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Editor’s Choice Articles

Editor’s Choice articles are based on recommendations by the scientific editors of MDPI journals from around the world. Editors select a small number of articles recently published in the journal that they believe will be particularly interesting to readers, or important in the respective research area. The aim is to provide a snapshot of some of the most exciting work published in the various research areas of the journal.

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13 pages, 9623 KiB  
Article
Effect of Tropical Marine Atmospheric Environment on Corrosion Behaviour of the 7B04-T74 Aluminium Alloy
by Ning Li, Weifang Zhang, Xiaojun Yan, Meng Zhang, Lu Han and Yikun Cai
Metals 2023, 13(5), 995; https://doi.org/10.3390/met13050995 - 21 May 2023
Cited by 4 | Viewed by 2238
Abstract
In this work, the effects of the tropical marine atmospheric environment on the corrosion behaviour of the 7B04-T74 aluminium alloy were systematically investigated by using accelerated testing, together with corrosion kinetic analysis, microstructure observation, product composition analysis, and potentiodynamic polarization curve tests. The [...] Read more.
In this work, the effects of the tropical marine atmospheric environment on the corrosion behaviour of the 7B04-T74 aluminium alloy were systematically investigated by using accelerated testing, together with corrosion kinetic analysis, microstructure observation, product composition analysis, and potentiodynamic polarization curve tests. The weight loss method was used for the corrosion kinetics analysis. The surface morphology and corrosion products transformation law were investigated by OM, SEM, EDS, and XPS. The electrochemical characteristics were studied using potentiodynamic polarization curves. The research indicated that the 7B04-T74 aluminium alloy has eminent corrosion resistance in the tropical marine atmospheric environment. Localized pitting corrosion occurred rapidly in the tropical marine atmosphere. In the later stage of corrosion, the corrosion of aluminium alloy did not become serious. Specifically, no obvious intergranular corrosion was found, which is related to the thermal treatment method. Corrosion products included Al(OH)3, Al2O3, and AlCl3, of which Al(OH)3 is the most notable. Full article
(This article belongs to the Special Issue Corrosion Prediction in Different Environment)
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Figure 1

Figure 1
<p>Schematic diagram of potentiodynamic polarization curve test.</p> Full article ">Figure 2
<p>Calculated corrosion depth (<b>a</b>) and corrosion rate (<b>b</b>) of the 7B04-T74 aluminium alloy.</p> Full article ">Figure 3
<p>The surface morphology of uncorroded (<b>a</b>), after 48 h (<b>b</b>) and 576 h (<b>c</b>) of the 7B04-T74 aluminium alloy.</p> Full article ">Figure 4
<p>The metallographic morphology of the uncorroded 7B04-T74 aluminium alloy.</p> Full article ">Figure 5
<p>The surface morphology of the 7B04-T74 aluminium alloy corroded for 48 h (<b>a</b>) and 576 h (<b>c</b>), where (<b>b</b>) is the enlarged image of (<b>a</b>,<b>d</b>) is the enlarged image of (<b>c</b>).</p> Full article ">Figure 6
<p>The EDS point scanning of the 7B04-T74 aluminium alloy corroded for 48 h.</p> Full article ">Figure 7
<p>The EDS surface scanning results of the 7B04-T74 aluminium alloy after 576 h of corrosion.</p> Full article ">Figure 8
<p>Total XPS spectrogram of the corrosion products of the 7B04-T74 aluminium alloy after 48 h of corrosion.</p> Full article ">Figure 9
<p>XPS spectrogram for (<b>a</b>) Al2p, (<b>b</b>) O1s, and (<b>c</b>) Cl2p of the corrosion products after 48 h.</p> Full article ">Figure 10
<p>Relative content of the corrosion products of the 7B04-T74 aluminium alloy under different corrosion times.</p> Full article ">Figure 11
<p>Potentiodynamic polarization curves after different corrosion times.</p> Full article ">Figure 12
<p>The corrosion mechanism of the 7B04-T74 aluminium alloy in the tropical marine atmosphere.</p> Full article ">
11 pages, 1840 KiB  
Article
Corrosion of Copper in a Tropical Marine Atmosphere Rich in H2S Resulting from the Decomposition of Sargassum Algae
by Mahado Said Ahmed, Mounim Lebrini, Benoit Lescop, Julien Pellé, Stéphane Rioual, Olivia Amintas, Carole Boullanger and Christophe Roos
Metals 2023, 13(5), 982; https://doi.org/10.3390/met13050982 - 19 May 2023
Cited by 6 | Viewed by 2214
Abstract
The atmospheric corrosion of copper exposed in Martinique (Caribbean Sea) for 1 year was reported. This island suffered the stranding of sargassum algae, which decompose and release toxic gases such as hydrogen sulfide (H2S) or ammonia (NH3). Four sites [...] Read more.
The atmospheric corrosion of copper exposed in Martinique (Caribbean Sea) for 1 year was reported. This island suffered the stranding of sargassum algae, which decompose and release toxic gases such as hydrogen sulfide (H2S) or ammonia (NH3). Four sites in Martinique (France) more or less impacted by sargassum algae strandings were selected. The corrosion rate was studied via mass loss determination. The morphology and properties of the corrosion products were determined using Scanning Electron Microscopy (SEM) coupled with energy-dispersive X-ray spectroscopy (EDS) and X-ray diffraction (XRD). The samples were exposed for up to 12 months. The mass loss results after 1-year exposure were from 4.8 µm for the least impacted site to 325 µm for the site most affected by sargassum algae. This very high value proves that the presence of sargassum algae caused a significant degradation of copper. The morphological structures and properties of the corrosion products obtained at the impacted and non-impacted sites differed significantly. In the absence of sargassum algae, classical corrosion products of copper were reported such as Cu2O and Cu2Cl(OH)3. In the sites near the stranding of the sargassum algae, the CuS product is the main corrosion product obtained, but copper hydroxylsulfate is created. Full article
(This article belongs to the Section Corrosion and Protection)
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Figure 1
<p>Positions of the different sites studied on the map of Martinique (<b>a</b>), exposition desks (<b>b</b>).</p> Full article ">Figure 2
<p>Monthly average concentration of H<sub>2</sub>S (<b>a</b>) and chloride (<b>b</b>) in different sites during 1 year.</p> Full article ">Figure 3
<p>Thickness loss of copper exposed in 4 sites: Diamant/Vert pré (<b>a</b>) and Vauclin/Frégate est (<b>b</b>) for 12 months.</p> Full article ">Figure 4
<p>SEM images after 3, 6, 9 and 12 months of exposure in Vert pré (<b>a</b>–<b>d</b>), Diamant (<b>e</b>–<b>h</b>), Vauclin (<b>i</b>–<b>l</b>) and Fregate est (<b>m</b>–<b>p</b>). The scale bar corresponds to 500 μm.</p> Full article ">Figure 5
<p>XRD diffractograms of copper samples exposed for 12 months in Vert pré (<b>a</b>), Diamant (<b>b</b>), Vauclin (<b>c</b>) and Frégate est (<b>d</b>).</p> Full article ">
13 pages, 80879 KiB  
Article
Corrosion Behavior of the AZ31 Mg Alloy in Neutral Aqueous Solutions Containing Various Anions
by Duyoung Kwon, Hien Van Pham, Pungkeun Song and Sungmo Moon
Metals 2023, 13(5), 962; https://doi.org/10.3390/met13050962 - 16 May 2023
Cited by 1 | Viewed by 1861
Abstract
This work demonstrates the corrosion behavior of the AZ31 Mg alloy as a function of an immersion time of 48 h in 0.1 M HCl, H2SO4, H3PO4 and HF solutions, in which pH was adjusted to [...] Read more.
This work demonstrates the corrosion behavior of the AZ31 Mg alloy as a function of an immersion time of 48 h in 0.1 M HCl, H2SO4, H3PO4 and HF solutions, in which pH was adjusted to 6 to exclude the contribution of hydrogen ions (H+) and hydroxide ions (OH−). In situ observations, open circuit potential (OCP), weight changes and AC impedance measurements were performed with an immersion time of 48 h and the morphologies and chemical compositions of the surface products after 48 h of immersion were analyzed by SEM, EDS and XPS. In the chloride ion (Cl−)-containing solution, the corrosion of the AZ31 Mg alloy initiated locally and propagated discontinuously over the surface with immersion time. The OCP value of the AZ31 Mg alloy showed an initial increase from −1.51 VAg/AgCl to −1.47 VAg/AgCl after about 5 h of immersion and then a decrease to −1.51 VAg/AgCl due to corrosion initiation. In the F−-containing solution, after 48 h of immersion, the OCP showed an extremely large value of −0.6 VAg/AgCl, while the relatively lower values of −1.52 VAg/AgCl, −1.59 VAg/AgCl were seen in the solutions containing SO42− and PO43, respectively. In the sulfate ion (SO42−)-containing neutral aqueous solution, needle-like surface films were formed and there were no changes in the weight of the AZ31 Mg alloy with immersion time. In the phosphate ion (PO43−)-containing neutral aqueous solution, a vigorous gas evolution occurred, together with the formation of black surface films with cracks, and a high corrosion rate of −13.8018 × 10−3 g·cm−2·day−1 was obtained. In the fluoride ion (F−)-containing neutral aqueous solution, a surface film with crystalline grains of MgF2 was formed and the weight of the AZ31 Mg alloy increased continuously with immersion time. In conclusion, the corrosion of the AZ31 Mg alloy occurred uniformly in neutral phosphate solution but locally in chloride solution. No corrosion was observed in either the neutral sulfate or fluoride solutions. Full article
(This article belongs to the Special Issue Corrosion and Protection of Lightweight Engineering Materials: Mg Alloys, Al Alloys, Ti Alloys and Other Related Metals)
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<p>Photographs of the AZ31 Mg alloy surfaces during immersion for 48 h in 0.1 M HCl, H<sub>2</sub>SO<sub>4</sub>, H<sub>3</sub>PO<sub>4</sub> and HF solutions at 20 ± 0.5 °C. The solution pH was adjusted to 6 by adding 10 M NaOH solution.</p> Full article ">Figure 2
<p>The OCP of the AZ31 Mg alloy specimens during immersion for 48 h in 0.1 M HCl, H<sub>2</sub>SO<sub>4</sub>, H<sub>3</sub>PO<sub>4</sub> and HF solutions at 20 ± 0.5 °C. The solution pH was adjusted to 6 by adding 10 M NaOH solution.</p> Full article ">Figure 3
<p>Changes in weight of the AZ31 Mg alloy specimens with immersion time in 0.1 M HCl, H<sub>2</sub>SO<sub>4</sub>, H<sub>3</sub>PO<sub>4</sub> and HF solutions at 20 ± 0.5 °C. The solution pH was adjusted to 6 by adding 10 M NaOH solution.</p> Full article ">Figure 4
<p>Nyquist plots obtained during the immersion of the AZ31 Mg alloy in (<b>a</b>) 0.1 M HCl, (<b>b</b>) H<sub>2</sub>SO<sub>4</sub>, (<b>c</b>) H<sub>3</sub>PO<sub>4</sub> and (<b>d</b>) HF solutions at 20 ± 0.5 °C. The solution pH was adjusted to 6 by adding 10 M NaOH solution.</p> Full article ">Figure 5
<p>Plot of the surface film resistance of the AZ31 Mg alloy specimens (<a href="#metals-13-00962-f003" class="html-fig">Figure 3</a>) with immersion time.</p> Full article ">Figure 6
<p>SEM images (<b>a</b>–<b>o</b>) and EDS results (<b>p</b>) of the AZ31 Mg alloy surfaces before (<b>a</b>,<b>b</b>) and after immersion for 48 h in 0.1 M HCl (<b>c</b>–<b>f</b>), H<sub>2</sub>SO<sub>4</sub> (<b>g</b>–<b>i</b>), H<sub>3</sub>PO<sub>4</sub> (<b>j</b>–<b>l</b>) and HF (<b>m</b>–<b>o</b>) solutions, in which pH was adjusted to 6 by adding 10 M NaOH solution.</p> Full article ">Figure 7
<p>XPS spectra of the AZ31 Mg alloy specimens after immersion for 48 h in (<b>a</b>) 0.1 M HCl, (<b>b</b>) H<sub>2</sub>SO<sub>4</sub>, (<b>c</b>) H<sub>3</sub>PO<sub>4</sub> and (<b>d</b>) HF solutions, in which pH was adjusted to 6 by adding 10 M NaOH solution.</p> Full article ">Figure 7 Cont.
<p>XPS spectra of the AZ31 Mg alloy specimens after immersion for 48 h in (<b>a</b>) 0.1 M HCl, (<b>b</b>) H<sub>2</sub>SO<sub>4</sub>, (<b>c</b>) H<sub>3</sub>PO<sub>4</sub> and (<b>d</b>) HF solutions, in which pH was adjusted to 6 by adding 10 M NaOH solution.</p> Full article ">
14 pages, 4704 KiB  
Article
Electrochemical Corrosion Behavior of 310S Stainless Steel in Hot Concentrated Tap Water
by Wen Xian, Zhong Yin, Lele Liu and Moucheng Li
Metals 2023, 13(4), 713; https://doi.org/10.3390/met13040713 - 5 Apr 2023
Cited by 3 | Viewed by 1641
Abstract
The corrosion behavior of 310S stainless steel was investigated in synthetic tap water and Ca2+ and Mg2+-free solutions with different concentration ratios at 80 °C using electrochemical measurement techniques and surface analysis methods. The main purpose was to obtain the [...] Read more.
The corrosion behavior of 310S stainless steel was investigated in synthetic tap water and Ca2+ and Mg2+-free solutions with different concentration ratios at 80 °C using electrochemical measurement techniques and surface analysis methods. The main purpose was to obtain the electrochemical corrosion characteristics under carbonate scale conditions. The specimens displayed a spontaneous passivation state in the solutions with or without Ca2+ and Mg2+ ions. With the enlargement of the concentration ratio of synthetic tap water from 1 to 10 times, the polarization resistance under free corrosion conditions and the pitting potential decreased by about 48% and 327 mV, respectively. The pitting tendency increased with increasing concentration ratio of tap water. The carbonate scales deposited from the synthetic tap water solutions were mainly composed of CaCO3, which led to a slight increase in the polarization resistance and the pitting potential and decrease in the passive current density. Full article
(This article belongs to the Special Issue Applications of Electrochemistry in Corrosion Science and in Practice)
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<p>Open circuit potential curves of specimens immersed in the synthetic tap water solutions with different concentration ratios from 1 to 10 times for 60 min.</p> Full article ">Figure 2
<p>Impedance spectra (<b>a</b>) Nyquist and (<b>b</b>) Bode for specimens in the synthetic tap water solutions with different concentration ratios from 1 to 10 times. Symbols: experimental data; lines: fitted values.</p> Full article ">Figure 3
<p>Polarization curves for specimens in the synthetic tap water solutions with different concentration ratios from 1 to 10 times.</p> Full article ">Figure 4
<p>Open circuit potential curves of specimens in the Ca<sup>2+</sup> and Mg<sup>2+</sup>-free solutions with different concentration ratios from 1 to 10 times for 60 min.</p> Full article ">Figure 5
<p>Impedance spectra (<b>a</b>) Nyquist and (<b>b</b>) Bode for specimens in the Ca<sup>2+</sup> and Mg<sup>2+</sup>-free solutions with different concentration ratios from 1 to 10 times. Symbols: experimental data; lines: fitted values.</p> Full article ">Figure 6
<p>Polarization curves for specimens in the Ca<sup>2+</sup> and Mg<sup>2+</sup>-free solutions with different concentration ratios from 1 to 10 times.</p> Full article ">Figure 7
<p>Typical SEM-SE morphologies (×500) of specimen surfaces after the polarization measurement in the synthetic tap water solutions with different concentration ratios: (<b>a</b>,<b>b</b>) 1, (<b>c</b>,<b>d</b>) 6, and (<b>e</b>,<b>f</b>) 8 times.</p> Full article ">Figure 8
<p>SEM-EDS mapping of elements for the area marked by the blue rectangle at point A in <a href="#metals-13-00713-f007" class="html-fig">Figure 7</a>f.</p> Full article ">Figure 9
<p>XRD patterns for the specimen surfaces after the polarization measurement in the solutions with the concentration ratio 1, 6 and 8 times.</p> Full article ">Figure 10
<p>Variation of corrosion potential with the concentration ratio for specimens in the two solutions.</p> Full article ">Figure 11
<p>Equivalent circuit models for the corrosion system of 310S stainless steel in different concentrated solutions [<a href="#B11-metals-13-00713" class="html-bibr">11</a>].</p> Full article ">Figure 12
<p>Pitting potentials of specimens in the two solutions with different concentration ratios.</p> Full article ">
38 pages, 7189 KiB  
Review
High-Entropy Alloy Coatings Deposited by Thermal Spraying: A Review of Strengthening Mechanisms, Performance Assessments and Perspectives on Future Applications
by Rakesh Bhaskaran Nair, Raunak Supekar, Seyyed Morteza Javid, Wandong Wang, Yu Zou, André McDonald, Javad Mostaghimi and Pantcho Stoyanov
Metals 2023, 13(3), 579; https://doi.org/10.3390/met13030579 - 13 Mar 2023
Cited by 20 | Viewed by 5331
Abstract
Thermal spray deposition techniques have been well-established, owing to their flexibility in addressing degradation due to wear and corrosion issues faced due to extreme environmental conditions. With the adoption of these techniques, a broad spectrum of industries is experiencing continuous improvement in resolving [...] Read more.
Thermal spray deposition techniques have been well-established, owing to their flexibility in addressing degradation due to wear and corrosion issues faced due to extreme environmental conditions. With the adoption of these techniques, a broad spectrum of industries is experiencing continuous improvement in resolving these issues. To increase industrial-level implementation, state-of-the-art advanced materials are required. High-entropy alloys (HEAs) have recently gained considerable attention within the scientific community as advanced materials, mainly due to their exceptional properties and desirable microstructural features. Unlike traditional material systems, high-entropy alloys are composed of multi-component elements (at least five elements) with equimolar or nearly equimolar concentrations. This allows for a stable microstructure that is associated with high configurational entropy. This review article provides a critical assessment of different strengthening mechanisms observed in various high-entropy alloys developed by means of deposition techniques. The wear, corrosion, and oxidation responses of these alloys are reviewed in detail and correlated to microstructural and mechanical properties and behavior. In addition, the review focused on material design principles for developing next-generation HEAs that can significantly benefit the aerospace, marine, oil and gas, nuclear sector, etc. Despite having shown exceptional mechanical properties, the article describes the need to further evaluate the tribological behavior of these HEAs in order to show proof-of-concept perspectives for several industrial applications in extreme environments. Full article
(This article belongs to the Special Issue Wear and Corrosion Behavior of High-Entropy Alloy)
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<p>Differences between low-, medium-, and high-entropy alloys based on their configurational entropy.</p> Full article ">Figure 2
<p>The number of publications of thermal spray high-entropy alloy coatings per year until December 2022 (exported with Scopus analysis tool).</p> Full article ">Figure 3
<p>Backscattered scanning electron microscope images of (<b>a</b>) mechanically alloyed and (<b>b</b>) gas-atomized feedstocks of AlCoCrFeNiTi powders. Coating microstructure of the AlCoCrFeNiTi HEA coatings for (<b>c</b>) mechanically blended, (<b>d</b>) mechanically alloyed, and (<b>e</b>) gas atomized feedstock powders. The different microstructural features with respect to different feedstock powders are clearly visible in the figure, with the gas-atomized powder coatings showing homogeneity compared to the other two coatings [<a href="#B14-metals-13-00579" class="html-bibr">14</a>].</p> Full article ">Figure 4
<p>Setup of HEA synthesis with RF-ICP. Well-mixed powder bed is directly synthesized by RF-ICP under 1100 W with 5000–8000 K and becomes HEAs within 40 s [<a href="#B17-metals-13-00579" class="html-bibr">17</a>].</p> Full article ">Figure 5
<p>Rapid synthesis of alloys. (<b>a</b>) The temperature profile of CuNi alloy using the RF-ICP synthesis. The OM images demonstrate the evolution of the porosity of CuNi alloy during the HEA synthesis process. (<b>b</b>–<b>e</b>) OM images of synthesized FeCoNi MEA, FeCoNiCu MEA, FeCoNiCuAl HEA, and FeCoNiCuTi HEA [<a href="#B17-metals-13-00579" class="html-bibr">17</a>].</p> Full article ">Figure 6
<p>Backscattered scanning electron micrographs of atmospheric plasma spraying (APS) of AlCoCrFeNi HEA coatings. (<b>a</b>) shows a high-magnified image highlighting the nanoscale precipitates within the splats. (<b>b</b>) represents SEM images with different distinct regions (black, white, and gray phases) along with different splats (oxides, thick splats, thin splats, and partial and unmelted particles) [<a href="#B34-metals-13-00579" class="html-bibr">34</a>].</p> Full article ">Figure 7
<p>Schematic representation of how varying particle size ranges expose to the temperature and oxidizing atmosphere during in-flight [<a href="#B37-metals-13-00579" class="html-bibr">37</a>].</p> Full article ">Figure 8
<p>(<b>a</b>) A cross-sectional backscattered SEM image of AlCoCrCuFeNi coated on an Mg substrate showing the as-sprayed layer and a laser-remelted layer. (<b>b</b>) A high-resolution backscattered SEM image showing the epitaxial growth of columnar dendrites at the laser-remelted layer [<a href="#B53-metals-13-00579" class="html-bibr">53</a>].</p> Full article ">Figure 9
<p>Comparative assessment of different strengthening mechanisms for thermal-sprayed HEA coatings. The precipitation strengthening seems to have a higher influence on the microhardness compared to all other strengthening mechanisms [<a href="#B16-metals-13-00579" class="html-bibr">16</a>,<a href="#B38-metals-13-00579" class="html-bibr">38</a>,<a href="#B47-metals-13-00579" class="html-bibr">47</a>,<a href="#B51-metals-13-00579" class="html-bibr">51</a>,<a href="#B56-metals-13-00579" class="html-bibr">56</a>,<a href="#B57-metals-13-00579" class="html-bibr">57</a>,<a href="#B58-metals-13-00579" class="html-bibr">58</a>,<a href="#B59-metals-13-00579" class="html-bibr">59</a>,<a href="#B60-metals-13-00579" class="html-bibr">60</a>,<a href="#B61-metals-13-00579" class="html-bibr">61</a>,<a href="#B62-metals-13-00579" class="html-bibr">62</a>,<a href="#B63-metals-13-00579" class="html-bibr">63</a>,<a href="#B64-metals-13-00579" class="html-bibr">64</a>,<a href="#B65-metals-13-00579" class="html-bibr">65</a>,<a href="#B66-metals-13-00579" class="html-bibr">66</a>,<a href="#B67-metals-13-00579" class="html-bibr">67</a>,<a href="#B68-metals-13-00579" class="html-bibr">68</a>,<a href="#B69-metals-13-00579" class="html-bibr">69</a>,<a href="#B70-metals-13-00579" class="html-bibr">70</a>,<a href="#B71-metals-13-00579" class="html-bibr">71</a>,<a href="#B72-metals-13-00579" class="html-bibr">72</a>,<a href="#B73-metals-13-00579" class="html-bibr">73</a>,<a href="#B74-metals-13-00579" class="html-bibr">74</a>,<a href="#B75-metals-13-00579" class="html-bibr">75</a>,<a href="#B76-metals-13-00579" class="html-bibr">76</a>,<a href="#B77-metals-13-00579" class="html-bibr">77</a>,<a href="#B78-metals-13-00579" class="html-bibr">78</a>,<a href="#B79-metals-13-00579" class="html-bibr">79</a>,<a href="#B80-metals-13-00579" class="html-bibr">80</a>].</p> Full article ">Figure 10
<p>Schematic illustrating the classifications of material properties for quantifying the mechanical performance for extreme environmental applications.</p> Full article ">Figure 11
<p>Summary of different thermal-sprayed HEA coatings investigated for wear behavior.</p> Full article ">Figure 12
<p>Wear morphologies of AlCoCrFeNiTi HEA coatings tested at room temperature and elevated temperature (secondary electron mode). The presence of contrast oxide regions was more profound when tested at room temperature compared to that at elevated temperature [<a href="#B14-metals-13-00579" class="html-bibr">14</a>].</p> Full article ">Figure 13
<p>Optical microscopy images of alumina counter balls after sliding at 5 N load for (<b>a</b>) unpolished, (<b>b</b>) polished cold-sprayed HEA coatings, (<b>c</b>) unpolished and (<b>d</b>) polished flame-sprayed AlCoCrFeMo HEA coatings [<a href="#B95-metals-13-00579" class="html-bibr">95</a>].</p> Full article ">Figure 14
<p>Backscattered SEM images of (<b>a</b>) cold-sprayed and (<b>b</b>) flame-sprayed AlCoCrFeMo HEA coatings [<a href="#B59-metals-13-00579" class="html-bibr">59</a>].</p> Full article ">Figure 15
<p>(<b>a</b>) Potentiodynamic polarization curves of APS AlCoCrFeNi HEA coatings and SS 316L and (<b>b</b>,<b>c</b>) depict the corroded HEA surfaces [<a href="#B68-metals-13-00579" class="html-bibr">68</a>].</p> Full article ">Figure 16
<p>Electrochemical corrosion studies representing potentiodynamic polarization curves and Nyquist plot of flame-sprayed and cold-sprayed AlCoCrFeMo HEA coatings. The figure indicates that the cold-sprayed HEA coatings showed better corrosion performance compared to flame-sprayed HEA coatings [<a href="#B59-metals-13-00579" class="html-bibr">59</a>].</p> Full article ">Figure 17
<p>A comparative assessment of corrosion current density for thermally sprayed HEA coatings with other HEA coatings fabricated using different methods under 3.5 wt% NaCl solution [<a href="#B59-metals-13-00579" class="html-bibr">59</a>,<a href="#B96-metals-13-00579" class="html-bibr">96</a>,<a href="#B97-metals-13-00579" class="html-bibr">97</a>,<a href="#B99-metals-13-00579" class="html-bibr">99</a>,<a href="#B102-metals-13-00579" class="html-bibr">102</a>,<a href="#B103-metals-13-00579" class="html-bibr">103</a>,<a href="#B104-metals-13-00579" class="html-bibr">104</a>,<a href="#B105-metals-13-00579" class="html-bibr">105</a>]. The thermal-sprayed HEA coatings showed lower corrosion rates compared to stainless steel 316L and other HEA coatings. GTAC is gas tungsten arc cladding, CS is cold spraying, FS is flame spraying, HVOF is high-velocity oxy-fuel, APS is air plasma spraying, and LC is laser cladding.</p> Full article ">Figure 18
<p>Typical seals and their locations in gas turbine engines [<a href="#B118-metals-13-00579" class="html-bibr">118</a>].</p> Full article ">
17 pages, 5148 KiB  
Article
Stress Corrosion Cracking Mechanisms of UNS S32205 Duplex Stainless Steel in Carbonated Solution Induced by Chlorides
by Ulises Martin and David M. Bastidas
Metals 2023, 13(3), 567; https://doi.org/10.3390/met13030567 - 11 Mar 2023
Cited by 6 | Viewed by 3925
Abstract
Herein, the chloride-induced stress corrosion cracking (SCC) mechanisms of UNS S32205 duplex stainless steel (DSS) reinforcing bars in alkaline and carbonated solutions are studied. Electrochemical monitoring and mechanical properties were tested using linear polarization resistance and electrochemical impedance spectroscopy, coupled with the slow [...] Read more.
Herein, the chloride-induced stress corrosion cracking (SCC) mechanisms of UNS S32205 duplex stainless steel (DSS) reinforcing bars in alkaline and carbonated solutions are studied. Electrochemical monitoring and mechanical properties were tested using linear polarization resistance and electrochemical impedance spectroscopy, coupled with the slow strain rate tensile test (SSRT) to evaluate the SCC behavior and unravel the pit-to-crack mechanisms. Pit initiation and crack morphology were identified by fractographic analysis, which revealed the transgranular (TG) SCC mechanism. HCO3− acidification enhanced the anodic dissolution kinetics, thus promoting a premature pit-to-crack transition, seen by the decrease in the maximum phase angle in the Bode plot at low frequencies (≈ 1 Hz) for the carbonated solution. The crack propagation rate for the carbonated solution increased by over 100% compared to the alkaline solution, coinciding with the lower phase angle from the Bode plots, as well as with the lower charge transfer resistance. Pit initiation was found at the TiN nonmetallic inclusion inside the ferrite phase cleavage facet, which developed TG-SCC. Full article
(This article belongs to the Special Issue Corrosion and Protection of Stainless Steels)
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<p>Microstructural characterization of as-received UNS S32205 reinforcement: (<b>a</b>) micrograph of the rolling direction ×50, and (<b>b</b>) XRD pattern and alloy phase fraction (γ–phase and α–phase).</p> Full article ">Figure 2
<p>Stress/strain curves of UNS S32205 reinforcement as a function of the chloride content: (<b>a</b>) simulated concrete pore solution (SCPS, pH 12.6), and (<b>b</b>) carbonated solution (CBS, pH 9.1).</p> Full article ">Figure 3
<p>Linear polarization resistance measurements of UNS S32205 reinforcement during slow strain rate test (SSRT): (<b>a</b>) <span class="html-italic">E</span><sub>corr</sub>, and (<b>b</b>) <span class="html-italic">i</span><sub>corr</sub>.</p> Full article ">Figure 4
<p>Nyquist plots of UNS S32205 reinforcement during slow strain rate test (SSRT): (<b>a</b>) SCPS 0 wt.% Cl<sup>–</sup>, (<b>b</b>) SCPS 4 wt.% Cl<sup>–</sup>, (<b>c</b>) SCPS 8 wt.% Cl<sup>–</sup>, (<b>d</b>) CBS 0 wt.% Cl<sup>–</sup>, (<b>e</b>) CBS 4 wt.% Cl<sup>–</sup>, and (<b>f</b>) CBS 8 wt.% Cl<sup>–</sup>.</p> Full article ">Figure 5
<p>Electric equivalent circuit (EEC) with two time constants.</p> Full article ">Figure 6
<p>Micrographs of UNS S32205 reinforcement after slow strain rate test (SSRT) immersed in SCPS: 0 wt.% Cl<sup>–</sup> (<b>a</b>) rupture surface ×50 (<b>b</b>) microvoids and coalescence of dimples ×1300, (<b>c</b>) ductile overload areas ×1000; 4 wt.% Cl<sup>–</sup> (<b>d</b>) rupture surface ×50, (<b>e</b>) ductile overload area ×600, (<b>f</b>) brittle fracture inside the ferrite cleavage facets ×22,800 ; 8 wt.% Cl<sup>–</sup> (<b>g</b>) rupture surface ×50, (<b>h</b>) cracks inside the ferrite cleavage facets ×4500, and (<b>i</b>) crack nucleation due to inclusions ×12,300.</p> Full article ">Figure 7
<p>Micrographs of UNS S32205 reinforcement after slow strain rate test (SSRT) immersed in CBS: 0 wt.% Cl<sup>–</sup> (<b>a</b>) rupture surface ×50 (<b>b</b>) microvoids and coalescence of dimples ×550, (<b>c</b>) broken inclusions ×11,200; 4 wt.% Cl<sup>–</sup> (<b>d</b>) rupture surface ×50, (<b>e</b>) brittle fracture mode ×1200, (<b>f</b>) microcracks inside the ferrite cleavage facets ×10,300; 8 wt.% Cl<sup>–</sup> (<b>g</b>) rupture surface ×50, (<b>h</b>) cracks inside the ferrite cleavage facets ×3100, and (<b>i</b>) crack nucleation due to inclusions ×7600.</p> Full article ">Figure 8
<p>Crack propagation rate and stress monitoring UNS S32205 reinforcement during SSRT: (<b>a</b>) simulated concrete pore solution (SCPS, pH 12.6), and (<b>b</b>) carbonated solution (CBS, pH 9.1).</p> Full article ">Figure 9
<p>Bode plots of UNS S32205 reinforcement during SSRT: (<b>a</b>) simulated concrete pore solution (SCPS, pH 12.6), and (<b>b</b>) carbonated solution (CBS, pH 9.1).</p> Full article ">Figure 10
<p>EDX analysis of Ti-based nonmetallic inclusion in UNS S32205 reinforcement after failure immersed in carbonated solution (CBS, pH 9.1) contaminated with 8 wt.% Cl<sup>–</sup>: (<b>a</b>) SEM micrograph ×7600, and (<b>b</b>) EDX spectrum.</p> Full article ">
21 pages, 14270 KiB  
Article
Effect of Laser Shock Peening on the Stress Corrosion Cracking of 304L Stainless Steel
by Young-Ran Yoo, Seung-Heon Choi and Young-Sik Kim
Metals 2023, 13(3), 516; https://doi.org/10.3390/met13030516 - 3 Mar 2023
Cited by 8 | Viewed by 2519
Abstract
Storage canisters used in nuclear power plants operating in seaside areas—where the salt content in the atmosphere is high—may be susceptible to chloride-induced stress corrosion cracking (CISCC). Chloride-induced stress corrosion cracking is one of the ways in which dry storage canisters made of [...] Read more.
Storage canisters used in nuclear power plants operating in seaside areas—where the salt content in the atmosphere is high—may be susceptible to chloride-induced stress corrosion cracking (CISCC). Chloride-induced stress corrosion cracking is one of the ways in which dry storage canisters made of stainless steel can degrade. Stress corrosion cracking depends on the microstructure and residual stress, and it is therefore very important to improve the surface properties of materials. Laser shock peening both greatly deforms the material surface and refines grains, and it generates compressive residual stress in the deep part from the surface of the material. This study focused on the effect of laser shock peening on the stress corrosion cracking of 304L stainless steel. The laser shock peening was found to induce compressive residual stress from the surface to a 1 mm depth, and the SCC properties were evaluated by a U-bend test. The results showed that the SCC resistance of laser-peened 304L stainless steel in a chloride environment was enhanced, and that it was closely related to grain size, the pitting potential of the cross section, and residual stress. Full article
(This article belongs to the Special Issue Feature Paper Collection of “Current Challenges in Corrosion Research”)
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Figure 1

Figure 1
<p>(<b>a</b>) Corrosion Cell for U-Bend SCC Test; (<b>b</b>) Setup of U-Bend Specimen; (<b>c</b>) Insulation on the Sides of U-Bend Specimen.</p> Full article ">Figure 2
<p>Effect of laser shock peening on the crack times of 304LB by U−bend SCC test at 155 °C, 42% MgCl<sub>2</sub>: (<b>a</b>) total crack time; (<b>b</b>) crack propagation rates.</p> Full article ">Figure 3
<p>Surface appearance of 304LB by laser shock peening: (<b>a</b>) 304LB; (<b>b</b>) 304LB−L−NC; (<b>c</b>) 304LB−L−WC.</p> Full article ">Figure 4
<p>Photographs of the cross section of 304LB after U-bend SCC test (OM, ×50, 155 °C, 42% MgCl<sub>2</sub>): (<b>a</b>) 304LB; (<b>b</b>) 304LB−L−NC; (<b>c</b>) 304LB−L−WC.</p> Full article ">Figure 5
<p>Crack morphologies of 304LB after U-bend SCC test (OM, ×200, 155 °C, 42% MgCl<sub>2</sub>, Areas A, B, and C in <a href="#metals-13-00516-f005" class="html-fig">Figure 5</a>): (<b>a</b>) 304LB; (<b>b</b>) 304LB−L−NC; (<b>c</b>) 304LB−L−WC.</p> Full article ">Figure 6
<p>Cracking mode of 304LB after U-bend SCC test (SEM, ×400, 155 °C, 42% MgCl<sub>2</sub>, top, middle, and bottom areas in <a href="#metals-13-00516-f005" class="html-fig">Figure 5</a> and <a href="#metals-13-00516-f006" class="html-fig">Figure 6</a>): (<b>a</b>) 304LB; (<b>b</b>) 304LB−L−NC; (<b>c</b>) 304LB−L−WC.</p> Full article ">Figure 7
<p>Effect of laser shock peening on the total crack times of welded 304LW by U−bend SCC test at 155 °C, 42% MgCl<sub>2</sub>: (<b>a</b>) crack time; (<b>b</b>) crack propagation rates.</p> Full article ">Figure 8
<p>Surface appearance of 304LW by LP: (<b>a</b>) 304LW; (<b>b</b>) 304LW−L−NC; (<b>c</b>) 304LW−L−WC.</p> Full article ">Figure 9
<p>Photographs of the cross section of 304LW after U-bend SCC test (OM, ×50, 155 °C, 42% MgCl<sub>2</sub>): (<b>a</b>) 304LW; (<b>b</b>) 304LW−L−NC; (<b>c</b>) 304LW−L−WC.</p> Full article ">Figure 10
<p>Crack morphologies of 304LW after U-bend SCC test (OM, ×200, 155 °C, 42% MgCl<sub>2</sub>, Areas A, B, and C in <a href="#metals-13-00516-f010" class="html-fig">Figure 10</a>): (<b>a</b>) 304LW; (<b>b</b>) 304LW−L−NC; (<b>c</b>) 304LW−L−WC.</p> Full article ">Figure 11
<p>Cracking mode of 304LW after U-bend SCC test (SEM, ×400, 155 °C, 42% MgCl<sub>2</sub>, top, middle, and bottom areas in <a href="#metals-13-00516-f010" class="html-fig">Figure 10</a> and <a href="#metals-13-00516-f011" class="html-fig">Figure 11</a>): (<b>a</b>) 304LW; (<b>b</b>) 304LW−L−NC; (<b>c</b>) 304LW−L−WC.</p> Full article ">Figure 12
<p>Relationship between the crack propagation rate and average grain size: (<b>a</b>) total crack propagation rate, (<b>b</b>) net crack propagation rate.</p> Full article ">Figure 13
<p>Relationship between crack propagation rate and intergranular corrosion: (<b>a</b>,<b>b</b>) total crack propagation, (<b>a’</b>,<b>b’</b>) net crack propagation rate.</p> Full article ">Figure 14
<p>Relationship between the crack propagation rate and pitting potential: (<b>a</b>,<b>b</b>) total crack propagation, (<b>a’</b>,<b>b’</b>) net crack propagation rate.</p> Full article ">Figure 15
<p>Relationship between the crack propagation rate and residual stress: (<b>a</b>) total crack propagation, (<b>b</b>) net crack propagation rate.</p> Full article ">Figure 16
<p>EBSD results of the cracking area of non-peened 304L base metal after SCC test: (<b>a</b>) SEM image, (<b>b</b>) band contrast, (<b>c</b>) inverse pole figure (IPF) coloring.</p> Full article ">Figure 17
<p>EBSD results of the cracking area of laser-peened 304L base metal after SCC test: (<b>a</b>) SEM image, (<b>b</b>) band contrast, (<b>c</b>) inverse pole figure (IPF) coloring.</p> Full article ">Figure 18
<p>Proposed model of crack initiation and propagation of 304L stainless steel by laser shock peening: (<b>a</b>) non-peening; (<b>b</b>) laser shock peening.</p> Full article ">
29 pages, 15823 KiB  
Review
Corrosion Behavior of High Entropy Alloys and Their Application in the Nuclear Industry—An Overview
by Tianrun Li, Debin Wang, Suode Zhang and Jianqiang Wang
Metals 2023, 13(2), 363; https://doi.org/10.3390/met13020363 - 10 Feb 2023
Cited by 13 | Viewed by 3221
Abstract
With multiple principal components, high entropy alloys (HEAs) have aroused great interest due to their unique microstructures and outstanding properties. Recently, the corrosion behavior of HEAs has become a scientific hotspot in the area of material science and engineering, and HEAs can exhibit [...] Read more.
With multiple principal components, high entropy alloys (HEAs) have aroused great interest due to their unique microstructures and outstanding properties. Recently, the corrosion behavior of HEAs has become a scientific hotspot in the area of material science and engineering, and HEAs can exhibit good protection against corrosive environments. A comprehensive understanding of the corrosion mechanism of HEAs is important for further design of HEAs with better performance. This paper reviews the corrosion properties and mechanisms of HEAs (mainly Cantor alloy and its variants) in various environments. More crucially, this paper is focused on the influences of composition and microstructure on the evolution of the corrosion process, especially passive film stability and localized corrosion resistance. The corrosion behavior of HEAs as structural materials in nuclear industry applications is emphasized. Finally, based on this review, the possible perspectives for scientific research and engineering applications of HEAs are proposed. Full article
(This article belongs to the Special Issue Feature Paper Collection of “Current Challenges in Corrosion Research”)
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Figure 1
<p>An overview of corrosion behavior and mechanism of HEAs regarding their elements (red area), microstructure (green area), and applications (blue area). (<b>a</b>) Passive film stability and pitting behavior for various alloys; (<b>b</b>) corrosion mechanism diagram of EHEA in H<sub>2</sub>SO<sub>4</sub> solution; (<b>c</b>) SEM results of the cross section of HEAs after LBE corrosion [<a href="#B25-metals-13-00363" class="html-bibr">25</a>,<a href="#B26-metals-13-00363" class="html-bibr">26</a>,<a href="#B27-metals-13-00363" class="html-bibr">27</a>].</p> Full article ">Figure 2
<p>The polarization and EIS performance of FeCoNiCr<span class="html-italic"><sub>x</sub></span> (<span class="html-italic">x</span> = 0, 0.5, 1) HEAs. Potentiodynamic curves for FeCoNiCr<span class="html-italic"><sub>x</sub></span> in (<b>a</b>) H<sub>2</sub>SO<sub>4</sub> aqueous solution (0.5 M) and (<b>b</b>) NaCl aqueous solution (3.5 wt.%); (<b>c</b>–<b>e</b>) the Nyquist, Bode curves and electrical equivalent circuit for FeCoNiCr<span class="html-italic"><sub>x</sub></span> in 0.5 M H<sub>2</sub>SO<sub>4</sub> aqueous solution [<a href="#B32-metals-13-00363" class="html-bibr">32</a>].</p> Full article ">Figure 3
<p>The electrochemical polarization performance, corrosion morphologies, and XPS results for Ti-containing HEAs. (<b>a</b>) Potentiodynamic polarization curves of AlCoCrFeNiTi<span class="html-italic"><sub>x</sub></span> (<span class="html-italic">x</span> = 0, 0.2, 0.4, 0.6, 0.8, and 1.0) coatings and AISI1045 steel substrate; (<b>b</b>) potentiodynamic polarization plots of Al<sub>2−<span class="html-italic">x</span></sub>CoCrFeNiTi<span class="html-italic"><sub>x</sub></span>; corrosion morphologies of Al<sub>2−<span class="html-italic">x</span></sub>CoCrFeNiTi<span class="html-italic"><sub>x</sub></span> HEAs after the potentiodynamic polarization in 3.5 wt.% NaCl aqueous solution: (<b>c</b>) Ti<sub>0</sub>Al<sub>2</sub>, (<b>d</b>) Ti<sub>0.2</sub>Al<sub>1.8</sub>, (<b>e</b>) Ti<sub>0.5</sub>Al<sub>1.5</sub>; XPS results of the surface films upon (<b>f</b>) Ti<sub>0.5</sub>Al<sub>1.5</sub>, and (<b>g</b>) Ti<sub>1.0</sub>Al<sub>1.0</sub> HEAs [<a href="#B9-metals-13-00363" class="html-bibr">9</a>,<a href="#B34-metals-13-00363" class="html-bibr">34</a>].</p> Full article ">Figure 4
<p>The polarization behavior and the surface chemistry for FeCoCrNiMo<span class="html-italic"><sub>x</sub></span> (<span class="html-italic">x</span> = 0, 0.1, 0.3, 0.6) HEAs. (<b>a</b>) Potentiodynamic polarization plots of FeCoCrNiMo<span class="html-italic"><sub>x</sub></span> in 1 M NaCl aqueous solution; the content percentages of Fe<sub>(o<span class="html-italic">x</span>)(hyd)</sub>, Cr<sub>(ox)</sub>, Cr<sub>(hyd)</sub> and Mo<sub>(ox)</sub> within the surface layers of Mo<sub>0</sub>, Mo<sub>0.1</sub> and Mo<sub>0.6</sub> alloys: (<b>b</b>) the surface of the passive layers, and (<b>c</b>) the 1 nm depth of the passive layers; (<b>d</b>) potentiodynamic polarization curves and (<b>e</b>) Motto–Schottky plots after polarization for 4 h at 0.4 V<sub>SCE</sub> of FeCoCrNiMo<span class="html-italic"><sub>x</sub></span> in 0.5 M H<sub>2</sub>SO<sub>4</sub> [<a href="#B38-metals-13-00363" class="html-bibr">38</a>,<a href="#B41-metals-13-00363" class="html-bibr">41</a>].</p> Full article ">Figure 5
<p>The polarization behaviors, surface chemistry, and pitting corrosion behaviors of Al<span class="html-italic"><sub>x</sub></span>(CoCrFeNi)<sub>1−<span class="html-italic">x</span></sub> HEAs. (<b>a</b>) Chemical compositions vs. number of samples (spots 1 to 10); (<b>b</b>) XRD diffraction patterns of Al–containing HEA samples; (<b>c</b>) polarization plots of the ten HEA samples; #1–10 represent the Al<span class="html-italic"><sub>x</sub></span>(CoCrFeNi)<sub>1−<span class="html-italic">x</span></sub> HEAs in (<b>a</b>); (<b>d</b>) Al 2p, (<b>e</b>) Cr 2p3/2 spectra of #2, #4, #6, and #8 HEA samples after immersion in 3.5 wt.% NaCl aqueous solution for 24 h; (<b>f</b>) enlarged current fluctuation of potentiodynamic polarization curves for Al<sub>0.3</sub>CoCrFeNi, red boxes presents the enlarged current fluctuations; (<b>g</b>) corrosion micrograph after the polarization test at 0.6 V<sub>SCE</sub> [<a href="#B25-metals-13-00363" class="html-bibr">25</a>,<a href="#B46-metals-13-00363" class="html-bibr">46</a>].</p> Full article ">Figure 6
<p>The effect of the addition of N on the polarization behaviors, pitting corrosion performance, and nucleation and growth processes of CoCrFeMnNiN<span class="html-italic"><sub>x</sub></span>. (<b>a</b>) The potentiodynamic polarization plots for CoCrFeMnNi and CoCrFeMnNiN in 0.1 M H<sub>2</sub>SO<sub>4</sub> aqueous solution; (<b>b</b>) potentiodynamic polarization plots for CoCrFeMnNi and CoCrFeMnNiN<span class="html-italic"><sub>x</sub></span> in 3.5 wt.% NaCl aqueous solution; (<b>c</b>) corrosion morphologies containing enlarged image (inset) after potentiostatic polarization; variation of nucleation rates of (<b>d</b>) Cr<sub>2</sub>O<sub>3</sub> and (<b>e</b>) Fe<sub>2</sub>O<sub>3</sub> within the passive films of CoCrFeMnNiN<span class="html-italic"><sub>x</sub></span> [<a href="#B33-metals-13-00363" class="html-bibr">33</a>,<a href="#B58-metals-13-00363" class="html-bibr">58</a>].</p> Full article ">Figure 7
<p>The design of corrosion resistant HEAs by using integrated computational materials engineering (ICME), and corresponding characterizations of the film chemistry and electrochemical behaviors. (<b>a</b>) Pseudo-binary phase diagram of the Ni<sub>59−<span class="html-italic">x</span></sub>Cr<span class="html-italic"><sub>x</sub></span>Fe<sub>20</sub>Ru<sub>13</sub>Mo<sub>6</sub>W<sub>2</sub> alloy system; (<b>b</b>) CALPHAD-calculated <span class="html-italic">E</span>-pH plots for Ni<sub>38</sub>Cr<sub>21</sub>Fe<sub>20</sub>Ru<sub>13</sub>Mo<sub>6</sub>W<sub>2</sub> in 1 kg H<sub>2</sub>O at 25 °C and 1 atm.; (<b>c</b>) potentiodynamic polarization curve of the Ni<sub>38</sub>Cr<sub>21</sub>Fe<sub>20</sub>Ru<sub>13</sub>Mo<sub>6</sub>W<sub>2</sub> HEA in 0.1 M Na<sub>2</sub>SO<sub>4</sub> solution with different pH; (<b>d</b>–<b>f</b>) APT characterization of the passive film formed on the HEA; (<b>g</b>–<b>i</b>) the polarization curve (upper) and the elemental dissolution rates (lower) of the HEA in de-aerated 2 M H<sub>2</sub>SO<sub>4</sub>; (<b>j</b>) elemental dissolution rates during a potentiostatic polarization at 0.6 V<sub>SCE</sub> [<a href="#B63-metals-13-00363" class="html-bibr">63</a>,<a href="#B64-metals-13-00363" class="html-bibr">64</a>,<a href="#B65-metals-13-00363" class="html-bibr">65</a>].</p> Full article ">Figure 8
<p>The microstructure, polarization behaviors, and corrosion failure morphologies of dual-phase Al<span class="html-italic"><sub>x</sub></span>CoCrFeNi (<span class="html-italic">x</span> = 0.3, 0.5, and 0.7) HEAs. (<b>a</b>) The elemental distribution by EDS mappings in Al<span class="html-italic"><sub>x</sub></span>CoCrFeNi, in which the brown colour represents Al, the violet colour represents Ni, the yellow colour represents Ni, the blue colour represents Cr, and the green colour represents Fe; (<b>b</b>) potentiodynamic polarization curves in the 3.5 wt.% NaCl aqueous solution, in which the black curve is Al<sub>0.3</sub>CoCrFeNi, the red curve is Al<sub>0.5</sub>CoCrFeNi, the green curve is Al<sub>0.7</sub>CoCrFeNi; and SEM micrographs after potentiodynamic polarization for the Al<sub>0.3</sub>CoCrFeNi (<b>c<sub>1</sub></b>), Al<sub>0.5</sub>CoCrFeNi (<b>c<sub>2</sub></b>) and Al<sub>0.7</sub>CoCrFeNi (<b>c<sub>3</sub></b>), the white arrows represent the corrosion pits and selective corrosion of BCC phase. [<a href="#B44-metals-13-00363" class="html-bibr">44</a>].</p> Full article ">Figure 9
<p>The effect of homogenization on the microstructure characterizations, electrochemical polarizations, and surface corroded morphologies for Al<sub>0.7</sub>CoCrFeNi HEAs. The surface morphology images, phase maps, and inverse pole-figures (IPF) of (<b>a<sub>1</sub></b>–<b>a<sub>3</sub></b>) the as-forged and (<b>b<sub>1</sub></b>–<b>b<sub>3</sub></b>) the as-equilibrated Al<sub>0.7</sub>CoCrFeNi, the red areas are A1 phases while the blue areas are A2/B2 phases in (<b>a<sub>2</sub></b>,<b>b<sub>2</sub></b>), the different colours in (<b>a<sub>3</sub></b>,<b>b<sub>3</sub></b>) represent the grains with different oriented; the elemental distribution line-scanning by TEM-EDS of (<b>c<sub>1</sub></b>) the as-forged and (<b>c<sub>2</sub></b>) the as-equilibrated Al<sub>0.7</sub>CoCrFeNi, the different colours of the lines represent different elements; potentiodynamic polarization plots of (<b>d<sub>1</sub></b>) the as-forged and (<b>d<sub>2</sub></b>) as-equilibrated Al<sub>0.7</sub>CoCrFeNi in 3.5 wt.% NaCl solution, in which the black curve is Al<sub>0.3</sub>CoCrFeNi, the red curve is Al<sub>0.5</sub>CoCrFeNi, the green curve is Al<sub>0.7</sub>CoCrFeNi; (<b>e<sub>1</sub></b>) the contact potential difference (V<sub>CPD</sub>) map and line scanning of the as-forged Al<sub>0.7</sub>CoCrFeNi, the net-like A2 precipitates with a lower V<sub>CPD</sub> distributed inside the bright B2 matrix; (<b>e<sub>2</sub></b>) V<sub>CPD</sub> maps and line profiles for the as-equilibrated Al<sub>0.7</sub>CoCrFeNi at the phase boundary, the blue arrows are the line scanning sites, the bright areas are the phases with high V<sub>CPD</sub>, and vice versa, the red dotted lines represent the average values of different phases; surface morphologies of (<b>f<sub>1</sub></b>) the as-forged and (<b>f<sub>2</sub></b>) the as-equilibrated Al<sub>0.7</sub>CoCrFeNi after the potentiodynamic polarization tests, the white arrows represent the selective corrosion of phases. [<a href="#B17-metals-13-00363" class="html-bibr">17</a>].</p> Full article ">Figure 10
<p>SEM morphologies and 3D characterizations of the FeCoCrNiMo<span class="html-italic"><sub>x</sub></span> (<span class="html-italic">x</span> = 0, 0.1, 0.3, 0.6) alloy after the anodic polarization in 1 M NaCl aqueous solution: (<b>a<sub>1</sub></b>–<b>c<sub>1</sub></b>) FeCoCrNiMo<sub>0</sub> alloy, (<b>a<sub>2</sub></b>–<b>c<sub>2</sub></b>) FeCoCrNiMo<sub>0.1</sub> alloy in which the red circles are the regions of the anodic dissolution, (<b>a<sub>3</sub></b>–<b>c<sub>3</sub></b>) FeCoCrNiMo<sub>0.3</sub> alloy, and (<b>a<sub>4</sub></b>–<b>c<sub>4</sub></b>) FeCoCrNiMo<sub>0.6</sub> alloy. (<b>a<sub>1</sub></b>–<b>a<sub>4</sub></b>) and (<b>b<sub>1</sub></b>–<b>b<sub>4</sub></b>) SEM and 3D images after anodic polarization for 600 s; (<b>c<sub>1</sub></b>–<b>c<sub>4</sub></b>) 3D images after anodic polarization for 1800s; (<b>d</b>) topographical, (<b>e</b>) potential distribution mapping of the FeCoCrNiMo<sub>0.3</sub> alloy and (<b>f</b>) line-scanning results of the white line in (<b>e</b>) [<a href="#B41-metals-13-00363" class="html-bibr">41</a>].</p> Full article ">Figure 11
<p>(<b>a<sub>1</sub></b>) SEM morphology of the eutectic AlCoCrFeNi<sub>2.1</sub> specimen surface; (<b>a<sub>2</sub></b>) line-scanning of eutectic AlCoCrFeNi<sub>2.1</sub> immersed in 1 wt.% NaCl aqueous solution at different time interval obtained at the dashed line; (<b>a<sub>3</sub></b>) corrosion morphology of the specimen after immersion for 5 days; 3D image (<b>a<sub>4</sub></b>) and line profile (<b>a<sub>5</sub></b>) of the corroded specimen showing preferential dissolution of B2 phase. TEM characterizations of corrosion product film upon specimen at 2 h: (<b>b<sub>1</sub></b>) BF-TEM morphology; (<b>b<sub>2</sub></b>,<b>b<sub>3</sub></b>) the magnified TEM images of films upon FCC and BCC phases; (<b>b<sub>4</sub></b>,<b>b<sub>5</sub></b>) elemental distribution of films upon FCC and BCC phases; (<b>b<sub>6</sub></b>,<b>b<sub>7</sub></b>) HRTEM images of film on FCC and BCC phase; (<b>c<sub>1</sub></b>–<b>c<sub>4</sub></b>) corrosion morphologies in different regions of EHEA in 0.5 M H<sub>2</sub>SO<sub>4</sub> aqueous solution after being immersed for different times: 5 min and 8 h. The schematic of corrosion mechanism of EHEA: (<b>d<sub>1</sub></b>) micro-galvanic corrosion effect; (<b>d<sub>2</sub></b>) short-term self-protection; (<b>d<sub>3</sub></b>) the dissolution of surface BCC phase and the dissolution of FCC lamellas; (<b>d<sub>4</sub></b>) localized dissolution of lamellar region [<a href="#B27-metals-13-00363" class="html-bibr">27</a>,<a href="#B80-metals-13-00363" class="html-bibr">80</a>].</p> Full article ">Figure 12
<p>(<b>a</b>) the SEM morphology of the FeCrNiCoNb<sub>0.5</sub> EHEA; (<b>b</b>) the low magnification TEM morphology of FCC and Laves phases, and the corresponding FFT images in the inset; (<b>c</b>) the HRTEM image of the boundary of the phases; (<b>d<sub>1</sub></b>–<b>d<sub>5</sub></b>) XRD results and (<b>e<sub>1</sub></b>–<b>e<sub>5</sub></b>) the SEM images of the lamellar structure for the samples with different sizes of lamellar (note that scale bars are equal to 2 μm in (<b>e<sub>1</sub></b>–<b>e<sub>5</sub></b>)). (<b>f<sub>1</sub></b>–<b>f<sub>5</sub></b>) potentiodynamic polarization plots of samples with different sizes of lamellar [<a href="#B75-metals-13-00363" class="html-bibr">75</a>,<a href="#B81-metals-13-00363" class="html-bibr">81</a>].</p> Full article ">Figure 13
<p>(<b>a<sub>1</sub></b>) SEM image, XRD pattern, and EBSD phase map; (<b>a<sub>2</sub></b>) HAADF morphology and (<b>a<sub>3</sub></b>) EDS elemental distribution of the dual-phase Al<sub>0.7</sub>CoCrFeNi HEA; SEM micrographs of the dual-phase Al<sub>0.7</sub>CoCrFeNi HEA after being immersed in LBE at (<b>b<sub>1</sub></b>) 350 °C and (<b>b<sub>2</sub></b>) 500 °C; (<b>c<sub>1</sub></b>) optical micrograph and XRD pattern of the Al<sub>0.4</sub>CoCrFeNi HEA; (<b>c<sub>2</sub></b>) BSE morphology of intergranular precipitates; (<b>c<sub>3</sub></b>) TEM-EDS maps of the intergranular precipitates; (<b>d<sub>1</sub></b>) the LBE ingress depth at grain boundaries of Al<sub>0.3</sub>CoCrFeNi and Al<sub>0.4</sub>CoCrFeNi HEAs; (<b>d<sub>2</sub></b>) BSE image of the Al<sub>0.4</sub>CoCrFeNi after being immersed in LBE at 500 °C for 500 h [<a href="#B19-metals-13-00363" class="html-bibr">19</a>,<a href="#B96-metals-13-00363" class="html-bibr">96</a>].</p> Full article ">Figure 14
<p>SEM and EDS line-scanning of the cross section of (<b>a</b>) AlCrFeNiCu, (<b>b</b>) AlCrFeNiTi and (<b>c</b>) AlCrFeNiNb after being immersed in 10<sup>−6</sup> wt.% oxygen containing molten Pb at 600 °C for 1000 h, the red arrows represent the positions of the line-scanning, the blue arrow represent the transitional layer, the different colours of the lines represent different elements, and the dotted lines represent the boundary of the different layers [<a href="#B26-metals-13-00363" class="html-bibr">26</a>].</p> Full article ">Figure 15
<p>SEM morphologies of the FeNiCrCuAl HEA: (<b>a</b>) low magnification image, (<b>b</b>) high magnification image, and the yellow arrows show the phase distribution; (<b>c</b>) the weight gain curves vs. exposure period at 550 °C and 25 MPa in SCW for different materials; the line-scanning elemental distribution of the oxide films (<b>d</b>) upon the FeCr-rich BCC phase after being immersed at 550 °C and 25 MPa for 100 h, (<b>e</b>) upon the FeCr-rich BCC phase after being immersed at 550 °C and 25 MPa for 1000 h, and (<b>f</b>) upon the NiAl-rich phase after being immersed at 550 °C and 25 MPa for 100 h, the dotted lines represent the boundary of the different layers [<a href="#B101-metals-13-00363" class="html-bibr">101</a>].</p> Full article ">Figure 16
<p>Proposed future work and perspectives for the corrosion fields of HEAs, including design composition for highly corrosion resistant HEAs, customized specific microstructures for highly corrosion resistant HEAs, applications in special environments, and applications of surface coating technology. The red part represents the composition design for highly corrosion resistant HEAs, the different colours of the lines represent different elements; the green part represents the customize specific microstructure for highly corrosion resistant HEAs, the blue part represents the application of HEAs in special environment, and the brown represents the application of HEAs of surface coating technology [<a href="#B26-metals-13-00363" class="html-bibr">26</a>,<a href="#B63-metals-13-00363" class="html-bibr">63</a>].</p> Full article ">
13 pages, 13119 KiB  
Article
Effect of Superhydrophobic Surface on Corrosion Resistance of Magnesium-Neodymium Alloy in Artificial Hand Sweat
by Changyang Liu, Jiapeng Sun and Guosong Wu
Metals 2023, 13(2), 219; https://doi.org/10.3390/met13020219 - 24 Jan 2023
Cited by 5 | Viewed by 1714
Abstract
A superhydrophobic surface can endow metals with some intriguing characteristics such as self-cleaning behavior. In this study, a simple solution-immersion method based on the concept of predesigned corrosion is developed to enhance the corrosion resistance of a magnesium-neodymium alloy. The Mg alloy is [...] Read more.
A superhydrophobic surface can endow metals with some intriguing characteristics such as self-cleaning behavior. In this study, a simple solution-immersion method based on the concept of predesigned corrosion is developed to enhance the corrosion resistance of a magnesium-neodymium alloy. The Mg alloy is directly soaked in potassium dihydrogen phosphate solution with the addition of ultrasound, and a layer of rough but dense coating is uniformly formed on the Mg-Nd alloy after the immersion process, which is mainly composed of MgHPO4∙3H2O. A superhydrophobic surface with an average wetting angle of 150.5° and a sliding angle of about 4.5° can be obtained on the Mg alloy by further chemical surface modification with perfluorodecyltriethoxysilane. This superhydrophobic surface has an interesting self-cleaning effect as well as good corrosion resistance in artificial hand sweat. In brief, this study provides a feasible way to prepare a superhydrophobic surface on the Mg-Nd alloy and reveals the effect of a superhydrophobic surface on the corrosion behavior of the Mg-Nd alloy, offering new technical insights into the corrosion protection of magnesium alloys. Full article
(This article belongs to the Special Issue Corrosion and Protection of Lightweight Engineering Materials: Mg Alloys, Al Alloys, Ti Alloys and Other Related Metals)
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<p>(<b>a</b>) Schematic diagram of the fabrication process; (<b>b</b>) temperature and pH value evolution during the process; (<b>c</b>) surface appearance of the treated and untreated samples.</p> Full article ">Figure 2
<p>XRD patterns of the untreated sample and the sample treated by uniform corrosion.</p> Full article ">Figure 3
<p>(<b>a</b>,<b>b</b>) SEM images of surface morphology of the coated sample after PFDS-modification; (<b>c</b>,<b>d</b>) SEM images of cross-section morphology of the coated sample after PFDS-modification.</p> Full article ">Figure 4
<p>EDS mapping of the cross-section of the coated sample after PFDS-modification: (<b>a</b>) SEM image of the selected region, (<b>b</b>–<b>e</b>) maps showing the distribution of the elements.</p> Full article ">Figure 5
<p>XPS spectra of the PFDS-modified surface: (<b>a</b>) survey spectrum and high-resolution spectra of (<b>b</b>) C 1s, (<b>c</b>) F 1s and (<b>d</b>) Si 2p.</p> Full article ">Figure 6
<p>(<b>a</b>) Macro-wettability of the investigated samples; (<b>b</b>) water contact angle of the investigated samples; (<b>c</b>) sliding angle of the hydrophobic samples.</p> Full article ">Figure 7
<p>(<b>a</b>) Self-cleaning behavior of all the investigated samples; (<b>b</b>) schematic diagram of the test.</p> Full article ">Figure 8
<p>(<b>a</b>) Polarization curves of all the investigated samples after immersion for 1 h; (<b>b</b>) E<sub>corr</sub> and I<sub>corr</sub> calculated from the corresponding polarization curves.</p> Full article ">Figure 9
<p>(<b>a</b>) Bode plots of the sample after immersion for 1 h: impedance versus frequency; (<b>b</b>) Bode plots of the sample after immersion for 1 h: phase angle versus frequency; (<b>c</b>) Nyquist plots of the sample after immersion for 1 h; (<b>d</b>) equivalent electrical circuits models for the investigated samples.</p> Full article ">Figure 10
<p>Optical and SEM images of surface morphology of (<b>a</b>) MA, (<b>b</b>) MA/CM, (<b>c</b>) MA/UC and (<b>d</b>) MA/UC/CM after immersion for 24 h.</p> Full article ">Figure 11
<p>(<b>a</b>,<b>b</b>) Schematic diagram of the coating formation during ultrasonic-assisted chemical conversion; (<b>c</b>) superhydrophobic surface on Mg-Nd alloy; (<b>d</b>) anti-corrosion mechanism in artificial hand sweat solution.</p> Full article ">
37 pages, 17666 KiB  
Article
Effect of σ-Phase on the Strength, Stress Relaxation Behavior, and Corrosion Resistance of an Ultrafine-Grained Austenitic Steel AISI 321
by Vladimir I. Kopylov, Aleksey V. Nokhrin, Natalia A. Kozlova, Mikhail K. Chegurov, Mikhail Yu. Gryaznov, Sergey V. Shotin, Nikolay V. Melekhin, Nataliya Yu. Tabachkova, Ksenia E. Smetanina and Vladimir N. Chuvil’deev
Metals 2023, 13(1), 45; https://doi.org/10.3390/met13010045 - 24 Dec 2022
Cited by 6 | Viewed by 2764
Abstract
This paper reported the results of research into the effect of Equal Channel Angular Pressing (ECAP) temperature and 1-h annealing temperature on mechanical properties, stress-relaxation resistance, and corrosion resistance of austenitic steel AISI 321L with strongly elongated thin δ-ferrite particles in its microstructure. [...] Read more.
This paper reported the results of research into the effect of Equal Channel Angular Pressing (ECAP) temperature and 1-h annealing temperature on mechanical properties, stress-relaxation resistance, and corrosion resistance of austenitic steel AISI 321L with strongly elongated thin δ-ferrite particles in its microstructure. The formation of α′-martensite and fragmentation of austenite grains takes place during ECAP. Ultrafine-grained (UFG) steels demonstrate increased strength. However, we observed a reduced Hall–Petch coefficient as compared with coarse-grained (CG) steels due to the fragmentation of δ-ferrite particles. UFG steel specimens were found to have 2–3 times higher stress-relaxation resistance as compared with CG steels. For the first time, the high stress-relaxation resistance of UFG steels was shown to stem from a internal stress-relaxation mechanism, i.e., the interaction of lattice dislocations with non-equilibrium grain boundaries. Short-time 1-h annealing of UFG steel specimens at 600–800 °C was found to result in the nucleation of σ-phase nanoparticles. These nanoparticles affect the grain boundary migration, raise strength, and stress-relaxation resistance of steel but reduce the corrosion resistance of UFG steel. Lower corrosion resistance of UFG steel was shown to be related to the formation of α′-martensite during ECAP and the nucleation of σ-phase particles during annealing. Full article
(This article belongs to the Special Issue Environmentally-Assisted Degradation of Metals and Alloys)
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<p>Deformation scheme during ECAP (<b>a</b>) and general view of specimens obtained by ECAP at 150 °C (<b>b</b>).</p> Full article ">Figure 2
<p>Specimens (<b>a</b>,<b>b</b>) and test procedures (<b>c</b>,<b>d</b>). General view of specimens prepared for tension (<b>a</b>) and stress-relaxation (<b>b</b>) tests. Stress-relaxation test technique: (<b>c</b>) typical load and σ<sub>i</sub>(t) relaxation curves, (<b>d</b>) ∆σ<sub>i</sub>(σ) typical curve.</p> Full article ">Figure 3
<p>Microstructure of steel in its initial state: (<b>a</b>,<b>b</b>) δ-phase particles in steel in its initial state ((<b>a</b>)—optical microscopy; (<b>b</b>)—SEM); (<b>c</b>,<b>d</b>) microstructure of austenite grains (TEM).</p> Full article ">Figure 4
<p>Macrostructure (<b>a</b>,<b>b</b>) and microstructure (<b>c</b>–<b>f</b>) of UFG steel specimens: (<b>a</b>,<b>b</b>) macrostructure of steel specimens after the first ECAP cycle at 150 °C (<b>a</b>) and 450 °C (<b>b</b>); TEM-microphotographs and electron diffraction patterns of a steel microstructure after ECAP (N = 4) at 150 °C (<b>c</b>–<b>e</b>) and 450 °C (<b>f</b>). Nanotwins in (<b>f</b>) are marked with dash lines. (<b>b</b>) shows microscopic shear bands.</p> Full article ">Figure 5
<p>Results of XRD phase analysis of steel specimens in their initial state (line (1)) and after ECAP (lines (2)–(9)): (2) N = 1 at 150 °C; (3) N = 2 at 150 °C; (4) N = 3 at 150 °C; (5) N = 4 at 150 °C; (6) N = 1 at 450 °C; (7) N = 2 at 450 °C; (8) N = 3 at 450 °C; (9) N = 4 at 450 °C.</p> Full article ">Figure 6
<p>Dependences of mean grain sizes on 1-h annealing temperature for UFG steel specimens subjected to ECAP at T<sub>ECAP</sub> = 450 °C (<b>a</b>) and SEM images of microstructure after annealing at different temperatures: (<b>b</b>) 750 °C (ECAP, N = 4 at 450 °C); (<b>c</b>,<b>d</b>) 750 °C (ECAP, N = 3 at 450 °C); (<b>e</b>) 900 °C (ECAP, N = 3 at 450 °C); (<b>f</b>) 900 °C (ECAP, N = 4 at 450 °C).</p> Full article ">Figure 7
<p>Nucleation of σ-phase particles in UFG steels (N = 4, T<sub>ECAP</sub> = 450 °C) after heating to 800 °C and holding for 60 min. TEM-microphotographs (<b>a</b>) and electron diffraction patterns (<b>b</b>). Intensive particle nucleation regions are marked with dash lines.</p> Full article ">Figure 8
<p>Results of investigations of mechanical properties of UFG steel (T<sub>ECAP</sub> = 450 °C): (<b>a</b>) dependencies of mean grain sizes and mechanical properties of steel on the number of ECAP cycles; (<b>b</b>) dependence of the yield strength on grain size on σ<sub>y</sub>–d<sup>−1/2</sup> axes.</p> Full article ">Figure 9
<p>Dependencies of the macroelasticity stress (<b>a</b>) and yield strength (<b>b</b>) on temperatures during 1 h annealing of UFG steel.</p> Full article ">Figure 10
<p>Stress–strain tension curves for CG and UFG steel specimens at RT.</p> Full article ">Figure 11
<p>Fractographic analysis of fractures of steel specimens after tension tests at RT: (<b>a</b>,<b>b</b>) coarse-grained steel, (<b>c</b>,<b>d</b>) UFG steel (N = 4, T<sub>ECAP</sub> = 150 °C), (<b>e</b>,<b>f</b>) UFG steel (N = 4, T<sub>ECAP</sub> = 450 °C). (<b>a</b>,<b>b</b>): Zone 1—a fibrous fracture zone; Zone 2—a radial zone; Zone 3—a cut zone; (<b>d</b>): a fibrous zone consisting of a set of pits and featuring viscous fracture.</p> Full article ">Figure 12
<p>Results of stress-relaxation tests: (<b>a</b>) stress-relaxation curves for CG and UFG steel specimens; (<b>b</b>) stress-relaxation curves for UFG steel specimens (ECAP, N = 1, 150 °C) after annealing at different temperatures.</p> Full article ">Figure 13
<p>Results of electrochemical investigations of CG and UFG steel specimens: (<b>a</b>) Tafel curves lg(<span class="html-italic">i</span>)–<span class="html-italic">E</span>; (<b>b</b>) results of DLEPR tests.</p> Full article ">Figure 14
<p>Surfaces of CG steel specimens (<b>a</b>,<b>c</b>) and UFG steel specimens (N = 4, 450 °C) (<b>b</b>,<b>d</b>) after DLEPR tests (<b>a</b>,<b>b</b>) and after tests in a boiling acid solution (<b>c</b>,<b>d</b>).</p> Full article ">Figure 15
<p>The effect of annealing temperature on IGC resistance in UFG steel. The results of investigations of resistance against IGC by DLEPR: lg(<span class="html-italic">i</span>)–<span class="html-italic">E</span> curves for UFG steel after ECAP at 150 °C (<b>a</b>,<b>b</b>), 450 °C (<b>c</b>,<b>d</b>) and annealing at 600, 700, 800 °C, 1 h.</p> Full article ">Figure 16
<p>Contributions of microstructure parameters to yield strength in CG and UFG steel.</p> Full article ">Figure 17
<p>Dependences of stress-relaxation magnitude on stress applied on ln(∆σ<sub>i</sub>) − 1−σ/σ<sub>b</sub> axes: (<b>a</b>) comparison of CG and UFG steels (analysis of the data presented in <a href="#metals-13-00045-f012" class="html-fig">Figure 12</a>a); (<b>b</b>) effect of annealing temperatures on relaxation curves for UFG steel (analysis of the data presented in <a href="#metals-13-00045-f012" class="html-fig">Figure 12</a>b).</p> Full article ">Figure A1
<p>TEM images of the microstructure of UFG steel AISI 321L (ECAP, N = 4 at 450 °C) at room temperature (<b>a</b>) and after in situ heating: (<b>b</b>) 300 °C, 1 h; (<b>c</b>) 500 °C, 0 h; (<b>d</b>) 500 °C, 0.5 h; (<b>e</b>) 500 °C, 1 h; (<b>f</b>) 600 °C, 0 h; (<b>g</b>) 600 °C, 0.5 h; (<b>h</b>) 600 °C, 1 h; (<b>i</b>) 700 °C, 0 h; (<b>j</b>) 700 °C, 0.5 h; (<b>k</b>) 700 °C, 1 h; (<b>l</b>) 800 °C, 0 h; (<b>m</b>) 800 °C, 0.5 h; (<b>n</b>) 800 °C, 1 h.</p> Full article ">Figure A1 Cont.
<p>TEM images of the microstructure of UFG steel AISI 321L (ECAP, N = 4 at 450 °C) at room temperature (<b>a</b>) and after in situ heating: (<b>b</b>) 300 °C, 1 h; (<b>c</b>) 500 °C, 0 h; (<b>d</b>) 500 °C, 0.5 h; (<b>e</b>) 500 °C, 1 h; (<b>f</b>) 600 °C, 0 h; (<b>g</b>) 600 °C, 0.5 h; (<b>h</b>) 600 °C, 1 h; (<b>i</b>) 700 °C, 0 h; (<b>j</b>) 700 °C, 0.5 h; (<b>k</b>) 700 °C, 1 h; (<b>l</b>) 800 °C, 0 h; (<b>m</b>) 800 °C, 0.5 h; (<b>n</b>) 800 °C, 1 h.</p> Full article ">Figure A1 Cont.
<p>TEM images of the microstructure of UFG steel AISI 321L (ECAP, N = 4 at 450 °C) at room temperature (<b>a</b>) and after in situ heating: (<b>b</b>) 300 °C, 1 h; (<b>c</b>) 500 °C, 0 h; (<b>d</b>) 500 °C, 0.5 h; (<b>e</b>) 500 °C, 1 h; (<b>f</b>) 600 °C, 0 h; (<b>g</b>) 600 °C, 0.5 h; (<b>h</b>) 600 °C, 1 h; (<b>i</b>) 700 °C, 0 h; (<b>j</b>) 700 °C, 0.5 h; (<b>k</b>) 700 °C, 1 h; (<b>l</b>) 800 °C, 0 h; (<b>m</b>) 800 °C, 0.5 h; (<b>n</b>) 800 °C, 1 h.</p> Full article ">Figure A2
<p>The steel AISI 321L workpieces after ECAP at different temperatures.</p> Full article ">Figure A3
<p>XRD curves for specimens cut out from transverse (<b>a</b>) and longitudinal (<b>b</b>) sections of workpieces obtained by ECAP at different temperatures.</p> Full article ">
15 pages, 5943 KiB  
Article
A Survey on the Oxidation Behavior of a Nickel-Based Alloy Used in Natural Gas Engine Exhaust Valve Seats
by José Henrique Alano, Renato Luiz Siqueira, Claudio Beserra Martins Júnior, Rodrigo Silva, Guilherme dos Santos Vacchi and Carlos Alberto Della Rovere
Metals 2023, 13(1), 49; https://doi.org/10.3390/met13010049 - 24 Dec 2022
Cited by 2 | Viewed by 2147
Abstract
This study reports the oxidation behavior of a Ni-based alloy used in the manufacture of valve seats for automotive engine exhaust systems. Isothermal thermogravimetric analyses were carried out at temperatures of 660, 740, 860, and 900 °C under an oxygen atmosphere for up [...] Read more.
This study reports the oxidation behavior of a Ni-based alloy used in the manufacture of valve seats for automotive engine exhaust systems. Isothermal thermogravimetric analyses were carried out at temperatures of 660, 740, 860, and 900 °C under an oxygen atmosphere for up to 1 h. At 660 and 740 °C, only one stage was observed during the whole time studied. At this stage, the oxide layer was formed mainly by NiO + Cr2O3, following a linear oxidation law with a rate constant (Kl) on the order of magnitude of 10−6 kg/m2s and an apparent activation energy (Ea) of ~47 kJ/mol. At 860 and 900 °C, an identical first stage was observed with a transition to a different stage. In the second stage, the oxidation layer was composed of Cr2O3, and a parabolic oxidation law was followed with a rate constant (Kp) on the order of 10−8 kg2/m4s and Ea of ~128 kJ/mol. Moreover, the Ni-based alloy formed a dense and compact oxide layer after oxidation, with no apparent cavities, pores, or microcracks. Characterization techniques such as Scanning Electron Microscopy (SEM), Energy-Dispersive X-ray Spectroscopy (EDS), Fourier Transform Infrared Spectroscopy (FTIR), and Raman Spectroscopy were carried out to characterize the formed oxide layer. Full article
(This article belongs to the Special Issue Composition Design and Damage Mechanism of Crystal Superalloys)
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<p>Exhaust valve seats: (<b>a</b>) as received from the supplier and (<b>b</b>) samples prepared for thermogravimetric analyses.</p> Full article ">Figure 2
<p>Experimental procedure adopted for isothermal thermogravimetric analysis.</p> Full article ">Figure 3
<p>BSE SEM micrographs of the Ni-based alloy: (<b>a</b>,<b>b</b>) chemical contrast of the microstructure; and (<b>c</b>) EDS spectra of the different phases.</p> Full article ">Figure 4
<p>Mass variation per unit surface area as a function of the temperature for the Ni-based alloy.</p> Full article ">Figure 5
<p>Mass variation per unit surface area as a function of the exposure time for the Ni-based alloy at 660, 740, 860, and 900 °C.</p> Full article ">Figure 6
<p>Mass variation per unit surface area as a function of the exposure time for the Ni-based alloy at 660 °C: (<b>a</b>) log <span class="html-italic">y</span> versus log <span class="html-italic">t</span> to determine the oxidation rate index (<span class="html-italic">n</span>) and (<b>b</b>) <span class="html-italic">y</span> versus <span class="html-italic">t</span> to determine the linear oxidation rate constant (<span class="html-italic">K</span><sub><span class="html-italic">l</span></sub>).</p> Full article ">Figure 7
<p>Log of the mass variation per unit surface area as a function of the exposure time for the Ni-based alloy at 860 °C to determine the oxidation rate index (<span class="html-italic">n</span>) for the two different oxidation stages.</p> Full article ">Figure 8
<p>Mass variation per unit surface area as a function of the exposure time for the Ni-based alloy at 860 °C to determine the linear (<span class="html-italic">K</span><sub><span class="html-italic">l</span></sub>) and parabolic (<span class="html-italic">K</span><sub><span class="html-italic">p</span></sub>) oxidation rate constants: (<b>a</b>) first stage and (<b>b</b>) second stage.</p> Full article ">Figure 9
<p>ln of the linear (<span class="html-italic">K</span><sub><span class="html-italic">l</span></sub>) and parabolic (<span class="html-italic">K</span><sub><span class="html-italic">p</span></sub>) oxidation rate constants as a function of the inverse exposure time (<span class="html-italic">K</span><sup>−1</sup>) to obtain the activation energy (<span class="html-italic">E<sub>a</sub></span>) for the two different oxidation stages.</p> Full article ">Figure 10
<p>Ni-based alloy characterizations before and after the oxidation analyses at 660, 740, 860, and 900 °C: (<b>a</b>) FTIR and (<b>b</b>) FT-Raman.</p> Full article ">Figure 11
<p>SEM micrographs of the Ni-based alloy after the oxidation test: (<b>a</b>,<b>b</b>) 660 °C; (<b>c</b>,<b>d</b>) 740 °C; (<b>e</b>,<b>f</b>) 860 °C; and (<b>g,h</b>) 900 °C.</p> Full article ">Figure 12
<p>Cr and Ni contents in the oxide layers formed on Ni-based alloy as a function of the temperature.</p> Full article ">Figure 13
<p>SEM micrographs of the nickel-based alloy after the oxidation analysis at 900 °C with emphasis on the oxide layer formed: (<b>a</b>) lower and (<b>b</b>) higher magnification.</p> Full article ">
18 pages, 8364 KiB  
Article
Challenges and Latest Developments in Diffusion Bonding of High-Magnesium Aluminium Alloy (Al-5056/Al-5A06) to Stainless Steels
by Amir A. Shirzadi, Chengcong Zhang, Muhammad Zeeshan Mughal and Peiyun Xia
Metals 2022, 12(7), 1193; https://doi.org/10.3390/met12071193 - 13 Jul 2022
Cited by 5 | Viewed by 3273
Abstract
The aim of this work was to investigate the challenges associated with bonding Al-Mg alloys and develop a new method for bonding these alloys to steels. During an extensive R&D project, over 80 attempts, using 11 methods, were made to bond Al-6 wt.% [...] Read more.
The aim of this work was to investigate the challenges associated with bonding Al-Mg alloys and develop a new method for bonding these alloys to steels. During an extensive R&D project, over 80 attempts, using 11 methods, were made to bond Al-6 wt.% Mg alloy (Al-5056/Al-5A06) to two types of stainless steels (heat-resistant 1Cr18Ni9Ti and conventional 316). Wide ranges of temperature (500 °C to 580 °C), pressure (0.5 MPa to 10 MPa) and time (1 min to 2 h) were used when direct diffusion bonding of these alloys. Then, effects of using various interlayers and brazing foils were investigated. The interlayers used in this work were gallium, pure titanium, copper and aluminium foils, aluminium 6061 alloy sheets, aluminium-silicon brazing foils, zinc and zinc alloy foils as well as an active brazing foil (known as Incusil-ABA containing silver, copper, indium and titanium). Several complex and multi-stage processes, using up to 3 different interlayers in the same joint, were also developed and assessed. Examination and assessment of the bonded samples, including failed attempts, paved the way of developing new methods for bonding these dissimilar materials. A number of samples with tensile strengths from 200 MPa to 226 MPa were made by using complex combinations of 2 or 3 interlayers and triple-stage bonding cycles. The highest recorded bond strength was 226 MPa in the as-bonded condition. This value is above the measured yield strength (134 MPa) and about 93% of the measured ultimate strength (243 MPa) of the parent Al-Mg alloy after it was subjected to the same bonding cycle. Since the use of complex processes was not feasible for bonding large components, a simpler and more practical bonding cycle was also developed in the project. Using the simpler process, joints with tensile strengths around 90 MPa could be made. This article also sheds light on the difficulties associated with brazing and soldering aluminium alloys with a high magnesium content. Full article
(This article belongs to the Special Issue Advances in Technology and Applications of Diffusion Bonding)
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<p>Disintegration of Al-Mg alloy during solid-state bonding it to stainless steel at 560 °C. Almost no bond formation occurred despite using a temperature close to the melting point of the Al-Mg alloy.</p> Full article ">Figure 2
<p>The sudden melting of Al-Mg alloy due to migration of Mg to the surface and its eutectic reaction with liquid gallium.</p> Full article ">Figure 3
<p>Phase diagrams of Ga-Mg (left) and Al-Ga (right) show eutectic reactions may occur well below the melting point of the Al-Mg alloy (solidus 568 °C) used in this project.</p> Full article ">Figure 4
<p>A graphite crucible was used to heat the Al-Mg and Cu-cladded stainless steel. The bonding pressure was applied by the graphite plug placed on top of the assembled sample.</p> Full article ">Figure 5
<p>Decomposition of the Al-Mg alloy in the presence of copper just below 520 °C is consistent with the low temperature eutectic reactions seen in the corresponding phase diagram.</p> Full article ">Figure 6
<p>Elemental mapping of the reaction zone of the sample shown in <a href="#metals-12-01193-f005" class="html-fig">Figure 5</a>. The bias affinity of Mg to Cu led to low temperature melting at the joint interface and the decomposition of the Al-Mg alloy.</p> Full article ">Figure 7
<p>High concentration of Mg up to 12 wt.% was found in six randomly selected points in the reaction zone of the sample shown in <a href="#metals-12-01193-f005" class="html-fig">Figure 5</a>.</p> Full article ">Figure 8
<p>Magnesium from the Al-Mg alloy had a long-range diffusion into thin pure aluminium coating.</p> Full article ">Figure 9
<p>Up to 11 wt.% Mg was detected on the fracture surface of a bonded sample despite using pure aluminium coating as a diffusion barrier.</p> Full article ">Figure 10
<p>A plate of Al-6061 was inserted between Al-Mg alloy and steel to create a buffer layer and prevent migration of Mg into joint interface.</p> Full article ">Figure 11
<p>Diffusion bonding of Al-Mg to Al-6061 in vacuum resulted in decomposition of Al-Mg when bonding temperature exceeded ~550 °C, while Al 6061 remained intact.</p> Full article ">Figure 12
<p>EDX mapping shows Mg diffuses up to 0.5 mm deep into the Al-6061 interlayer.</p> Full article ">Figure 13
<p>Specimen setups for triple (top left), double (top right) and single stage (bottom) bonding cycles developed in this work; the greater the complexity of the bonding process, the higher the bond strength.</p> Full article ">Figure 14
<p>Bonds with tensile strengths above the yield point of Al-Mg (5A06) alloy were achieved using complex and multi-stage processes. Note that the relative elongations are not comparable directly—see the text for more detail.</p> Full article ">Figure 15
<p>Cross-section (SEM) and hardness profile of bonded Al-Mg (5A06) to stainless steel (1Cr18Ni9Ti).</p> Full article ">Figure 16
<p>SEM micrographs show joint continuity in a bonded Al-Mg (5A06) to stainless steel (1Cr18Ni9Ti) sample. The presence of Al-6061, as a buffer layer, prevented the migration of Mg toward the surface of stainless steel.</p> Full article ">Figure 17
<p>Migration and high concentration of Mg at joint interfaces occur when diffusion bonding or brazing Mg-rich aluminium alloys.</p> Full article ">Figure 18
<p>Distribution of main elements around an Al-6061/steel joint. Smooth transitions in the concentration profiles were achieved using the optimum bonding temperature and time developed in this work.</p> Full article ">
13 pages, 6920 KiB  
Article
Effect of Process Control Agent on Microstructures and High-Temperature Oxidation Behavior of a Nickel-Based ODS Alloy
by Zhe Mao, Jing Li, Shi Liu and Liangyin Xiong
Metals 2022, 12(6), 1029; https://doi.org/10.3390/met12061029 - 17 Jun 2022
Viewed by 2039
Abstract
Two nickel-based oxide-dispersion-strengthened (ODS) alloys supplemented with different amounts of process control agent (PCA) were prepared. The microstructures including grains and nanometric oxides and the subsequent oxidation behavior of these ODS alloys were investigated. It was found that the distribution of nanometric oxides [...] Read more.
Two nickel-based oxide-dispersion-strengthened (ODS) alloys supplemented with different amounts of process control agent (PCA) were prepared. The microstructures including grains and nanometric oxides and the subsequent oxidation behavior of these ODS alloys were investigated. It was found that the distribution of nanometric oxides in the nickel-based ODS alloy is uniform and the grains are refined by adding a proper amount of PCA in the mechanical milling, while the blocking effect on the diffusion of active elements Y, Al and Ti among powders takes place with an excessive amount of PCA, resulting in the precipitation of large-size oxides in local areas of the alloy. After oxidation in air at 1000 °C for 200 h, the oxide scales on the surface of both nickel-based ODS alloys are composed of Cr2O3. As Y-rich oxide particles are precipitated in the matrix, the thickness of the oxide scale is significantly reduced compared with non-ODS alloys. However, due to the influence of grain boundaries on the diffusion of elements, the oxide scale on the surface of an alloy with finer grain size is thicker. The oxidation resistance of ODS alloys strongly depends on the exact manufacturing process. Full article
(This article belongs to the Special Issue High Temperature Corrosion or Oxidation of Metals and Alloys)
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<p>(<b>a</b>) TEM image and size distribution of secondary phase oxides and (<b>b</b>) EBSD image and size distribution of grains in the as-HIP-treated S1.</p> Full article ">Figure 2
<p>(<b>a</b>) TEM image and size distribution of secondary phase oxides and (<b>b</b>) EBSD image and size distribution of grains in as-HIP-treated S2. The fine grain regions with secondary phase oxides are marked with red circles.</p> Full article ">Figure 3
<p>SEM secondary electron image and corresponding element distribution mapping of a TEM foil of as-HIP-treated (<b>a</b>) S1 and (<b>b</b>) S2.</p> Full article ">Figure 4
<p>Oxidation kinetics curves of two ODS alloys.</p> Full article ">Figure 5
<p>XRD patterns of S1 and S2 oxidized in air at 1000 °C for 200 h.</p> Full article ">Figure 6
<p>Cross-sectional SEM micrograph and element distribution of oxide scales on S1 after oxidation in air at 1000 °C for 200 h.</p> Full article ">Figure 7
<p>Cross-sectional SEM micrograph at (<b>a</b>) low magnification and (<b>b</b>) high magnification and element distribution of oxide scale on S2 after oxidation in air at 1000 °C for 200 h.</p> Full article ">Figure 8
<p>TEM image and size distribution of secondary phase oxides in (<b>a</b>) S1 and (<b>b</b>) S2 after oxidation in air at 1000 °C for 200 h.</p> Full article ">Figure 9
<p>Schematic diagram of chromia scale formation on the surface of (<b>a</b>) S1 and (<b>b</b>) S2.</p> Full article ">
13 pages, 3661 KiB  
Article
Effect of Austenitizing Temperature on the Work Hardening Behavior of Air-Hardening Steel LH800
by Xiang Luo, Zhenli Mi, Yanxin Wu, Yonggang Yang, Haitao Jiang and Kuanhui Hu
Metals 2022, 12(6), 1026; https://doi.org/10.3390/met12061026 - 16 Jun 2022
Cited by 2 | Viewed by 2040
Abstract
In this paper, we present the effect of austenitizing temperature on the work hardening behavior of air-hardening steel LH800 by evaluating the influence of austenitizing temperature on microstructure evolution and mechanical properties, using Hollomon, Differential Crussard–Jaoul (DC-J), and Modified C-J (M [...] Read more.
In this paper, we present the effect of austenitizing temperature on the work hardening behavior of air-hardening steel LH800 by evaluating the influence of austenitizing temperature on microstructure evolution and mechanical properties, using Hollomon, Differential Crussard–Jaoul (DC-J), and Modified C-J (MC-J) work hardening models. The results reveal that with an increase in austenitizing temperature, there is an increase in the percentage of martensite, along with an increase in the strength and hardness of the LH800 steel; on the other hand, there is a decrease in the plasticity. Austenitized at 825 °C, LH800 steel exhibits its highest strength and good plasticity, with a tensile strength of 897 MPa and an elongation of 13.6%. The comparison between the three strain hardening models revealed that the Hollomon model was the finest fit for the experimental data utilized and could illustrate the work hardening behavior of LH800 steel most suitably. This model manifests a two-stage work hardening mechanism; the first stage is related to the plastic deformation of ferrite phase, while the second stage deals with the co-deformation of ferrite and martensite/bainite phase. As austenitizing temperature increases, the work hardening ability of LH800 steel diminishes at each stage, the transition strain decreases, and the plastic deformation of martensite starts earlier. Full article
(This article belongs to the Special Issue Development and Performance Optimization of High-Strength Steels)
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<p>Dilatation vs. temperature curve of LH800 steel.</p> Full article ">Figure 2
<p>Schematic diagram of the heat treatment process of LH800 steel.</p> Full article ">Figure 3
<p>Initial microstructure of LH800 steel: (<b>a</b>) cold-rolled state and (<b>b</b>) as-supplied state. F: ferrite; P: pearlite; RD: rolling direction; ND: normal direction.</p> Full article ">Figure 4
<p>Engineering stress-strain curves of LH800 steel in cold-rolled and as-supplied state.</p> Full article ">Figure 5
<p>The microstructure of LH800 steel held at different temperatures for 15 min followed by air quenching: (<b>a</b>) 750 °C; (<b>b</b>) 775 °C; (<b>c</b>) 800 °C; (<b>d</b>) 825 °C. F: ferrite; M: martensite; GB: granular bainite; RD: rolling direction; ND: normal direction.</p> Full article ">Figure 6
<p>Mechanical properties of LH800 steel at different temperatures for 15 min followed by air quenching. (<b>a</b>) Engineering stress–strain curves of LH800. (<b>b</b>) Changes in elongation of LH800. (<b>c</b>) Changes in yield strength, tensile strength, and Vickers hardness of LH800. (<b>d</b>) The relationship between yield strength and Vickers hardness. (UE: uniform elongation; TE: total elongation; YS: yield stress; UTS: tensile stress; HV: Vickers hardness).</p> Full article ">Figure 7
<p>(<b>a</b>) The ln<span class="html-italic">σ</span> vs. ln<span class="html-italic">ε</span> plots for the Hollomon model of LH800 steel. The slope of the line segment is <span class="html-italic">n</span>; <span class="html-italic">ε</span><sub>t</sub> is the transition strain between each deformation stage. (<b>b</b>) Effect of austenitizing temperature on the work hardening behavior of each stage and the transition strain in the Hollomon model of LH800 steel.</p> Full article ">Figure 8
<p>The ln(d<span class="html-italic">σ</span>/d<span class="html-italic">ε</span>) vs. ln<span class="html-italic">ε</span> plots for the D<sub>C-J</sub> model of LH800 steel at different austenitizing temperatures: (<b>a</b>) 750 °C; (<b>b</b>) 775 °C; (<b>c</b>) 800 °C; (<b>d</b>) 825 °C.</p> Full article ">Figure 9
<p>(<b>a</b>) The ln(d<span class="html-italic">σ</span>/d<span class="html-italic">ε</span>) vs. ln<span class="html-italic">σ</span> plots for the M<sub>C-J</sub> model of the work hardening behavior of LH800 steel. (<b>b</b>) Effect of austenitizing temperature on work hardening ability of LH800 steel in the M<sub>C-J</sub> model. (<b>c</b>) The correlation coefficients (R<sup>2</sup>) of different models vary with austenitizing temperature.</p> Full article ">
12 pages, 6416 KiB  
Article
Structure, Mechanical Properties and Friction Characteristics of the Al-Mg-Sc Alloy Modified by Friction Stir Processing with the Mo Powder Addition
by Tatiana Kalashnikova, Evgeny Knyazhev, Denis Gurianov, Andrey Chumaevskii, Andrey Vorontsov, Kirill Kalashnikov, Natalya Teryukalova and Evgeny Kolubaev
Metals 2022, 12(6), 1015; https://doi.org/10.3390/met12061015 - 15 Jun 2022
Cited by 4 | Viewed by 2063
Abstract
In this study, samples of Al-Mg-Sc alloy were investigated after friction stir processing with the addition of Mo powder. Holes were drilled into 5 mm-thick aluminum alloy sheets into which Mo powder was added at percentages of 5, 10, and 15 wt%. The [...] Read more.
In this study, samples of Al-Mg-Sc alloy were investigated after friction stir processing with the addition of Mo powder. Holes were drilled into 5 mm-thick aluminum alloy sheets into which Mo powder was added at percentages of 5, 10, and 15 wt%. The workpieces with different powder contents were then subjected to four passes of friction stir processing. Studies have shown that at least three tool passes are necessary and sufficient for a uniform Mo powder distribution in the stir zone, but the number of required passes is higher with an increase in the Mo content. Due to the temperature specifics of the processing, no intermetallic compounds are formed in the stir zone, and Mo is distributed as separate particles of different sizes. The average ultimate strength of the composite materials after four passes is approximately 387 MPa in the stir zone, and the relative elongation of the material changes from 15 to 24%. The dry sliding friction test showed that the friction coefficient of the material decreases with the addition of 5 wt% Mo, but with a further increase in Mo content, returns to the original material values. Full article
(This article belongs to the Special Issue Mechanical Properties of Metals Welding Joints)
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<p>Scheme of the production of a composite based on an Al-Mg-Sc alloy with the addition of Mo powder via friction stir processing. ω—tool rotation speed; P—tool loading force; d—diameter of a hole for the powder addition.</p> Full article ">Figure 2
<p>Al-Mg-Sc alloy macrostructure after FSP with 5 wt% Mo after 1–4 passes (samples 5.1–5.4). AS—advancing side, RS—retreating side.</p> Full article ">Figure 3
<p>CT images of the processed area of Al-Mg-Sc alloy with 5 wt% Mo.</p> Full article ">Figure 4
<p>Al-Mg-Sc alloy macrostructure after FSP with 10 wt% Mo after 1–4 passes (samples 10.1–10.4). AS—advancing side, RS—retreating side.</p> Full article ">Figure 5
<p>CT images of the processed area of Al-Mg-Sc alloy with 10 wt% Mo.</p> Full article ">Figure 6
<p>Al-Mg-Sc alloy macrostructure after FSP with 15 wt% Mo after 1–4 passes (samples 15.1–15.4). AS–advancing side, RS–retreating side.</p> Full article ">Figure 7
<p>CT images of the processed area of Al-Mg-Sc alloy with 15 wt% Mo.</p> Full article ">Figure 8
<p>Light-field TEM images of the stir zone of samples (<b>a</b>) 5.4, (<b>b</b>) 10.4 and (<b>c</b>) 15.4.</p> Full article ">Figure 9
<p>Light-field TEM images of the stir zone and chemical element distribution maps of samples 15.1 (<b>a</b>–<b>d</b>) and 15.4 (<b>e</b>–<b>h</b>).</p> Full article ">Figure 10
<p>Microhardness profiles of Al-Mg-Sc alloy with Mo addition after four passes.</p> Full article ">Figure 11
<p>Static tensile test results of samples with 5 wt% Mo (<b>a</b>), 10 wt% Mo (<b>b</b>), and 15 wt% Mo (<b>c</b>), as well as a comparison of the stress–strain curves of the samples after four passes (<b>d</b>).</p> Full article ">Figure 11 Cont.
<p>Static tensile test results of samples with 5 wt% Mo (<b>a</b>), 10 wt% Mo (<b>b</b>), and 15 wt% Mo (<b>c</b>), as well as a comparison of the stress–strain curves of the samples after four passes (<b>d</b>).</p> Full article ">Figure 12
<p>Friction coefficients (μ) of samples after four passes obtained in dry sliding friction tests.</p> Full article ">
25 pages, 8807 KiB  
Article
Biofilm Development on Carbon Steel by Iron Reducing Bacterium Shewanella putrefaciens and Their Role in Corrosion
by Sachie Welikala, Saad Al-Saadi, Will P. Gates, Christopher Panter and R. K. Singh Raman
Metals 2022, 12(6), 1005; https://doi.org/10.3390/met12061005 - 12 Jun 2022
Cited by 5 | Viewed by 2696
Abstract
Microscopic, electrochemical and surface characterization techniques were used to investigate the effects of iron reducing bacteria (IRB) biofilm on carbon steel corrosion for 72 and 168 h under batch conditions. The organic nutrient availability for the bacteria was varied to evaluate biofilms formed [...] Read more.
Microscopic, electrochemical and surface characterization techniques were used to investigate the effects of iron reducing bacteria (IRB) biofilm on carbon steel corrosion for 72 and 168 h under batch conditions. The organic nutrient availability for the bacteria was varied to evaluate biofilms formed under nutritionally rich, as compared to nutritionally deficient, conditions. Focused ion beam-scanning electron microscopy (FIB-SEM) was used to investigate the effect of subsurface biofilm structures on the corrosion characteristics of carbon steel. Hydrated biofilms produced by IRB were observed under environmental scanning electron microscope (ESEM) with minimal surface preparation, and the elemental composition of the biofilms was investigated using energy dispersive spectroscopy (EDX). Attenuated total reflectance-Fourier transform infrared spectroscopy (ATR-FTIR) was used to provide information on the organic and inorganic chemical makeup of the biofilms. Electrochemical techniques employed for assessing corrosion, by open circuit potential, linear polarization and potentiodynamic polarization tests indicated no significant difference in the corrosion resistance for carbon steel in IRB-inoculated, compared to the abiotic solutions of common Postgate C after 72 and 168 h. However, the steel was found to be more susceptible to corrosion when the yeast extract was removed from the biotic environment for the 168 h test. In the absence of yeast nutrient, it is postulated that IRB received energy by transforming the protective film of Fe3+ into more soluble Fe2+ products. Full article
(This article belongs to the Special Issue Advances in Corrosion and Protection of Materials)
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<p>Growth of the IRB, <span class="html-italic">Shewanella putrefaciens</span> in B7 medium.</p> Full article ">Figure 2
<p>Quantification of total iron (Fe<sup>2+</sup> and Fe<sup>3+</sup>) in solution after 72 and 168 h in modified Postgate C medium in the presence of the IRB and carbon steel specimen.</p> Full article ">Figure 3
<p>ESEM images of hydrated biofilm formed on carbon steel by a pure IRB culture following 72 h exposure to Postgate C medium, (<b>a</b>,<b>b</b>) reduced Fe(II) oxide species, (<b>c</b>) sites of pitting attack under loose cap of corrosion product and (<b>d</b>) pit initiated at a MnS inclusion site. (<b>e</b>) IR spectra of IRB biofilm formed on carbon steel upon to exposure to pure IRB culture in Postgate C medium and abiotic Postgate C for a period of 72 h.</p> Full article ">Figure 4
<p>(<b>a</b>,<b>b</b>) FIB-SEM cross-sections of the biofilm formed on carbon steel and corrosion pitting due to pure culture of IRB in Postgate C medium following a 72 h exposure.</p> Full article ">Figure 5
<p>(<b>a</b>) ESEM and (<b>b</b>) SEM images of biofilm formed on carbon steel by a pure IRB culture (168 h exposure time to Postgate C medium). (<b>c</b>) EDS analysis of the biofilm corresponding to the ESEM image. (<b>d</b>) FIB-SEM images of the cross-section and (<b>e</b>) initial layer of fibrous EPS material produced by IRB.</p> Full article ">Figure 6
<p>(<b>a</b>) Polarization resistance (<b>a</b>) of carbon steel exposed to a pure IRB culture and abiotic conditions in Postgate C medium, and (<b>b</b>) a potentiodynamic polarization and (<b>b</b>) scan of carbon steel after 72 h exposure with and without IRB.</p> Full article ">Figure 7
<p>ESEM analysis of hydrated biofilm formed on carbon steel by a pure IRB culture following 72 h exposure to modified Postgate C medium: appearance of uniform corrosion (<b>a</b>) across the surface of carbon steel, a hollow iron oxide shell (<b>b</b>) at the surface of a corrosion pit, and needle (<b>c</b>) and globular (<b>d</b>) forms of reduced iron oxyhydroxide compound. (<b>e</b>) EDS spectra of the biofilm. (<b>f</b>) IR spectra of IRB biofilm formed on carbon steel exposed to pure IRB culture in modified Postgate C medium for a period of 72 h.</p> Full article ">Figure 7 Cont.
<p>ESEM analysis of hydrated biofilm formed on carbon steel by a pure IRB culture following 72 h exposure to modified Postgate C medium: appearance of uniform corrosion (<b>a</b>) across the surface of carbon steel, a hollow iron oxide shell (<b>b</b>) at the surface of a corrosion pit, and needle (<b>c</b>) and globular (<b>d</b>) forms of reduced iron oxyhydroxide compound. (<b>e</b>) EDS spectra of the biofilm. (<b>f</b>) IR spectra of IRB biofilm formed on carbon steel exposed to pure IRB culture in modified Postgate C medium for a period of 72 h.</p> Full article ">Figure 8
<p>FIB-SEM images of the cross-sections (<b>a</b>,<b>b</b>) of biofilm formed two areas of a carbon steel surface by pure culture of IRB in modified Postgate C medium following a 72 h exposure a detailed view (<b>c</b>) of the corroded area of the cross-section identified in <a href="#metals-12-01005-f008" class="html-fig">Figure 8</a>b at higher magnification showing channels that could allow electrolyte transport.</p> Full article ">Figure 8 Cont.
<p>FIB-SEM images of the cross-sections (<b>a</b>,<b>b</b>) of biofilm formed two areas of a carbon steel surface by pure culture of IRB in modified Postgate C medium following a 72 h exposure a detailed view (<b>c</b>) of the corroded area of the cross-section identified in <a href="#metals-12-01005-f008" class="html-fig">Figure 8</a>b at higher magnification showing channels that could allow electrolyte transport.</p> Full article ">Figure 9
<p>ESEM images of hydrated biofilm formed on carbon steel due to a pure IRB culture following 168 h exposure in the modified Postgate C medium: (<b>a</b>) overall covering (inorganic and organic components), (<b>b</b>) an iron oxyhydroxide shell covered in adhesive EPS material and (<b>c</b>) SEM image of precipitates of an iron phosphate species. (<b>d</b>) EDS analysis of the layered, plate-like deposit in <a href="#metals-12-01005-f009" class="html-fig">Figure 9</a>c.</p> Full article ">Figure 10
<p>(<b>a</b>) Change in OCP of carbon steel with time of exposure to a modified Postgate C medium with pure IRB culture and without the culture (abiotic). (<b>b</b>) Potentiodynamic polarization scans of carbon steel after 72 h exposure in modified Postgate C medium with and without IRB. (<b>c</b>) Potentiodynamic polarization scans of carbon steel after 168 h exposure in modified Postgate C medium with and without IRB.</p> Full article ">Figure 10 Cont.
<p>(<b>a</b>) Change in OCP of carbon steel with time of exposure to a modified Postgate C medium with pure IRB culture and without the culture (abiotic). (<b>b</b>) Potentiodynamic polarization scans of carbon steel after 72 h exposure in modified Postgate C medium with and without IRB. (<b>c</b>) Potentiodynamic polarization scans of carbon steel after 168 h exposure in modified Postgate C medium with and without IRB.</p> Full article ">Figure 11
<p>ESEM images of hydrated biofilm formed on carbon steel by a pure IRB culture following 72 h exposure to IRB and Postgate C without organic constituents: (<b>a</b>) amorphous corrosion product deposits containing carbon, oxygen phosphorous and iron. (<b>b</b>) SEM image of iron oxyhyrdroxide shell. FIB cross-section was carried out across the area indicated by the red line.</p> Full article ">Figure 12
<p>FIB-SEM cross-sectional analysis of biofilm formed on carbon steel by pure culture of IRB in inorganic medium following a 72 h exposure: (<b>a</b>) overview of the biofilm cross-section, and (<b>b</b>,<b>c</b>) the extent of cracking attack observed under the IRB biofilm at higher magnification.</p> Full article ">Figure 13
<p>ESEM images of hydrated biofilm formed on carbon steel by a pure IRB culture following 168 h exposure in inorganic medium: (<b>a</b>) pits covered under corrosion deposit, (<b>b</b>) micro-pits in an area of low corrosion rate, and (<b>c</b>) SEM image showing the location where the cross-sectional FIB milling was performed.</p> Full article ">Figure 14
<p>FIB-SEM cross-section of specific features of carbon steel exposed to the pure culture of IRB in the inorganic medium (i.e., Postgate C without organic components) for 168 h: (<b>a</b>) an overview of the cross-section (the circled area marks undercutting pitting attack making its way back towards the outer surface, EDS was performed in the area in the rectangle), (<b>b</b>) a close up view of the undercutting pit, and (<b>c</b>) a close up view of the side walls of the pit where biofilm is loosely attached in this area allowing electrolyte transport.</p> Full article ">Figure 15
<p>FIB-SEM cross-section of more common features of carbon steel exposed to the pure culture of IRB in the inorganic medium (i.e., Postgate C without organic components) for 168 h: (<b>a</b>) overview of a cross-section under an area of a corrosion deposit, (<b>b</b>) undercutting pitting has occurred under the biofilm, and (<b>c</b>) localised corrosion penetration and biofilm.</p> Full article ">Figure 16
<p>(<b>a</b>,<b>b</b>) Potentiodynamic polarization scans of carbon steel following 72 and 168 h exposure, with and without the IRB, to inorganic Postgate C solution.</p> Full article ">
15 pages, 6999 KiB  
Article
Effect of Two-Step High Temperature Treatment on Phase Transformation and Microstructure of V-Bearing Bainitic Steel
by Bo Lv, Dongxin Yin, Dongyun Sun, Zhinan Yang, Xiaoyan Long and Zeliang Liu
Metals 2022, 12(6), 983; https://doi.org/10.3390/met12060983 - 7 Jun 2022
Cited by 1 | Viewed by 1753
Abstract
The effects of VC precipitation on phase transformation, microstructure, and mechanical properties were studied by controlling two-step isothermal treatment, i.e., austenization followed by intercritical transformation. The results show that the bainite transformation time of 950 °C–860 °C treatment and 950 °C–848 °C treatment [...] Read more.
The effects of VC precipitation on phase transformation, microstructure, and mechanical properties were studied by controlling two-step isothermal treatment, i.e., austenization followed by intercritical transformation. The results show that the bainite transformation time of 950 °C–860 °C treatment and 950 °C–848 °C treatment is shorter than that of 950 °C single-step treatment. This is related to the isothermal ferrite transformation in the intercritical transformation range. The formation of ferrite nuclei increases the density of medium temperature bainite nucleation sites and decrease the bainite nucleation activation energy. At the same time, a large number of VC particles are precipitated. The additional VC particles provide numbers of preferential nucleation sites. The toughness of the specimen treated at 950~870 °C is improved, which is related to the large proportion of high angle grain boundaries. High angle grain boundaries can hinder crack propagation or change the direction of crack propagation. The specimen treated at 950 °C–848 °C exhibits large proportion of low angle grain boundaries, which is beneficial for the strength improvement. Full article
(This article belongs to the Special Issue High Performance Bainitic Steels)
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<p>Thermal expansion curve.</p> Full article ">Figure 2
<p>Schematic diagram of heat treatment process.</p> Full article ">Figure 3
<p>(<b>a</b>) V component in the austenite phase with the increase of the austenitization temperature; (<b>b</b>) and (<b>c</b>) Relationship between VC volume fraction and time.</p> Full article ">Figure 4
<p>(<b>a</b>) Temperature–expansion curves at different austenitization temperatures; (<b>b</b>) Enlarged figure of (<b>a</b>) located at the transformation point.</p> Full article ">Figure 5
<p>(<b>a</b>) Time–expansion curves at different temperatures; (<b>b</b>) changes in expansion amount under different processes; (<b>c</b>) isothermal curves of different processes; and (<b>d</b>) bainite transformation rates under different heat treatment process.</p> Full article ">Figure 6
<p>SEM micrographs of specimens with different heat treatment processes. (<b>a</b>) 950 °C; (<b>b</b>) 950 °C–890 °C; (<b>c</b>) 950 °C–870 °C; (<b>d</b>) 950 °C–860 °C; and (<b>e</b>) 950 °C–848 °C. (BF: bainitic ferrite; F-RA: film retained austenite; B-RA: blocky retained austenite).</p> Full article ">Figure 7
<p>EDS analysis of specimen treated at 950 °C–870 °C and 950 °C–848 °C (<b>a</b>,<b>b</b>).</p> Full article ">Figure 8
<p>TEM micrographs of specimens with different heat treatment processes. (<b>a</b>,<b>b</b>) 950 °C; (<b>c</b>,<b>d</b>) 950 °C–870 °C; and (<b>e</b>,<b>f</b>) 950 °C–848 °C.</p> Full article ">Figure 8 Cont.
<p>TEM micrographs of specimens with different heat treatment processes. (<b>a</b>,<b>b</b>) 950 °C; (<b>c</b>,<b>d</b>) 950 °C–870 °C; and (<b>e</b>,<b>f</b>) 950 °C–848 °C.</p> Full article ">Figure 9
<p>TEM map scanning images of different heat treatment processes (<b>a</b>,<b>d</b>) 950 °C; (<b>b</b>) 950 °C–870 °C; and (<b>c</b>) 950 °C–848 °C.</p> Full article ">Figure 9 Cont.
<p>TEM map scanning images of different heat treatment processes (<b>a</b>,<b>d</b>) 950 °C; (<b>b</b>) 950 °C–870 °C; and (<b>c</b>) 950 °C–848 °C.</p> Full article ">Figure 10
<p>(<b>a</b>) High magnification TEM image of 950–848 process; (<b>b</b>) High resolution micrographs image of 950 °C–848 °C process.</p> Full article ">Figure 11
<p>Engineering stress–strain curves of tested steel under different heat treatment processes.</p> Full article ">Figure 12
<p>XRD patterns of the specimens after high-temperature two-step isothermal treatment.</p> Full article ">Figure 13
<p>IPF maps ((<b>a</b>) 950 °C; (<b>b</b>) 950 °C–870 °C; and (<b>c</b>) 950 °C–848 °C) and misorientation angle distributions (<b>d</b>) of the samples after high-temperature two-step isothermal processes.</p> Full article ">Figure 13 Cont.
<p>IPF maps ((<b>a</b>) 950 °C; (<b>b</b>) 950 °C–870 °C; and (<b>c</b>) 950 °C–848 °C) and misorientation angle distributions (<b>d</b>) of the samples after high-temperature two-step isothermal processes.</p> Full article ">
10 pages, 2383 KiB  
Article
A New Strategy for Dissimilar Material Joining between SiC and Al Alloys through Use of High-Si Al Alloys
by Yongjing Yang, Ayan Bhowmik, Jin Lee Tan, Zehui Du and Wei Zhou
Metals 2022, 12(5), 887; https://doi.org/10.3390/met12050887 - 23 May 2022
Cited by 1 | Viewed by 2467
Abstract
Joining metals and ceramics plays a crucial role in many engineering applications. The current research aims to develop a simple and convenient approach for dissimilar material joining between SiC and Al alloys. In this work, Al alloys with Si contents varying from 7 [...] Read more.
Joining metals and ceramics plays a crucial role in many engineering applications. The current research aims to develop a simple and convenient approach for dissimilar material joining between SiC and Al alloys. In this work, Al alloys with Si contents varying from 7 wt.% to 50 wt.% were bonded with SiC at a high temperature of 1100 °C by a pressure-less bonding process in a vacuum furnace, and shear tests were carried out to study the bonding strength. When using low-Si Al alloys to bond with SiC, the bonding strength was very low. The bonding strength of Al/SiC joints increased significantly through the use of high-Si Al alloys with 30 wt.% and 50 wt.% Si. The shear strength achieved (28.8 MPa) is far higher than those reported previously. The remarkable improvement in bonding strength is attributed to the suppression of brittle interfacial products and reduced thermal stresses. This research provides a new strategy for joining between SiC and a wide range of Al alloys through the use of high-Si Al alloys as the interlayers. Full article
(This article belongs to the Special Issue Advances in Welding, Joining and Surface Coating Technology)
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<p>Photos showing (<b>a</b>) placement of 0.7 g Al alloy on SiC block before heating and (<b>b</b>) the sample after the completion of the bonding process.</p> Full article ">Figure 2
<p>Schematic graph of the shear test jig.</p> Full article ">Figure 3
<p>Typical load–displacement curve of (<b>a</b>) A356/SiC joint, (<b>b</b>) Al4047/SiC joint, (<b>c</b>) Al-30Si/SiC joint, and (<b>d</b>) Al-50Si/SiC joint; (<b>e</b>) comparison of the load–displacement curves of the four different Al/SiC joints.</p> Full article ">Figure 4
<p>Plot showing the effect of Si content on the bonding strength of the Al/SiC joint (error bar represents standard deviation in tested samples).</p> Full article ">Figure 5
<p>(<b>a</b>) Backscattered electron micrograph of Al4047/SiC joint showing interfacial layer and microcracks (arrowed) in the layer; (<b>b</b>) XRD pattern of Al4047/SiC joint; (<b>c</b>) SEM image showing the rectangular area for EDS in the Al4047/SiC joint; and (<b>d</b>) corresponding EDS mapping results of the area in (<b>c</b>). (The white dashed lines denote the reaction layer in the interfacial area).</p> Full article ">Figure 6
<p>(<b>a</b>) Backscattered electron micrograph of Al-50Si/SiC joint showing no obvious interfacial reaction layer or microcracks; (<b>b</b>) XRD pattern of Al-50Si/SiC joint.</p> Full article ">Figure 7
<p>Illustration of the strategy for joining between SiC and a wide range of Al alloys through use of high-Si Al alloys.</p> Full article ">
31 pages, 6813 KiB  
Article
The Holes of Zn Phosphate and Hot Dip Galvanizing on Electrochemical Behaviors of Multi-Coatings on Steel Substrates
by Thiago Duarte, Yuri A. Meyer and Wislei R. Osório
Metals 2022, 12(5), 863; https://doi.org/10.3390/met12050863 - 18 May 2022
Cited by 31 | Viewed by 2793
Abstract
The aim of this investigation is focused on the evaluation of distinctive coatings commonly applied in the automotive industry. The resulting corrosion behavior is analyzed by using electrochemical impedance spectroscopy (EIS), equivalent circuit (EC) and potentiodynamic polarization curves. The novelty concerns a comparison [...] Read more.
The aim of this investigation is focused on the evaluation of distinctive coatings commonly applied in the automotive industry. The resulting corrosion behavior is analyzed by using electrochemical impedance spectroscopy (EIS), equivalent circuit (EC) and potentiodynamic polarization curves. The novelty concerns a comparison between tricationic phosphate (TCP), cataphoretic electrodeposition (CED) of an epoxy layer, TCP + CED and HDG (hot-dip galvanized) + TCP + CED multi-coatings. Both the naturally deposited and defect-induced damage (incision) coatings are examined. The experimental impedance parameters and corrosion current densities indicate that multi-coating system (HDG + TCP + CED layers) provides better protection. Both planar and porous electrode behaviors are responsible to predict the corrosion mechanism of the majority of samples examined. Although induced-damage samples reveal that corrosion resistances decreased up to 10×, when compared with no damaged samples, the same trend of the corrosion protection is maintained, i.e., TCP < CED < TCP + CED < HDG + TCP + CED. It is also found that the same trend verified by using electrochemical parameters is also observed when samples are subjected under salt spray condition (500 h). It is also found that porous electrode behavior is not a deleterious aspect to corrosion resistance. It is more intimately associated with initial thickness coating, while corrosion resistance is associated with adhesion of the CED layer on TCP coating. The results of relative cost-to-efficiency to relative coating density ratios are associated with fact that a CED coating is necessary to top and clear coating applications and the TCP + CED and the HDG/TCP + CED coating systems exhibit the best results. Full article
(This article belongs to the Special Issue Investigations on Mechanical Properties and Corrosion Resistance of Alloys and Compounds)
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<p>A schematic representation of the panel displacing the area where the samples for EIS were withdrawn.</p> Full article ">Figure 2
<p>Experimental potentiodynamic polarization curves of three distinct coating systems (TCP, CED and TCP +CED) in an 0.5 M NaCl solution, platinum plate as counter electrode and SCE as a reference electrode, where: (<b>a</b>) the samples have no incisions (damage-induced samples), and (<b>b</b>) the samples with incisions expose the steel substrate.</p> Full article ">Figure 3
<p>Moduli of the imaginary parts of the impedances with frequencies of the coating system samples without incisions (damages induced): (<b>a</b>) TCP + CED sample, (<b>b</b>) only CED coating and (<b>c</b>) TCP coating; and when incisions are provoked: (<b>d</b>) TCP + CED, (<b>e</b>) CED and (<b>f</b>) TCP coating systems, respectively. These results are obtained when a 0.5 M NaCl solution at environmental temperature is considered.</p> Full article ">Figure 4
<p>Experimental results of EIS in Bode and Bode-phase representations of the TCP, CED and TCP + CED coating systems in 0.5 M NaCl with: (<b>a</b>) without incision and (<b>b</b>) with incision (damage exposing the CRS substrate).</p> Full article ">Figure 5
<p>Nyquist plots of the TCP, CED and TCP + CED coating systems in a 0.5 M NaCl solution without incision depicted in two distinct scale ranges: (<b>a</b>) evidencing the completed depressed semi arcs of the CED and TCP + CED coating system samples and (<b>b</b>) detailing the lowest semi arc of the TCP samples. Straight line forming 45° characterizes porous electrode behavior.</p> Full article ">Figure 6
<p>Nyquist plots of the TCP, CED and TCP + CED coating systems in a 0.5 M NaCl solution with incision depicted in two distinct scale ranges: (<b>a</b>) evidencing the completed depressed semi arcs and (<b>b</b>) detailing semi arcs in a low frequency domain, characterizing the straight line of 45° (porous electrode behavior).</p> Full article ">Figure 7
<p>Schematic representation with an equivalent circuit proposed: (<b>a</b>) substrate and TCP coating; (<b>b</b>) substrate TCP + CED coating systems; and (<b>c</b>) complete/reorganized equivalent circuit. The resulting microstructures and microconstituents of a mild steel (similar to CRS) is agreed as in the previously reported investigation [<a href="#B27-metals-12-00863" class="html-bibr">27</a>].</p> Full article ">Figure 8
<p>(<b>a</b>) Typical SEM micrographs of induced-damage incision coating system, evidencing (<b>b</b>) substrate (bottom), (<b>c</b>) the CED layer (on top), and (<b>d</b>,<b>e</b>) the TCP layer depicting Zn phosphate crystal particles.</p> Full article ">Figure 9
<p>Experimental results of the EDX (energy dispersive x-ray analysis) of (<b>a</b>) the TCP + CED coating system sample showing an EDX analysis of: (<b>b</b>) steel substrate, (<b>c</b>) the CED layer and (<b>d</b>) the TCP layer.</p> Full article ">Figure 10
<p>Typical XRD patterns of (<b>a</b>) the TCP coating system sample, (<b>b</b>) the CED sample and (<b>c</b>) the TCP + CED coating system. The JCPDS numbers are 37−0465 and 29−1427 for hopeite and phosphophyllite and 44−1415 and 34−1266 for lepidocrocite and akaganeite, respectively.</p> Full article ">Figure 11
<p>Results of (<b>a</b>) potentiodynamic polarization curves; (<b>b</b>) moduli of the imaginary parts of the impedances with frequencies; (<b>c</b>) Bode and Bode-phase plots; (<b>d</b>) Nyquist representations; and (<b>e</b>) details of Nyquist plots in a minor range of analysis of the four distinctive coating systems, i.e., the HDG, HDG/TCP, HDG/CED and HDG/TCP + CED samples, examined in 0.5 M NaCl solution at an environmental temperature. All samples contain incisions to damage the coating systems.</p> Full article ">Figure 12
<p>Schematic representation of a multi-coating system with the substrate being firstly coated with an HDG layer followed by other layers.</p> Full article ">Figure 13
<p>Typical resulting photographs of the coating systems examined: (<b>a</b>) no immersed sample showing scribe marks as ASTM B117; (<b>b</b>) the TCP + CED sample; (<b>c</b>) CED; (<b>d</b>) TCP sample coating; (<b>e</b>) HDG/TCP + CED; (<b>f</b>) HDG/CED; and (<b>g</b>) HDG/TCP samples exposed in NaCl solution during 500 h (~21 days), according to ASTM B 117. Samples without an HDG layer clearly exhibit red rust, while only the HDG/TCP sample has no depicted white rust formation.</p> Full article ">Figure 14
<p>(<b>a</b>) Variation of coating weight (CW) and (<b>b</b>) the coating thickness (CT) and CW-to-CT ratio of the all examined coating systems involving two distinct substrate, i.e., CRS and HDG. (<b>c</b>) correlation of η with the coating density of the examined samples.</p> Full article ">
13 pages, 3133 KiB  
Article
An Improved Approach to Manufacture Carbon Nanotube Reinforced Magnesium AZ91 Composites with Increased Strength and Ductility
by Samaneh Nasiri, Guang Yang, Erdmann Spiecker and Qianqian Li
Metals 2022, 12(5), 834; https://doi.org/10.3390/met12050834 - 13 May 2022
Cited by 5 | Viewed by 2346
Abstract
Multiwalled carbon nanotubes (MWCNTs) are decorated with Pt nanoparticles by a “layer-by-layer” approach using poly (sodium 4-styrene sulfonate) (PSS) and poly (diallyl dimethylammonium chloride) (PDDA). Transmission electron microscopy (TEM) images and Energy Dispersive X-ray (EDX) analysis of the samples confirm Pt deposition on [...] Read more.
Multiwalled carbon nanotubes (MWCNTs) are decorated with Pt nanoparticles by a “layer-by-layer” approach using poly (sodium 4-styrene sulfonate) (PSS) and poly (diallyl dimethylammonium chloride) (PDDA). Transmission electron microscopy (TEM) images and Energy Dispersive X-ray (EDX) analysis of the samples confirm Pt deposition on surfaces of CNTs. Dispersibility and dispersion stability of MWCNTs in the solvents are enhanced when MWCNTs are coated with Pt nanoparticles. Mg AZ91 composites reinforced with MWCNTs are then produced by a melt stirring process. Compression tests of the composites show that adding 0.05% wt Pt-coated MWCNTs in AZ91 improves the composite’s mechanical properties compared to the pure AZ91 and pristine MWCNT/AZ91. Fracture surface analysis of the composite using a scanning electron microscope (SEM) shows individual pulled out MWCNTs in the case of the Pt-coated MWCNT/AZ91 composites. This finding can be attributed to the uniform dispersion of Pt-coated MWCNTs in Mg due to the improved wettability of Pt-coated MWCNTs in Mg melts. The study of the pull-out behaviour of pristine and Pt-coated CNTs from an Mg matrix using molecular dynamics simulation supports this interpretation. Full article
(This article belongs to the Special Issue Advances in High Strength–Ductility Synergy Materials)
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<p>Schematic drawing of the melt stirring furnace used in this study.</p> Full article ">Figure 2
<p>TEM image of overview of Pt coated MWCNTs (containing 30 wt% of Pt), containing Pt-coated MWCNTs and polymer residue with Pt clusters; (<b>a</b>) Overview TEM image of Pt coated MWCNTs; (<b>b</b>) TEM image of Pt coating on the surface of individual MWCNTs; (<b>c</b>) EDX analysis of selected area on the surface.</p> Full article ">Figure 3
<p>Dispersion ability and stability of Pt coated CNTs and pristine CNTs. Same weight (1.14 mg) of Pt coated CNT (coating ratio has been considered) and pristine CNTs in same amount of ethanol (20 mL). Sonicated for 15 min.</p> Full article ">Figure 4
<p>Typical stress-strain curves of pure AZ91, 0.05 wt% uncoated MWCNT/AZ91 composite, and 0.05 wt% Pt-MWCNT/AZ91 composite.</p> Full article ">Figure 5
<p>Comparison of (<b>a</b>) compression strain at failure (CSF), (<b>b</b>) yield strength (YS), (<b>c</b>) density, and (<b>d</b>) ultimate compressive strength (UCS) between pure AZ91, 0.05 wt% MWCNT/AZ91 composite, and 0.05 wt% Pt-coated MWCNT/AZ91 composite.</p> Full article ">Figure 6
<p>Fracture surface of (<b>a</b>) Pt coated CNT/AZ91 composite and (<b>b</b>) raw CNT/AZ91 composite after compression test. Arrow points at the individual CNT pulled out of the fracture surface after failure.</p> Full article ">Figure 7
<p>MD simulation of CNT pull-out from Mg matrix; (<b>a</b>) pristine CNT in Mg matrix, (<b>b</b>) Pt-coated in Mg matrix; inset pictures show the cross section of the configuration during the pull-out simulation. Blue, red, and yellow atoms at the insets represent Mg, C, and Pt atoms, respectively.</p> Full article ">
14 pages, 4127 KiB  
Article
Real-Time Quality Monitoring of Laser Cladding Process on Rail Steel by an Infrared Camera
by Pornsak Srisungsitthisunti, Boonrit Kaewprachum, Zhigang Yang and Guhui Gao
Metals 2022, 12(5), 825; https://doi.org/10.3390/met12050825 - 11 May 2022
Cited by 9 | Viewed by 4692
Abstract
Laser cladding is considered to be a highly complex process to set up and control because it involves several parameters, such as laser power, laser scanning speed, powder flow rate, powder size, etc. It has been widely studied for metal-part coating and repair [...] Read more.
Laser cladding is considered to be a highly complex process to set up and control because it involves several parameters, such as laser power, laser scanning speed, powder flow rate, powder size, etc. It has been widely studied for metal-part coating and repair due to its advantage in controllable deposited materials on a small target substrate with low heat-affected distortion. In this experiment, laser cladding of U75V and U20Mn rail steels with Inconel 625 powder was captured by an infrared camera with image analysis software to monitor the laser cladding process in order to determine the quality of the cladded substrates. The cladding temperature, thermal gradient, spot profile, and cooling rate were determined from infrared imaging of the molten pool. The results showed that cladding temperature and molten pool’s spot closely related to the laser cladding process condition. Infrared imaging provided the cooling rate from a temperature gradient which was used to correctly predict the microhardness and microstructure of the HAZ region. This approach was able to effectively detect disturbance and identify geometry and microstructure of the cladded substrate. Full article
(This article belongs to the Special Issue High Performance Bainitic Steels)
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<p>Experimental setup for process monitoring of laser cladding.</p> Full article ">Figure 2
<p>Online monitoring of laser-cladding experiment’s equipment.</p> Full article ">Figure 3
<p>Laser-cladding process and infrared image of the laser hot spot on the substrate taken by an IR camera.</p> Full article ">Figure 4
<p>Image-analysis steps with LabVIEW to determine the molten pool’s size from infrared image.</p> Full article ">Figure 5
<p>Parametric study of laser cladding on (<b>a</b>,<b>b</b>) U75V and (<b>c</b>,<b>d</b>) U20Mn rail steels. For all cases, the laser cladded four times (four layers) on the substrates. (<b>a</b>,<b>c</b>) Effect of scanning speed and (<b>b</b>,<b>d</b>) effect of laser power.</p> Full article ">Figure 6
<p>SEM images from cross-section of cladded rail steels at different conditions. For all cases, the laser cladded four times (four layers) on the substrates.</p> Full article ">Figure 7
<p>Temperature converted from infrared images during laser cladding for different laser powers and scanning speeds.</p> Full article ">Figure 8
<p>Laser hot-spot profiles during laser cladding on U75V substrate for different laser powers and scanning speeds.</p> Full article ">Figure 9
<p>Temperature profile of laser hot spots plotted along the dashed lines from <a href="#metals-12-00825-f008" class="html-fig">Figure 8</a>.</p> Full article ">Figure 10
<p>Microstructure of the interface zone between the cladded material and the heat-affected zone of substrate on U75V rail steels with different scanning speeds.</p> Full article ">Figure 11
<p>Correlation between IR image width (molten pool’s width) and clad width from experiment.</p> Full article ">Figure 12
<p>Correlation between IR image width (W<sub>IR</sub>) and actual clad width (W<sub>Clad</sub>) from experiment.</p> Full article ">Figure 13
<p>Laser cladding on U75V with laser power 500 W scan speed 1 mm/s with 10-layer deposition resulted in porosities. (<b>a</b>) Cladded substrate, (<b>b</b>) substrate’s cross-section, (<b>c</b>) IR image during laser cladding, (<b>d</b>,<b>e</b>) SEM images at the interface region and (<b>f</b>) SEM at the HAZ region.</p> Full article ">
14 pages, 9661 KiB  
Article
Microstructure Characteristics and Wear Performance of a Carburizing Bainitic Ferrite + Martensite Si/Al-Rich Gear Steel
by Yanhui Wang, Qingsong He, Qian Yang, Dong Xu, Zhinan Yang and Fucheng Zhang
Metals 2022, 12(5), 822; https://doi.org/10.3390/met12050822 - 10 May 2022
Cited by 4 | Viewed by 2433
Abstract
In this paper, a new low-carbon alloy gear steel is designed via Si/Al alloying. The carburizing and austempering, at a temperature slightly higher than the martensitic transformation point (Ms) of the surface and much lower than the Ms of the core, for different [...] Read more.
In this paper, a new low-carbon alloy gear steel is designed via Si/Al alloying. The carburizing and austempering, at a temperature slightly higher than the martensitic transformation point (Ms) of the surface and much lower than the Ms of the core, for different times, were carried out on the newly designed gear steel. After heat treatment, a series of different microstructures (superfine bainitic ferrite + retained austenite, superfine bainitic ferrite + martensite + retained austenite, and martensite + retained austenite) were obtained on the surface, whilst the low-carbon lath martensitic microstructure was obtained in the core. The microstructure of the surface was examined using optical microscopy (OM), scanning electron microscopy (SEM), and transmission electron microscopy (TEM). The phase composition was analyzed using X-ray diffraction (XRD). The hardness and wear resistance of the surface as well as the hardness distribution of carburizing layer of the samples with different microstructures were studied. The results show that the Si/Al-rich gear steel, after carburizing and austempering at 200 °C for 8 h, not only has excellent mechanical properties but also has high wear resistance, which meets the technical requirements of heavy-duty gear steel. The research work in this paper can provide a data reference for the application of carburized steel with mixed microstructures of bainitic ferrite and martensite in the design of heavy-duty gear. Full article
(This article belongs to the Special Issue High Performance Bainitic Steels)
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<p>Schematic diagrams of carburizing and heat treatment processes.</p> Full article ">Figure 2
<p>The optical micrographs of carburized surface austempered at 200 °C: (<b>a</b>) 2 h, (<b>b</b>) 4 h, (<b>c</b>) 8 h, (<b>d</b>) 12 h, (<b>e</b>) 24 h, and (<b>f</b>) 48 h.</p> Full article ">Figure 3
<p>The optical micrographs of carburized surface austempered at 230 °C: (<b>a</b>) 2 h, (<b>b</b>) 4 h, (<b>c</b>) 8 h, (<b>d</b>) 12 h, (<b>e</b>) 24 h, and (<b>f</b>) 48 h.</p> Full article ">Figure 4
<p>The SEM images of the carburized surface austempered at 200 °C: (<b>a</b>) 2 h, (<b>b</b>) 8 h, and (<b>c</b>) 48 h. Notes: BF—bainitic ferrite, RA—retained austenite, and M—martensite.</p> Full article ">Figure 5
<p>The SEM images of the carburized surface austempered at 230 °C: (<b>a</b>) 2 h and (<b>b</b>) 48 h. Notes: BF—bainitic ferrite, RA—retained austenite, and M—martensite.</p> Full article ">Figure 6
<p>The TEM micrographs of the carburized surface after isothermal treatment: (<b>a</b>) 200 °C × 8 h, (<b>b</b>) 200 °C × 48 h, (<b>c</b>) 230 °C × 8 h, and (<b>d</b>) 230 °C × 48 h. Notes: BF—bainitic ferrite, RA—retained austenite, and M—martensite.</p> Full article ">Figure 7
<p>XRD patterns of carburized surface austempered at 200 °C (<b>a</b>) and 230 °C (<b>b</b>).</p> Full article ">Figure 8
<p>Carbon content distribution curve of the specimen after carburizing.</p> Full article ">Figure 9
<p>Vickers hardness distributions of carburized layer austempered at 200 °C (<b>left</b>) and 230 °C (<b>right</b>) for different times.</p> Full article ">Figure 10
<p>Relationship curves between weight loss and wear time of the carburized surface austempered at 200 °C.</p> Full article ">Figure 11
<p>Wear morphology of carburized surface austempered at 200 °C: (<b>a</b>) 2 h, (<b>b</b>) 8 h, (<b>c</b>) 48 h.</p> Full article ">Figure 12
<p>XRD patterns of the surface at 200 °C for 8 h before and after wear.</p> Full article ">Figure 13
<p>Optical micrographs of the transition layer (<b>a</b>) and core (<b>b</b>) of the experimental steel, after carburizing and austempering at 200 °C for 8 h.</p> Full article ">
21 pages, 5807 KiB  
Article
Observing the Effect of Grain Refinement on Crystal Growth of Al and Mg Alloys during Solidification Using In-Situ Neutron Diffraction
by Abdallah Elsayed, Francesco D’Elia, Comondore Ravindran and Dimitry Sediako
Metals 2022, 12(5), 793; https://doi.org/10.3390/met12050793 - 4 May 2022
Cited by 4 | Viewed by 2214
Abstract
The present research uses in-situ neutron diffraction to examine the effect of grain refinement on grain growth during solidification of Al-5 wt.% Cu and Mg-5 wt.% Zn alloys. The alloys were grain refined through additions of Al-5Ti-1B and Zr, respectively. The in-situ neutron [...] Read more.
The present research uses in-situ neutron diffraction to examine the effect of grain refinement on grain growth during solidification of Al-5 wt.% Cu and Mg-5 wt.% Zn alloys. The alloys were grain refined through additions of Al-5Ti-1B and Zr, respectively. The in-situ neutron diffraction experiments were carried out by heating the alloys to temperatures above the liquidus and subsequently cooling in 5 or 10 °C temperature steps to temperatures below solidus, while being irradiated by thermal neutrons. With the addition of grain refiners, grain size reductions of 92% were observed for both the Al-5 wt.% Cu and Mg-5 wt.% Zn alloys. The refined and unrefined Al-5 wt.% Cu alloys contained α-Al with Al2Cu along the grain boundary regions. Differences in Al2Cu morphology were observed in the grain refined alloys. The Mg-5 wt.% Zn alloy contained MgZn intermetallic phases with primary Mg. The refined Mg-5 wt.% Zn-0.7 wt.% Zr alloy contained Mg, MgZn and Zn2Zr phases. In-situ neutron diffraction enabled quantification of individual plane solid fraction growth for the α-Al and Al2Cu phases in the Al-Cu alloys, and for α-Mg in the Mg alloys. For the unrefined Al-5 wt.% Cu, the coarse microstructure resulted in a rapid solid fraction rise at temperatures just below liquidus followed by a gradual increase in solid fraction until the sample was fully solid. The grain-refined Al-5 wt.% Cu alloys showed a columnar to equiaxed microstructure transition and a more gradual growth in fraction solid throughout solidification. For the Mg-5 wt.% Zn alloy, the more packed (0002) and (101¯1) α-Mg plane intensities grew at a slower rate than the (101¯0) plane intensity, resulting in an irregular grain structure. With the addition of the Zr grain refiner, the Mg-5 wt.% Zn-0.7 wt.% Zr alloy had (101¯0), (0002) and (101¯1) planes intensities all increasing at similar rates, especially at the early stages of solidification. FactSage™ (version 6.4, Montréal, QC, Canada) equilibrium solidification models followed the fraction solid curves developed by tracking the fastest growing planes of the Mg alloys. Full article
(This article belongs to the Special Issue Material Properties—Superalloys, Ferrous and Lightweight Alloys and Metal Matrix Composites)
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<p>Grain structures of unrefined (<b>a</b>) Al-Cu and (<b>b</b>) Mg-Zn prior to in-situ neutron diffraction experiments.</p> Full article ">Figure 2
<p>Grain structures of (<b>a</b>) Al-Cu, (<b>b</b>) Al-Cu-02Ti, (<b>c</b>) Al-Cu-05Ti, (<b>d</b>) Mg-Zn and (<b>e</b>) Mg-Zn-Zr after in-situ neutron diffraction experiments.</p> Full article ">Figure 3
<p>SEM images of (<b>a</b>) Al-Cu, (<b>b</b>) Al-Cu-05Ti, (<b>c</b>) Mg-Zn and (<b>d</b>) Mg-Zn-Zr alloys after in-situ neutron diffraction experiments.</p> Full article ">Figure 4
<p>In-situ neutron diffraction plots of (<b>a</b>) Al-Cu and (<b>b</b>) Mg-Zn at different temperatures (°C, select isotherms omitted for clarity).</p> Full article ">Figure 5
<p>In-situ neutron diffraction plot of (<b>a</b>) Al-Cu alloy <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <mrow> <mn>111</mn> </mrow> <mo>)</mo> </mrow> </mrow> </semantics></math> plane and (<b>b</b>) Al<sub>2</sub>Cu <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <mrow> <mn>310</mn> </mrow> <mo>)</mo> </mrow> </mrow> </semantics></math> plane in Al-Cu alloy at different temperatures (°C, select isotherms omitted for clarity).</p> Full article ">Figure 6
<p>In-situ neutron diffraction plot of Mg-Zn-Zr alloy <math display="inline"><semantics> <mrow> <mrow> <mo>(</mo> <mrow> <mn>10</mn> <mover accent="true"> <mn>1</mn> <mo>¯</mo> </mover> <mn>0</mn> </mrow> <mo>)</mo> </mrow> </mrow> </semantics></math> plane at different temperatures (°C, select isotherms omitted for clarity).</p> Full article ">Figure 7
<p>Fraction solid (%) versus temperature of (<b>a</b>) α-Al and Al<sub>2</sub>Cu for Al-Cu, (<b>b</b>) α-Al for Al-Cu-02Ti, and (<b>c</b>) α-Al and Al<sub>2</sub>Cu for Al-Cu-05Ti.</p> Full article ">Figure 8
<p>Fraction solid (%) versus temperature of (<b>a</b>) α-Mg for Mg-Zn and (<b>b</b>) α-Mg for Mg-Zn-Zr.</p> Full article ">Figure 9
<p>Grain size versus 1/Q for Al and Mg alloys.</p> Full article ">
8 pages, 5218 KiB  
Article
Electrochemical Corrosion Resistance of Mg Alloy ZK60 in Different Planes with Respect to Extrusion Direction
by G. Keerthiga, Dandapani Vijayshankar, MJNV Prasad, Mirco Peron, Jafar Albinmousa and RK Singh Raman
Metals 2022, 12(5), 782; https://doi.org/10.3390/met12050782 - 30 Apr 2022
Cited by 6 | Viewed by 2258
Abstract
The electrochemical corrosion resistance of a Mg-Zn-Zr alloy, ZK60, in different planes with respect to the extrusion direction was investigated in 3.5 wt% NaCl. The motivation of this study lies in the influence of extrusion on the grain size, texture and precipitation characteristics [...] Read more.
The electrochemical corrosion resistance of a Mg-Zn-Zr alloy, ZK60, in different planes with respect to the extrusion direction was investigated in 3.5 wt% NaCl. The motivation of this study lies in the influence of extrusion on the grain size, texture and precipitation characteristics of magnesium alloys, and the profound role of these characteristics in the corrosion resistance of the alloys. Corrosion resistance was found to be considerably superior in the plane transverse to the extrusion direction (TD) than in the extrusion direction (ED) or normal to the extrusion direction (ND). The difference in the corrosion resistance was attributed to the variations in microstructural features in the TD, ED and ND directions. Full article
(This article belongs to the Special Issue Investigations on Mechanical Properties and Corrosion Resistance of Alloys and Compounds)
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<p>Schematic depiction of different orientations (ED, ND and TD) with respect to the extrusion direction of a ZK60 extruded bar.</p> Full article ">Figure 2
<p>(<b>a</b>–<b>c</b>) Optical micrographs of the microstructure of ZK60 alloy in different planes (i.e., TD, ND and ED, respectively) with respect to the extrusion direction, and (<b>d</b>–<b>f</b>) the corresponding grain size distributions.</p> Full article ">Figure 3
<p>XRD spectrum of the extruded ZK60 alloy.</p> Full article ">Figure 4
<p>Potentiodynamic polarisation behaviour of the extruded ZK60 alloy in different planes with respect to the extrusion direction (i.e., ED, ND and TD) in 3.5 wt% NaCl.</p> Full article ">Figure 5
<p>Comparison of corrosion current density of the heat-treated (450 °C/1 h) ED (i.e., ED + HT) with TD, ND and ED alloy in 3.5 wt% NaCl. Error bars are the standard deviation of the replicated measurements.</p> Full article ">Figure 6
<p>SEM images of the microstructure of (<b>a</b>) ED and (<b>b</b>) ED + HT.</p> Full article ">
3 pages, 162 KiB  
Editorial
Recent Developments in Medium and High Manganese Steels
by Colin P. Scott
Metals 2022, 12(5), 743; https://doi.org/10.3390/met12050743 - 27 Apr 2022
Cited by 2 | Viewed by 2396
Abstract
A huge amount of intellectual effort is currently being devoted to the study of medium and high manganese steels due to the diverse and impressive mechanical properties that can be achieved with these steels [...] Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
13 pages, 5018 KiB  
Article
Investigation on Corrosion Resistance Properties of 17-4 PH Bound Metal Deposition As-Sintered Specimens with Different Build-Up Orientations
by Pietro Forcellese, Tommaso Mancia, Michela Simoncini and Tiziano Bellezze
Metals 2022, 12(4), 588; https://doi.org/10.3390/met12040588 - 30 Mar 2022
Cited by 9 | Viewed by 3633
Abstract
Additive manufacturing is a promising and emerging technology that can transform the global manufacturing and logistics by cutting costs and times of production. Localized corrosion resistance properties of 0°, 45°, and 90° build-up orientations of 17-4 PH as-sintered samples, manufactured by means of [...] Read more.
Additive manufacturing is a promising and emerging technology that can transform the global manufacturing and logistics by cutting costs and times of production. Localized corrosion resistance properties of 0°, 45°, and 90° build-up orientations of 17-4 PH as-sintered samples, manufactured by means of Bound Metal Deposition (BMD), have been investigated by electrochemical and morphological investigations. The cyclic potentiodynamic polarization curves and the open circuit potential monitoring, together with potential drop analysis, revealed that the BMD localized corrosion resistance properties were lowered if compared to a wrought 17-4 PH: a characteristic anodic behavior and many drops in potential were recorded for BMD, whilst the wrought specimens presented a typical passive behavior with pitting corrosion. Morphological investigations by scanning electron microscopy and energy-dispersive X-ray analysis revealed the presence of porosities and defects, especially for the 90° build-up orientation, and inclusions of SiO2. The 45° build-up orientation showed the best corrosion resistance properties among all the BMD specimens, even though defects and porosities were observed, suggesting that their morphology and geometry influenced the overall corrosion behavior. Full article
(This article belongs to the Special Issue Corrosion and Protection of Stainless Steels)
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<p>Build-up orientations: (<b>a</b>) 0°; (<b>b</b>) 45°; (<b>c</b>) 90°.</p> Full article ">Figure 2
<p>The 17-4 PH BMD as-sintered plate specimens: (<b>a</b>) 0°; (<b>b</b>) 45°; (<b>c</b>) 90°.</p> Full article ">Figure 3
<p>SEM images of the surface area of 17-4 PH BMD specimens: (<b>a</b>) 0°; (<b>b</b>) 45°; (<b>c</b>) 90°. Red arrows point at porosities and defects.</p> Full article ">Figure 4
<p>SEM-EDX acquisition of a SiO<sub>2</sub> inclusion on a 90° specimen.</p> Full article ">Figure 5
<p>CPP curve of 17-4 PH wrought stainless steel sample in NaCl 0.35 wt%. Arrows point at the direction of the scan.</p> Full article ">Figure 6
<p>CPP curves of 17-4 PH BMD stainless-steel samples 0° (red), 45° (blue), and 90° (green) build-up orientations. Arrows point at the direction of the scan.</p> Full article ">Figure 7
<p>OM images of 17-4 PH BMD specimens: (<b>a</b>,<b>b</b>) 0°; (<b>c</b>) 45°; (<b>d</b>,<b>e</b>) 90°. Red lines were signed after CPP to underline the border of the exposed area to the testing solution.</p> Full article ">Figure 8
<p>OCP of 0°, 45°, and 90° build-up orientations and wrought 17-4 PH specimens: (<b>a</b>) 108 h OCP monitoring; (<b>b</b>) enlargement of the first 6 h of OCP monitoring.</p> Full article ">Figure 9
<p>OM image of corrosion products on 90° sample after the OCP investigation.</p> Full article ">Figure 10
<p>Potential drops analysis of the OCP curves: (<b>a</b>) 0°; (<b>b</b>) 45°; (<b>c</b>) wrought.</p> Full article ">
24 pages, 2796 KiB  
Review
A Review of Corrosion under Insulation: A Critical Issue in the Oil and Gas Industry
by Qing Cao, Thunyaluk Pojtanabuntoeng, Marco Esmaily, Sebastian Thomas, Michael Brameld, Ayman Amer and Nick Birbilis
Metals 2022, 12(4), 561; https://doi.org/10.3390/met12040561 - 25 Mar 2022
Cited by 28 | Viewed by 12514
Abstract
Corrosion under insulation (CUI) is defined as any form of external corrosion that occurs on the underlying metal beneath insulated equipment, due to water ingress through the insulation layer. This type of corrosion is frequently observed in oil and gas production, where insulated [...] Read more.
Corrosion under insulation (CUI) is defined as any form of external corrosion that occurs on the underlying metal beneath insulated equipment, due to water ingress through the insulation layer. This type of corrosion is frequently observed in oil and gas production, where insulated piping is prevalent, and has historically remained a predominant materials integrity issue. The prediction and direct visualisation of CUI are challenging tasks because of the coverage of the insulation layer(s) and any external jacketing or cladding. Several factors, including the local/ambient environment, system design, and the piping installation process, can influence how CUI initiates and propagates. In this review, CUI background, CUI monitoring, and CUI mitigation strategies are discussed. Full article
(This article belongs to the Special Issue Predicting and Preventing Corrosion of Metals in Different Environments)
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<p>Schematic of a typical insulated pipe, which consists of an external cladding/jacketing, insulation material, a protective coating layer, and a steel pipe itself.</p> Full article ">Figure 2
<p>Schematic of “non-contact” or “distance-insulation” system [<a href="#B47-metals-12-00561" class="html-bibr">47</a>].</p> Full article ">Figure 3
<p>(<b>a</b>) Schematic of metal ions (Fe<sup>3+</sup>, Ca<sup>2+</sup>, K<sup>+,</sup> and Mg<sup>2+</sup>) that leach out from the insulation and (<b>b</b>) X-ray diffraction characterisation of corrosion products formed upon the steel pipe during its interaction with moist insulation, analysed by Cao et al. [<a href="#B49-metals-12-00561" class="html-bibr">49</a>].</p> Full article ">Figure 4
<p>CUI failure in a refinery plant: (<b>a</b>) pipework fracture detail with a huge hole and (<b>b</b>) the installation of walkway bracket touching the pipework. Reprinted with permission from Ref. [<a href="#B23-metals-12-00561" class="html-bibr">23</a>]. Copyright 2013 Elsevier.</p> Full article ">Figure 5
<p>Effect of temperature on steel corrosion in water [<a href="#B6-metals-12-00561" class="html-bibr">6</a>].</p> Full article ">Figure 6
<p>Four major steps involved in a CUI inspection process: (<b>a</b>) visual inspection, (<b>b</b>) partial or complete removal of insulation, (<b>c</b>) ultrasonic testing and (<b>d</b>) reinstallation of insulation material. Reprinted with authors’ permission from Ref. [<a href="#B64-metals-12-00561" class="html-bibr">64</a>].</p> Full article ">Figure 7
<p>Water detection using an infrared thermal imaging technique. The bright area in the thermal image reveals the presence of water. Examples of obtained thermal images taken during the daytime (<b>a</b>,<b>c</b>); and taken during the night time (<b>b</b>,<b>d</b>). Reprinted with permission from Ref. [<a href="#B66-metals-12-00561" class="html-bibr">66</a>]. Copyright 2012 Society of Photo-Optical Instrumentation Engineers (SPIE).</p> Full article ">Figure 7 Cont.
<p>Water detection using an infrared thermal imaging technique. The bright area in the thermal image reveals the presence of water. Examples of obtained thermal images taken during the daytime (<b>a</b>,<b>c</b>); and taken during the night time (<b>b</b>,<b>d</b>). Reprinted with permission from Ref. [<a href="#B66-metals-12-00561" class="html-bibr">66</a>]. Copyright 2012 Society of Photo-Optical Instrumentation Engineers (SPIE).</p> Full article ">Figure 8
<p>Radiography technique to image the wall thickness loss using C-arm equipment. (<b>a</b>) Operator holds the C-arm machine to check potential piping damage; (<b>b</b>) Radiography images showing pipe wall thickness loss. Reprinted with authors’ permission from Ref. [<a href="#B64-metals-12-00561" class="html-bibr">64</a>].</p> Full article ">Figure 9
<p>Schematic of eddy current used to inspect the occurrence of corrosion under insulation on a conductive steel pipe.</p> Full article ">Figure 10
<p>The procedure of employing electrochemical noise data as input to build a random forest model for the identification of CUI types. (<b>a</b>) Electrochemical noise data collected from an experimental CUI rig. (<b>b</b>) Recurrence plot produced from electrochemical noise data input. (<b>c</b>) Steps for training a random forest model to identify the form of CUI. Reprinted with permission from Ref. [<a href="#B94-metals-12-00561" class="html-bibr">94</a>]. Copyright 2020 Springer Nature.</p> Full article ">Figure 11
<p>The spectrogram of reflected signals collected from a corrosion sensor, which indicates the location of the corrosion initiation part of a pipe. Reprinted with permission from Ref. [<a href="#B98-metals-12-00561" class="html-bibr">98</a>] CorrosionRadar.</p> Full article ">Figure 12
<p>Apparatus sketch for CUI detection equipment that consists of an electrically conductive excitation coil and an electrical power source. Reprinted with authors’ permission from Ref. [<a href="#B105-metals-12-00561" class="html-bibr">105</a>].</p> Full article ">Figure 13
<p>Schematic of thermal sprayed aluminium (TSA) coating process, which involves projecting small molten particles onto a prepared surface and achieving a continuous coating with strong surface adhesion.</p> Full article ">Figure 14
<p>Coating degradation in the form of chained pitting. Reprinted with permission from Ref. [<a href="#B109-metals-12-00561" class="html-bibr">109</a>] Copyright 2013 Springer Nature.</p> Full article ">Figure 15
<p>CUI mitigation using cathodic protection, (<b>a</b>,<b>b</b>) a proof-of-concept that cathodic protection using sacrificial Zn can mitigate CUI on mild steel, in which (<b>a</b>) represents mild steel surface morphologies in the absence of Zn protection, (<b>b</b>) represents mild steel surface morphologies with Zn protection and (<b>c</b>) represents the achievable cathodic protection area (represented by the Cu plated region) on steel test plates under different Zn to MS contact modes (0 cm, 2 cm, and 4 cm contact distances) and different Zn to MS area ratios (1:10 and 1:40). Samples and data were analysed by Cao et al. [<a href="#B121-metals-12-00561" class="html-bibr">121</a>].</p> Full article ">
15 pages, 4157 KiB  
Article
Effect of Diffusion on Dissimilar Welded Joint between Al0.8CoCrFeNi High-Entropy Alloy and S235JR Structural Steel
by Ionelia Voiculescu, Victor Geanta, Elena Violeta Stefanescu, George Simion and Elena Scutelnicu
Metals 2022, 12(4), 548; https://doi.org/10.3390/met12040548 - 24 Mar 2022
Cited by 9 | Viewed by 3003
Abstract
This research focused on the investigation of the metallurgical behavior of the Al0.8CoCrFeNi high-entropy alloy and S235JR structural steel, welded with (Ni, Fe)-rich filler metal, by the Gas Tungsten Arc Welding (GTAW) method. The electric arc and the welding pool were [...] Read more.
This research focused on the investigation of the metallurgical behavior of the Al0.8CoCrFeNi high-entropy alloy and S235JR structural steel, welded with (Ni, Fe)-rich filler metal, by the Gas Tungsten Arc Welding (GTAW) method. The electric arc and the welding pool were protected against the contamination with gases from the environment, by employing high-purity Ar 4.8 inert gas that plays an important role in reducing the oxidation effects and the development of cracks in the weld and the adjacent areas. The microstructure and microhardness analysis did not reveal the existence of fragile phases, cracks, inadequate penetration, or other imperfections, showing an appropriate adhesion between the deposited metal and the substrates. At the interface between the Ni-rich weld metal and the high-entropy alloy, a higher hardness (448 HV0.2) than in the base material (358 HV0.2) was measured. Energy-dispersive X-ray analysis (EDS), performed at the interface between the weld metal and the base materials, did not show significant modifications of Co, Fe, and Cr percentages. However, during the investigation, significant variations in Al and Ni concentrations were observed, caused by the fast diffusion of chemical elements, and the development of hard (Ni, Al)-rich compounds. In some areas of the deposited metal, located at a distance of about 10 µm from the interface, the percentages of Ni and Al were higher than in the high-entropy alloy base material, being around 41% by weight Ni and over 13% by weight Al, while the concentrations of the Co, Cr, and Fe elements proportionally decreased (i.e., approximately 14% by weight Co, 12% by weight Cr, and 17% by weight Fe). The development of Ni3Al and NiAl compounds was also noticed, whose formation was determined by the local chemical concentration and the temperature reached in the vicinity of the diffusion zone. The XRD analysis showed a group of X-ray peaks in the Al0.8CrFeCoNi alloy that corresponded to both α-type—BCC and FCC phases. The crystallite size of the high-entropy alloy investigated was found to be 22.05 nm. Despite the diffusion phenomenon, if filler materials and process parameters are appropriately selected, quality joints of high-entropy alloys and structural steels can be carried out under good welding conditions. Full article
(This article belongs to the Special Issue High-Productivity Welding of Metals and Alloys)
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<p>Aspect of the welded joint between HEA and S235JR substrate (<b>a</b>) and cross-sectional sample prepared for metallographic analysis (<b>b</b>).</p> Full article ">Figure 2
<p>Optical microstructure of specific areas of the welded joint: BM—base material; HAZ—Heat-Affected Zone; FL—Fusion Line; WD—Weld Deposit.</p> Full article ">Figure 3
<p>XRD diagram for Al<sub>0.8</sub>CrFeCoNi high-entropy alloy.</p> Full article ">Figure 4
<p>SEM microstructure of HEA sample: (<b>a</b>) longitudinal and (<b>b</b>) transversal cross-sections; (<b>c</b>) EDS analysis on gamma phase (GP) and (<b>d</b>) alpha phase (AP).</p> Full article ">Figure 5
<p>Distribution of the main chemical elements through the fusion line between HEA and weld metals: (<b>a</b>) HAZ—Spot 1 to 3; FL—Spot 4; WD—Spot 5 to 7; (<b>b</b>) chemical elements modification on the cross direction on the fusion line.</p> Full article ">Figure 6
<p>Elements distribution maps on fusion line between HEA and weld deposit.</p> Full article ">Figure 7
<p>Transition from weld metal to S235 JR: (<b>a</b>) fusion line between WD and unalloyed steel; (<b>b</b>) distribution of the main chemical elements in this area.</p> Full article ">Figure 8
<p>Chemical composition analysis near the fusion line between weld and S235 JR steel.</p> Full article ">Figure 9
<p>Hardness profile in the regions of the welded joint.</p> Full article ">
31 pages, 102999 KiB  
Review
The Influence of Precipitation, High Levels of Al, Si, P and a Small B Addition on the Hot Ductility of TWIP and TRIP Assisted Steels: A Critical Review
by Barrie Mintz and Abdullah Qaban
Metals 2022, 12(3), 502; https://doi.org/10.3390/met12030502 - 16 Mar 2022
Cited by 14 | Viewed by 2790
Abstract
The hot ductility of Transformation Induced Plasticity (TRIP) and Twinning Induced Plasticity (TWIP) steels is reviewed, concentrating on the likelihood of cracking occurring on continuous casting during the straightening operation. In this review, the influence of high levels of Al, Si, P, Mn [...] Read more.
The hot ductility of Transformation Induced Plasticity (TRIP) and Twinning Induced Plasticity (TWIP) steels is reviewed, concentrating on the likelihood of cracking occurring on continuous casting during the straightening operation. In this review, the influence of high levels of Al, Si, P, Mn and C on their hot ductility will be discussed as well as the important role B can play in improving their hot ductility. Of these elements, Al has the worst influence on ductility but a high Al addition is often needed in both TWIP and TRIP steels. AlN precipitates are formed often as thin coatings covering the austenite grain surfaces favouring intergranular failure and making them difficult to continuous cast without cracks forming. Furthermore, with TWIP steels the un-recrystallised austenite, which is the state the austenite is when straightening, suffers from excessive grain boundary sliding, so that the ductility often decreases with increasing temperature, resulting in the RA value being below that needed to avoid cracking on straightening. Fortunately, the addition of B can often be used to remedy the deleterious influence of AlN. The influence of precipitation hardeners (Nb, V and Ti based) in strengthening the room temperature yield strength of these TWIP steels and their influence on hot ductility is also discussed. Full article
(This article belongs to the Special Issue Continuous Casting and Hot Ductility of Advanced High-Strength Steels)
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Figure 1
<p>Schematic diagram showing the normal sequence of stacking (111) planes, ABCABCABC in the fcc crystal to that of missing out a plane so that the sequence changes to ABCBCABC, the stacking fault line being hcp [<a href="#B9-metals-12-00502" class="html-bibr">9</a>].</p> Full article ">Figure 2
<p>Mechanism of retained austenite formation during heat treatment of TRIP and DP steels [<a href="#B15-metals-12-00502" class="html-bibr">15</a>].</p> Full article ">Figure 3
<p>Schematic representation of a continuous casting machine.</p> Full article ">Figure 4
<p>A 2-D computerised strand temperature model predicting the thermal history during continuous casting of a 240 mm thick strand cast at 1 m/min [<a href="#B20-metals-12-00502" class="html-bibr">20</a>].</p> Full article ">Figure 5
<p>Thermal schedule used to generate the thermal condition of the billet surface in the continuous casting process: T<sub>m</sub> is melting point, T<sub>min</sub> and T<sub>max</sub> are lowest and highest temperatures respectively. T<sub>u</sub> is the temperature at the straightener and ΔT is the undercooling step [<a href="#B20-metals-12-00502" class="html-bibr">20</a>].</p> Full article ">Figure 6
<p>Hot ductility curve for a 0.4% C plain C-Mn steel tested at strain rate of 3 × 10<sup>−4</sup> s<sup>−1</sup> having no micro-alloying precipitates present showing all the different regions that are possible in the trough on cooling down through the austenitic temperature range, region (a) Deformation Induced ferrite (DIF) between Ae3 (the transformation temperature below which ferrite first starts to form under equilibrium conditions) and the Ar3 (the transformation temperature below which the steel first starts to form ferrite under non equilibrium conditions of cooling), region (b) Un-recrystallised γ, region (c) Recrystallised γ [<a href="#B17-metals-12-00502" class="html-bibr">17</a>].</p> Full article ">Figure 7
<p>Influence of particle size on the RA value for C-Mn-Al-Ti and C-Mn-Nb-Al steels [<a href="#B19-metals-12-00502" class="html-bibr">19</a>].</p> Full article ">Figure 8
<p>Influence of grain size on the RA value of steels, where D<sub>o</sub> is the original grain size before deformation for steels having 0.15% C, 1.4% Mn in which precipitation hardening is not contributing to the RA value [<a href="#B17-metals-12-00502" class="html-bibr">17</a>] and for TWIP steels having 0.6% C [<a href="#B21-metals-12-00502" class="html-bibr">21</a>].</p> Full article ">Figure 9
<p>Schematic diagram showing (<b>a</b>) how the width of the ductility trough could be controlled by the dynamic recrystallization (DRX), (<b>b</b>) how increasing the strain rate reduces the depth and width of the trough where ε<sub>c1</sub>, ε<sub>f1</sub> and TD<sub>1</sub> refer to the lower strain rate. ε<sub>c2</sub>, ε<sub>f2</sub> and TD<sub>2</sub> refer to the higher strain rate (<b>c</b>) how refining the grain size reduces the depth and width of the trough, where ε<sub>c1</sub>, ε<sub>f1</sub> and TD<sub>1</sub> refer to the coarser grain size and ε<sub>c2</sub>, ε<sub>f2</sub> and TD<sub>2</sub> refer to the finer grain size and (<b>d</b>) the influence of precipitation in increasing depth and width of the trough where ε<sub>c1</sub>, ε<sub>f1</sub> and TD<sub>1</sub> refer to trough without precipitation and ε<sub>c2</sub>, ε<sub>f2</sub> and TD<sub>2</sub>, the trough with precipitation. TD is the temperature on cooling when the changeover occurs from DRX of the ϒ, to the presence of a thin film of deformation induced ferrite surrounding the unrecrystallised ϒ.</p> Full article ">Figure 10
<p>Example of delayed fracture in a deep drawn Fe-22% Mn-0.6% C, TWIP steel cup (top, left). Right, Top row: Suppression of delayed fracture by 1.5% Al addition in deep drawn Fe-xMn-0.6% C, TWIP steel cups having 15, 16 and 17% Mn. Right, bottom row: Cups showing cracks in the same steels without an Al addition due to delayed fracture [<a href="#B5-metals-12-00502" class="html-bibr">5</a>].</p> Full article ">Figure 11
<p>Various forms of AlN precipitation at the boundaries in as-cast high Al, TWIP steels (0.6% C, 18% Mn). Thin films of AlN on the dendritic surface of a high Al, TWIP steel (<b>a</b>) 1.6% Al, 0.007% N, (<b>b</b>) 1.4% Al, 0.004% N and very low S, 0.002% S (<b>c</b>) Very coarse AlN precipitates situated at the austenite grain boundaries [<a href="#B26-metals-12-00502" class="html-bibr">26</a>,<a href="#B27-metals-12-00502" class="html-bibr">27</a>].</p> Full article ">Figure 11 Cont.
<p>Various forms of AlN precipitation at the boundaries in as-cast high Al, TWIP steels (0.6% C, 18% Mn). Thin films of AlN on the dendritic surface of a high Al, TWIP steel (<b>a</b>) 1.6% Al, 0.007% N, (<b>b</b>) 1.4% Al, 0.004% N and very low S, 0.002% S (<b>c</b>) Very coarse AlN precipitates situated at the austenite grain boundaries [<a href="#B26-metals-12-00502" class="html-bibr">26</a>,<a href="#B27-metals-12-00502" class="html-bibr">27</a>].</p> Full article ">Figure 12
<p>“Rock candy” fracture in a 1.05% Al containing TRIP steel [<a href="#B28-metals-12-00502" class="html-bibr">28</a>]. The composition of the TRIP steel was 0.15% C, 2.45% Mn, 0.025% Nb, 0.005% S, 0.0065% N.</p> Full article ">Figure 13
<p>Hot ductility curves for TWIP steels having different products of [Al][N], from 0.61 to 35 × 10<sup>−3</sup>. Al additions varied from 0.047 to 1.5% and N from 0.004 to 0.0023%. The base composition of the steels was 0.6% C, 18% Mn and 0.006% S [<a href="#B30-metals-12-00502" class="html-bibr">30</a>,<a href="#B31-metals-12-00502" class="html-bibr">31</a>].</p> Full article ">Figure 14
<p>The hot ductility of TWIP steel at different Al content for the base composition: 0.6% C, 0.008% S, 18% Mn and 0.01% N [<a href="#B32-metals-12-00502" class="html-bibr">32</a>]. The top curve is for a TWIP steel free of Al and shows the big improvement in overall ductility when DRX occurs.</p> Full article ">Figure 15
<p>Influence of Mn content on the ductility of high Al containing TWIP steels. The base composition of the steels was 0.6% C, 0.007% P, 0.008% S and 1.5% Al [<a href="#B32-metals-12-00502" class="html-bibr">32</a>].</p> Full article ">Figure 16
<p>Hot ductility curves for a TWIP steel having the composition 0.45% C, 22% Mn, 1.5% Al and 1.5% Si with 0.02% Ti for two different cooling rates, the tensile specimens were cast in metallic and sand molds [<a href="#B33-metals-12-00502" class="html-bibr">33</a>].</p> Full article ">Figure 17
<p>Thermo-Calc precipitation predictions for a TWIP steel with a composition of 0.61% C, 18.0% Mn, 0.003% S, 0.062% Ti, 1.54% Al and 0.007% N. Ti:N ratio of 7:1.</p> Full article ">Figure 18
<p>(<b>a</b>) MnS inclusions acting as a nucleus for AlN precipitation in TWIP steels having S in the range 0.01–0.023% and (<b>b</b>) a similar steel with a low S, content of 0.003% [<a href="#B49-metals-12-00502" class="html-bibr">49</a>].</p> Full article ">Figure 19
<p>Hot ductility curves for three TWIP steels having 0.003, 0.01 and 0.023% S with base composition: 0.6% C, 18.3% Mn, 1.5% Al and 0.009% N [<a href="#B49-metals-12-00502" class="html-bibr">49</a>].</p> Full article ">Figure 20
<p>P addition of 0.054% P improving the hot ductility of a low alloy Cr-Mo steel [<a href="#B53-metals-12-00502" class="html-bibr">53</a>].</p> Full article ">Figure 21
<p>Influence of P on the hot ductility of high C (0.6%), low alloy steels: (<b>a</b>) hot ductility curves for as-cast C-Mn-Nb-Al and plain C-Mn steels at two levels of P, 0.007% and 0.045% (<b>b</b>) low melting point iron phosphide phase at prior austenite grain boundaries in the low P plain C-Mn steel. The base composition of the steels was 0.6% C, 2% Si, 0.8% Mn and 0.025% Al with and without 0.03% Nb [<a href="#B56-metals-12-00502" class="html-bibr">56</a>].</p> Full article ">Figure 21 Cont.
<p>Influence of P on the hot ductility of high C (0.6%), low alloy steels: (<b>a</b>) hot ductility curves for as-cast C-Mn-Nb-Al and plain C-Mn steels at two levels of P, 0.007% and 0.045% (<b>b</b>) low melting point iron phosphide phase at prior austenite grain boundaries in the low P plain C-Mn steel. The base composition of the steels was 0.6% C, 2% Si, 0.8% Mn and 0.025% Al with and without 0.03% Nb [<a href="#B56-metals-12-00502" class="html-bibr">56</a>].</p> Full article ">Figure 22
<p>Influence of P on hot ductility curves of a TWIP steel. Steels had a base composition of 0.6% C, 0.3% Si, 18.2% Mn, 0.005% S, 1.5% Al, 0.01% Ti, 0.007% N with a B addition of 0.0026% B and P additions of 0.007, 0.019, 0.037 and 0.074%, steels, 1–4, respectively [<a href="#B59-metals-12-00502" class="html-bibr">59</a>].</p> Full article ">Figure 23
<p>Influence of Si on the hot ductility curves of plain 0.1% C, 1.2% Mn steels (Al free) having 0.011% N. Curves move to higher temperatures with increasing Si level [<a href="#B60-metals-12-00502" class="html-bibr">60</a>].</p> Full article ">Figure 24
<p>Influence of Si on the hot ductility of medium C steels (0.5% C, 0.01% N), again curve moves to higher temperature with increase in Si content [<a href="#B61-metals-12-00502" class="html-bibr">61</a>].</p> Full article ">Figure 25
<p>A comparison of the room temperature yield strength increases in TWIP steels for cold rolled and annealed strips as a function of the microalloying additions, Ti, Nb and V [<a href="#B62-metals-12-00502" class="html-bibr">62</a>].</p> Full article ">Figure 26
<p>Hot ductility curves for V containing high Al TWIP steels 1–5 having respectively, 0.05, 0.11, 0.29, 0.5 and 0.75% V [<a href="#B22-metals-12-00502" class="html-bibr">22</a>].</p> Full article ">Figure 27
<p>Fine VC precipitation in a 0.3% V containing TWIP steel both at the grain boundaries and within the matrix [<a href="#B22-metals-12-00502" class="html-bibr">22</a>]. A Ni grid was used to support the replica.</p> Full article ">Figure 28
<p>The beneficial influence of B on the hot ductility of a high V (0.5% V), TWIP steel [<a href="#B22-metals-12-00502" class="html-bibr">22</a>].</p> Full article ">Figure 29
<p>Hot ductility curves for high Al TWIP steels with the same V content, 0.011% V, (a) recrystallised austenite (b) un-recrystallised austenite. Compositions of TWIP steels were: Recrystallised top curve –21% Mn, 0.56% C, 1.3% Si, 1.50% Al, 0.011% V and 0.012% N [<a href="#B33-metals-12-00502" class="html-bibr">33</a>]. Unrecrystallised bottom curve –18% Mn, 0.61% C, 0.2% Si, 1.54% Al, 0.011% V and 0.007% N [<a href="#B22-metals-12-00502" class="html-bibr">22</a>].</p> Full article ">Figure 30
<p>Hot ductility curves for high Al TWIP steels having a variety of microalloying elements, Nb, Nb-V, Ti, V, Ti-B and Ti-B-V. Steels had the base composition 0.6% C, 18% Mn, 1.5% Al and the cooling rate and strain rate were 60 K/min and 3 × 10<sup>−3</sup> s<sup>−1</sup>, respectively [<a href="#B22-metals-12-00502" class="html-bibr">22</a>,<a href="#B26-metals-12-00502" class="html-bibr">26</a>,<a href="#B27-metals-12-00502" class="html-bibr">27</a>,<a href="#B65-metals-12-00502" class="html-bibr">65</a>,<a href="#B66-metals-12-00502" class="html-bibr">66</a>].</p> Full article ">Figure 31
<p>Schematic types of hot ductility curves for high Mn TWIP steels (<b>a</b>) No DRX, no fine matrix precipitation, the curve is relevant to straightening (<b>b</b>) GBS at the low and high temperature ends of the straightening temperature range but DRX in intermediate temperature range, curve relevant to hot forming (<b>c</b>) Separation of curve in (<b>b</b>) into regions of GBS and DRX and drawing a straight line relationship for continued GBS in the temperature range in which DRX is occurring [<a href="#B21-metals-12-00502" class="html-bibr">21</a>].</p> Full article ">Figure 31 Cont.
<p>Schematic types of hot ductility curves for high Mn TWIP steels (<b>a</b>) No DRX, no fine matrix precipitation, the curve is relevant to straightening (<b>b</b>) GBS at the low and high temperature ends of the straightening temperature range but DRX in intermediate temperature range, curve relevant to hot forming (<b>c</b>) Separation of curve in (<b>b</b>) into regions of GBS and DRX and drawing a straight line relationship for continued GBS in the temperature range in which DRX is occurring [<a href="#B21-metals-12-00502" class="html-bibr">21</a>].</p> Full article ">Figure 32
<p>Influence of Nb on the hot ductility of high Al, Ti-B containing TWIP steels. The composition of the steels (1–6) is given in <a href="#metals-12-00502-t004" class="html-table">Table 4</a>. Cooling rate and strain rate were 60 K/min and 3 × 10<sup>−3</sup> s<sup>−1</sup>, respectively [<a href="#B66-metals-12-00502" class="html-bibr">66</a>].</p> Full article ">Figure 33
<p>Coarse TiN precipitates in 0.1% Ti, containing B steel, TWIP steel free of Nb, tested at a 1000 °C, Steel 4 in <a href="#metals-12-00502-t004" class="html-table">Table 4</a> [<a href="#B27-metals-12-00502" class="html-bibr">27</a>].</p> Full article ">Figure 34
<p>Nb-Ti carbo-nitrides found in a B containing high Al, TWIP tested at 1000 °C, steel 1 in <a href="#metals-12-00502-t004" class="html-table">Table 4</a>. The precipitates varied considerably in size (~80 nm in average size) but always gave similar composition. Composition of steel was 0.6% C, 18% Mn, 1.51% Al, 0.033% Nb, 0.075% Ti and 0.011% N. Cooling rate and strain rate were 60 K/min and 3 × 10<sup>−3</sup> s<sup>−1</sup>, respectively [<a href="#B66-metals-12-00502" class="html-bibr">66</a>].</p> Full article ">
14 pages, 14185 KiB  
Article
Corrosion Behavior of Gravity Cast and High-Pressure Die-Cast AM60 Mg Alloys with Ca and Y Addition
by Hui Yu, Xin Yang, Wei Yu, Youngmin Kim, Shaoming Kang, Lixin Huang, Bongsun You, Chunhai Liu, Kwangseon Shin and Weimin Zhao
Metals 2022, 12(3), 495; https://doi.org/10.3390/met12030495 - 15 Mar 2022
Viewed by 2336
Abstract
In this study, the microstructure and related corrosion behavior of AM60 alloys with/without Ca and Y addition upon gravity casting (GC) and high-pressure die-casting (HPDC) are investigated by means of SEM/EDS characterization, immersion/salt spray test, hydrogen releasing, as well as electrochemistry examination. When [...] Read more.
In this study, the microstructure and related corrosion behavior of AM60 alloys with/without Ca and Y addition upon gravity casting (GC) and high-pressure die-casting (HPDC) are investigated by means of SEM/EDS characterization, immersion/salt spray test, hydrogen releasing, as well as electrochemistry examination. When utilizing GC, AM60 alloy with Ca and Y addition (named AZXW6000 alloy) has better corrosion resistance compared with AM60 alloy. Thanks to grain refinement and second phase networks introduced in HPDC, the anti-corrosion properties of the AM alloy seem much better than those of the GC counterpart. The corrosion mechanism of both GCed and HPDCed AM60-based alloys is also investigated in detail. The inspiration from present work can provide more thinking for developing high corrosion-resistant Mg alloys using different casting methods. Full article
(This article belongs to the Special Issue Corrosion and Protection of Lightweight Engineering Materials: Mg Alloys, Al Alloys, Ti Alloys and Other Related Metals)
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<p>Solidification calculation of (<b>a</b>) AM60 alloy and (<b>b</b>) AMXW6000 alloy by Pandat<sup>TM</sup>; (<b>c</b>) XRD patterns of AM60 and AMXW6000 alloys under GC and HPDC.</p> Full article ">Figure 2
<p>SEM images of (<b>a</b>) GC AM60 alloy, (<b>b</b>) GC AMXW6000 alloy, (<b>c</b>) HPDC AM60 alloy, (<b>d</b>) HPDC AMXW6000 alloy, and corresponding (<b>e</b>) EDS results in (<b>a</b>–<b>d</b>), respectively.</p> Full article ">Figure 3
<p>SEM observation of (<b>a</b>–<b>e</b>) GC AM60 alloy and (<b>f</b>–<b>j</b>) GC AMXW6000 alloy by immersion after 1 h, 3 h, 12 h, 2 days, 8 days, respectively.</p> Full article ">Figure 4
<p>SEM images of HPDCed AM60 alloy under various immersion time: (<b>a</b>) 3 h; (<b>b</b>) 12 h; (<b>c</b>) 2 days; (<b>d</b>) 8 days; SEM with EDS mapping of HPDCed AM60 alloy after immersion for (<b>e</b>) 12 h; (<b>f</b>) 2 days; (<b>g</b>) 8 days, respectively.</p> Full article ">Figure 5
<p>SEM images of HPDCed AMXW6000 alloy under various immersion time: (<b>a</b>) 3 h; (<b>b</b>) 12 h; (<b>c</b>) 2 days; (<b>d</b>) 8 days; SEM with EDS mapping of HPDCed AM6000 alloy after immersion for (<b>e</b>) 12 h; (<b>f</b>) 2 days; (<b>g</b>) 8 days, respectively.</p> Full article ">Figure 6
<p>Cross sectional SEM images of HPDCed (<b>a</b>) AM60 alloy and (<b>b</b>) AMXW6000 alloy after 4 days immersion.</p> Full article ">Figure 7
<p>Schematic diagram of corrosion behavior in (<b>a</b>) HPDCed sample and (<b>b</b>) GCed alloys.</p> Full article ">Figure 8
<p>H<sub>2</sub> release volume (<b>a</b>,<b>c</b>) and calculated evolution rate (<b>b</b>,<b>d</b>) of AM60 and AMXW6000 alloy in terms of time in GC and HPDC state, respectively.</p> Full article ">Figure 9
<p>Corrosion rate examined by weight loss tests of GC and HPDC samples: (<b>a</b>) IT and (<b>b</b>) SST results, respectively.</p> Full article ">Figure 10
<p>Macro-morphology of corroded surface &amp; side under IT and SST after 5 and 15 days.</p> Full article ">Figure 11
<p>Linear polarization curves of (<b>a</b>) as-received GC and HPDC samples, (<b>b</b>) HPDCed AM60 and (<b>c</b>) AMXW6000 alloys, for different values of corrosion time.</p> Full article ">Figure 12
<p>Nyquist plots and Bode impedance plots of (<b>a</b>) as-received GC and HPDC samples, (<b>b</b>,<b>c</b>) HPDCed AM60 and AMXW6000 alloy with different corrosion times, respectively.</p> Full article ">Figure 13
<p>Equivalent electric circuit diagrams of EIS from <a href="#metals-12-00495-f012" class="html-fig">Figure 12</a>.</p> Full article ">
10 pages, 29014 KiB  
Article
Investigation on Blood Compatibility of Cu/Ti Metal Coating Prepared via Various Bias Voltages and Copper Content
by Qiong Hu, Hengquan Liu, Fei Gao, Xi Yang, Junfeng Li, Ren Liu, Zexuan Liu and Dongfang Wang
Metals 2022, 12(3), 435; https://doi.org/10.3390/met12030435 - 1 Mar 2022
Cited by 1 | Viewed by 2219
Abstract
Surface modification of some metal coatings is usually used to improve the blood compatibility of biomaterials; however, some aspects of the bological properties of metal coatings cannot be adjusted via the content of each component. In this work, Cu/Ti metal coatings with various [...] Read more.
Surface modification of some metal coatings is usually used to improve the blood compatibility of biomaterials; however, some aspects of the bological properties of metal coatings cannot be adjusted via the content of each component. In this work, Cu/Ti metal coatings with various amounts of copper content were prepared by the physical vapor deposition (PVD) method, and the influence of deposition bias was further investigated. Phase structure, element composition and surface morphology were investigated by X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS) and scanning electron microscopy, respectively. The hemolysis ratio, platelet adhesion and protein adsorption were applied to evaluate the blood compatibility. The results show that a Cu/Ti coating of uniform quality can be obtained; the dispersion of the deposition and copper content is regulated by the number of copper sheets, but the deposition bias does not obviously affect the copper content of the Cu/Ti coating. The hemolysis rate of the Cu/Ti coating is less than 0.4%, the degree of platelet adhesion is significantly reduced on Cu/Ti coatings compared to control samples, and the contact angle of all coatings is greater than that of pure titanium. The largest adsorption capacity of BSA was found on the coating with the deposition bias voltage of −40 V. The number of copper flakes is increased, and the adsorption of FIB on the Cu/Ti coating surface is reduced. Therefore, Cu/Ti coatings prepared via this deposition method have potential for applications to regulate blood compatibility and surface performance. Full article
(This article belongs to the Special Issue Advances in Stability of Metallic Implants)
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<p>XRD pattern of Cu/Ti coating: (<b>a</b>) different bias voltage, (<b>b</b>) different number of copper sheets.</p> Full article ">Figure 2
<p>XPS spectra of 2# and 4# samples: (<b>a</b>) full spectrum, (<b>b</b>) 2# Cu2p, (<b>c</b>) 2# Ti2p, (<b>d</b>) 4# Cu2p, (<b>e</b>) 4# tip.</p> Full article ">Figure 3
<p>The SEM images of various Cu/Ti composite coatings: (<b>a</b>) 1#, (<b>b</b>) 2#, (<b>c</b>) 3#, (<b>d</b>) 4#, (<b>e</b>) 2#, (<b>f</b>) 5#, (<b>g</b>) 6#.</p> Full article ">Figure 4
<p>The contact angle and surface energy of various samples: (<b>a</b>) the contact angle of 1#–4#, (<b>b</b>) the contact angle of 2#, 5# and 6#, (<b>c</b>) free energy of 1#–4#, (<b>d</b>) free energy of 2#, 5# and 6#.</p> Full article ">Figure 5
<p>Platelet adhesion of Cu/Ti composite coatings: (<b>a</b>) Ti, (<b>b</b>) 1#, (<b>c</b>) 2#, (<b>d</b>) 3#, (<b>e</b>) 4#, (<b>f</b>) 2#, (<b>g</b>) 5#, (<b>h</b>) 6#.</p> Full article ">Figure 6
<p>The platelet activation. (<b>a</b>) Different bias voltage, (<b>b</b>) different copper content.</p> Full article ">Figure 7
<p>Surface fibrinogen and albumin adsorption capacity on various surfaces (indicated by OD value). (<b>a</b>) Different bias voltage, (<b>b</b>) different copper content.</p> Full article ">
20 pages, 9493 KiB  
Article
Effect of Intercritical Annealing Parameters and Starting Microstructure on the Microstructural Evolution and Mechanical Properties of a Medium-Mn Third Generation Advanced High Strength Steel
by Kazi M. H. Bhadhon, Xiang Wang, Elizabeth A. McNally and Joseph R. McDermid
Metals 2022, 12(2), 356; https://doi.org/10.3390/met12020356 - 18 Feb 2022
Cited by 13 | Viewed by 2912
Abstract
A prototype medium-Mn TRIP steel (0.2 C–6 Mn–1.7 Si–0.4 Al–0.5 Cr (wt %)) with a cold-rolled tempered martensite (CR) and martensitic (M) starting microstructures was subjected to continuous galvanizing line (CGL) compatible heat treatments. It was found that the M starting microstructures achieved [...] Read more.
A prototype medium-Mn TRIP steel (0.2 C–6 Mn–1.7 Si–0.4 Al–0.5 Cr (wt %)) with a cold-rolled tempered martensite (CR) and martensitic (M) starting microstructures was subjected to continuous galvanizing line (CGL) compatible heat treatments. It was found that the M starting microstructures achieved greater than 0.30 volume fraction of retained austenite and target 3G properties (UTS × TE ≥ 24,000 MPa%) using an intercritical annealing temperature (IAT) of 675 °C with an IA holding time of 60–360 s, whereas the CR microstructure required an IAT of 710 °C and annealing times of 360 s or greater to achieve comparable fractions of retained austenite and target 3G properties. This was attributed to the rapid austenite reversion kinetics for the M starting microstructures and rapid C partitioning from the C supersaturated martensite, providing chemical and mechanical stability to the retained austenite, thereby allowing for a gradual deformation-induced transformation of retained austenite to martensite—the TRIP effect—and the formation of nano-scale planar faults in the retained austenite (TWIP effect), such that a high work-hardening rate was maintained to elongation of greater than 0.20. Overall, it was concluded that the prototype steel with the M starting microstructure is a promising candidate for CGL processing for 3G AHSS properties. Full article
(This article belongs to the Special Issue Recent Developments in Medium and High Manganese Steels)
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<p>Schematic diagram of the heat treatment profile.</p> Full article ">Figure 2
<p>(<b>a</b>) Load–unload–reload test; (<b>b</b>) enlarged portion of (<b>a</b>) showing process for determining <span class="html-italic">σ<sub>F</sub></span> and <span class="html-italic">σ<sub>R</sub></span>.</p> Full article ">Figure 3
<p>SEM micrographs of (<b>a</b>) as-received CR samples with tempered martensite and (<b>b</b>) heat-treated M samples with martensite; C = carbides, TM = tempered martensite, M = martensite, TT = through thickness, TD = transverse direction.</p> Full article ">Figure 4
<p>Montage of TEM micrographs showing carbide distribution in (<b>a</b>) CR and (<b>b</b>) M starting microstructures.</p> Full article ">Figure 5
<p>Retained austenite volume fraction as a function of intercritical annealing temperature and holding time for (<b>a</b>) CR and (<b>b</b>) M starting microstructures.</p> Full article ">Figure 6
<p>Estimated fresh martensite volume fraction as a function of holding time at 675 and 710 °C IATs for (<b>a</b>) CR and (<b>b</b>) M starting microstructures.</p> Full article ">Figure 7
<p>SEM micrographs of CR samples annealed at 710 °C for (<b>a</b>) 120 s, (<b>b</b>) 240 s, and (<b>c</b>) 600 s; C = carbide, F = ferrite, M = martensite, and A = retained austenite.</p> Full article ">Figure 8
<p>SEM micrographs of M samples annealed at 675 °C for (<b>a</b>) 120 s, (<b>b</b>) 360 s, (<b>c</b>) 600 s; C = carbide, F = ferrite, M = martensite, and A = retained austenite.</p> Full article ">Figure 8 Cont.
<p>SEM micrographs of M samples annealed at 675 °C for (<b>a</b>) 120 s, (<b>b</b>) 360 s, (<b>c</b>) 600 s; C = carbide, F = ferrite, M = martensite, and A = retained austenite.</p> Full article ">Figure 9
<p>TEM results for the M 675 °C + 120 s sample: (<b>a</b>) Montage of bright field (BF) TEM micrographs; (<b>b</b>) dark field (DF) TEM corresponding to <math display="inline"><semantics> <mrow> <mo>〈</mo> <mn>110</mn> <mo>〉</mo> </mrow> </semantics></math>γ; (<b>c</b>) SAD patterns corresponding to <math display="inline"><semantics> <mrow> <mfenced close="]" open="["> <mrow> <mn>100</mn> </mrow> </mfenced> <msup> <mi>α</mi> <mo>′</mo> </msup> <mn>1</mn> <mo>∥</mo> <mfenced close="]" open="["> <mrow> <mn>111</mn> </mrow> </mfenced> <msup> <mi>α</mi> <mo>′</mo> </msup> <mn>2</mn> <mo>∥</mo> <mfenced close="]" open="["> <mrow> <mn>110</mn> </mrow> </mfenced> <mi>γ</mi> </mrow> </semantics></math>.</p> Full article ">Figure 10
<p>Montage of BF TEM micrographs showing carbide distribution in (<b>a</b>) CR 710 °C + 120 s and (<b>b</b>) M 675 °C + 120 s samples.</p> Full article ">Figure 11
<p>Carbide area fraction as a function of microstructure/heat treatment.</p> Full article ">Figure 12
<p>(<b>a</b>) Engineering stress vs. engineering strain; (<b>b</b>) true stress vs. true strain curves; work hardening rate vs. true strain curves for selected (<b>c</b>) CR and (<b>d</b>) M starting microstructures.</p> Full article ">Figure 13
<p>Retained austenite transformation kinetics for selected CR and M samples.</p> Full article ">Figure 14
<p>(<b>a</b>) Bright field (BF) TEM; (<b>b</b>) dark field (DF) TEM corresponding to <math display="inline"><semantics> <mrow> <mo>〈</mo> <mn>110</mn> <mo>〉</mo> </mrow> </semantics></math>γ; (<b>c</b>) SAD patterns corresponding to <math display="inline"><semantics> <mrow> <mfenced close="]" open="["> <mrow> <mn>100</mn> </mrow> </mfenced> <msup> <mi>α</mi> <mo>′</mo> </msup> <mn>1</mn> <mo>∥</mo> <mfenced close="]" open="["> <mrow> <mn>111</mn> </mrow> </mfenced> <msup> <mi>α</mi> <mo>′</mo> </msup> <mn>2</mn> <mo>∥</mo> <mfenced close="]" open="["> <mrow> <mn>110</mn> </mrow> </mfenced> <mi>γ</mi> </mrow> </semantics></math> for M 675 °C + 120 s sample at <span class="html-italic">ε</span> = 0.10.</p> Full article ">Figure 15
<p>R (<span class="html-italic">ε</span>) versus true strain for the M 675 °C + 120 s sample.</p> Full article ">
15 pages, 7246 KiB  
Article
Role of Hot Rolling in Microstructure and Texture Development of Strip Cast Non-Oriented Electrical Steel
by Haitao Jiao, Xinxiang Xie, Xinyi Hu, Longzhi Zhao, Raja Devesh Kuma Misra, Dejia Liu, Yanchuan Tang and Yong Hu
Metals 2022, 12(2), 354; https://doi.org/10.3390/met12020354 - 18 Feb 2022
Cited by 4 | Viewed by 3652
Abstract
In this study, the effect of the hot-cold rolling process on the evolution of the microstructure, texture and magnetic properties of strip-cast non-oriented electrical steel was investigated by introducing hot rolling with different reductions. The results indicate that hot rolling with an appropriate [...] Read more.
In this study, the effect of the hot-cold rolling process on the evolution of the microstructure, texture and magnetic properties of strip-cast non-oriented electrical steel was investigated by introducing hot rolling with different reductions. The results indicate that hot rolling with an appropriate reduction, such as the 20% used in this study, increases the shear bands and {100} deformed microstructure in the cold roll sheet. As a result, in our study, enhanced η and Cube recrystallization texture and the improved magnetic induction were obtained. However, hot rolling with excessive reduction (36–52%) decreased the shear bands and increased the α-oriented deformation microstructure with low stored energy. It enhanced the α recrystallization texture and weakened the η texture, resulting in a decrease in the magnetic induction. In addition, hot rolling promoted the precipitation of supersaturated solid solution elements in the as-cast strip, thereby affecting the subsequent microstructure evolution and the optimization of its magnetic properties. Full article
(This article belongs to the Special Issue Texture, Microstructure and Properties of Electrical Steels)
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<p>Schematic diagram of twin-roll strip casting, rolling, and annealing process.</p> Full article ">Figure 2
<p>Microstructure of (<b>a</b>) as-cast Fe-2.8%Si strip and hot-rolled strip with reduction of (<b>b</b>) 20%, (<b>c</b>) 36%, and (<b>d</b>) 52%, respectively.</p> Full article ">Figure 3
<p>Macro-texture (φ2 = 45° ODFs) of (<b>a</b>) as-cast strip and hot-rolled strip with reductions of (<b>b</b>) 20%, (<b>c</b>) 36%, and (<b>d</b>) 52%, respectively.</p> Full article ">Figure 4
<p>Cold-rolled microstructure of sheet subjected to (<b>a</b>) processing route A, (<b>b</b>) route B, (<b>c</b>) route C, and (<b>d</b>) route D, respectively.</p> Full article ">Figure 5
<p>Macro-texture (φ<sub>2</sub> = 45° ODFs) of cold-rolled sheet subjected to (<b>a</b>) processing route A, (<b>b</b>) route B, (<b>c</b>) route C, and (<b>d</b>) route D, respectively.</p> Full article ">Figure 6
<p>Orientation intensities along (<b>a</b>) α-fiber, (<b>b</b>) γ-fiber, and (<b>c</b>) λ-fiber textures in cold-rolled sheet subjected to different processing routes.</p> Full article ">Figure 7
<p>Recrystallized microstructure of sheet subjected to different hot-rolling and cold-rolling process: (<b>a</b>) route A by direct cold rolling; (<b>b</b>) route B by 20% reduction hot rolling; (<b>c</b>) route B by 36% reduction hot rolling; (<b>d</b>) route B by 52% reduction hot rolling.</p> Full article ">Figure 8
<p>Precipitates in the final annealed micro-structure: (<b>a</b>) sample processed by route A; (<b>b</b>) sample processed by route B; (<b>c</b>) sample processed by route D.</p> Full article ">Figure 9
<p>Recrystallization texture (φ<sub>2</sub> = 45° and φ<sub>2</sub> = 0° ODFs) of sheet subjected to different processing routes: (<b>a</b>) route A; (<b>b</b>) route B; (<b>c</b>) route C; (<b>d</b>) route D.</p> Full article ">Figure 10
<p>Intensity of specific texture components in the annealed samples subjected to different rolling processes: (<b>a</b>) η-fiber; (<b>b</b>) α-fiber; and (<b>c</b>) λ-fiber texture.</p> Full article ">Figure 11
<p>EBSD inverse pole figure (IPF) map and orientation distribution image (OIM) of partially recrystallized microstructure: (<b>a</b>) sample processed by route A; (<b>b</b>) sample processed by route B; (<b>c</b>) sample processed by route D.</p> Full article ">Figure 12
<p>Magnetic properties of annealed sheet subjected to different processing routes: (<b>a</b>) magnetic induction <span class="html-italic">B</span><sub>50</sub>; (<b>b</b>) core loss <span class="html-italic">P</span><sub>15/50</sub>.</p> Full article ">
13 pages, 3792 KiB  
Article
Use of Porous Titanium Trabecular as a Bone Defect Regenerator: In Vivo Study
by Ana Isabel Torres Pérez, Mariano Fernández Fairén, Ángel Antonio Torres Pérez and Javier Gil Mur
Metals 2022, 12(2), 327; https://doi.org/10.3390/met12020327 - 12 Feb 2022
Cited by 2 | Viewed by 2253
Abstract
The application of porous materials is increasingly being used in orthopaedic surgery due to its good results. Bone growth within the pores results in excellent mechanical fixation with the bone, as well as good bone regeneration. The pores, in addition to being colonised [...] Read more.
The application of porous materials is increasingly being used in orthopaedic surgery due to its good results. Bone growth within the pores results in excellent mechanical fixation with the bone, as well as good bone regeneration. The pores, in addition to being colonised by bone, produce a decrease in the modulus of elasticity that favours the transfer of loads to the bone. This research shows the results of an experimental study where we have created critical osteoperiosteal defects of 10 mm on rabbit’s radius diaphysis. In one group of 10 rabbits (experimental group) we have implanted a bioactive porous titanium cylinder, and in another group we have allowed spontaneous regeneration (control group). Mechanical tests were performed to assess the material. Image diagnostic techniques (X-ray, scanner and 3D scan: there are no references on the literature with the use of CT-scan in bone defects) and histological and histomorphometric studies post-op and after 3, 6 and 12 months after the surgery were performed. All the control cases went through a pseudoarthrosis. In 9 of the 10 cases of the experimental group complete regeneration was observed, with a normal cortical-marrow structure established at 6 months, similar to normal bone. Titanium trabecular reached a bone percentage of bone inside the implant of 49.3% on its surface 3 months post-op, 75.6% at 6 months and 81.3% at 12 months. This porous titanium biomaterial has appropriate characteristics to allow bone ingrowth, and it can be proposed as a bone graft substitute to regenerate bone defects, as a scaffold, or as a coating to achieve implant osteointegration. Full article
(This article belongs to the Special Issue Titanium Alloys: Microstructure, Mechanical Properties and Biomedical Applications)
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<p>Surgical technique.</p> Full article ">Figure 2
<p>Porous titanium obtained by PM.</p> Full article ">Figure 3
<p>X-rays during the follow-up. (<b>A</b>) 0 months, (<b>B</b>) 2 months, (<b>C</b>) 4 months, (<b>D</b>) 6 months.</p> Full article ">Figure 4
<p>(<b>A</b>) CT-Scan helicoidally images to estimate restauration of the radius continuity. (<b>B</b>) 3D Reconstruction of the CT-scan helicoidally images to estimate restauration of the radius continuity.</p> Full article ">Figure 5
<p>Defect bone growth progression in the titanium trabecular group (circles, blue) and control group (squares, green).</p> Full article ">Figure 6
<p>Cortical-marrow pattern of regenerated bone.</p> Full article ">Figure 7
<p>Normal bone structure with osteocytes inside the lacunas and canalicular mesh (Giemsa stain).</p> Full article ">Figure 8
<p>Transmission electron microscopy image, where mature laminar bone can be seen inside the pores of the metallic structure. Image on the <b>left</b>: 35×. Image on the <b>right</b>: segmentation performed to achieve the histomorphometric study.</p> Full article ">Figure 9
<p>Scanning electron microscopy images: (<b>A</b>) Position on the implant inside de radius, bone integration at 3 months; (<b>B</b>) Bone regeneration inside and around the implant at 12 months; and (<b>C</b>) Direct union bone and implant.</p> Full article ">Figure 9 Cont.
<p>Scanning electron microscopy images: (<b>A</b>) Position on the implant inside de radius, bone integration at 3 months; (<b>B</b>) Bone regeneration inside and around the implant at 12 months; and (<b>C</b>) Direct union bone and implant.</p> Full article ">
12 pages, 4228 KiB  
Article
Rejuvenation-to-Relaxation Transition Induced by Elastostatic Compression and Its Effect on Deformation Behavior in a Zr-Based Bulk Metallic Glass
by Jingxian Cui, Qiang Luo, Siyi Di, Zhengguo Zhang and Baolong Shen
Metals 2022, 12(2), 282; https://doi.org/10.3390/met12020282 - 4 Feb 2022
Cited by 9 | Viewed by 2529
Abstract
The effect of uniaxial elastostatic compression on the deformation behavior of the Zr41.2Ti13.8Cu12.5Ni10Be22.5 (Vit1) bulk metallic glass (BMG) was reported. The as-cast alloy was pre-compressed under various time (20, 40 and 60 h) at [...] Read more.
The effect of uniaxial elastostatic compression on the deformation behavior of the Zr41.2Ti13.8Cu12.5Ni10Be22.5 (Vit1) bulk metallic glass (BMG) was reported. The as-cast alloy was pre-compressed under various time (20, 40 and 60 h) at a preloading level of 87% of its yield strength. It was found that elastostatic compression can lead to structural rejuvenation or relaxation depending on the pre-compression time. Elastostatic compression, for 40 h, increased the free volume and improved the plasticity of the BMGs from 1.4% to 3.4%, but preloading for 60 h decreased the free volume and worsened the plasticity. In addition, the heterogeneous structure evolution during creep deformation has been analyzed by the Maxwell-Voigt model with two Kelvin units, revealing that more (less) defects with larger size are activated after elastostatic compression treatment for 40 h (60 h). This work sheds new light on the correlation between heterogeneous structure and plasticity/creep behaviors of Zr-based BMGs. Full article
(This article belongs to the Section Entropic Alloys and Meta-Metals)
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<p>The representative engineering stress-strain curves of the AC (as-cast) and pre-compression treated (T20, T40, T60) Vit1 BMG samples.</p> Full article ">Figure 2
<p>SEM images of the fracture samples of (<b>a</b>) AC, (<b>b</b>) T20, (<b>c</b>) T40 and (<b>d</b>) T60 Vit1 BMG samples, which were compressed to failure.</p> Full article ">Figure 3
<p>(<b>a</b>) DSC traces corresponding to the AC and T20, T40 and T60 Vit1 BMG samples (<span class="html-italic">T</span><sub>g</sub>—glass transition temperature; <span class="html-italic">T<sub>x</sub></span>,<sub>1</sub>, <span class="html-italic">T<sub>x</sub></span><sub>,2</sub>, <span class="html-italic">T<sub>x</sub></span><sub>,3</sub>—crystallization onset temperatures). (<b>b</b>) The relaxation enthalpy for the AC and T20, T40 and T60 Vit1 BMG samples.</p> Full article ">Figure 4
<p>The typical load-displacement (<span class="html-italic">P</span>-<span class="html-italic">h</span>) curves of (<b>a</b>) AC, (<b>b</b>)T20, (<b>c</b>) T40 and (<b>d</b>) T60 Vit1 BMG samples at different loading rates under a maximum load of 50 mN.</p> Full article ">Figure 5
<p>(<b>a</b>) The Maxwell-Voigt model used for analyzing creep curves. (<b>b</b>) The typically fitting creep curve of AC Vit1 BMG samples at different loading rates under a maximum load of 50 mN (<span class="html-italic">R<sup>2</sup></span>—correlation coefficient).</p> Full article ">Figure 5 Cont.
<p>(<b>a</b>) The Maxwell-Voigt model used for analyzing creep curves. (<b>b</b>) The typically fitting creep curve of AC Vit1 BMG samples at different loading rates under a maximum load of 50 mN (<span class="html-italic">R<sup>2</sup></span>—correlation coefficient).</p> Full article ">Figure 6
<p>The relaxation spectrum of the (<b>a</b>) AC, (<b>b</b>) T20, (<b>c</b>) T40 Vit1 BMG samples measured at different loading rates with a maximum load of 50 mN, (<b>d</b>) the relaxation spectrum of the AC and pre-compressed samples measured under the loading rate of 0.5 mN/s based on the anelastic part of the creep displacement curves.</p> Full article ">
17 pages, 8727 KiB  
Article
Effects of Thermomechanical Processing on Hydrogen Embrittlement Properties of UltraHigh-Strength TRIP-Aided Bainitic Ferrite Steels
by Tomohiko Hojo, Yutao Zhou, Junya Kobayashi, Koh-ichi Sugimoto, Yoshito Takemoto, Akihiko Nagasaka, Motomichi Koyama, Saya Ajito and Eiji Akiyama
Metals 2022, 12(2), 269; https://doi.org/10.3390/met12020269 - 31 Jan 2022
Cited by 1 | Viewed by 2617
Abstract
The effects of thermomechanical processing on the microstructure and hydrogen embrittlement properties of ultrahigh-strength, low-alloy, transformation-induced plasticity (TRIP)-aided bainitic ferrite (TBF) steels were investigated to apply to automobile forging parts such as engine and drivetrain parts. The hydrogen embrittlement properties were evaluated by [...] Read more.
The effects of thermomechanical processing on the microstructure and hydrogen embrittlement properties of ultrahigh-strength, low-alloy, transformation-induced plasticity (TRIP)-aided bainitic ferrite (TBF) steels were investigated to apply to automobile forging parts such as engine and drivetrain parts. The hydrogen embrittlement properties were evaluated by conducting conventional tensile tests after hydrogen charging and constant load four-point bending tests with hydrogen charging. The 0.4 mass%C-TBF steel achieved refinement of the microstructure, improved retained austenite characteristics, and strengthening, owing to thermomechanical processing. This might be attributed to dynamic and static recrystallizations during thermomechanical processing in TBF steels. Moreover, the hydrogen embrittlement resistances were improved by the thermomechanical processing in TBF steels. This might be caused by the refinement of the microstructure, an increase in the stability of the retained austenite, and low hydrogen absorption of the thermomechanically processed TBF steels. Full article
(This article belongs to the Special Issue Mechanical Properties and Microstructure of Forged Steel)
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<p>Thermomechanical processing and heat treatment diagrams of steels A, B and C. <span class="html-italic">R</span> represents reduction ratio. W. Q. indicates water quenching.</p> Full article ">Figure 2
<p>(<b>a</b>,<b>c</b>,<b>e</b>,<b>g</b>,<b>i</b>,<b>k</b>) Inverse pole figure (IPF) and (<b>b</b>,<b>d</b>,<b>f</b>,<b>h</b>,<b>j</b>,<b>l</b>) phase maps of steels (<b>a</b>,<b>b</b>,<b>g</b>,<b>h</b>) A, (<b>c</b>,<b>d</b>,<b>i</b>,<b>j</b>) B, and (<b>e</b>,<b>f</b>,<b>k</b>,<b>l</b>) C (<b>a</b>–<b>f</b>) without and (<b>g</b>–<b>l</b>) with hot forging. <span class="html-italic">R</span> represents reduction ratio. BCC and FCC denote body-centered-cubic and face-centered-cubic, respectively.</p> Full article ">Figure 3
<p>Nominal stress–strain curves of (<b>a</b>,<b>b</b>,<b>c</b>) conventional and (<b>d</b>,<b>e</b>,<b>f</b>) hot-forged steels (<b>a</b>,<b>d</b>) A, (<b>b</b>,<b>e</b>) B, and (<b>c</b>,<b>f</b>) C with and without hydrogen. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">Figure 4
<p>Variations in (<b>a</b>,<b>b</b>) tensile strength (<span class="html-italic">TS</span>), yield strength (<span class="html-italic">YS</span>), (<b>c</b>,<b>d</b>) total elongation (<span class="html-italic">TEl</span>), and uniform elongation (<span class="html-italic">UEl</span>) as a function of carbon content in conventional and hot-forged steels A, B, and C (<b>a</b>,<b>c</b>) without and (<b>b</b>,<b>d</b>) with hydrogen. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">Figure 5
<p>Fracture surfaces of (<b>a</b>–<b>f</b>) conventional and (<b>g</b>–<b>l</b>) hot-forged steels (<b>a</b>,<b>d</b>,<b>g</b>,<b>j</b>) A, (<b>b</b>,<b>e</b>,<b>h</b>,<b>k</b>) B, and (<b>c</b>,<b>f</b>,<b>i</b>,<b>l</b>) C (<b>d</b>-<b>f</b>,<b>j</b>-<b>l</b>)with and (<b>a</b>–<b>c</b>,<b>g</b>-<b>i</b>) without hydrogen after tensile tests. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">Figure 6
<p>Variations in hydrogen embrittlement susceptibility (<span class="html-italic">HES</span>) as functions of (<b>a</b>) yield strength (<span class="html-italic">YS</span>) and (<b>b</b>) tensile strength (<span class="html-italic">TS</span>) in conventional and hot-forged steels A, B, and C. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">Figure 7
<p>Hydrogen desorption curves of conventional and hot-forged steels A, B, and C charged with hydrogen by a 3% NaCl + 5 g/L NH<sub>4</sub>SCN solution at a current density of 10 A/m<sup>2</sup> for 48 h. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">Figure 8
<p>Applied stress–time to fracture (<span class="html-italic">σ</span><sub>A</sub>–<span class="html-italic">t</span><sub>f</sub>) curves of conventional and hot-forged steels A, B, and C. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">Figure 9
<p>Relationship between delayed fracture strength (<span class="html-italic">DFS</span>) and tensile strength (<span class="html-italic">TS</span>) of conventional and hot-forged steels A, B, and C. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">Figure 10
<p>Fracture surfaces of (<b>a</b>,<b>b</b>) conventional steel B and (<b>c</b>,<b>d</b>) conventional and (<b>e</b>,<b>f</b>) hot-forged steel C after four-point bending tests, in which arrows represent crack initiation region. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">Figure 11
<p>Illustrations of microstructure evolution behavior of hot-forged steel A. (<b>a</b>) Austenite at annealing temperature, (<b>b</b>) final microstructure of conventional TBF steel, (<b>c</b>) microstructure after hot forging, (<b>d</b>) final microstructure of hot-forged TBF steel, (<b>e</b>) TTT diagram, respectively. <span class="html-italic">R</span> represents reduction ratio. α, α<sub>b</sub>, α<sub>bf</sub>, γ, and γ<sub>R</sub> denote ferrite, bainite, bainitic ferrite, austenite, and retained austenite, respectively. The TTT diagram (<b>e</b>) is the time transition temperature diagram.</p> Full article ">Figure 12
<p>Illustrations of microstructure evolution behavior of hot-forged steel C. (<b>a</b>) Austenite at annealing temperature, (<b>b</b>) final microstructure of conventional TBF steel, (<b>c</b>) microstructure after hot forging, (<b>d</b>) final microstructure of hot-forged TBF steel, (<b>e</b>) TTT diagram, respectively. <span class="html-italic">R</span> represents reduction ratio. α, α<sub>b</sub>, α<sub>bf</sub>, γ, and γ<sub>R</sub> denote ferrite, bainite, bainitic ferrite, austenite, and retained austenite, respectively. The TTT diagram (<b>e</b>) is the time transition temperature diagram.</p> Full article ">Figure 13
<p>Variations in hydrogen embrittlement susceptibility (<span class="html-italic">HES</span>) as functions of (<b>a</b>) initial volume fraction (<span class="html-italic">f</span><sub>γ0</sub>), (<b>b</b>) initial carbon concentration (<span class="html-italic">C</span><sub>γ0</sub>), and (<b>c</b>) total carbon content (<span class="html-italic">f</span><sub>γ0</sub> × <span class="html-italic">C</span><sub>γ0</sub>) in retained austenite in steels A, B, and C. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">Figure 14
<p>Variations in diffusible hydrogen concentration (<span class="html-italic">H</span><sub>C</sub>) as a function of initial volume fraction of retained austenite (<span class="html-italic">f</span><sub>γ0</sub>) in steels A, B, and C. <span class="html-italic">R</span> represents reduction ratio.</p> Full article ">
17 pages, 3815 KiB  
Article
Gallic Acid as a Potential Green Corrosion Inhibitor for Aluminum in Acidic Solution
by Przemysław Kwolek, Kamil Dychtoń, Barbara Kościelniak, Andrzej Obłój, Agnieszka Podborska and Marek Wojnicki
Metals 2022, 12(2), 250; https://doi.org/10.3390/met12020250 - 28 Jan 2022
Cited by 15 | Viewed by 4456
Abstract
Plant extracts are intensively studied as green corrosion inhibitors of aluminum. Because these extracts are complex systems, the influence of their individual constituents on the corrosion of aluminum should be determined. In this work, gallic acid was tested for the first time as [...] Read more.
Plant extracts are intensively studied as green corrosion inhibitors of aluminum. Because these extracts are complex systems, the influence of their individual constituents on the corrosion of aluminum should be determined. In this work, gallic acid was tested for the first time as a corrosion inhibitor of aluminum in orthophosphoric acid aqueous solution. So far, its potential inhibiting properties in acidic solutions were only suggested based on promising results obtained for various plant extracts. Evaluation of the potential inhibiting properties of gallic acid was performed using electrochemical methods. The corrosion potential, polarization curves, and impedance spectra of aluminum in 0.5 M orthophosphoric acid, at T = 303 K, were determined. The corrosion potential, corrosion current density, and corrosion rate of aluminum in orthophosphoric acid were equal to −1.151 V vs. Ag|AgCl (3M KCl) reference electrode, 36 μA∙cm−2 and 0.39 mm∙year−1, respectively. These values did not change with the addition of gallic acid. The results obtained show that gallic acid does not inhibit aluminum corrosion. UV-Vis absorption spectra of gallic acid solutions and quantum mechanical calculations show that this organic compound did not adsorb onto the aluminum surface under the studied conditions. Full article
(This article belongs to the Special Issue Electrochemical Processes at Metallic Electrodes—Corrosion and Protection)
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<p>Gallic acid molecule.</p> Full article ">Figure 2
<p>Open circuit potential of aluminum immersed in deaerated 0.5 M H<sub>3</sub>PO<sub>4</sub> aqueous solution as a function of initial concentration of gallic acid and exposition time; agitation rate 300 rpm, <span class="html-italic">T</span> = 303 K.</p> Full article ">Figure 3
<p>Polarization curves of aluminum: (<b>a</b>) as a function of the initial concentration of gallic acid; (<b>b</b>) an example showing how corrosion current density was determined; the red line indicates the approximation of the linear part of the polarization curve. Experiments were performed in deaerated 0.5 M H<sub>3</sub>PO<sub>4</sub> aqueous solutions with an agitation rate of 300 rpm, <span class="html-italic">T</span> = 303 K.</p> Full article ">Figure 4
<p>Scanning electron microscopy (SEM) micrographs of corroded aluminum: (<b>a</b>) <sub>CGA</sub> = 0 mM; (<b>b</b>) <sub>CGA</sub> = 38 mM. Experiments were performed in deaerated 0.5 M H<sub>3</sub>PO<sub>4</sub> aqueous solutions with an agitation rate of 300 rpm, <span class="html-italic">T</span> = 303 K.</p> Full article ">Figure 5
<p>Impedance spectra of aluminum: (<b>a</b>) recorded at OCP and subsequent anodic polarization; (<b>b</b>) recorded at OCP and subsequent cathodic polarization. Continuous lines indicate an approximation of these spectra using appropriate electrical equivalent circuit. Experiments were performed in a deaerated 0.5 M H<sub>3</sub>PO<sub>4</sub> aqueous solution with an agitation rate of 300 rpm, <span class="html-italic">T</span> = 303 K.</p> Full article ">Figure 6
<p>Electrical equivalent circuits used for approximation of impedance spectra: (<b>a</b>) recorded at open circuit potential; (<b>b</b>) recorded at η = +0.1 V; (<b>c</b>) recorded at η = −0.1 V.</p> Full article ">Figure 7
<p>Absorption spectra of aqueous solutions of gallic acid: (<b>a</b>) measured experimentally as a function of pH, <sub>CGA</sub> = 65 μM; (<b>b</b>) obtained using quantum mechanical calculations.</p> Full article ">Figure 8
<p>Electrical charge distribution in gallic acid molecule; C, O, and H atoms are depicted as grey, red, and white balls, respectively.</p> Full article ">Figure 9
<p>Speciation diagram for Al-H<sub>3</sub>PO<sub>4</sub>-GA-H<sub>2</sub>O system. Calculations were performed for 0.5 M H<sub>3</sub>PO<sub>4</sub> aqueous solution. Total aluminum concentration of 5 mM was obtained from corrosion rate, assuming <span class="html-italic">t</span> = 3 h, the total gallic acid concentration was 38 mM. Stability constants for complex ions are from [<a href="#B40-metals-12-00250" class="html-bibr">40</a>,<a href="#B46-metals-12-00250" class="html-bibr">46</a>].</p> Full article ">Figure 10
<p>Impedance spectra of aluminum in 0.1 M HCl aqueous solution recorded at OCP and subsequent cathodic polarization. Continuous lines indicate an approximation of these spectra using appropriate electrical equivalent circuit. Experiments were performed in deaerated, agitated solutions with an agitation rate of 300 rpm, <span class="html-italic">T</span> = 303 K.</p> Full article ">
10 pages, 11230 KiB  
Perspective
A Prospective Way to Achieve Ballistic Impact Resistance of Lightweight Magnesium Alloys
by Abdul Malik, Faisal Nazeer and Yangwei Wang
Metals 2022, 12(2), 241; https://doi.org/10.3390/met12020241 - 27 Jan 2022
Cited by 10 | Viewed by 3736
Abstract
The ballistic impact resistance of lightweight magnesium alloys is an eye-catching material for the military and aerospace industries, which can decrease the cost of a project and the fuel consumption. The shockwave mitigation ability of a magnesium alloy is 100 times stronger than [...] Read more.
The ballistic impact resistance of lightweight magnesium alloys is an eye-catching material for the military and aerospace industries, which can decrease the cost of a project and the fuel consumption. The shockwave mitigation ability of a magnesium alloy is 100 times stronger than an aluminum alloy; nonetheless, ballistic impact resistance has still not been achieved against blunt and API projectiles. The major obstacles are the low hardness, low mechanical strength, basal texture and strain hardening ability under loading along the normal direction of the sheet. The high yield strength and ultimate strength can be achieved for a specific loading condition (tensile or compression) by adjusting the texture in magnesium alloys. The projectile impact along the normal direction in a strong basal-textured magnesium alloy can only produce a slip-induced deformation or minor twinning activity. Here, we propose a practical technique that can be valuable for altering the texture from c-axes//ND to c-axes//ED or TD, and can produce high strain hardening and high strength through a twinning and de-twinning activity. Subsequently, it can improve the ballistic impact resistance of magnesium alloys. The effect of the technique on the evolution of the microstructure and possible anticipated deformation mechanisms after ballistic impact is proposed and discussed. Full article
(This article belongs to the Special Issue Development and Applications of New Lightweighting Metal Technologies: High Strength Al and Mg Alloys)
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<p>The macro images of deformed specimens: (<b>a</b>–<b>i</b>) <span class="html-italic">AZ31 Mg</span> alloy at given conditions against <span class="html-italic">2017Al</span> projectile [<a href="#B27-metals-12-00241" class="html-bibr">27</a>] (Copyright 2022 Elsevier); (<b>j</b>–<b>m</b>) front and rear view, perforation channel and backing plate images of <span class="html-italic">Mg–Zn–Zr Mg</span> alloy (thickness 30 mm) against standard 7.62 mm soft steel core projectile under standard velocity 830 m/s [<a href="#B17-metals-12-00241" class="html-bibr">17</a>] (Copyright 2022 Elsevier).</p> Full article ">Figure 2
<p>Schematic illustration of pre-compression process and texture of <span class="html-italic">Mg</span> alloys.</p> Full article ">Figure 3
<p>The different <span class="html-italic">Mg</span> alloys subjected to pre-straining under different loading conditions [<a href="#B35-metals-12-00241" class="html-bibr">35</a>,<a href="#B38-metals-12-00241" class="html-bibr">38</a>,<a href="#B39-metals-12-00241" class="html-bibr">39</a>].</p> Full article ">Figure 4
<p>Schematic illustration of the anticipated deformation mechanism after the ballistic impact of (<b>a</b>) non-pre-compressed (<b>b</b>) pre-compressed <span class="html-italic">Mg</span> alloy.</p> Full article ">Figure 5
<p>Schematic illustration of texture, strength, elongation, slips, and deformation mechanism in differently processed <span class="html-italic">Mg</span> alloys.</p> Full article ">
24 pages, 14256 KiB  
Article
Structural and Tribological Studies of “(TiC + WC)/Hardened Steel” PMMC Coating Deposited by Air Pulsed Plasma
by Yuliia Chabak, Vasily Efremenko, Vadym Zurnadzhy, Viktor Puchý, Ivan Petryshynets, Bohdan Efremenko, Victor Fedun, Kazumichi Shimizu, Iurii Bogomol, Volodymyr Kulyk and Dagmar Jakubéczyová
Metals 2022, 12(2), 218; https://doi.org/10.3390/met12020218 - 24 Jan 2022
Cited by 24 | Viewed by 3290
Abstract
The deposition of a thin (several tens of microns) protective coating in atmospheric conditions is a challenging task for surface engineering. The structural features and tribological properties of a particle-reinforced metal matrix composite coating synthesized on middle-carbon steel by air pulse-plasma treatments were [...] Read more.
The deposition of a thin (several tens of microns) protective coating in atmospheric conditions is a challenging task for surface engineering. The structural features and tribological properties of a particle-reinforced metal matrix composite coating synthesized on middle-carbon steel by air pulse-plasma treatments were studied in the present work. The 24–31 µm thick coating of “24 vol.% (TiC + WC)/Hardened steel matrix” was produced by 10 plasma pulses generated by an electro-thermal axial plasma accelerator equipped with a consumable cathode of novel design (low-carbon steel tube filled with “TiC/WC + Epoxy resin” mixture). The study included optical microscopy (OM), scanning electron microscopy (SEM), energy-dispersive X-ray spectroscopy (EDX), X-ray diffraction (XRD, microhardness measurements, and dry “Ball-on-Plate” testing. The carbides were directly plasma-transferred to the substrate (steel of AISI 4145H grade) from the cathode without substantial melting. The hard (500–1044 HV) coating matrix consisted of 57 vol.% austenite (1.43 wt.% C) and 43 vol.% plate martensite was formed via carbon enrichment of steel from plasma flow. Additionally, a minor amount of oxide phases (TiO2, WO2, WO3) were dispersed in the matrix. As compared to substrate, the coating had a lower coefficient of friction; its volumetric wear was decreased by 4.4 times when sliding against hardened steel ball and by 16 times when sliding against SiC ball. Full article
(This article belongs to the Special Issue Cavitation Erosion, Abrasive and Sliding Wear Behaviour of Metal-Based Structures)
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<p>The principal scheme of EAPA and its electrical installation: 1, 7—steel shells, 2—cathode, 3—paper reinforced bakelite, 4—copper mount, 5—sheet copper cylinder, 6—consumable part of the cathode, 8—EAPA edge, 9—the target (the substrate) (ICG—impulse current generator, IVG—impulse voltage generator).</p> Full article ">Figure 2
<p>The solid model of the consumable part of the EAPA cathode.</p> Full article ">Figure 3
<p>The consumable part of the EAPA cathode: the microstructure of (<b>a</b>) a steel tube and (<b>b</b>) a filler, (<b>c</b>) carbide size distributions for the filler and the coating, and (<b>d</b>) general view of the ready cathode before PPD (pulsed-plasma deposition) and the eroded cathode after deposition of five coatings (50 pulses). (<b>e</b>) The microstructure of a substrate (steel of AISI 4145H grade).</p> Full article ">Figure 4
<p>The scheme of the volumetric wear calculation using the wear track’s profile.</p> Full article ">Figure 5
<p>The coating microstructure: (<b>a</b>,<b>d</b>) cross-sectional direction, (<b>b</b>,<b>c</b>) a view from above. ((<b>a</b>)—SEM (secondary electron image—SEI); (<b>b</b>–<b>d</b>)—OM, (<b>e</b>)—SEM (back-scattered electron image—BSEI)).</p> Full article ">Figure 6
<p>The coating microhardness: distribution of the values for the matrix and carbide particles.</p> Full article ">Figure 7
<p>XRD pattern of the coating.</p> Full article ">Figure 8
<p>Elemental EDX-mapping of the coating: (<b>a</b>) SEM(BSEI), distribution of (<b>b</b>) Ti, (<b>c</b>) W, (<b>d</b>) Fe, (<b>e</b>) C, (<b>f</b>) O.</p> Full article ">Figure 9
<p>The EDX-profiles of C, W, Ti and Fe within the coating: (<b>a</b>) across WC and TiC carbides, (<b>b</b>) across carbide-free matrix and “carbide network” matrix.</p> Full article ">Figure 10
<p>EDX-spectra for the indicative points shown in <a href="#metals-12-00218-f008" class="html-fig">Figure 8</a>: (<b>a</b>) WC carbide (point 004), (<b>b</b>) TiC carbide (point 005), (<b>c</b>) “carbide network”-matrix (point 013), (<b>d</b>) carbide-free matrix (point 008).</p> Full article ">Figure 11
<p>Variation of a mean CoF value during the wear test: (<b>a</b>) sliding against 100Cr6 ball, (<b>b</b>) sliding against SiC ball.</p> Full article ">Figure 12
<p>3D images of the wear tracks for different sliding couples: (<b>a</b>) 100Cr6-substrate (steel of AISI 4145H grade), (<b>b</b>) SiC-substrate (steel of AISI 4145H grade), (<b>c</b>) 100Cr6-coating, (<b>d</b>) SiC-coating. Dry-sliding wear characteristics of the specimens: (<b>e</b>) wear track profiles, and (<b>f</b>) volumetric wear (The values of a profile height in Figure (<b>a</b>–<b>d</b>) are given in µm).</p> Full article ">Figure 13
<p>Worn surface of the counter balls: (<b>a</b>) 100Cr6-substrate (steel of AISI 4145H grade), (<b>b</b>) SiC-substrate (steel of AISI 4145H grade), (<b>c</b>) 100Cr6-coating, (<b>d</b>) SiC-coating.(<b>e</b>) Mean surface area of the wear spot on the ball after sliding test.</p> Full article ">Figure 14
<p>Optical micrographs of the wear debris for different sliding couples: (<b>a</b>) 100Cr6-substrate (steel of AISI 4145H grade), (<b>b</b>) SiC-substrate (steel of AISI 4145H grade), (<b>c</b>) 100Cr6-coating, (<b>d</b>) SiC-coating.</p> Full article ">Figure 15
<p>Oxide films on the worn surface of the specimens: (<b>a</b>) 100Cr6-substrate (steel of AISI 4145H grade), (<b>b</b>) SiC-substrate (steel of AISI 4145H grade), (<b>c</b>) 100Cr6-coating, (<b>d</b>) SiC-coating.</p> Full article ">
15 pages, 10670 KiB  
Article
The Microstructure and Mechanical Properties of 5083, 6005A and 7N01 Aluminum Alloy Gas Metal Arc-Welded Joints for High-Speed Train: A Comparative Study
by Laijun Wu, Biao Yang, Xiaohui Han, Guolong Ma, Bingxiao Xu, Yuhang Liu, Xiaoguo Song and Caiwang Tan
Metals 2022, 12(2), 213; https://doi.org/10.3390/met12020213 - 24 Jan 2022
Cited by 19 | Viewed by 4754
Abstract
This study aimed to conduct a comparative study on the microstructure and mechanical performance of 5083, 6005A and 7N01 Al joints used in China Railway High-speed (CRH) trains. We connected 10 mm-thick plates by three-layer and three-pass gas metal arc welding (GMAW). The [...] Read more.
This study aimed to conduct a comparative study on the microstructure and mechanical performance of 5083, 6005A and 7N01 Al joints used in China Railway High-speed (CRH) trains. We connected 10 mm-thick plates by three-layer and three-pass gas metal arc welding (GMAW). The results indicated that 6005A and 7N01 Al joints were more sensitive to grain boundary liquation in the partially melted zone (PMZ) than 5083 Al joins. Besides, recrystallization was obtained in heat-affected zones (HAZ). The 5083 Al joints experienced the most severe recrystallization and the grain size changed from 6.32 (BM) to 32.44 (HAZ) μm duo to intracrystalline strain induced by cold-rolled processes. The 7N01 Al alloys experienced the lowest extent of recrystallization and the grain size increased from 5.32 (BM) to 22.31 (HAZ) μm. Furthermore, significant precipitate evolution in the HAZ was observed. Original thin β” precipitates dissolved in HAZ of 6005A Al joints and transformed to the softer β phase. However, less precipitation transition was examined in 5083 and 7N01 Al joints. The precipitates’ evolution produced a softening region in HAZ of 6005A joints where the hardness was only 55 HV. The microhardness profile of the other two Al joints was less affected. The tensile strength of 5083, 6005A, and 7N01 Al alloy joints reached 323, 206 and 361 MPa, respectively. The 5083 Al and 6005A Al joints failed at HAZ near the fusion line while 7N01 Al joints failed at the fusion zone owing to the high strength of the base metal. The liquation, coarse grains by recrystallization, and precipitate evolution all decreased local strength, resulting in the fracture at HAZ. Full article
(This article belongs to the Special Issue Advanced Welding Technology in Metals II)
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<p>The schematic of gas metal arc welding (GMAW) of aluminum alloy. (<b>a</b>)welding diagram; (<b>b</b>) three-layer and three-pass welding layer arrangement.</p> Full article ">Figure 2
<p>Geometry of the tensile specimens (mm).</p> Full article ">Figure 3
<p>The weld appearance and cross section of welding joints: (<b>a</b>,<b>b</b>) for 5083; (<b>c</b>,<b>d</b>) for 6005A; (<b>e</b>,<b>f</b>) for 7N01.</p> Full article ">Figure 4
<p>The microstructure of different aluminum alloy welded joints: 5083: (<b>a</b>) fusion zone (FZ), (<b>b</b>) partially melted zone (PMZ), (<b>c</b>) heat-affected zone (HAZ); 6005A: (<b>d</b>) FZ, (<b>e</b>) PMZ, (<b>f</b>) HAZ; 7N01: (<b>g</b>) FZ, (<b>h</b>) PMZ, (<b>i</b>) HAZ.</p> Full article ">Figure 5
<p>Energy-dispersive X-ray spectroscopy (EDS) scanning results of different aluminum alloy welded joints: (<b>a</b>) 5083; (<b>b</b>) 6005A; (<b>c</b>) 7N01.</p> Full article ">Figure 6
<p>The differential scanning calorimetry (DSC) curves of different aluminum alloy joints: (<b>a</b>) 5083; (<b>b</b>) 6005A; (<b>c</b>) 7N01.</p> Full article ">Figure 7
<p>Electron backscatter diffraction detector (EBSD) images of aluminum alloy welding joints: 5083: (<b>a</b>) base metal (BM); (<b>d</b>) HAZ; 6005A: (<b>b</b>) BM; (<b>e</b>) HAZ; 7N01: (<b>c</b>) BM; (<b>f</b>) HAZ.</p> Full article ">Figure 8
<p>The grain size of different aluminum alloy welded joints: 5083: (<b>a</b>) BM; (<b>b</b>) HAZ; 6005A: (<b>c</b>) BM; (<b>d</b>) HAZ; 7N01: (<b>e</b>) BM; (<b>f</b>) HAZ.</p> Full article ">Figure 9
<p>EDS results of different aluminum alloy welded joints: 5083: (<b>a</b>) BM; (<b>d</b>) HAZ; 6005A: (<b>b</b>) BM; (<b>e</b>) HAZ; 7N01: (<b>c</b>) BM, (<b>f</b>) HAZ.</p> Full article ">Figure 10
<p>Bright-field transmission electron microscopy (TEM) micrograph of different aluminum alloy welded joints: 5083: (<b>a</b>) BM; (<b>d</b>) HAZ; 6005A: (<b>b</b>) BM; (<b>e</b>) HAZ; 7N01: (<b>c</b>) BM, (<b>f</b>) HAZ.</p> Full article ">Figure 11
<p>Microhardness profile of different aluminum alloy welded joints: (<b>a</b>) 5083; (<b>b</b>) 6005A; (<b>c</b>) 7N01.</p> Full article ">Figure 12
<p>The displacement-strength curves of 5083, 6005A and 7N01 Al alloys welded joints.</p> Full article ">Figure 13
<p>Fracture position of different aluminum alloy welded joints: (<b>a</b>) 5083; (<b>b</b>) 6005A; (<b>c</b>) 7N01.</p> Full article ">Figure 14
<p>The fractured surface of different Al alloy joints: (<b>a</b>) and (<b>d</b>) for 5083; (<b>b</b>) and (<b>e</b>) for 6005A; (<b>c</b>) and (<b>f</b>) for 7N01.</p> Full article ">Figure 15
<p>Aluminum alloys pillow beam in China Railway High-Speed trains.</p> Full article ">
15 pages, 1232 KiB  
Review
Viewpoints on Technological Aspects of Advanced High-Strength Bainitic Steels
by Lucia Morales-Rivas
Metals 2022, 12(2), 195; https://doi.org/10.3390/met12020195 - 21 Jan 2022
Cited by 8 | Viewed by 4133
Abstract
The development of advanced high-strength bainitic steels has been preceded and linked to different metallurgical advances, both in the field of fundamental materials science and in technological fields closer to the production and final application. The diversity and abundance of documents in literature [...] Read more.
The development of advanced high-strength bainitic steels has been preceded and linked to different metallurgical advances, both in the field of fundamental materials science and in technological fields closer to the production and final application. The diversity and abundance of documents in literature has favored the co-existence of extensive terminology in the context of advanced high-strength steels and bainitic steels. In this work, the concept of advanced high-strength bainitic steels is briefly revisited from a wide perspective, with the aim of highlighting the main limitations and challenges for further development of these microstructures. Full article
(This article belongs to the Special Issue Advanced High-Strength Bainitic Steels)
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<p>Secondary electron SEM images of Nital-etched metallographic surface of: (<b>a</b>) granular CFB obtained by continuous cooling, as reported in adapted from [<a href="#B17-metals-12-00195" class="html-bibr">17</a>]; (<b>b</b>) nanostructured bainite from a steel with 1C-2.5Si wt.% (among other elements) treated at 250 °C during 16 h, as reported adapted from [<a href="#B18-metals-12-00195" class="html-bibr">18</a>]; (<b>c</b>) nanostructured bainite from a steel with 0.7C-1.4Si wt.% (among other elements) treated at 220 °C during supplementary time, as reported in [<a href="#B19-metals-12-00195" class="html-bibr">19</a>]. BF stands for bainitic ferrite.</p> Full article ">Figure 2
<p>Mechanical performance of conventional steels and AHSS, adapted from [<a href="#B35-metals-12-00195" class="html-bibr">35</a>,<a href="#B36-metals-12-00195" class="html-bibr">36</a>,<a href="#B37-metals-12-00195" class="html-bibr">37</a>].</p> Full article ">Figure 3
<p>Schematic illustration of advanced high-strength bainitic steels (within the dashed blue line) vs. C content and size of the cross-section of the part.</p> Full article ">
12 pages, 6292 KiB  
Article
Performance of a Nozzle to Control Bath Level Oscillations and Turbulence of the Metal-Flux Interface in Slab Molds
by María Guadalupe González-Solórzano, Rodolfo Davila Morales, Javier Guarneros, Carlos Rodrigo Muñiz-Valdés and Alfonso Nájera Bastida
Metals 2022, 12(1), 140; https://doi.org/10.3390/met12010140 - 12 Jan 2022
Cited by 3 | Viewed by 1863
Abstract
The characterization of the turbulent flow of liquid steel in a slab mold using a commercial nozzle was carried out through physical experiments and mathematical models. Six ultrasonic sensors were located at each side of the nozzle to obtain real-time plotting of the [...] Read more.
The characterization of the turbulent flow of liquid steel in a slab mold using a commercial nozzle was carried out through physical experiments and mathematical models. Six ultrasonic sensors were located at each side of the nozzle to obtain real-time plotting of the bath levels during the experimental time. An ultrasonic transducer located in the mold, 20 mm below the meniscus, determines the velocities and the turbulent variables along with the distance from the narrow face to the position of the nozzle’s outer wall. These data, together with the mathematical simulations, demonstrated a high correlation of bath level oscillations and the time-dependent behavior of the discharging jets. The flow inside the mold shows low-frequency non-symmetric patterns without a severe turbulent in the meniscus. The source of this instability is the partial opening of the slide valve gate used to control the mass flow of liquid. Full article
(This article belongs to the Special Issue Advanced Tundish Metallurgy and Clean Steel Technology)
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<p>The nozzle and the experimental setup. (<b>a</b>) Geometric features of the industrial nozzle (mm). (<b>b</b>) The mold model equipped with meniscus levels sensors and the ultrasonic transducer to measure turbulence. (<b>c</b>) Interior wall of the nozzle.</p> Full article ">Figure 2
<p>Tracer mixing in the mold model: (<b>a</b>) 0.5 s after injection, (<b>b</b>) 1.5 s after injection, (<b>c</b>) 4.5 s after injection. (<b>d</b>) Numerical velocity field.</p> Full article ">Figure 3
<p>Variations of meniscus level and their covariances: (<b>a</b>) Sensors 1 and 6, (<b>b</b>) Sensors 2 and 5, (<b>c</b>) Sensors 3 and 4. (<b>d</b>) Covariances of bath level fluctuations.</p> Full article ">Figure 4
<p>Velocities calculated along a line located 20 mm below the meniscus by the mathematical model: (<b>a</b>) right port, (<b>b</b>) left port. (<b>c</b>) Comparison of meniscus velocities with critical velocities to induce powder entrainment.</p> Full article ">Figure 5
<p>Numerical velocity fields in the central plane of the mold at different times ((<b>a</b>) 20 s; (<b>b</b>) 80 s; (<b>c</b>) 140 s; (<b>d</b>) 200 s; (<b>e</b>) 260 s; (<b>f</b>) 300 s) are indicated in each image.</p> Full article ">Figure 6
<p>Numerical velocity fields in the nozzle: (<b>a</b>) longitudinal plane axis of symmetry, (<b>b</b>) transverse plane axis of symmetry. (<b>c</b>) Flow in the central plane of the internal axis of symmetry through the right port. (<b>d</b>) Flow through the outer port plane of the right port. (<b>e</b>) Flow in the central plane of the internal axis of symmetry through the left port. (<b>f</b>) Flow through the outer port plane of the left port.</p> Full article ">Figure 7
<p>Turbulence features. (<b>a</b>) PDF of velocity in point L33. (<b>b</b>) PDF of velocity in point L60. (<b>c</b>) measured velocity profile along the line located 20 mm below the meniscus.</p> Full article ">Figure 8
<p>Autocorrelation of velocities: (<b>a</b>) at point L33, (<b>b</b>) at point L60.</p> Full article ">
17 pages, 8878 KiB  
Article
The Research on Recrystallization Behaviors and Mechanism of a Medium-Density Ni-Based Alloy
by Kai Feng, Xiaxu Huang, Rui Wang, Wenli Xue, Yilei Fu and Zhaoxin Li
Metals 2022, 12(1), 137; https://doi.org/10.3390/met12010137 - 11 Jan 2022
Cited by 2 | Viewed by 2365
Abstract
Revealing the recrystallization behavior and mechanism of this new alloy is of great significance to subsequent research. In this study, the Ni-36.6W-15Co ternary medium heavy alloy was solution-treated at 1100–1200 °C for different lengths of time. The grain size change, microstructure and texture [...] Read more.
Revealing the recrystallization behavior and mechanism of this new alloy is of great significance to subsequent research. In this study, the Ni-36.6W-15Co ternary medium heavy alloy was solution-treated at 1100–1200 °C for different lengths of time. The grain size change, microstructure and texture evolution as well as twin development during recrystallization annealing were analyzed using SEM, EBSD and TEM techniques. The study found that complete recrystallization occurs at 1150 °C/60 min. In addition, it takes a longer amount of time for complete recrystallization to occur at 1100 °C. The value of the activation energy Q1 of the studied alloys is 701 kJ/mol and the recrystallization process is relatively slow. By comparing the changes of microstructure and texture with superalloys, it is found that the recrystallization mechanism of the studied alloy is different from that of the superalloy. The development of annealing twins has a great influence on the recrystallization behavior and mechanism. The results show that the twin mechanism is considered as the dominant recrystallization mechanism of the studied alloy, although the formation and development of sub-grains appear in the early stage of recrystallization. Full article
(This article belongs to the Special Issue Thermomechanical Processing, Microstructure Evolution and Mechanical Properties of Alloys)
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<p>The secondary electron image of sample in (<b>a</b>) rolled state, and (<b>b</b>) is a large version of the red box in (<b>a</b>).</p> Full article ">Figure 2
<p>The microstructure of the Ni-36.6W-15Co sample after annealing at (<b>a</b>) 1100 °C, (<b>b</b>) 1150 °C and (<b>c</b>) 1200 °C for 60 min.</p> Full article ">Figure 3
<p>The microstructure of the Ni-36.6W-15Co sample after annealing for (<b>a</b>) 10 min, (<b>b</b>) 30 min, (<b>c</b>) 60 min and (<b>d</b>) 120 min at 1150 °C.</p> Full article ">Figure 4
<p>Relationship between annealing temperature and average grain size.</p> Full article ">Figure 5
<p>Grain and grain boundary network of Ni-36.6W-15Co after annealing for (<b>a</b>) 10 min, (<b>b</b>) 30 min, (<b>c</b>) 60 min and (<b>d</b>) 120 min at 1150 °C.</p> Full article ">Figure 6
<p>Grain size distribution maps Ni-36.6W-15Co after annealing for (<b>a</b>) 10 min, (<b>b</b>) 30 min, (<b>c</b>) 60 min, (<b>d</b>) 120 min at 1150 °C and (<b>e</b>) line chart of average grain size.</p> Full article ">Figure 7
<p>Relationship between annealing time and hardness at 1150 °C.</p> Full article ">Figure 8
<p>Orientation imaging maps of Ni-36.6W-15Co after annealing for (<b>a</b>) 10 min, (<b>b</b>) 30 min, (<b>c</b>) 60 min and (<b>d</b>) 120 min at 1150 °C.</p> Full article ">Figure 8 Cont.
<p>Orientation imaging maps of Ni-36.6W-15Co after annealing for (<b>a</b>) 10 min, (<b>b</b>) 30 min, (<b>c</b>) 60 min and (<b>d</b>) 120 min at 1150 °C.</p> Full article ">Figure 9
<p>Grain boundary maps of Ni-36.6W-15Co after annealing for (<b>a</b>) 10 min, (<b>b</b>) 30 min, (<b>c</b>) 60 min and (<b>d</b>) 120 min at 1150 °C.</p> Full article ">Figure 10
<p>Variations of density of low-CSL boundary with annealing time.</p> Full article ">Figure 11
<p>The pole figure (PF) of Ni-36.6W-15Co in (<b>a</b>) rolled, and after annealing for (<b>b</b>) 10 min, (<b>c</b>) 30 min, (<b>d</b>) 60 min and (<b>e</b>) 120 min at 1150 °C.</p> Full article ">Figure 12
<p>The orientation distribution function (ODF) of Ni-36.6W-15Co in (<b>a</b>) rolled, and after annealing for (<b>b</b>) 10 min, (<b>c</b>) 30 min, (<b>d</b>) 60 min and (<b>e</b>) 120 min at 1150 °C.</p> Full article ">Figure 13
<p>The TEM bright field photograph of Ni-36.6W-15Co in (<b>a</b>) rolled, and after annealing for (<b>b</b>) 10 min, (<b>c</b>) 30 min and (<b>d</b>) 120 min at 1150 °C.</p> Full article ">Figure 14
<p>Comparisons of the experimental and predicted average grain sizes.</p> Full article ">Figure 15
<p>The TEM bright field photograph of Ni-36.6W-15Co after annealing at 1150 °C for 10 min, (<b>a</b>) selected area diffraction patters (SADP) of the twin and (<b>b</b>) is HRTEM image of the stacking fault.</p> Full article ">
46 pages, 6639 KiB  
Review
Recent Developments and Future Challenges in Incremental Sheet Forming of Aluminium and Aluminium Alloy Sheets
by Tomasz Trzepieciński, Sherwan Mohammed Najm, Valentin Oleksik, Delia Vasilca, Imre Paniti and Marcin Szpunar
Metals 2022, 12(1), 124; https://doi.org/10.3390/met12010124 - 9 Jan 2022
Cited by 26 | Viewed by 6334
Abstract
Due to a favourable strength-to-density ratio, aluminium and its alloys are increasingly used in the automotive, aviation and space industries for the fabrication of skins and other structural elements. This article explores the opportunities for and limitations of using Single- and Two Point [...] Read more.
Due to a favourable strength-to-density ratio, aluminium and its alloys are increasingly used in the automotive, aviation and space industries for the fabrication of skins and other structural elements. This article explores the opportunities for and limitations of using Single- and Two Point Incremental Sheet Forming techniques to form sheets from aluminium and its alloys. Incremental Sheet Forming (ISF) methods are designed to increase the efficiency of processing in low- and medium-batch production because (i) it does not require the production of a matrix and (ii) the forming time is much higher than in conventional methods of sheet metal forming. The tool in the form of a rotating mandrel gradually sinks into the sheet, thus leading to an increase in the degree of deformation of the material. This article provides an overview of the published results of research on the influence of the parameters of the ISF process (feed rate, tool rotational speed, step size), tool path strategy, friction conditions and process temperature on the formability and surface quality of the workpieces. This study summarises the latest development trends in experimental research on, and computer simulation using, the finite element method of ISF processes conducted in cold forming conditions and at elevated temperature. Possible directions for further research are also identified. Full article
(This article belongs to the Special Issue Challenges and Achievements in Metal Forming)
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<p>Main methods of incremental sheet forming.</p> Full article ">Figure 2
<p>(<b>a</b>) SPIF, (<b>b</b>) counter tool TPIF, (<b>c</b>) partial die TPIF, (<b>d</b>) full die TPIF: 1—forming tool, 2—blank holder, 3—backing plate, 4—workpiece, 5—rig frame, 6—counter tool, 7—partial die, 8—full die.</p> Full article ">Figure 3
<p>Process parameters of WJISF (reprinted with permission from [<a href="#B64-metals-12-00124" class="html-bibr">64</a>]; Copyright © 2022, Springer-Verlag London Limited).</p> Full article ">Figure 4
<p>Comparison between ISF and WJISF (reprinted with permission from [<a href="#B64-metals-12-00124" class="html-bibr">64</a>]; Copyright © 2022, Springer-Verlag London Limited).</p> Full article ">Figure 5
<p>Laminated tools with various shapes and parts formed using different WJ trajectories (reprinted with permission from [<a href="#B68-metals-12-00124" class="html-bibr">68</a>]; copyright © 2022, Springer-Verlag London Limited).</p> Full article ">Figure 6
<p>Electromagnetic incremental sheet forming technology: (<b>a</b>) the initial position of the coil and positions after rotation through (<b>b</b>) 90°, (<b>c</b>) 180°, (<b>d</b>) 270°, and (<b>e</b>) the region of plastic deformation of workpiece (reprinted with permission from [<a href="#B84-metals-12-00124" class="html-bibr">84</a>]; copyright © 2022 The Author(s). Published by Elsevier Ltd.).</p> Full article ">Figure 7
<p>Schematic of the Tension under Cyclic Bending and Compression test concept (reprinted with permission from [<a href="#B96-metals-12-00124" class="html-bibr">96</a>]; copyright © 2022 Elsevier B.V. All rights reserved.).</p> Full article ">Figure 8
<p>The three strategies used (<b>a</b>) A strategy, (<b>b</b>) B strategy and (<b>c</b>) C strategy (reprinted with permission from [<a href="#B106-metals-12-00124" class="html-bibr">106</a>]; copyright © 2022 The Society of Manufacturing Engineers. Published by Elsevier Ltd. All rights reserved).</p> Full article ">Figure 9
<p>Scheme of local springback in SPIF (reprinted with permission from [<a href="#B130-metals-12-00124" class="html-bibr">130</a>]; copyright © 2022, German Academic Society for Production Engineering (WGP)).</p> Full article ">Figure 10
<p>The ANN structure for modelling the effect of process parameters on the surface roughness of components: (<b>a</b>) two outputs, (<b>b</b>) one output (reprinted from [<a href="#B148-metals-12-00124" class="html-bibr">148</a>]; copyright © 2022, The Authors, this is an open access article distributed under the terms of the Creative Commons CC BY license, which permits unrestricted use, distribution and reproduction in any medium, provided the original work is properly cited).</p> Full article ">Figure 11
<p>Potential surface problem in non-symmetrical part (reprinted with permission from [<a href="#B155-metals-12-00124" class="html-bibr">155</a>]; copyright © 2022, Elsevier B.V. Published by Elsevier B.V. All rights reserved).</p> Full article ">Figure 12
<p>An in situ springback compensation method (reprinted with permission from [<a href="#B157-metals-12-00124" class="html-bibr">157</a>]; copyright © 2022, Published by Elsevier Ltd. on behalf of CIRP).</p> Full article ">Figure 13
<p>Schematic outline of the pillow effect (reprinted with permission from [<a href="#B163-metals-12-00124" class="html-bibr">163</a>]; copyright © 2022, Elsevier B.V. All rights reserved).</p> Full article ">Figure 14
<p>Schematic diagram of the pillow effect (reprinted with permission from [<a href="#B188-metals-12-00124" class="html-bibr">188</a>]; copyright © 2022, Springer-Verlag London).</p> Full article ">Figure 15
<p>Roughness for different speeds of tool rotation (reprinted with permission from [<a href="#B195-metals-12-00124" class="html-bibr">195</a>]; copyright © 2022, Elsevier B.V. All rights reserved).</p> Full article ">Figure 16
<p>The ORB tool (reprinted from [<a href="#B203-metals-12-00124" class="html-bibr">203</a>], this is an open access article distributed under the terms of the Creative Commons CC-BY license, which permits unrestricted use, distribution and reproduction in any medium, provided the original work is properly cited).</p> Full article ">Figure 17
<p>Effect of SPIF parameters on (<b>a</b>) the arithmetical mean roughness of the drawpiece surface Ra and (<b>b</b>) the ten point roughness of a drawpiece surface Rz (reprinted from [<a href="#B148-metals-12-00124" class="html-bibr">148</a>]), this is an open access article distributed under the terms of the Creative Commons CC BY license, which permits unrestricted use, distribution and reproduction in any medium, provided the original work is properly cited.</p> Full article ">Figure 18
<p>(<b>a</b>) Equipment for electrically assisted hot ISF (reprinted with permission from [<a href="#B40-metals-12-00124" class="html-bibr">40</a>]; copyright © 2022, Elsevier Ltd. All rights reserved) and (<b>b</b>) the experimental equipment used in the experiments: A—base structure, B—tool holder, C—workpiece, D—blank holder, E—forming tool, F—Anode, G—cathode (reprinted with permission from [<a href="#B228-metals-12-00124" class="html-bibr">228</a>]; copyright © 2022, Springer-Verlag London Ltd., part of Springer Nature).</p> Full article ">Figure 19
<p>(<b>a</b>) Distribution of surface textures on the forming tool tip (reprinted with permission from [<a href="#B222-metals-12-00124" class="html-bibr">222</a>]; copyright © 2022, Elsevier Ltd. All rights reserved) and (<b>b</b>) surface fish scales (reprinted with permission from [<a href="#B241-metals-12-00124" class="html-bibr">241</a>]; copyright © 2022, Elsevier B.V. All rights reserved).</p> Full article ">
11 pages, 8506 KiB  
Article
Transmission Electron Microscopy Study on the Precipitation Behaviors of Laser-Welded Ferritic Stainless Steels and Their Implications on Intergranular Corrosion Resistance
by Niklas Sommer, Clementine Warres, Tarek Lutz, Martin Kahlmeyer and Stefan Böhm
Metals 2022, 12(1), 86; https://doi.org/10.3390/met12010086 - 4 Jan 2022
Cited by 6 | Viewed by 2519
Abstract
The intergranular corrosion susceptibility of ferritic stainless-steel weldments is strongly dependent on chromium carbide precipitation phenomena. Hence, stabilization is widely used to mitigate the aforementioned precipitation. In contrast, stabilization has proved ineffective to fully prevent intergranular corrosion due to segregation of unreacted chromium [...] Read more.
The intergranular corrosion susceptibility of ferritic stainless-steel weldments is strongly dependent on chromium carbide precipitation phenomena. Hence, stabilization is widely used to mitigate the aforementioned precipitation. In contrast, stabilization has proved ineffective to fully prevent intergranular corrosion due to segregation of unreacted chromium during solid-state heat-treatments. To analyze the precipitation behavior of 17 wt.-% chromium ferritic stainless steels during laser welding, sheets of unstabilized and titanium-stabilized ferritic stainless steels were welded in a butt joint configuration and characterized with special consideration of precipitation behavior by means of transmission electron microscopy. While unstabilized ferritic stainless steels exhibit pronounced chromium precipitate formation at grain boundaries, titanium-stabilization leads to titanium precipitates without adjacent chromium segregation. However, corrosion tests reveal three distinctive corrosion mechanisms within the investigated ferritic stainless steels based on their inherent precipitation behaviors. In light of the precipitation formation, it is evident that immersion in sulfuric acid media leads to the dissolution of either grain boundaries or the grain boundary vicinity. As a result, the residual mechanical strength of the joint is substantially degraded. Full article
(This article belongs to the Special Issue Corrosion and Protection of Stainless Steels)
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<p>Schematic detailing (<b>a</b>) sample extraction location and (<b>b</b>) tensile testing geometry. adapted from [<a href="#B16-metals-12-00086" class="html-bibr">16</a>] under CC-BY-license. All dimensions in millimeters.</p> Full article ">Figure 2
<p>Microstructural evolution of laser-welds. Light microscopy images of (<b>a</b>) AISI 430 and (<b>b</b>) AISI 430Ti. (<b>c</b>,<b>d</b>) STEM-images of investigated grain boundary regions following FIB-extraction.</p> Full article ">Figure 3
<p>(<b>a</b>) HAADF-image of investigated grain boundary within weld metal of AISI 430 with EDS-maps for (<b>b</b>) iron and (<b>c</b>) chromium. (<b>d</b>) quantitative EDS-linescan extracted from image (<b>a</b>).</p> Full article ">Figure 4
<p>(<b>a</b>) HAADF-image of investigated grain boundary region within weld metal of AISI 430Ti with EDS-maps for (<b>b</b>) titanium and (<b>c</b>) chromium. (<b>d</b>) shows an overlay of Ti-EDS-map and HAADF image, (<b>e</b>) represents results of the EDS-line scan.</p> Full article ">Figure 5
<p>Optical micrographs of (<b>a</b>) AISI 430 and (<b>b</b>) AISI 430Ti following exposure to 16% boiling sulfuric acid for a duration of 20 h.</p> Full article ">Figure 6
<p>(<b>a</b>) SEM overview of the AISI 430Ti welding bead following exposure to 16% boiling sulfuric acid for a duration of 20 h, (<b>b</b>) detailed SEM image of a precipitate on top of the welding bead and (<b>c</b>) EDS linescan across the identified precipitate.</p> Full article ">Figure 7
<p>Stress-strain curves depicting the residual mechanical properties of laser-welded AISI 430Ti specimens following exposure to boiling 16% boiling sulfuric acid for a duration of 20 h. All specimens were extracted from a single weld seam following corrosion testing.</p> Full article ">
12 pages, 3199 KiB  
Article
Microstructure and Strengthening Mechanisms in an HSLA Steel Subjected to Tempforming
by Anastasiia Dolzhenko, Alexander Pydrin, Sergey Gaidar, Rustam Kaibyshev and Andrey Belyakov
Metals 2022, 12(1), 48; https://doi.org/10.3390/met12010048 - 27 Dec 2021
Cited by 15 | Viewed by 3345
Abstract
An effect of tempforming on the microstructure, the carbide precipitation, and the strengthening mechanisms of high-strength low-alloyed steel has been analyzed. The quenched steel was subjected to 1 h tempering at a temperature of 873 K, 923 K, or 973 K followed by [...] Read more.
An effect of tempforming on the microstructure, the carbide precipitation, and the strengthening mechanisms of high-strength low-alloyed steel has been analyzed. The quenched steel was subjected to 1 h tempering at a temperature of 873 K, 923 K, or 973 K followed by plate rolling at the same temperature. Tempforming resulted in the formation of an ultrafine grained lamellar-type microstructure with finely dispersed carbides of (Nb,V)C, Fe3C and Cr23C6. A decrease in tempforming temperature resulted in a reduction of the transverse grain size from 950 nm to 350 nm. Correspondingly, the size of Fe3C/Cr23C6 particles decreased from 90 nm to 40 nm while the size of (Nb,V)C particles decreased from 17 nm to 4 nm. Refining the tempformed microstructure with a decrease in thetempforming temperature provided an increase in the yield strength from 690 MPa to 1230 MPa. Full article
(This article belongs to the Section Metal Casting, Forming and Heat Treatment)
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<p>Microstructures developed in an HSLA steel subjected to tempforming at 873 K (<b>a</b>), 923 K (<b>b</b>) and 973 K (<b>c</b>) along with grain/subgrain misorientation distributions and sections of orientation distribution functions at φ<sub>2</sub> = 0°. Colors correspond to the crystallographic direction along the normal direction (ND). The black and white lines indicate high-angle boundaries (θ ≥ 15°) and low-angle sub-boundaries (2° ≤ θ &lt; 15°), respectively. Arrows on the misorientation distributions indicate an average sub-boundary misorientation.</p> Full article ">Figure 2
<p>Dispersed carbides in an HSLA steel subjected to tempforming at 873 K.</p> Full article ">Figure 3
<p>Particles of Cr<sub>23</sub>C<sub>6</sub> (<b>a</b>) and (V,Nb)C (<b>b</b>) in an HSLA steel tempformed at 923 K. The number in (<b>a</b>) indicates the sub-boundary misorientation in degrees.</p> Full article ">Figure 4
<p>Particles of Cr<sub>23</sub>C<sub>6</sub> (<b>a</b>) and (Nb,V)C (<b>b</b>) in an HSLA steel tempformed at 973 K.</p> Full article ">Figure 5
<p>Relationship between the subgrain size (d) and the dislocation density inside subgrains (<span class="html-italic">ρ</span>) in an HSLA steel subjected to tempforming at different temperatures.</p> Full article ">Figure 6
<p>Effect of deformation conditions on the transverse grain size in an HSLA steel subjected to tempforming (black symbols) and low carbon steel after large strain compression (open symbols) [<a href="#B19-metals-12-00048" class="html-bibr">19</a>].</p> Full article ">Figure 7
<p>X-Ray diffraction of powder residuals obtained by a method of electrolysis from the tempformed specimens.</p> Full article ">Figure 8
<p>Tensile stress-elongation curves of an HSLA steel subjected to tempforming at indicated temperatures (T<sub>TF</sub>).</p> Full article ">
8 pages, 4053 KiB  
Article
Effect of V and Ti on the Oxidation Resistance of WMoTaNb Refractory High-Entropy Alloy at High Temperatures
by Shuaidan Lu, Xiaoxiao Li, Xiaoyu Liang, Wei Yang and Jian Chen
Metals 2022, 12(1), 41; https://doi.org/10.3390/met12010041 - 25 Dec 2021
Cited by 18 | Viewed by 3350
Abstract
Alloying with V and Ti elements effectively improves the strength of WMoTaNb refractory high entropy alloys (RHEAs) at elevated temperatures. However, their effects on the oxidation resistance of WMoTaNb RHEAs are unknown, which is vitally important to their application at high temperatures. In [...] Read more.
Alloying with V and Ti elements effectively improves the strength of WMoTaNb refractory high entropy alloys (RHEAs) at elevated temperatures. However, their effects on the oxidation resistance of WMoTaNb RHEAs are unknown, which is vitally important to their application at high temperatures. In this work, the effect of V and Ti on the oxidation behavior of WMoTaNb RHEA at 1000 °C was investigated using a thermogravimetric system, X-ray diffraction and scanning electron microscopy. The oxidation of all alloys was found to obey a power law passivating oxidation at the early stage. The addition of V aggravates the volatility of V2O5, MoO3 and WO3, and leads to disastrous internal oxidation. The addition of Ti reduces the mass gain in forming the full coverage of passivating scale and prolongs the passivation duration of alloys. Full article
(This article belongs to the Special Issue Advanced Refractory Alloys)
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<p>BSE micrographs and XRD patterns (<b>d</b>) of as-cast (<b>a</b>) WMoTaNb, (<b>b</b>) WMoTaNbV and (<b>c</b>) WMoTaNbTi alloys.</p> Full article ">Figure 2
<p>The oxidation kinetics of WMoTaNb, WMoTaNbV and WMoTaNbTi RHEAs during isothermal exposure to dry air at 1000 °C.</p> Full article ">Figure 3
<p>XRD patterns of the oxide scales formed on three alloys after different oxidation times at 1000 °C.</p> Full article ">Figure 4
<p>Surface morphologies, cross sections and EDS results of (<b>a</b>,<b>d</b>) WMoTaNb, (<b>b</b>,<b>e</b>) WMoTaNbV and (<b>c</b>,<b>f</b>) WMoTaNbTi RHEAs after 0.5 h of oxidation at 1000 °C.</p> Full article ">
20 pages, 6696 KiB  
Article
The Effects of Deteriorated Boundary Conditions on Horizontally Framed Miter Gates
by Guillermo A. Riveros, Felipe J. Acosta, Christine M. Lozano and Eileen Glynn
Metals 2022, 12(1), 37; https://doi.org/10.3390/met12010037 - 24 Dec 2021
Cited by 2 | Viewed by 3515
Abstract
The U.S. navigable infrastructure is a system of waterways dependent upon hydraulic steel structures (HSS) to facilitate the passage of ships and cargo. The system is linear in the sense that if one HSS is impassable, the entire river system is halted at [...] Read more.
The U.S. navigable infrastructure is a system of waterways dependent upon hydraulic steel structures (HSS) to facilitate the passage of ships and cargo. The system is linear in the sense that if one HSS is impassable, the entire river system is halted at that point. The majority of the HSS in this system were built in the first half of the 20th Century, and over seventy percent of them are past or near past their design life. Miter gates are critical HSS components within the system and many are showing signs of structural distress from continued operation past their design life. Common distress features include shear cracking within the pintle socket, partially missing Quoin blocks, fatigue fracturing, and bolt failure in the pintle region. This article focuses on gaining a fundamental understanding of the consequences of quoin block deterioration on a miter gate. The work was conducted by developing a computational model of a miter gate with different levels of quoin block deterioration. This model was validated using analytical solutions. The deterioration results demonstrated that the miter gate thrust diaphragm and quoin post experienced changes in their limit states due to deterioration. The results also demonstrated that the miter gate could overcome up to 10% of quoin block deterioration. Full article
(This article belongs to the Section Computation and Simulation on Metals)
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<p>Map of the inland waterways along with a cost comparison of intercity transportation [<a href="#B1-metals-12-00037" class="html-bibr">1</a>,<a href="#B2-metals-12-00037" class="html-bibr">2</a>,<a href="#B3-metals-12-00037" class="html-bibr">3</a>].</p> Full article ">Figure 2
<p>Total nationwide lock unavailability [<a href="#B15-metals-12-00037" class="html-bibr">15</a>].</p> Full article ">Figure 3
<p>Vessel moving through lock gates, from downstream to upstream.</p> Full article ">Figure 4
<p>Horizontal miter gate (<b>a</b>) the miter gate is in the closed position and (<b>b</b>) the miter gate is in the open position [<a href="#B23-metals-12-00037" class="html-bibr">23</a>].</p> Full article ">Figure 5
<p>Downstream view of a miter gate with components.</p> Full article ">Figure 6
<p>(<b>a</b>) Top view of the tapered end section and (<b>b</b>) front view of miter gate leaf with force distribution.</p> Full article ">Figure 7
<p>Distress features in horizontal miter gates caused by Quoin block deterioration, (<b>A</b>) out of plane deformation near Quoin block, (<b>B</b>) cracks propagate around bolts, (<b>C</b>) crack at the pintle, (<b>D</b>) crack above pintle.</p> Full article ">Figure 8
<p>Main structural elements of the lower gate, Holt Lock. (<b>a</b>) miter gate leaf with 16 girders, (<b>b</b>) pintle bolt, (<b>c</b>) pintle socket, (<b>d</b>) pintle ball.</p> Full article ">Figure 9
<p>Graphic illustration showing geometry and displacement boundary conditions for the method validation experiments.</p> Full article ">Figure 10
<p>Three charts showing agreement in numerical and analytical solutions for the L&amp;D No. 27. (<b>A</b>) Girder 3 near the top of the gate, (<b>B</b>) Girder 8 near the middle, (<b>C</b>) Girder 12 at the bottom.</p> Full article ">Figure 11
<p>Vertical stress paths along the thrust diaphragm.</p> Full article ">Figure 12
<p>Horizontal Paths along the thrust diaphragm aligned with Girders 9, 10, 11, 12, 13, 14, and 15; and with the centerline of Panel 12.</p> Full article ">Figure 13
<p>Horizontal stresses in thrust diaphragm along vertical paths 1, 2, 3, and 4, on plots (<b>A</b>), (<b>B</b>), (<b>C</b>), and (<b>D</b>), respectively.</p> Full article ">Figure 13 Cont.
<p>Horizontal stresses in thrust diaphragm along vertical paths 1, 2, 3, and 4, on plots (<b>A</b>), (<b>B</b>), (<b>C</b>), and (<b>D</b>), respectively.</p> Full article ">Figure 14
<p>Calculated vertical stresses along vertical paths 1, 2, 3, and 4, on plots (<b>A</b>), (<b>B</b>), (<b>C</b>), and (<b>D</b>), respectively.</p> Full article ">Figure 14 Cont.
<p>Calculated vertical stresses along vertical paths 1, 2, 3, and 4, on plots (<b>A</b>), (<b>B</b>), (<b>C</b>), and (<b>D</b>), respectively.</p> Full article ">Figure 15
<p>(<b>A</b>–<b>G</b>) Illustration of the horizontal stress captured along these paths.</p> Full article ">Figure 15 Cont.
<p>(<b>A</b>–<b>G</b>) Illustration of the horizontal stress captured along these paths.</p> Full article ">Figure 16
<p>(<b>A</b>–<b>G</b>) Calculated vertical stresses along seven horizontal paths within the thrust diaphragm.</p> Full article ">Figure 16 Cont.
<p>(<b>A</b>–<b>G</b>) Calculated vertical stresses along seven horizontal paths within the thrust diaphragm.</p> Full article ">
18 pages, 7624 KiB  
Article
Effect of Ni Doping on the Embrittlement of Liquid Zinc at Σ5 Fe Austenite Grain Boundary
by Chengfa Ding, Wangjun Peng, Zheng Ma, Yan Zhao, Huaxiang Teng and Guangxin Wu
Metals 2022, 12(1), 27; https://doi.org/10.3390/met12010027 - 23 Dec 2021
Cited by 2 | Viewed by 2876
Abstract
In this study, first-principles computational tensile tests have been performed for the Σ5 symmetrically tilted grain boundaries of the face-centered cubic (fcc) Fe to investigate the effects of Zn and Zn-Ni doping on the boundary energy and electronic structure. The obtained results indicate [...] Read more.
In this study, first-principles computational tensile tests have been performed for the Σ5 symmetrically tilted grain boundaries of the face-centered cubic (fcc) Fe to investigate the effects of Zn and Zn-Ni doping on the boundary energy and electronic structure. The obtained results indicate that the mismatch between the sizes of Zn and Fe atoms at the Zn-doped grain boundary causes its expansion, which increases the lengths of Fe-Fe bonds, leading to their weakening, and reduces the overall boundary strength. After the Zn doping of the Fe grain boundary, Zn atoms form covalent bonds with Fe atoms, that decreases the charge density of Fe-Fe bonds and their strength. Meanwhile, the strength of the newly formed Fe-Zn covalent bonds oriented at a certain angle with respect to the grain boundary direction is very low. The breakage of Fe-Fe bonds that occurs under tensile loading rapidly decreases the boundary strength. Finally, after the Zn-Ni co-doping of the Fe grain boundary, Ni atoms form metallic bonds with Fe atoms, thus increasing both the charge density of Fe-Fe bonds (as compared with that of the Fe-Fe bonds at the Zn-doped grain boundary). Full article
(This article belongs to the Section Computation and Simulation on Metals)
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<p>Analysis region of electron backscattering diffraction (EBSD).</p> Full article ">Figure 2
<p>EBSD grain boundary analysis results: (<b>a</b>) morphology of grain, (<b>b</b>) phase proportion, (<b>c</b>) coincidence lattice GB, and (<b>d</b>) misorientation of grain boundary at positions 1 and 2 in (<b>c</b>).</p> Full article ">Figure 3
<p>Fcc Fe Σ5 [001] (210) 53.6° symmetric tilt grain boundary model, GB (Grain Boundary) is grain boundary surface of the GB model, a total contains two layers of GB surface, the size of initial super cell of GB model is 3.44900 × 7.71210 × 15.42420 Å3.</p> Full article ">Figure 4
<p>Grain boundary model based on the fcc Fe Σ5 [001] (210) 53.6° symmetric tilt boundary model of <a href="#metals-12-00027-f003" class="html-fig">Figure 3</a>: (<b>a</b>) Undoped GB, (<b>b</b>) Zn-doped GB, and (<b>c</b>) Zn-Ni co-doped GB.</p> Full article ">Figure 5
<p>GB model of Ni atom at different sites: (<b>a</b>) The central position of the Ni atom, (<b>b</b>) Position of Ni atom on edge.</p> Full article ">Figure 6
<p>Interlayer spacing of grain boundary before and after doping Zn and Ni: (<b>a</b>) Undoped GB, (<b>b</b>) Zn-doped GB, and (<b>c</b>) Zn-Ni co-doped GB.</p> Full article ">Figure 7
<p>Relationship of total energy with strain of different GB.</p> Full article ">Figure 8
<p>Stress-strain curve calculated by first-principles.</p> Full article ">Figure 9
<p>Grain boundaries under different strains: (<b>a</b>) Undoped GB, (<b>b</b>) Zn-doped GB, and (<b>c</b>) Zn-Ni co-doped GB.</p> Full article ">Figure 10
<p>Chemical bond structures of Fe-Fe, Fe-Zn and Ni-Fe bonds in different GB models: (<b>a</b>) Undoped GB, (<b>b</b>) Zn-doped GB, and (<b>c</b>) Zn-Ni co-doped GB, unit: Å.</p> Full article ">Figure 11
<p>Variation of bond length of Fe-Fe, Fe-Zn and Ni-Fe bonds in the GB model at different strain: (<b>a</b>) Fe(2)-Fe(–2) bond of undoped GB, Zn-doped GB, and Zn-Ni co-doped GB, (<b>b</b>) Fe(1)-Fe(3) bond of undoped GB, Zn(1)-Fe(3) bond of Zn-doped GB and Zn-Ni co-doped GB, (<b>c</b>) Fe(1)-Fe(2) bond of undoped GB, Zn(1)-Fe(2) bond of Zn-doped GB, Ni(1)-Fe(2) bond Zn-Ni co-doped GB.</p> Full article ">Figure 12
<p>Charge distribution of Fe(2)-Fe(–2) bond on (002) plane under different strains: (<b>a</b>) Undoped GB, (<b>b</b>) Zn-doped GB, and (<b>c</b>) Zn-Ni co-doped GB.</p> Full article ">Figure 13
<p>Charge distribution of (<b>a</b>) Fe(1)-Fe(2) bond and (<b>b</b>) Fe(1)-Fe(3) bond on (002) plane of undoped GB at 0% stain.</p> Full article ">Figure 14
<p>Charge distribution of (<b>a</b>) Zn(1)-Fe(2) bond and (<b>b</b>) Zn(1)-Fe(3) bond on (002) plane of Zn-doped GB at 0% stain.</p> Full article ">Figure 15
<p>Charge distribution of (<b>a</b>) Ni(1)-Fe(2) bond and (<b>b</b>) Ni(1)-Fe(3) bond on (002) plane of Zn-Ni co-doped GB at 0% stain.</p> Full article ">Figure 16
<p>Partial Density of States: (<b>a</b>) Fe(2) atom in undoped GB, (<b>b</b>) and (<b>c</b>) Fe(2) and Zn(1) atom in Zn-doped GB, and (<b>d</b>) Fe(2) atom in Zn-Ni co-doped GB.</p> Full article ">Figure 17
<p>(<b>a</b>) Differential charge of Fe(1) atom on (100) plane of undoped GB, (<b>b</b>) Differential charge of Zn(1) atom on (100) plane of Zn-doped GB, (<b>c</b>) Differential charge of Ni(1) atom on (100) plane of Zn-Ni co-doped GB, and (<b>d</b>) Differential charge of Zn(1) atom on (100) plane of Zn-Ni co-doped GB.</p> Full article ">
13 pages, 6862 KiB  
Article
Kinetic and Metallography Study of the Oxidation at 1250 °C of {Co+Ni}-Based Superalloys Containing Ti to Form MC Carbides
by Patrice Berthod, Synthia Annick Ozouaki Wora, Lionel Aranda, Ghouti Medjahdi and Erwan Etienne
Metals 2022, 12(1), 10; https://doi.org/10.3390/met12010010 - 22 Dec 2021
Cited by 5 | Viewed by 2372
Abstract
Six conventionally cast chromium-rich titanium-containing alloys based on cobalt and nickel with various Co/Ni ratios were considered. They were tested in oxidation in air at 1250 °C for 70 h in a thermo-balance. The mass gain curves were exploited to specify different types [...] Read more.
Six conventionally cast chromium-rich titanium-containing alloys based on cobalt and nickel with various Co/Ni ratios were considered. They were tested in oxidation in air at 1250 °C for 70 h in a thermo-balance. The mass gain curves were exploited to specify different types of kinetic constants as well as several parameters characterizing the oxide spallation occurring during cooling. The obtained results show that, the higher the Ni content, the slower the mass gain and the better the quality of the protective external chromia scale. Secondly, no dependence of the oxide spallation characteristics on the Co content was clearly noted. Globally, the isothermal oxidation behavior becomes better when Ni is more and more present at the expense of Co. Titanium seems to be playing a particular role in the process of oxidation. It notably leads to the presence of an external thin TiO2 continuous scale beyond the chromia scale. The thermogravimetry records were numerically treated to determine the parabolic constant and the chromia volatilization constant. The values of these constants evidenced a double tendency: chromia growth acceleration and chromia volatilization slow-down. These trends are to be confirmed and further investigated. Full article
(This article belongs to the Special Issue Hot Oxidation and Corrosion of High Performance Metallic Alloys)
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<p>The as-cast microstructures of the six alloys: 5Co0NiTi (<b>a</b>), 4Co1NiTi (<b>b</b>), 3Co2NiTi (<b>c</b>), 2Co3NiTi (<b>d</b>), 1Co4NiTi (<b>e</b>) and 0Co5NiTi (<b>f</b>); SEM/BSE micrographs.</p> Full article ">Figure 2
<p>The mass variation curves plotted versus time to observe the mass gain evolution for all the six alloys during the isothermal stage.</p> Full article ">Figure 3
<p>Diffractogram acquired on one side of the oxidized 5Co0NiTi sample.</p> Full article ">Figure 4
<p>Diffractogram acquired on one side of the oxidized 3Co2NiTi sample.</p> Full article ">Figure 5
<p>Diffractogram acquired on one side of the oxidized 2Co3NiTi sample.</p> Full article ">Figure 6
<p>Diffractogram acquired on one side of the oxidized 0Co5NiTi sample.</p> Full article ">Figure 7
<p>X-map obtained on the oxidized surface zone of the 3Co2NiTi alloy (evidencing the TiO<sub>2</sub>, CoCr<sub>2</sub>O<sub>4</sub> and Cr<sub>2</sub>O<sub>3</sub> zones).</p> Full article ">Figure 8
<p>X-map obtained on the oxidized surface zone of the 2Co3NiTi alloy (evidencing the TiO<sub>2</sub>, Cr<sub>2</sub>O<sub>3</sub> and denuded zones).</p> Full article ">Figure 9
<p>Cross-sectional observation of the external oxides and of the subsurfaces affected by oxidation (Co-richest alloys); white dotted arrows: line scans corresponding to the concentration profiles presented in <a href="#metals-12-00010-f011" class="html-fig">Figure 11</a>; (<b>a</b>) and (<b>A</b>): 5Co0NiTi, (<b>b</b>) and (<b>B</b>): 4Co1NiTi, (<b>c</b>) and (<b>C</b>): 3Co2NiTi.</p> Full article ">Figure 10
<p>Cross-sectional observation of the external oxides and the subsurfaces affected by oxidation (Ni-richest alloys); white dotted arrows: line scans corresponding to the concentration profiles presented in <a href="#metals-12-00010-f011" class="html-fig">Figure 11</a>; (<b>a</b>) and (<b>A</b>): 2Co3NiTi, (<b>b</b>) and (<b>B</b>): 1Co4NiTi, (<b>c</b>) and (<b>C</b>): 0Co5NiTi.</p> Full article ">Figure 11
<p>Concentration profiles for chromium acquired by EDS across the subsurface; (<b>a</b>): 5Co0NiTi, (<b>b</b>): 4Co1NiTi, (<b>c</b>): 3Co2NiTi, (<b>d</b>): 2Co3NiTi, (<b>e</b>): 1Co4NiTi, (<b>f</b>): 0Co5NiTi.</p> Full article ">
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