</b>) NaCl addition. (<b>a</b>,<b>b</b>) are secondary electrons images while (<b>c</b>,<b>d</b>) are back scattered electrons images.</p> Full article ">Figure 6
<p>X-ray diffraction (XRD) patterns of the coated samples following thermal treatments of 500 °C for 30 min (Red) and 520 °C for 10 min (Black).</p> Full article ">Figure 7
<p>SEM micrographs following coating and subsequent thermal treatment at either a lower temperature of 500 °C (<b>a</b>,<b>c</b>) or shorter duration of 10 min. (<b>b</b>,<b>d</b>), compared to the previously applied 520 °C/30 min. condition. (<b>a</b>,<b>b</b>) are secondary electrons images while (<b>c</b>,<b>d</b>) are back scattered electrons images.</p> Full article ">
12 pages, 2605 KiB  
Article
Ab Initio Calculations on Elastic Properties of IF Steel Matrix Phase at High Temperature Based on Lattice Expansion Theory
by Songyuan Ai, Mujun Long, Siyuan Zhang, Dengfu Chen, Zhihua Dong, Peng Liu, Yanming Zhang and Huamei Duan
Metals 2020, 10(2), 283; https://doi.org/10.3390/met10020283 - 21 Feb 2020
Cited by 6 | Viewed by 3241
Abstract
Elucidating the evolution law of the elastic properties of the matrix phase is of great significance for the control of steel properties and quality during continuous casting and subsequent heat treatment. In this paper, thermal expansion experiments and ab initio calculations are used [...] Read more.
Elucidating the evolution law of the elastic properties of the matrix phase is of great significance for the control of steel properties and quality during continuous casting and subsequent heat treatment. In this paper, thermal expansion experiments and ab initio calculations are used to study the elastic properties of the interstitial free (IF) steel matrix phase in different magnetic states and crystal structures. The results show that the bulk modulus B and the tetragonal shear elastic constant C’ for the entire temperature range decrease with increasing temperature, but C44 is the opposite. While from paramagnetic (PM) to ferromagnetic (FM) state, C’(C44) have changed ~188% (~27%), B increases by ~55% during the crystal structure change (fcc→bcc). With the FM to PM state, the Zener anisotropy parameter increases sharply, and Young’s modulus decreases significantly in the [001] direction; the maximum difference is ~76 GPa. The evolution rate of average Young’s modulus in single bcc-phase FM (fcc-phase PM) range reaches ~5.5(~5.6) × 10−2 GPa K−1. The research provides an effective method for ab initio calculation of the elastic properties of interstitial free and ultra-low carbon steels at high temperature, also furnishing a basis for the application of ab initio calculations to the high temperature performance of steel materials. Full article
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<p>Schematic diagram of the position for samples.</p> Full article ">Figure 2
<p>Expansion curve and linear thermal expansion coefficient (LTEC) curve of sample at heating rate of 5 K/min.</p> Full article ">Figure 3
<p>Microstructure of the IF steel after thermal experiment.</p> Full article ">Figure 4
<p>Wigner–Seitz radius w of single bcc-phase and fcc-phase as a function of temperature. Experimental data of lattice parameter measurements (open symbols) from References [<a href="#B33-metals-10-00283" class="html-bibr">33</a>,<a href="#B34-metals-10-00283" class="html-bibr">34</a>] are taken. Solid symbols show the dilation test temperature dependencies of the Wigner–Seitz radius for single bcc-phase and fcc-phase, from left to right, and their linear regressions are plotted with red dashed lines.</p> Full article ">Figure 5
<p>Variation of magnetic moment <span class="html-italic">μ</span> with temperature. The values of present results are plotted as solid symbols connected with lines and the available data from Reference [<a href="#B36-metals-10-00283" class="html-bibr">36</a>] are represented as open symbols.</p> Full article ">Figure 6
<p>Calculated temperature dependence of elastic moduli of bcc-phase FM, bcc-phase PM, and fcc-phase PM in comparison with published data of References [<a href="#B38-metals-10-00283" class="html-bibr">38</a>,<a href="#B39-metals-10-00283" class="html-bibr">39</a>,<a href="#B40-metals-10-00283" class="html-bibr">40</a>].</p> Full article ">Figure 7
<p>Calculated 3D characteristic surfaces (<b>a</b>–<b>d</b>) and the 2D section (110) of the Young’s modulus as a function of crystallographic direction in IF steel matrix phase. The values on the color scale and on the axes are in GPa. The Zener anisotropy parameter <span class="html-italic">A<sub>Z</sub></span> is plotted against temperature for the matrix phase, along with the available data (open symbols) from References [<a href="#B7-metals-10-00283" class="html-bibr">7</a>,<a href="#B38-metals-10-00283" class="html-bibr">38</a>,<a href="#B39-metals-10-00283" class="html-bibr">39</a>,<a href="#B40-metals-10-00283" class="html-bibr">40</a>]. (<b>a</b>) 1043 K bcc-phase FM, (<b>b</b>) 1043 K bcc-phase PM, (<b>c</b>) 1083 K bcc-phase PM, (<b>d</b>) 1083 K fcc-phase PM, (<b>e</b>) the 2D section (110) of the Young’s modlus, (<b>f</b>) <span class="html-italic">A<sub>Z</sub></span> of matrix phase.</p> Full article ">Figure 8
<p>Relationship of average Young’s modulus and Poisson’s ratio of matrix phase with temperature for different lattice and magnetic structures, along with the available data from References [<a href="#B40-metals-10-00283" class="html-bibr">40</a>,<a href="#B42-metals-10-00283" class="html-bibr">42</a>].</p> Full article ">
13 pages, 75165 KiB  
Article
Influence of Martensite Deformation on Cu Precipitation Strengthening
by Jaromir Dlouhy, Pavel Podany and Ján Džugan
Metals 2020, 10(2), 282; https://doi.org/10.3390/met10020282 - 21 Feb 2020
Cited by 5 | Viewed by 2781
Abstract
Cu precipitation strengthening was compared in steels after treatments with and without cold rolling. A 0.2% C steel containing up to 1.5% Cu was quenched and tempered. Cu precipitation took place during tempering and increased its yield strength (YS). Quenched and tempered samples [...] Read more.
Cu precipitation strengthening was compared in steels after treatments with and without cold rolling. A 0.2% C steel containing up to 1.5% Cu was quenched and tempered. Cu precipitation took place during tempering and increased its yield strength (YS). Quenched and tempered samples were compared with samples where cold rolling was performed between quenching and tempering. They exhibited significantly different mechanical properties. In addition, Cu alloying influenced the properties of each group of samples in different ways. The quenched and tempered samples exhibited behavior that is typical of precipitation hardening. Cu caused yield strength to increase with tempering temperature and time. The cold rolling of martensite reduced the maximal Cu-related strengthening and also eliminated its time and temperature dependence. Full article
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<p>Yield stress (YS) of samples without tempering and after tempering at different temperatures for 60 min.</p> Full article ">Figure 2
<p>Yield stress after different times of tempering at 400 °C for 0, 15, 30, 60, and 120 min.</p> Full article ">Figure 3
<p>Engineering stress–strain diagrams of samples of 0Cu, 1Cu, and 1.5Cu materials: (<b>a</b>) quenched and tempered samples (60 min); (<b>b</b>) quenched, rolled, and tempered samples (60 min).</p> Full article ">Figure 4
<p>Microstructures of the material 0Cu tempered for 60 min at 500 °C: (<b>a</b>) quenched and tempered (QT); (<b>b</b>) quenched, rolled, and tempered (QRT) (the rolling direction is horizontal in the image).</p> Full article ">Figure 5
<p>Microstructures of the samples after tempering at 500 °C for 1 h: (<b>a</b>) 0Cu QT; (<b>b</b>) 0Cu QRT; (<b>c</b>) 1Cu QT; (<b>d</b>) 1Cu QRT; (<b>e</b>) 1Cu QT; (<b>f</b>) 1Cu QRT.</p> Full article ">Figure 6
<p>TEM bright field image of the sample 1.5Cu QT 500 °C/60 min. There are dark elongated particles of carbides (c) and small round precipitates (p) within the ferrite matrix.</p> Full article ">Figure 7
<p>TEM image of a precipitate in the sample 1.5Cu QT 500 °C/60 min. It is probably two coalesced particles or two overlapping ones. Their structure corresponds to fcc Cu.</p> Full article ">Figure 8
<p>Difference ΔYS between the YS of Cu-alloyed materials and 0Cu samples tempered for 60 min. YS of 0Cu samples is the zero-line.</p> Full article ">Figure 9
<p>Difference ΔYS between YS of Cu-alloyed materials and 0Cu samples tempered at 400 °C. YS of 0Cu samples is the zero-line.</p> Full article ">Figure 10
<p>Inverse pole figure (IPF) maps for 0Cu and 1.5Cu materials tempered for 60 min at temperature 500 °C: (<b>a</b>) 0Cu QT; (<b>b</b>) 0Cu QRT; (<b>c</b>) 1.5Cu QT; (<b>d</b>) 1.5Cu QRT. The map step was 0.05 µm. The coloring indicates local lattice orientation (legend in (<b>d</b>)). The rolling direction is oriented vertically for the QRT samples.</p> Full article ">Figure 11
<p>Kernell average misorientation (KAM) maps for the 0Cu and 1.5Cu materials tempered for 60 min at temperature 500 °C: (<b>a</b>) 0Cu QT; (<b>b</b>) 0Cu QRT; (<b>c</b>) 1.5Cu QT; (<b>d</b>) 1.5Cu QRT. The map step was 0.05 µm. Each pixel is colored according to its average misorientation with its six neighbors (maps were acquired in a hexagonal pixel pattern). The legend is in (<b>d</b>).</p> Full article ">Figure 12
<p>Grain size measured from EBSD maps for grains separated by boundaries with misorientation &gt;1°: (<b>a</b>) 0Cu QT; (<b>b</b>) 0Cu QRT; (<b>c</b>) 1.5Cu QT; (<b>d</b>) 1.5Cu QRT. Map step was 0.05 µm.</p> Full article ">Figure 12 Cont.
<p>Grain size measured from EBSD maps for grains separated by boundaries with misorientation &gt;1°: (<b>a</b>) 0Cu QT; (<b>b</b>) 0Cu QRT; (<b>c</b>) 1.5Cu QT; (<b>d</b>) 1.5Cu QRT. Map step was 0.05 µm.</p> Full article ">
43 pages, 6971 KiB  
Review
Processing and Properties of Reversion-Treated Austenitic Stainless Steels
by Antti Järvenpää, Matias Jaskari, Anna Kisko and Pentti Karjalainen
Metals 2020, 10(2), 281; https://doi.org/10.3390/met10020281 - 21 Feb 2020
Cited by 90 | Viewed by 11283
Abstract
Strength properties of annealed austenitic stainless steels are relatively low and therefore improvements are desired for constructional applications. The reversion of deformation induced martensite to fine-grained austenite has been found to be an efficient method to increase significantly the yield strength of metastable [...] Read more.
Strength properties of annealed austenitic stainless steels are relatively low and therefore improvements are desired for constructional applications. The reversion of deformation induced martensite to fine-grained austenite has been found to be an efficient method to increase significantly the yield strength of metastable austenitic stainless steels without impairing much their ductility. Research has been conducted during thirty years in many research groups so that the features of the reversion process and enhanced properties are reported in numerous papers. This review covers the main variables and phenomena during the reversion processing and lists the static and dynamic mechanical properties obtained in laboratory experiments, highlighting them to exceed those of temper rolled sheets. Moreover, formability, weldability and corrosion resistant aspects are discussed and finally the advantage of refined grain structure for medical applications is stated. The reversion process has been utilized industrially in a very limited extent, but apparently, it could provide a feasible processing route for strengthened austenitic stainless steels. Full article
(This article belongs to the Special Issue Manufacturing and Application of Stainless Steels)
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<p>Reversion treatment of a metastable austenitic stainless-steel consisting of cold rolling and annealing stages. Important variables have been designated.</p> Full article ">Figure 2
<p>Examples of increase of the deformation-induced martensite fraction during cold rolling deformation in 201/201L, 304/304L and 301LN grades. The chemical compositions of the steels are in <a href="#metals-10-00281-t001" class="html-table">Table 1</a>.</p> Full article ">Figure 3
<p>Reversion in 301LN ASS occurred by the shear reversion (<b>a</b>), where dislocation free grains are formed by continuous recrystallization (arrows pointing out such subgrains) [<a href="#B29-metals-10-00281" class="html-bibr">29</a>], diffusional reversion (<b>b</b>) (reproduced from [<a href="#B84-metals-10-00281" class="html-bibr">84</a>], with permission from Elsevier, 2017) and diffusional reversion with chromium nitride precipitation (<b>c</b>,<b>d</b>) (reproduced from [<a href="#B22-metals-10-00281" class="html-bibr">22</a>], with permission from Springer Nature, 2007).</p> Full article ">Figure 4
<p>The effect of prior cold rolling reduction and annealing parameters on the reversion kinetics: Annealing temperature (<b>a</b>) and soaking time at 700 °C (<b>b</b>). Data collected from References [<a href="#B23-metals-10-00281" class="html-bibr">23</a>,<a href="#B40-metals-10-00281" class="html-bibr">40</a>] (<b>a</b>) and [<a href="#B23-metals-10-00281" class="html-bibr">23</a>,<a href="#B26-metals-10-00281" class="html-bibr">26</a>,<a href="#B88-metals-10-00281" class="html-bibr">88</a>] (<b>b</b>). Legend: Reduction (%CR), annealing temperature (°C) and time (s).</p> Full article ">Figure 5
<p>The temperature range for the reversion, dependent on the reversion mechanism. Various reversion treatments (numbered 1…5) are possible. The original idea of the reversion-temperature-time diagram is from Reference [<a href="#B39-metals-10-00281" class="html-bibr">39</a>].</p> Full article ">Figure 6
<p>Examples of reversed microstructures in 301LN with various grain size. Partially reversed (32% cold rolling, 750 °C for 0.1 s) with shear reversed (SR) fine and medium-size grains, a large DA grain and retained DIM (red colored phase) (<b>a</b>) and diffusionally reversed ultrafine and medium-size grains in 63% cold-rolled steel, annealed at 800 °C for 1 s (<b>b</b>) and at 900 °C for 1 s (<b>c</b>).</p> Full article ">Figure 7
<p>Grain growth of reversed grain size at 800 and 900 °C in 201L [<a href="#B67-metals-10-00281" class="html-bibr">67</a>], 204Cu [<a href="#B82-metals-10-00281" class="html-bibr">82</a>], 301LN [<a href="#B26-metals-10-00281" class="html-bibr">26</a>], 304L [<a href="#B17-metals-10-00281" class="html-bibr">17</a>,<a href="#B70-metals-10-00281" class="html-bibr">70</a>] and 304 [<a href="#B11-metals-10-00281" class="html-bibr">11</a>] ASSs.</p> Full article ">Figure 8
<p>Examples of engineering tensile stress–strain curves for a 301LN ASS with refined grain size after various reversion treatments. Data collected from References [<a href="#B87-metals-10-00281" class="html-bibr">87</a>,<a href="#B102-metals-10-00281" class="html-bibr">102</a>,<a href="#B103-metals-10-00281" class="html-bibr">103</a>]. Legend: Reduction (%CR), temperature (°C), duration (s).</p> Full article ">Figure 9
<p>Examples of the Hall–Petch relations presented for reversed structures. Data for annealed coarse-grained 301LN is included from [<a href="#B151-metals-10-00281" class="html-bibr">151</a>].</p> Full article ">Figure 10
<p>Yield strength-total elongation combinations for some reversion-treated steels listed in <a href="#metals-10-00281-t002" class="html-table">Table 2</a> compared to those of annealed 301LN and temper-rolled 301LN and 201L steel.</p> Full article ">Figure 11
<p>Examples of DIM evolution during tensile testing in various 301LN (M<sub>d30</sub> = 27 °C) structures [<a href="#B82-metals-10-00281" class="html-bibr">82</a>,<a href="#B85-metals-10-00281" class="html-bibr">85</a>,<a href="#B87-metals-10-00281" class="html-bibr">87</a>] and coarse-grained annealed 301LN with M<sub>d30</sub> = 37 °C from Reference [<a href="#B155-metals-10-00281" class="html-bibr">155</a>]. Legend: Reduction (%CR), temperature (°C), duration (s).</p> Full article ">Figure 12
<p>Strain hardening rate curves for selected AISI 301LN structures [<a href="#B85-metals-10-00281" class="html-bibr">85</a>,<a href="#B87-metals-10-00281" class="html-bibr">87</a>] (<b>a</b>) and the relation between fracture strain and the strain at the peak SHR (<b>b</b>). SHR = strain hardening rate.</p> Full article ">Figure 13
<p>Initial parts of stress–strain curves of various reversed structures of a 30LN steel. Data from References [<a href="#B29-metals-10-00281" class="html-bibr">29</a>,<a href="#B87-metals-10-00281" class="html-bibr">87</a>]. Legend: Reduction (%CR), temperature (°C), duration (s).</p> Full article ">Figure 14
<p>Fatigue life of reversion-treated 301LN in total strain-controlled tests compared to coarse-grained or annealed steel based on (<b>a</b>) strain amplitude and (<b>b</b>) mid-life stress amplitude. Data for 304 and 16.5Cr-7Mn-6.5N are from References [<a href="#B43-metals-10-00281" class="html-bibr">43</a>] and [<a href="#B74-metals-10-00281" class="html-bibr">74</a>], respectively.</p> Full article ">Figure 15
<p>Cyclic medium-life stress amplitude (<b>a</b>) and DIM fraction (<b>b</b>) vs number of cycles curves of coarse-grained and reversed fine-grained 301LN at 0.4% and 0.6% strain amplitudes. Data from References [<a href="#B87-metals-10-00281" class="html-bibr">87</a>,<a href="#B99-metals-10-00281" class="html-bibr">99</a>,<a href="#B102-metals-10-00281" class="html-bibr">102</a>].</p> Full article ">Figure 16
<p>Cross-section of a laser weld on a reversion-treated 301LN sheet, hardness profile and grain size distribution. The hardness profile is from Reference [<a href="#B199-metals-10-00281" class="html-bibr">199</a>].</p> Full article ">
13 pages, 3137 KiB  
Article
Effects of Positions and Angulations of Titanium Dental Implants in Biomechanical Performances in the All-on-Four Treatment: 3D Numerical and Strain Gauge Methods
by Aaron Yu-Jen Wu, Jui-Ting Hsu, Lih-Jyh Fuh and Heng-Li Huang
Metals 2020, 10(2), 280; https://doi.org/10.3390/met10020280 - 21 Feb 2020
Cited by 6 | Viewed by 4324
Abstract
In finite element (FE) simulations, the peak bone stresses were higher when loading with a cantilever extension (CE) than when loading without a CE by 33–49% in the cortical bone. In the in vitro experiments, the highest values of principal strain were all [...] Read more.
In finite element (FE) simulations, the peak bone stresses were higher when loading with a cantilever extension (CE) than when loading without a CE by 33–49% in the cortical bone. In the in vitro experiments, the highest values of principal strain were all within the range of the minimum principal strain, and those peak bone strains were 40–58% greater when loading with a CE than when loading without a CE (p < 0.001). This study investigated how varying the implanted position and angulation of anterior implants in the All-on-Four treatment influenced the biomechanical environment in the alveolar bone around the dental implants. Ten numerical simulations of FE models and three in vitro samples of All-on-Four treatment of dental implants were created to investigate the effects of altering the implanted position and angulation type of anterior implants. A single load of 100 N was applied in the molar region in the presence or absence of a CE of the denture. The 3D FE simulations analyzed the von-Mises stresses in the surrounding cortical bone and trabecular bone. For the in vitro tests, the principal bone strains were recorded by rosette strain gauges and statistically evaluated using the Mann–Whitney U test and the Kruskal–Wallis test. Loading in the presence of a CE of the denture induced the highest bone stress and strain, which were 53–97% greater in the FE simulation and 68–140% in the in vitro experiments (p < 0.008) than when loading without a CE. The bone stresses in the FE models of various implanted positions and angulation types of anterior implants were similar to those in the model of a typical All-on-Four treatment. In vitro tests revealed that the bone strains were significantly higher in the samples with various angulation types of anterior implants (p < 0.008). In the All-on-Four treatment of dental implants, the bone stress and strain were higher when the load was applied to the CE of dentures. Altering the position or angulation of the anterior dental implant in the All-on-Four treatment has no benefit in relieving the stress and strain of the bone around the dental implant. Full article
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<p>The ten designs (test type, position of anterior implants, titled type of anterior implant), including (<b>a</b>) Model Typical (typical model, incisor region, vertical implantation), (<b>b</b>) Model P1 (implanted-position test, central incisor region, vertical implantation), (<b>c</b>) Model P2 (implanted-position test, incisor-canine region, vertical implantation), (<b>d</b>) Model P3 (implanted-position test, canine region, vertical implantation), (<b>e</b>) Model T1 (titled-type test, incisor region, 30°-tilted implants toward the distal direction), (<b>f</b>) Model T2 (titled-type test, incisor region, 30°-tilted implants toward the mesial direction), (<b>g</b>–<b>j</b>) Model T3–Model T6 (titled-type test, incisor region, one vertical implant and one 30°-tilted implants toward the mesial or distal direction), were investigated for All-on-Four implant treatment. (<b>k</b>) The 3D computer-aided design (CAD) model of Model Typical.</p> Full article ">Figure 2
<p>The finite element (FE) mesh model of Model Typical. The red circles indicate the areas of boundary condition.</p> Full article ">Figure 3
<p>Three samples involving (<b>a</b>) Sample Typical (typical model, incisor region, vertical implantation), (<b>b</b>) Sample P2 (implanted-position test, incisor-canine region, vertical implantation), and (<b>c</b>) Sample T2 (titled-type test, incisor region, 30°-tilted implants toward the mesial direction) were used in the in vitro tests. (<b>d</b>) Four strain gauges were attached to the buccal side of the alveolar bone around the four implants. Two kinds of loads were applied to the denture with and without a CE. A, B, C, and D refer to the location of each individual implant.</p> Full article ">Figure 4
<p>(<b>a</b>) The loading machine. (<b>b</b>) The experimental sample in the All-on-Four treatment was fixed to the self-developed jig, which was then attached to the platform of the loading machine for applying a single occlusal force to the denture of the sample.</p> Full article ">Figure 5
<p>Distributions of von-Mises stresses in cortical bone in (<b>a</b>) Model Typical, (<b>b</b>) Model P1, (<b>c</b>) Model P2, and (<b>d</b>) Model P3. The arrow shows where the load was applied.</p> Full article ">Figure 6
<p>Distributions of von-Mises stresses in cortical bone in (<b>a</b>) Model Typical, (<b>b</b>) Model T1, (<b>c</b>) Model T2, (<b>d</b>) Model T3, (<b>e</b>) Model T4, (<b>f</b>) Model T5, and (<b>g</b>) Model T6.</p> Full article ">Figure 7
<p>Peak values of the minimum principal strain in (<b>a</b>) Sample Typical, (<b>b</b>) Sample P2, and (<b>c</b>) Sample T2 when loading with and without a CE. Data are median and interquartile-range values. An asterisk indicates a significant difference (<span class="html-italic">p</span> &lt; 0.008).</p> Full article ">
16 pages, 8152 KiB  
Review
Research Status and Progress of Welding Technologies for Molybdenum and Molybdenum Alloys
by Qi Zhu, Miaoxia Xie, Xiangtao Shang, Geng An, Jun Sun, Na Wang, Sha Xi, Chunyang Bu and Juping Zhang
Metals 2020, 10(2), 279; https://doi.org/10.3390/met10020279 - 20 Feb 2020
Cited by 21 | Viewed by 5298
Abstract
Owing to its potential application prospect in novel accident tolerant fuel, molybdenum alloys and their welding technologies have gained great importance in recent years. The challenges of welding molybdenum alloys come from two aspects: one is related to its powder metallurgy manufacturing process, [...] Read more.
Owing to its potential application prospect in novel accident tolerant fuel, molybdenum alloys and their welding technologies have gained great importance in recent years. The challenges of welding molybdenum alloys come from two aspects: one is related to its powder metallurgy manufacturing process, and the other is its inherent characteristics of refractory metal. The welding of powder metallurgy materials has been associated with issues such as porosity, contamination, and inclusions, at levels which tend to degrade the service performances of a welded joint. Refractory metals usually present poor weldability due to embrittlement of the fusion zone as a result of impurities segregation and the grain coarsening in the heat-affected zone. A critical review of the current state of the art of welding Mo alloys components is presented. The advantages and disadvantages of the various methods, i.e., electron-beam welding (EBW), tungsten-arc inert gas (TIG) welding, laser welding (LW), electric resistance welding (ERW), and brazing and friction welding (FW) in joining Mo and Mo alloys, are discussed with a view to imagine future directions. This review suggests that more attention should be paid to high energy density laser welding and the mechanism and technology of welding Mo alloys under hyperbaric environment. Full article
(This article belongs to the Special Issue Technology of Welding and Joining)
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<p>(<b>a</b>) Electron backscatter diffraction (EBSD) image of cross section of Mo joint achieved by laser welding, and (<b>b</b>) enlarged view of area (<b>a</b>) [<a href="#B10-metals-10-00279" class="html-bibr">10</a>].</p> Full article ">Figure 2
<p>(<b>a</b>) SEM image of the fracture of Mo joint with plenty of MoO<sub>2</sub> at grain boundary, and (<b>b</b>) TEM bright field image of the weld bead of the Mo joint [<a href="#B10-metals-10-00279" class="html-bibr">10</a>].</p> Full article ">Figure 3
<p>Typical porosity morphology in Mo joints produced by laser welding: (<b>a</b>) orthogonal experiment 1# parameter sample, (<b>b</b>) orthogonal experiment 2# parameter sample, and (<b>c</b>) orthogonal experiment 3# parameter sample. [<a href="#B11-metals-10-00279" class="html-bibr">11</a>].</p> Full article ">Figure 4
<p>Cross-sectional morphology of the welded joints produced by electron-beam welding (EBW) welding (<b>a</b>) pore defects diagram with improved parameters, (<b>b</b>) crystal cracks, (<b>c</b>) low plastic embrittlement crack, and (<b>d</b>) forming after optimized process. [<a href="#B22-metals-10-00279" class="html-bibr">22</a>].</p> Full article ">Figure 5
<p>Microstructure of weld joint (<b>a</b>) welding seam zone, (<b>b</b>) heat-affected zone, and (<b>c</b>) TZM martix. [<a href="#B24-metals-10-00279" class="html-bibr">24</a>].</p> Full article ">Figure 6
<p>SEM observation of 50Mo-50Re overlap joint achieved by laser welding: (<b>a</b>) cross section and and (<b>b</b>) fracture surface [<a href="#B29-metals-10-00279" class="html-bibr">29</a>].</p> Full article ">Figure 7
<p>Fractographs of (<b>a</b>) the parent metal, (<b>b</b>) EB weld joint, and (<b>c</b>) LGTHW Joint of tensile samples. Change in the morphology and the presence of sharp faceting in samples containing weld joints could be noticed [<a href="#B5-metals-10-00279" class="html-bibr">5</a>].</p> Full article ">Figure 8
<p>TEM micrographs showing (<b>a</b>) the presence of the precipitates within the matrix. Inset shows the magnified view; (<b>b</b>) Magnified view showing the presence of Mo-oxide phase in the weld region. (<b>c</b>) The presence of needle-shaped long oxides may be noticed [<a href="#B5-metals-10-00279" class="html-bibr">5</a>].</p> Full article ">Figure 9
<p>The effect of C addition on tensile strength of the Mo alloy joints [<a href="#B9-metals-10-00279" class="html-bibr">9</a>].</p> Full article ">Figure 10
<p>TEM images of (<b>a</b>) molybdenum oxide particles in the FZ of LW joint and (<b>b</b>) carbide particles in the FZ of SC-150 joint [<a href="#B9-metals-10-00279" class="html-bibr">9</a>].</p> Full article ">Figure 11
<p>The cross-sectional microstructures of the NS Mo alloy laser welding joints. Heat input: (<b>a</b>) 250 J/cm and (<b>b</b>) 3600 J/cm [<a href="#B12-metals-10-00279" class="html-bibr">12</a>].</p> Full article ">Figure 12
<p>Tracing of the metal elements in the weld bead of the Mo-0.03Ti-B joint: (<b>a</b>) schematic diagram of the location where Zr was added. (<b>b</b>) Analysis of Mo, Zr, and Ti on the cross section of the joint by EDS surface scanning. (<b>c</b>) Analysis of Zr in zone A in panel (<b>b</b>) by EDS surface scanning. [<a href="#B10-metals-10-00279" class="html-bibr">10</a>].</p> Full article ">Figure 13
<p>Reconstructed 3D transparent distribution of pores in the three joints achieved under (<b>a</b>) P = 1.2 kW; (<b>b</b>) average power = 1.2 kW, Amplitude = 300 W, frequency = 50 HZ, N = 2 cycles; and (<b>c</b>) average power = 1.2 kW, amplitude = 300 W, frequency = 150 HZ, N = 2 cycles, adding Zr, respectively. [<a href="#B11-metals-10-00279" class="html-bibr">11</a>].</p> Full article ">Figure 14
<p>(<b>a</b>) Strength-displacement curves of Mo-0.03Ti-B and pure Mo tube; (<b>b</b>) images of the joint after fracture testing; (<b>c</b>) hydrostatic test curve and hydraulic bursting test curve for the molybdenum joint of Mo-0.03Ti-B; and (<b>d</b>) images of joint after hydraulic bursting testing [<a href="#B10-metals-10-00279" class="html-bibr">10</a>].</p> Full article ">Figure 15
<p>(<b>a</b>) Fracture cross section of the joint without Ni, (<b>b</b>) fracture cross section of the joint with Ni, and (<b>c</b>) tensile results of the two joints [<a href="#B31-metals-10-00279" class="html-bibr">31</a>].</p> Full article ">Figure 16
<p>(<b>a</b>,<b>b</b>) The shape and microstructure of the nuggets welded using the electrodes shown in (<b>c</b>,<b>d</b>), respectively. The horizontal arrows show the interfaces between two workpieces in panels (<b>a</b>) and (<b>b</b>) [<a href="#B38-metals-10-00279" class="html-bibr">38</a>].</p> Full article ">Figure 17
<p>Scheme of friction welding in IME82 oil [<a href="#B45-metals-10-00279" class="html-bibr">45</a>].</p> Full article ">Figure 18
<p>Medium-size friction welded Mo-tube (OD: 150 mm, ID: 130mm) [<a href="#B46-metals-10-00279" class="html-bibr">46</a>].</p> Full article ">Figure 19
<p>Detail of the microstructure friction welded samples: Overview of (<b>a</b>) Mo; overview of (<b>b</b>) TZM; (<b>c</b>) zone (1) of panel (<b>a</b>); (<b>d</b>) zone (2) of panel (a); (<b>e</b>) zone (3) of panel (<b>a</b>); (<b>f</b>) zone (4) of panel (<b>a</b>); (<b>g</b>) zone (1) of panel (<b>b</b>); (<b>h</b>) zone (2) of panel (<b>b</b>); (<b>i</b>) zone (3) of panel (<b>b</b>); (<b>j</b>) zone (4) of panel (<b>b</b>). [<a href="#B48-metals-10-00279" class="html-bibr">48</a>].</p> Full article ">
12 pages, 5589 KiB  
Article
CMT-Based Wire Arc Additive Manufacturing Using 316L Stainless Steel: Effect of Heat Accumulation on the Multi-Layer Deposits
by Seung Hwan Lee
Metals 2020, 10(2), 278; https://doi.org/10.3390/met10020278 - 20 Feb 2020
Cited by 67 | Viewed by 8716
Abstract
CMT welding sources are garnering attention as alternative heat sources for wire arc additive manufacturing because of their low-heat input. A comprehensive experimental and numerical study on the multi-layer deposition of STS316L was performed to investigate effect of heat accumulation during the deposition. [...] Read more.
CMT welding sources are garnering attention as alternative heat sources for wire arc additive manufacturing because of their low-heat input. A comprehensive experimental and numerical study on the multi-layer deposition of STS316L was performed to investigate effect of heat accumulation during the deposition. The numerical model which is appropriate for WAMM was developed considering the characteristics of the CMT heat source for the first time. Using a high-speed camera, the transient behavior of the CMT arc was investigated, and applied to the heat source of the numerical model. The model was then used to analyze 10-layered deposits of STS316L, fabricated using CMT-based WAAM. During deposition, the temperature is measured using a pyrometer to analyze the microstructure, after which the cooling rate of each layer is estimated. The measured and simulated SDAS were compared. Based on the comparison, a guideline for the equation regarding the SDAS size and cooling rate was suggested. Full article
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<p>Schematic of the experimental setup.</p> Full article ">Figure 2
<p>Evolution of the arc captured with a high-speed camera during once cycle (17.25 ms from (<b>a</b>–<b>d</b>)) of the CMT-based WAAM process.</p> Full article ">Figure 3
<p>Finite element model for WAAM simulation: (<b>a</b>) shape of the symmetric model and (<b>b</b>) initial temperature of a new layer.</p> Full article ">Figure 4
<p>Macro cross-section photograph for (<b>a</b>) 0 s IPT and (<b>b</b>) 60 s IPT, (<b>c</b>) 10<sup>th</sup> layer (surface) for 60 s IPT, (<b>d</b>) layer boundary of 60 s IPT, (<b>e</b>) the enlarged micrograph of the surface area, and (<b>f</b>) the enlarged micrograph of the boundary.</p> Full article ">Figure 5
<p>SEM micrographs and SEM-EDS analysis (inset) of 1<sup>st</sup> (<b>a</b>) and 2<sup>nd</sup> (<b>b</b>) layers.</p> Full article ">Figure 6
<p>Relationship between the solidification type and the pseudo-binary phase diagram for austenitic stainless steel [<a href="#B21-metals-10-00278" class="html-bibr">21</a>].</p> Full article ">Figure 7
<p>Optical-microscope images of the microstructure with respect to the IPT for STS 316L (<b>a</b>) 9th layer for 0 s IPT; (<b>b</b>) 9th layer for 60 s IPT; (<b>c</b>) 7th layer for 0 s IPT; (<b>d</b>) 7th layer for 60 s IPT; (<b>e</b>) 2nd layer for 0 s IPT; (<b>f</b>) 2nd layer for 60 s IPT; (<b>g</b>) 1st layer for 0 s IPT; (<b>h</b>) 1st layer for 60 s IPT.</p> Full article ">Figure 8
<p>Comparison of the simulated and experimental results for the temperature history of 10- layer deposits at: (<b>a</b>) point ① under 0 s IPT and (<b>b</b>) point ① under 60 s IPT.</p> Full article ">Figure 9
<p>Relationship between the cooling rate and SDAS under IPTs of 0 s and 60 s: (<b>a</b>) Simulated cooling rate, and (<b>b</b>) measured (experiment) and simulated SDAS (simulation).</p> Full article ">
13 pages, 9529 KiB  
Article
Mechanical Properties, Thermal Stability and Microstructures of W-Re-ZrC Alloys Fabricated by Spark Plasma Sintering
by Shu Miao, Zhuoming Xie, Yan Lin, Qianfeng Fang, Jinhong Tan and Yunqiang Zhao
Metals 2020, 10(2), 277; https://doi.org/10.3390/met10020277 - 20 Feb 2020
Cited by 10 | Viewed by 3171
Abstract
Tungsten materials, used as friction stir welding tools, undergo severe plastic deformation and even collapse at high operating temperatures. In order to improve the low-temperature toughness and high-temperature strength, W-10wt.%Re-0.5wt.%ZrC alloys were processed by high-energy ball milling and subsequent spark plasma sintering. Single [...] Read more.
Tungsten materials, used as friction stir welding tools, undergo severe plastic deformation and even collapse at high operating temperatures. In order to improve the low-temperature toughness and high-temperature strength, W-10wt.%Re-0.5wt.%ZrC alloys were processed by high-energy ball milling and subsequent spark plasma sintering. Single solid-solution W-Re powders with typical body-centered cubic structures were achieved when the milling time increases to 50 h. The microhardness, tensile properties, thermal stability and microstructures of this sintered W-10wt.%Re-0.5wt.%ZrC alloys were investigated. Synergetic effects of the solute Re and nanosized dispersion particles induce improvements in low-temperature toughness and high-temperature strength. The alloy suffers ductile fracture at 300 °C, which is about 400 °C and 300 °C lower than that of the spark plasma sintered pure W and W-0.5wt.%ZrC, respectively. Besides, this W-10wt.%Re-0.5wt.%ZrC has a high ultimate tensile strength of 818 MPa and uniform elongation of ~ 8.1% at 300 °C. Moreover, the microstructures and hardness remain stable even after 1500 °C anneal. Based on a detailed microstructure analysis, the mechanisms for the enhanced strength, low-temperature ductility and high thermal stability are proposed and discussed. Grain boundary mobility is impeded by the kinetics constraint through dispersed particles pinning and solute Re atoms dragging, which leads to improved thermal stability. The formation of Zr-C-O particles is most probably attributed to ZrC particles capturing and interacting with impurity oxygen during sintering. Full article
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<p>The temperature profile of the spark plasma sintering (SPS) process.</p> Full article ">Figure 2
<p>(<b>a</b>) X-ray diffraction (XRD) patterns of milled W-10wt.% Re-0.5wt.% ZrC (WRZ) powders as a function of milling time. (<b>b</b>) The (110) diffraction peaks of 5-h milled and 50-h milled WRZ powders.</p> Full article ">Figure 3
<p>SEM images of the 10-h milled powders (<b>a</b>), 30-h milled powders (<b>b</b>), 40-h milled powders (<b>c</b>) and 50-h milled powders (<b>d</b>).</p> Full article ">Figure 4
<p>Optical micrograph images of the spark plasma sinter (SPS)ed WRZ: material A (<b>a</b>), material B (<b>b</b>), material C (<b>c</b>) and material D (<b>d</b>).</p> Full article ">Figure 5
<p>Energy-dispersive X-ray spectroscopy (EDS)-line scan analysis showing the profile of tungsten and rhenium elements on the etched surface of WRZ alloys (material A).</p> Full article ">Figure 6
<p>EDS analysis of the surface of WRZ alloys (material A).</p> Full article ">Figure 7
<p>Tensile behavior (<b>a</b>) of a SPSed WRZ (material D), the ultimate tensile strength (UTS) (<b>b</b>) and uniform elongation (UE) (<b>c</b>) of SPSed W, WRZ, W-0.5wt.%ZrC (WZC), W-0.2wt.%Zr-1.0wt.%Y<sub>2</sub>O<sub>3</sub> (WZY) and W-0.5wt.%ZrC-1wt.%Re (WRZC).</p> Full article ">Figure 8
<p>Optical micrograph images of as-sintered material D samples (<b>a</b>), annealed samples at 1500 °C (<b>b</b>), 1600 °C (<b>c</b>), 1700 °C (<b>d</b>), 1800 °C (<b>e</b>) and the average grain size and hardness (<b>f</b>) of annealed material D samples at various temperatures.</p> Full article ">Figure 9
<p>Transmission electron microscope (TEM) micrographs (<b>a</b>) and (<b>b</b>) of the material D sample and size distributions of intragranular (<b>c</b>) and intergranular (<b>d</b>) particles.</p> Full article ">Figure 10
<p>The TEM image of dislocation remaining of the material D sample after fast sintering (<b>a</b>), selected particles (<b>b</b>) and the high-resolution (HR)TEM and inset inverse fast Fourier transform (IFFT) (selected areas as indicated by arrows) images of corresponding particles in the blue circle (<b>c</b>) and in the yellow circle (<b>d</b>), respectively.</p> Full article ">
17 pages, 4946 KiB  
Article
A Comprehensive Study into the Boltless Connections of Racking Systems
by Rodoljub Vujanac, Nenad Miloradović, Snežana Vulović and Ana Pavlović
Metals 2020, 10(2), 276; https://doi.org/10.3390/met10020276 - 20 Feb 2020
Cited by 8 | Viewed by 7210
Abstract
In practice, structures of pallet racks are characterized by very wide options of beam-to-column connections. The up to date part of the standard Eurocode 3 considers details for the design of connections. However, experimental determination of the joint properties in steel pallet racks [...] Read more.
In practice, structures of pallet racks are characterized by very wide options of beam-to-column connections. The up to date part of the standard Eurocode 3 considers details for the design of connections. However, experimental determination of the joint properties in steel pallet racks is the most reliable process, since it takes into account an inability to develop a general analytical model for the design of these connections. In this paper, a test procedure for the behavior of beam-to-column connections is presented and the results are analyzed according to the procedure defined in the relevant design codes. With aim to avoid expensive experiments to determine structural properties of different types of connections, a polynomial model and a corresponding numerical model were developed to be used for simulating the experiment. After verification, the developed analytical and numerical model can be applied for investigation of various combinations of beam-to-column connections. Full article
(This article belongs to the Special Issue Advances in Design by Metallic Materials: Synthesis, Characterization, Simulation and Applications)
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<p>Parts of the racking system.</p> Full article ">Figure 2
<p>Moment-rotation characteristics.</p> Full article ">Figure 3
<p>Cantilever test set-up.</p> Full article ">Figure 4
<p>Arrangement of the parts of the sample.</p> Full article ">Figure 5
<p>Disposition of the measuring equipment.</p> Full article ">Figure 6
<p>Derivation of moment-rotation relationship.</p> Full article ">Figure 7
<p>Finite element model of elements of cantilever test: (<b>a</b>) Column; (<b>b</b>) Beam; (<b>c</b>) Beam-end connector; (<b>d</b>) Screw; (<b>e</b>) Loading transfer plate; (<b>f</b>) Boundary conditions and the applied load.</p> Full article ">Figure 8
<p>Diagram of the displacement.</p> Full article ">Figure 9
<p>Moment-rotation curves for the S80ML-R140L joint.</p> Full article ">Figure 10
<p>Comparison of the results of the tested joints S80ML-R140L: (<b>a</b>) Deformation of the tested sample (<b>b</b>) Displacement field in the <span class="html-italic">x</span> direction obtained by the numerical model.</p> Full article ">Figure 11
<p>Comparative average moment-rotation curves for the tested samples and numerical model.</p> Full article ">
11 pages, 5108 KiB  
Article
Evolution of Inclusions in Steelmaking Process of Rare Earth Steels Containing Arsenic with Alumina Crucibles
by Hongpo Wang, Peng Yu, Silu Jiang, Bin Bai, Lifeng Sun and Yu Wang
Metals 2020, 10(2), 275; https://doi.org/10.3390/met10020275 - 20 Feb 2020
Cited by 20 | Viewed by 3132
Abstract
In order to determine strategies for removing arsenic from rare earth arsenic-containing steels, the evolution of inclusions in the whole steelmaking process with alumina crucibles was investigated. It has been proven that adding lanthanum has a significant effect on both the existing state [...] Read more.
In order to determine strategies for removing arsenic from rare earth arsenic-containing steels, the evolution of inclusions in the whole steelmaking process with alumina crucibles was investigated. It has been proven that adding lanthanum has a significant effect on both the existing state and content of arsenic in steel. The content of arsenic steeply decreased after adding 0.148% lanthanum by generating La–S–As inclusions. The addition of 0.054% lanthanum did not dramatically affect the content of arsenic. Both 0.148% and 0.054% additions of lanthanum modified the existing Si–Mn–Al–O inclusions, making them first change to La-containing inclusions, and then change back to Si–Mn–Al–O inclusions. During this process, the compositions of inclusions changed from (SiO2–MnO)-rich to Al2O3-rich ones, owing to the reactions between lanthanum and alumina crucibles. The addition of 0.148% lanthanum resulted in a relatively severe reaction with the alumina crucible. This led to the decomposition of a part of the existing La–S–As inclusions and a slight increase in the arsenic content. Therefore, it is noted that choosing an appropriate holding time after adding rare earth elements to molten steel has a significant effect on the arsenic removal and saving the consumption of rare earth elements. Full article
(This article belongs to the Special Issue Inclusion/Precipitate Engineering in Steels)
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<p>Details of the sample preparation process.</p> Full article ">Figure 2
<p>Evolution of inclusions in molten steel. (<b>a</b>–<b>e</b>) With the initial La addition of 0.148%; (<b>f</b>–<b>j</b>) with the initial La addition of 0.054%.</p> Full article ">Figure 3
<p>Changes in concentrations of As, La, and acid-soluble Al in steel (Arsenic additions were made at 10 min, La additions at 30 min, and the points at 60 min represent the samples taken from ingots). (<b>a</b>) the 0.148% La-added steel; (<b>b</b>) the 0.054% La-added steel.</p> Full article ">Figure 4
<p>Morphologies, distribution, and compositions of inclusions on the profile of the ingot (backscattered electron images). (<b>a</b>) Top surface of the ingot; (<b>b</b>) an enlarged image of La<sub>2</sub>O<sub>2</sub>S and LaAlO<sub>3</sub> inclusions; (<b>c</b>) side surface of the ingot; (<b>d</b>) an enlarged image of La–S–As inclusions near the top surface; (<b>e</b>) energy dispersive spectrum (EDS) of point C; (<b>f</b>) bottom of the ingot.</p> Full article ">Figure 5
<p>Morphologies of the LaAlO<sub>3</sub>-containing layer at the side surface of the ingot (secondary electron image).</p> Full article ">Figure 6
<p>Morphologies, distribution, and compositions of inclusions on the profile of the ingot with the initial La addition of 0.054% (backscattered electron images). (<b>a</b>) Top surface of ingot; (<b>b</b>) corner between the top and side surfaces of ingot; (<b>c</b>) energy dispersive spectrum (EDS) of point D; (<b>d</b>) bottom of ingot; (<b>e</b>) side surface of ingot; (<b>f</b>) EDS of point E.</p> Full article ">Figure 7
<p>Gibbs free energy of La-related reactions per lanthanum or aluminum in molten steel.</p> Full article ">Figure 8
<p>Morphologies and element distribution on the profile of the side surface of the ingot. (<b>a</b>) Backscattered electron image of morphologies.</p> Full article ">Figure 9
<p>The existing state of arsenic in the steelmaking process and its distribution in the ingot.</p> Full article ">
14 pages, 3517 KiB  
Article
A New 3D Creep-Fatigue-Elasticity Damage Interaction Diagram Based on the Total Tensile Strain Energy Density Model
by Qiang Wang, Naiqiang Zhang and Xishu Wang
Metals 2020, 10(2), 274; https://doi.org/10.3390/met10020274 - 20 Feb 2020
Cited by 6 | Viewed by 3453
Abstract
Fatigue damage, creep damage, and their interactions are the critical factors in degrading the integrity of most high-temperature engineering structures. A reliable creep-fatigue damage interaction diagram is a crucial issue for the design and assessment of high-temperature components used in power plants. In [...] Read more.
Fatigue damage, creep damage, and their interactions are the critical factors in degrading the integrity of most high-temperature engineering structures. A reliable creep-fatigue damage interaction diagram is a crucial issue for the design and assessment of high-temperature components used in power plants. In this paper, a new three-dimensional creep-fatigue-elasticity damage interaction diagram was constructed based on a developed life prediction model for both high-temperature fatigue and creep fatigue. The total tensile strain energy density concept is adopted as a damage parameter for life prediction by using the elastic strain energy density and mean stress concepts. The model was validated by a great deal of data such as P91 steel at 550 °C, Haynes 230 at 850 °C, Alloy 617 at 850 and 950 °C, and Inconel 625 at 815 °C. The estimation values have very high accuracy since nearly all the test data fell into the scatter band of 2.0. Full article
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<p>Schematic illustration of total tensile strain energy density in different loading waveforms. (<b>a</b>): tension-hold creep-fatigue loading; (<b>b</b>): compression-hold creep-fatigue loading.</p> Full article ">Figure 2
<p>Damage interaction diagram of P91 steel in time fraction method.</p> Full article ">Figure 3
<p>The relation between elastic modulus and temperature of Inconel 625.</p> Full article ">Figure 4
<p>The cyclic stress strain relation of Inconel 625 at 815 °C.</p> Full article ">Figure 5
<p>The relation between fatigue life and total tensile strain energy density of Inconel 625.</p> Full article ">Figure 6
<p>Model predictions by (<b>a</b>) present model and (<b>b</b>) traditional energy-based model.</p> Full article ">Figure 6 Cont.
<p>Model predictions by (<b>a</b>) present model and (<b>b</b>) traditional energy-based model.</p> Full article ">Figure 7
<p>Predicted life/tested life versus (<b>a</b>) inelastic strain and (<b>b</b>) tested life.</p> Full article ">Figure 8
<p>The relation between creep failure density and applied stress.</p> Full article ">Figure 9
<p>Damage interaction diagrams (<b>a</b>) Alloy 617 at 850 °C and 950 °C, Inconel 625 at 815 °C and (<b>b</b>) P91 steel at 550 °C.</p> Full article ">Figure 10
<p>Total damage vs. creep-fatigue life.</p> Full article ">Figure 11
<p>Creep damage per cycle versus creep-fatigue life.</p> Full article ">
12 pages, 3359 KiB  
Article
In Situ Measurements for Plastic Zone Ahead of Crack Tip and Continuous Strain Variation under Cyclic Loading Using Digital Image Correlation Method
by Yan Zhao, Dianyin Hu, Meng Zhang, Wei Dai and Weifang Zhang
Metals 2020, 10(2), 273; https://doi.org/10.3390/met10020273 - 19 Feb 2020
Cited by 18 | Viewed by 3115
Abstract
Fatigue crack is one of the most common damage forms for aeronautical aluminum alloy. With crack propagation, the strain fields of the whole object surface and plastic zone (PZ) ahead of the crack tip are changing continuously. For most metallic materials, the behavior [...] Read more.
Fatigue crack is one of the most common damage forms for aeronautical aluminum alloy. With crack propagation, the strain fields of the whole object surface and plastic zone (PZ) ahead of the crack tip are changing continuously. For most metallic materials, the behavior of PZ around the crack tip and continuous strain variation play a vital role in crack propagation. In this work, the “continuous” strain information at and in front of the crack tip on the specimen surface was obtained quantitatively and the PZ size ahead of crack tip was in situ measured quantitatively with crack propagation by using the digital image correlation (DIC) method, which overcomes the difficulty for the in situ measurement of mechanical variables. Moreover, the method of specimen preparation was simplified by using a white matt paint with strong adhesion, but also resulted in a higher resolution being shown, even for such a large area. Furthermore, the experimental results of the PZ size from the proposed method had good agreement with the theoretical values, which overcomes the limitation that the conventional approaches only consider the quasi-static crack. Finally, the continuous strain variation behavior was analyzed from the experimental results in detail with the consideration of crack propagation. Full article
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<p>Principle of the digital image correlation (DIC) method used for in-plane strain mapping.</p> Full article ">Figure 2
<p>The schematic diagram of the geometric dimensions of the aluminum specimen.</p> Full article ">Figure 3
<p>Experimental setup of strain measurement using the DIC technology. (<b>a</b>) A photo of the test setup, (<b>b</b>) DIC processing system, (<b>c</b>) the digital image recorded by the Charge Coupled Device (CCD) camera, and (<b>d</b>)the region of interest (ROI) corresponding to <a href="#metals-10-00273-f003" class="html-fig">Figure 3</a>c.</p> Full article ">Figure 4
<p>(<b>a</b>) Crack propagation path and speckle image; (<b>b</b>) manufactured specimen; (<b>c</b>) the crack growth curve.</p> Full article ">Figure 5
<p>(<b>a</b>) Schematic illustration of the plastic zone ahead of the crack tip; (<b>b</b>) strain distribution along the crack plane under different loadings.</p> Full article ">Figure 6
<p>The measurement results of the PZ size ahead of the crack tip with crack propagation based on the DIC method.</p> Full article ">Figure 7
<p>The measurement results of the continuous strain curves at and in front of the crack tip were measured by the DIC method, and the (<b>a</b>–<b>h</b>) show the continue strain variation with crack propagation.</p> Full article ">
17 pages, 9119 KiB  
Article
Additive Manufacturing with Superduplex Stainless Steel Wire by CMT Process
by Malin Lervåg, Camilla Sørensen, Andreas Robertstad, Bård M. Brønstad, Bård Nyhus, Magnus Eriksson, Ragnhild Aune, Xiaobo Ren, Odd M. Akselsen and Ivan Bunaziv
Metals 2020, 10(2), 272; https://doi.org/10.3390/met10020272 - 19 Feb 2020
Cited by 56 | Viewed by 6065
Abstract
For many years, the oil and gas industry has utilized superduplex stainless steels due to their high strength and excellent corrosion resistance. Wire arc additive manufacturing (WAAM) was used with superduplex filler wire to create walls with different heat input. Due to the [...] Read more.
For many years, the oil and gas industry has utilized superduplex stainless steels due to their high strength and excellent corrosion resistance. Wire arc additive manufacturing (WAAM) was used with superduplex filler wire to create walls with different heat input. Due to the multiple heating and cooling cycles during layer deposition, brittle secondary phases may form such as intermetallic sigma (σ) phase. By inspecting deposited walls within wide range of heat inputs (0.40–0.87 kJ/mm), no intermetallic phases formed due to low inter-pass temperatures used, together with the high Ni content in the applied wire. Lower mechanical properties were observed with high heat inputs due to low ferrite volume fraction, precipitation of Cr nitrides and formation of secondary austenite. The walls showed good toughness values based on both Charpy V-notch and CTOD (crack tip opening displacement) testing. Full article
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<p>Schematic illustration of additive manufacturing of component.</p> Full article ">Figure 2
<p>Mechanical testing of additively manufactured walls. Horizontal tensile samples to the left, vertical tensile specimens in the middle and Charpy V and CTOD (crack tip opening displacement) samples to the right. The latter two specimen types have different notches.</p> Full article ">Figure 3
<p>Macrosections of additively manufactured walls: (<b>a</b>) 0.40 kJ/mm, (<b>b</b>) 0.54 kJ/mm and (<b>c</b>) 0.87 kJ/mm. The dashed red lines indicate boundaries between deposited layers.</p> Full article ">Figure 4
<p>Stress strain curves of samples with heat inputs of (<b>a</b>) 0.40 kJ/mm, (<b>b</b>) 0.54 kJ/mm and (<b>c</b>) 0.87 kJ/mm. <span class="html-italic">H1</span>/<span class="html-italic">H2</span> are horizontal and <span class="html-italic">V1</span>/<span class="html-italic">V2</span> are vertical samples.</p> Full article ">Figure 5
<p>Effect of heat input on yield (<span class="html-italic">R<sub>p0.2</sub></span>) and tensile (<span class="html-italic">R<sub>m</sub></span>) strength for (<b>a</b>) vertical samples and (<b>b</b>) horizontal samples.</p> Full article ">Figure 6
<p>Effect of heat input on reduction in area.</p> Full article ">Figure 7
<p>Effect of heat input on hardness in deposited layers (WM) and heat affected zone (HAZ).</p> Full article ">Figure 8
<p>Charpy V-notch toughness of walls and base metal.</p> Full article ">Figure 9
<p>CTOD fracture toughness of walls.</p> Full article ">Figure 10
<p>Nomenclature of solidification modes in stainless steels. Adapted from [<a href="#B46-metals-10-00272" class="html-bibr">46</a>]. The wire composition is represented by the red circle marker based on the Cr and Ni equivalents adopted by WRC-1992 (Equation (2)).</p> Full article ">Figure 11
<p>Micrographs with indications of epitaxial solidification (one example is shown inside black dotted lines); between first and second build layer (0.87 kJ/mm).</p> Full article ">Figure 12
<p>Effect of heat input on volume fraction of ferrite in different areas of deposited walls: (<b>a</b>) WM of top layer, (<b>b</b>) WM between 5th and 6th layer, (<b>c</b>) WM of 5th layer. The dashed red lines represent volume fraction of ferrite in the BM (48 vol.%).</p> Full article ">Figure 13
<p>Microstructure in top layer in; (<b>a</b>) Sample No. 1 (0.40 kJ/mm), (<b>b</b>) Sample No. 2 (0.54 kJ/mm) and (<b>c</b>) Sample No. 3 (0.87 kJ/mm). WAS—Widmanstätten austenite side-plate.</p> Full article ">Figure 14
<p>Microstructure in heat-affected zone. WAS—Widmanstätten austenite side-plate. The dashed black line represents the boundary between the HAZ and the unaffected BM.</p> Full article ">Figure 15
<p>Cr nitrides in ferrite in HAZ. The dashed black line represents the boundary between the HAZ and the unaffected BM.</p> Full article ">Figure 16
<p>Secondary austenite formed in the ferrite, (<b>a</b>) overview with dark islands of <span class="html-italic">γ</span><span class="html-italic"><sub>2</sub></span>, and (<b>b</b>) close-up with <span class="html-italic">γ</span><span class="html-italic"><sub>2</sub></span> inside the dotted red ellipse. Arrows point at different nucleation sites.</p> Full article ">
21 pages, 13116 KiB  
Article
Sheet Metal Forming Optimization Methodology for Servo Press Process Control Improvement
by Antonio Del Prete and Teresa Primo
Metals 2020, 10(2), 271; https://doi.org/10.3390/met10020271 - 19 Feb 2020
Cited by 12 | Viewed by 6260
Abstract
In sheet metal forming manufacturing operations the use of servo presses is gaining more interest due to the opportunity to improve process performance (quality, productivity, cost reduction, etc.). It is not yet clear how to proceed in the engineering process when this type [...] Read more.
In sheet metal forming manufacturing operations the use of servo presses is gaining more interest due to the opportunity to improve process performance (quality, productivity, cost reduction, etc.). It is not yet clear how to proceed in the engineering process when this type of operating machine is used to achieve the maximum possible potential of this technology. Recently, several press builders have developed gap- and straight-sided metal forming presses adopting the mechanical servo-drive technology. The mechanical servo-drive press offers the flexibility of a hydraulic press with the speed, accuracy and reliability of a mechanical press. Servo drive presses give the opportunity to improve the productivity of process conditions and improve the quality of stamped parts. Forming simulation and numerical optimization can be useful tools to define beforehand the optimal process parameter set-up in terms of servo press downward curve properties. This is done by carrying out a sensitivity analysis of the forming parameters having influence on said curve. The authors have developed a numerical methodology able to analyze the influence factors, for comparison with the degrees of freedom made available by the usage of a servo press, in terms of stroke profile management, to obtain an optimized process parameters combination. Full article
(This article belongs to the Special Issue Analysis and Design of Metal-Forming Processes)
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<p>Finite element (FE) model created for simulation of the sheet forming process.</p> Full article ">Figure 2
<p>Blank size.</p> Full article ">Figure 3
<p>Forming limit diagram (FLD) for RUN0.</p> Full article ">Figure 4
<p>% Thinning distribution for RUN0.</p> Full article ">Figure 5
<p>Comparison between: (<b>a</b>) the real product obtained by forming using a traditional mechanical press and (<b>b</b>) the numerical model.</p> Full article ">Figure 6
<p>Definition of output variables: X1, X2, Y1 and Y2 and related detail of X1.</p> Full article ">Figure 7
<p>Input timing definition on stepwise curve.</p> Full article ">Figure 8
<p>Input stroke definition on stepwise curve.</p> Full article ">Figure 9
<p>% Thinning distribution and X1, X2, Y1 and Y2 values for the RUN22; blank0.</p> Full article ">Figure 10
<p>% Thinning distribution and X1, X2, Y1 and Y2 values for the RUN20; blank0.</p> Full article ">Figure 11
<p>% Thinning distribution and X1, X2, Y1 and Y2 values for the RUN1; blank1.</p> Full article ">Figure 12
<p>% Thinning distribution and X1, X2, Y1 and Y2 values for the RUN11; blank1.</p> Full article ">Figure 13
<p>Reaction force curves comparison between RUN20 and RUN22; blank0.</p> Full article ">Figure 14
<p>Punch displacement with stepwise, in correspondence of S and T1. TT represents the last point of the curve with maximum value in the x axis; blank0.</p> Full article ">Figure 15
<p>Reaction force curves comparison between RUN1 and RUN11; blank1.</p> Full article ">Figure 16
<p>Punch displacement with stepwise, in correspondence of S and T1. TT represents the last point of the curve with maximum value in x axis; blank1.</p> Full article ">Figure 17
<p>MEP for max thick % output.</p> Full article ">Figure 18
<p>MEP for min thick % output.</p> Full article ">Figure 19
<p>MEP for reaction force output.</p> Full article ">Figure 20
<p>MEP for X1 output.</p> Full article ">Figure 21
<p>MEP for X2 output.</p> Full article ">Figure 22
<p>MEP for Y1 output.</p> Full article ">Figure 23
<p>MEP for Y2 output.</p> Full article ">Figure 24
<p>Linear correlation matrix for the blank0 plan.</p> Full article ">Figure 25
<p>Design explored with multi-island genetic algorithm (MIGA), blank0 plan.</p> Full article ">Figure 26
<p>Linear correlation matrix, blank 0.</p> Full article ">Figure 27
<p>Design explored with MIGA for the blank1 plan.</p> Full article ">Figure 28
<p>Displacement vs time comparison for the optimal combination (according to the regression model) for blank0 and blank1 plan.</p> Full article ">Figure 29
<p>Die reaction force comparison for both simulation plans.</p> Full article ">Figure 30
<p>%Thinning distribution and X1, X2, Y1 and Y2 for optimal combination for blank0.</p> Full article ">Figure 31
<p>%Thinning distribution and X1, X2, Y1 and Y2 for optimal combination for blank1.</p> Full article ">
27 pages, 12282 KiB  
Review
Cavitation Peening: A Review
by Hitoshi Soyama
Metals 2020, 10(2), 270; https://doi.org/10.3390/met10020270 - 19 Feb 2020
Cited by 76 | Viewed by 8128
Abstract
The most popular surface modification technology used to enhance the mechanical properties of metallic materials is shot peening. Shot peening improves fatigue life and strength by introducing local plastic deformation pits. However, the pits increase surface roughness, which is a disadvantage for fatigue [...] Read more.
The most popular surface modification technology used to enhance the mechanical properties of metallic materials is shot peening. Shot peening improves fatigue life and strength by introducing local plastic deformation pits. However, the pits increase surface roughness, which is a disadvantage for fatigue properties. Recently, cavitation peening, in which cavitation bubble collapse impacts are used, has been developed as an advanced surface modification technology. The advantage of cavitation peening is the lesser increase in surface roughness compared with shot peening, as no solid collisions occur in cavitation peening. In conventional cavitation peening, cavitation is generated by injecting a high-speed water jet into water. However, cavitation peening is different from water jet peening, in which water column impacts are used. In the present review, to avoid confusing cavitation peening and water jet peening, fundamentals and mechanisms of cavitation peening are described in comparison to water jet peening, and the effects and applications of cavitation peening are reviewed compared with the other peening methods. Full article
(This article belongs to the Special Issue Advanced Surface Modification Technologies)
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<p>Development of relative aggressive intensity of cavitating jet, changing over time.</p> Full article ">Figure 2
<p>Aspect of cavitation through Venturi tube [<a href="#B47-metals-10-00270" class="html-bibr">47</a>].</p> Full article ">Figure 3
<p>Aspect of vortex cavitation arising downstream of a butterfly valve [<a href="#B48-metals-10-00270" class="html-bibr">48</a>].</p> Full article ">Figure 4
<p>Visualization of impact at cavitation collapse by using photo elasticity [<a href="#B54-metals-10-00270" class="html-bibr">54</a>].</p> Full article ">Figure 5
<p>Aspect of cavitation through a nozzle [<a href="#B54-metals-10-00270" class="html-bibr">54</a>]: (<b>a</b>) Venturi-type nozzle and (<b>b</b>) step-type nozzle.</p> Full article ">Figure 6
<p>Aspects of submerged high-speed water jets, i.e., cavitating jet. (<b>a</b>) Schematic diagram; (<b>b</b>) Observation with flush lamp; (<b>c</b>) Observation with normal light.</p> Full article ">Figure 7
<p>Aspect of impinging cavitating jet observed by high-speed video camera.</p> Full article ">Figure 8
<p>Typical treatment area by a fixed cavitating jet (pure aluminum, nozzle diameter <span class="html-italic">d</span> = 2 mm, upstream pressure of nozzle <span class="html-italic">p</span><sub>1</sub> = 30 MPa, downstream pressure of nozzle <span class="html-italic">p</span><sub>2</sub> = 0.1 MPa, standoff distance <span class="html-italic">s</span> = 262 mm, exposure time <span class="html-italic">t</span> = 1 min) [<a href="#B3-metals-10-00270" class="html-bibr">3</a>].</p> Full article ">Figure 9
<p>Aspect of cavitating jet in water and in air [<a href="#B54-metals-10-00270" class="html-bibr">54</a>]: (<b>a</b>) cavitating jet in water and (<b>b</b>) cavitating jet in air.</p> Full article ">Figure 10
<p>Aspect of cavitation induced by vibratory apparatus [<a href="#B70-metals-10-00270" class="html-bibr">70</a>]. Reproduced with permission from The Japan Society of Mechanical Engineers.</p> Full article ">Figure 11
<p>Laser ablation (LA) and laser cavitation (LC) produced by a pulse laser. (<b>a</b>) Aspect of laser ablation (LA) and laser cavitation (LC); (<b>b</b>) Signal from polyvinylidene fluoride (PVDF) sensor; (<b>c</b>) Signal from submerged shockwave sensor.</p> Full article ">Figure 12
<p>Compressive residual stress introduced by several types of cavitating jets (stainless steel) [<a href="#B3-metals-10-00270" class="html-bibr">3</a>].</p> Full article ">Figure 13
<p>Peening intensity as a function of standoff distance [<a href="#B97-metals-10-00270" class="html-bibr">97</a>]: (<b>a</b>) 40 and (<b>b</b>) 60 MPa. Reproduced with permission from The Water Jet Technology Society of Japan.</p> Full article ">Figure 14
<p>Schematic diagram of aggressive intensity of the cavitating jet using several types of nozzles.</p> Full article ">Figure 15
<p>Effect of nozzle diameter on introduction of compressive residual stress compared with non-peened (NP) (stainless steel) [<a href="#B100-metals-10-00270" class="html-bibr">100</a>].</p> Full article ">Figure 16
<p>Effect of injection pressure on introduction of compressive residual stress (stainless steel) [<a href="#B100-metals-10-00270" class="html-bibr">100</a>].</p> Full article ">Figure 17
<p>Normalized aggressive intensity of a cavitating jet as a function of cavitation number [<a href="#B3-metals-10-00270" class="html-bibr">3</a>].</p> Full article ">Figure 18
<p>Classification map for cavitation peening and water jet peening. [<a href="#B95-metals-10-00270" class="html-bibr">95</a>]</p> Full article ">Figure 19
<p>Processing capability as a function of injection pressure [<a href="#B97-metals-10-00270" class="html-bibr">97</a>]. Reproduced with permission from The Water Jet Technology Society of Japan; (<b>a</b>) Water jet peening; (<b>b</b>) Cavitation peening.</p> Full article ">Figure 20
<p>Improvement in fatigue strength of stainless steel by cavitation peening (CP), shot peening (SP), laser peening (LP), and water jet peening (WJP) compared with non-peened (NP).</p> Full article ">Figure 21
<p>Aspect of cavitating jet injecting to the gear.</p> Full article ">Figure 22
<p>Typical aspect of plastic deformation pit induced by cavitation impact in comparison with ball indentation. (<b>a</b>) Cavitation impact; (<b>b</b>) Ball indentation.</p> Full article ">Figure A1
<p>Schematic diagram of radius of curvature and arc height.</p> Full article ">
16 pages, 6195 KiB  
Article
Influence of Submerged Entry Nozzle Port Blockage on the Meniscus Fluctuation Considering Various Operational Parameters
by Manish Kumar, Praveen Mishra and Apurba Kumar Roy
Metals 2020, 10(2), 269; https://doi.org/10.3390/met10020269 - 18 Feb 2020
Cited by 3 | Viewed by 2924
Abstract
The continuous casting process (CCP) is the most vital part of steelmaking. The flow pattern near the submerged entry nozzle (SEN) and mould greatly influence the quality of the slab produced. The present investigation was carried out to gain knowledge regarding the meniscus [...] Read more.
The continuous casting process (CCP) is the most vital part of steelmaking. The flow pattern near the submerged entry nozzle (SEN) and mould greatly influence the quality of the slab produced. The present investigation was carried out to gain knowledge regarding the meniscus fluctuation under different nozzle port blockage conditions by water model experiments. The experiments were carried out to study the effect of no blockage, 25% blockage, 50% blockage, and 75% blockage of the nozzle port on mould-level fluctuations. The result shows that when the liquid flow rate increases, the wave amplitude increases. In these experiments, the average and maximum meniscus fluctuations were measured while changing different variables such as the water flow rate, gas flow rate, and one-side percentage blockage of the SEN port while the other side was fully open. The observation shows that when the port size decreases, the fluid steel mixed from the obstructing side to the open side results in asymmetry. The average and maximum wave amplitude increases with decreasing submergence depth. It was observed that the maximum height of the standing waves in the mould continued rising on the non-blocked side of the SEN. Blockage increases from 25% to 75%, and with 75% blockage of the right side of the SEN port, the mould-level fluctuation at the left side of the mould was extreme, while that of the right side was relatively quiet. Full article
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<p>(<b>a</b>) Pictorial view of the experimental setup. (<b>b</b>) Experimental setup at the Department of Mechanical Engineering, Birla Institute of Technology, Mesra (Ranchi) India.</p> Full article ">Figure 1 Cont.
<p>(<b>a</b>) Pictorial view of the experimental setup. (<b>b</b>) Experimental setup at the Department of Mechanical Engineering, Birla Institute of Technology, Mesra (Ranchi) India.</p> Full article ">Figure 2
<p>(<b>a</b>) Geometrical mould and SEN dimension of the experimentation setup. (<b>b</b>) Diagrammatic representation of different port blockages.</p> Full article ">Figure 3
<p>Images extracted from MATLAB coding.</p> Full article ">Figure 4
<p>Comparison of the effect of blockage of SEN (one side only) on the maximum surface amplitude at 40 and 60 L/min water flow rates at a submergence depth of 125 mm.</p> Full article ">Figure 5
<p>Comparison of the effect of blockage of SEN (one side only) on the average surface amplitude at 40 and 60 L/min water flow rates at a submergence depth of 125 mm.</p> Full article ">Figure 6
<p>Comparison of effect of blockage of SEN on the maximum surface wave amplitude at 40 and 60 L/min water flow rates at a submergence depth of 175 mm.</p> Full article ">Figure 7
<p>Comparison of effect of blockage of SEN on the average surface wave amplitude at 40 and 60 L/min at a submergence depth of 175 mm.</p> Full article ">Figure 8
<p>Average surface wave profile in the mould for two different water flow rates from the SEN with both ports fully open (no blockage condition).</p> Full article ">Figure 9
<p>Average surface wave profiles in the mould for two different water flow rates from the SEN with 25% blockage of one side of the port.</p> Full article ">Figure 10
<p>Average surface wave profiles in the mould for two different water flow rates from the SEN with 50% blockage of one side of the port.</p> Full article ">Figure 11
<p>Average surface wave profiles in the mould for two different water flow rates from the SEN with 75% blockage one side of the port.</p> Full article ">Figure 12
<p>Average wave profile in the mould at 40 L/min from the SEN at no blockage, 25% blockage, 50% blockage, and 75% blockage of the port, respectively.</p> Full article ">Figure 13
<p>Average wave profile in the mould at 60 L/min from the SEN at no blockage, 25% blockage, 50% blockage, and 75% blockage of the port, respectively.</p> Full article ">Figure 14
<p>Average meniscus fluctuations at air injection with an immersion depth of 125 mm.</p> Full article ">Figure 15
<p>Maximum meniscus fluctuations at air injection with an immersion depth of 125 mm.</p> Full article ">Figure 16
<p>Effect of submergence depth on the maximum wave amplitude with water flow rates of 40 L/min and 60 L/min, respectively.</p> Full article ">
14 pages, 13000 KiB  
Article
Study on the Mechanical and Tribological Properties and the Mechanisms of Cr-Free Ni-Based Self-Lubricating Composites at a Wide Temperature Range
by Penglin Zhang, Gaopan Zhao, Wenzhen Wang, Bin Wang, Peiying Shi, Gang Qi and Gewen Yi
Metals 2020, 10(2), 268; https://doi.org/10.3390/met10020268 - 18 Feb 2020
Cited by 8 | Viewed by 2459
Abstract
A Cr-free Ni-based self-lubricating composites with MoS2 and Ag as lubricants were fabricated by the powder metallurgy method. The microstructures were examined. The mechanical properties and tribological behaviors of the composites were evaluated from room temperature to 800 °C. The fractography was [...] Read more.
A Cr-free Ni-based self-lubricating composites with MoS2 and Ag as lubricants were fabricated by the powder metallurgy method. The microstructures were examined. The mechanical properties and tribological behaviors of the composites were evaluated from room temperature to 800 °C. The fractography was observed and the fracture mechanisms were analyzed. The morphologies and the phase compositions of worn surfaces were determined and the wear mechanisms were elaborated. The results indicate that MoS2 did not completely decompose after sintering, and the NiMoAl-MoS2-Ag composite has the best tribological properties (0.22, 1.68 × 10−5) at 800 °C. The main wear mechanisms are micro-ploughing and plastic deformation. The improvement of tribological properties was attributed to the formation of the lubricating film consisting of NiO, Mo oxides, various molybdates, and Ag particles. The reactions resulting in these compositions are proposed. The mechanical properties degrade with the rise of temperature and the addition of lubricants. Both NiMoAl and NiMoAlAg alloys exhibit micro-void accumulation fracture while the composites with MoS2 reveal intergranular fracture. Full article
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<p>XRD patterns of the (<b>a</b>) mixed powders and (<b>b</b>) sintered composites.</p> Full article ">Figure 2
<p>BEI micrographs of composites: (<b>a</b>) N, (<b>b</b>) NM, (<b>c</b>) NA, and (<b>d</b>) NMA.</p> Full article ">Figure 3
<p>(<b>a</b>) Friction coefficients and (<b>b</b>) wear rates of Ni-based composites at different temperatures.</p> Full article ">Figure 4
<p>(<b>a</b>) Tensile strength and (<b>b</b>) compressive strength of Ni-based composites at different temperatures.</p> Full article ">Figure 5
<p>Worn surface SEM morphologies of (<b>a</b>) N, (<b>b</b>) NM, (<b>c</b>) NA, and (<b>d</b>) NMA composites tested at (<b>1</b>) RT, (<b>2</b>) 400 °C, and (<b>3</b>) 800 °C.</p> Full article ">Figure 6
<p>XRD patterns of the worn surface of a Ni-based composites test at (<b>a</b>) RT, (<b>b</b>) 400 °C, and (<b>c</b>) 800 °C.</p> Full article ">Figure 7
<p>Raman spectra inside and outside the wear scar at 800 °C: (<b>a</b>) N, (<b>b</b>) NM, (<b>c</b>) NA, and (<b>d</b>) NMA composites.</p> Full article ">Figure 8
<p>SEM of the Al<sub>2</sub>O<sub>3</sub> counter ball, mating with N self-lubricating composite: (<b>a</b>) RT, (<b>b</b>) 400 °C, and (<b>c</b>) 800 °C.</p> Full article ">Figure 9
<p>SEM of the Al<sub>2</sub>O<sub>3</sub> counter ball, mating with NMA self-lubricating composite: (<b>a</b>) RT, (<b>b</b>) 400 °C, and (<b>c</b>) 800 °C.</p> Full article ">Figure 10
<p>(<b>a</b>) EDS and (<b>b</b>) Raman spectra of the Al<sub>2</sub>O<sub>3</sub> counter ball, mating with NMA self-lubricating a composite at 800 °C.</p> Full article ">Figure 11
<p>Fractography of (<b>a</b>) N, (<b>b</b>) NA, and (<b>c</b>) NMA composites after tensile test at (<b>1</b>) RT, (<b>2</b>) 400 °C, and (<b>3</b>) 800 °C.</p> Full article ">
16 pages, 3367 KiB  
Article
Kinetic Models for the in Situ Reaction between Cu-Ti Melt and Graphite
by Lei Guo, Xiaochun Wen and Zhancheng Guo
Metals 2020, 10(2), 267; https://doi.org/10.3390/met10020267 - 18 Feb 2020
Viewed by 2118
Abstract
The in situ reaction method for preparing metal matrix composites has the advantages of a simple process, good combination of the reinforcing phase and matrix, etc. Based on the mechanism of forming TiCx particles via the dissolution reaction of solid carbon (C) [...] Read more.
The in situ reaction method for preparing metal matrix composites has the advantages of a simple process, good combination of the reinforcing phase and matrix, etc. Based on the mechanism of forming TiCx particles via the dissolution reaction of solid carbon (C) particles in Cu-Ti melt, the kinetic models for C particle dissolution reaction were established. The kinetic models of the dissolution reaction of spherical, cylindrical, and flat C source particles in Cu-Ti melt were deduced, and the expressions of the time for the complete reaction of C source particles of different sizes were obtained. The mathematical relationship between the degree of reaction of C source and the reaction time was deduced by introducing the shape factor. By immersing a cylindrical C rod in a Cu-Ti melt and placing it in a super-gravity field for the dissolution reaction, it was found that the super-gravity field could cause the precipitated TiCx particles to aggregate toward the upper part of the sample under the action of buoyancy. Therefore, the consuming rate of the C rod was significantly accelerated. Based on the flat C source reaction kinetic model, the relationship between the floating speed of TiCx particles in the Cu-Ti melt and the centrifugal velocity (or the coefficient of super-gravity G) was derived. It was proven that, when the centrifugal velocity exceeded a critical value, the super-gravity field could completely avoid the accumulation behavior of TiCx particles on the surface of the C source, thereby speeding up the formation reaction of TiCx. The goal of this study is to better understand and evaluate the generating process of TiCx particles, thus finding possible methods to increase the reaction efficiency Full article
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>The schematic diagram of the reaction mechanism of spherical solid C source particles in Cu-Ti melt.</p> Full article ">Figure 2
<p>The schematic diagram of reaction mechanism of cylindrical solid C source particles in Cu-Ti melt.</p> Full article ">Figure 3
<p>The schematic diagram of reaction mechanism of flat solid C source particles in Cu-Ti melt.</p> Full article ">Figure 4
<p>Sketch map of the centrifugal apparatus: 1 counterweight, 2 centrifugal axis, 3 conductive slipping, 4 thermocouple, 5 insulating layer, 6 temperature controller, 7 TiC<sub>x</sub> particles, 8 Cu-Ti melt, 9 graphite crucible.</p> Full article ">Figure 5
<p>Comparison of the size change in solid C rods dissolved by Cu-Ti melt under normal gravity and super-gravity conditions (<b>a</b>) image of the designed graphite crucible; (<b>b</b>) the image of the sample after heating treatment; (<b>c</b>) the cross-section morphology of the sample after super-gravity treatment with <span class="html-italic">G</span> = 1000; (<b>d</b>) the cross-section morphology of the sample after heating in normal gravity.</p> Full article ">Figure 6
<p>The correlation between the TiC<sub>x</sub> particle floating speed in the Cu-Ti melt and the centrifugal velocity (super gravity coefficient <span class="html-italic">G</span>).</p> Full article ">
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