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13 pages, 4392 KiB  
Article
Aerosol-Deposited 8YSZ Coating for Thermal Shielding of 3YSZ/CNT Composites
by Maria Wiśniewska, Grzegorz Kubicki, Mateusz Marczewski, Volf Leshchynsky, Luca Celotti, Mirosław Szybowicz and Dariusz Garbiec
Coatings 2024, 14(9), 1186; https://doi.org/10.3390/coatings14091186 - 13 Sep 2024
Viewed by 411
Abstract
High-temperature conditions are harmful for carbon nanotube-based (CNT-based) composites, as CNTs are susceptible to oxidation. On the other hand, adding CNTs to ceramics with low electrical conductivity, such as 3YSZ, is beneficial because it allows the production of complex-shaped samples with spark plasma [...] Read more.
High-temperature conditions are harmful for carbon nanotube-based (CNT-based) composites, as CNTs are susceptible to oxidation. On the other hand, adding CNTs to ceramics with low electrical conductivity, such as 3YSZ, is beneficial because it allows the production of complex-shaped samples with spark plasma sintering (SPS). A shielding coating system may be applied to prevent CNT oxidation. In this work, the 8YSZ (yttria-stabilized zirconia) thermal shielding coating system was deposited by aerosol deposition (AD) to improve the composite’s resistance to CNT degradation without the use of bond-coat sublayers. Additionally, the influence of the annealing process on the mechanical properties and microstructure of the composite was evaluated by nanoindentation, scratch tests, scanning electron microscopy (SEM), X-ray diffraction (XRD), flame tests, and light microscopy (LM). Annealing at 1200 °C was the optimal temperature for heat treatment, improving the coating’s mechanical strength (the first critical load increased from 0.84 N to 3.69 N) and promoting diffusion bonding between the compacted powder particles and the substrate. The deposited coating of 8YSZ increased the composite’s thermal resistance by reducing the substrate’s heating rate and preventing the oxidation of CNTs. Full article
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Figure 1

Figure 1
<p>Particle (<b>a</b>) size distribution and (<b>b</b>) morphology of 8YSZ feedstock powders (magnification 1000×).</p>
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<p>Photographs of (<b>a</b>) thin and (<b>b</b>) thick 3YSZ/CNT 1” coupon with the 8YSZ coating annealed at 1200 °C.</p>
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<p>X-ray diffraction spectra comparison of (<b>a</b>) AD process results and (<b>b</b>) results of 8YSZ coating as-sprayed and heat-treated at 1000 °C (HT 1000 °C), 1100 °C (HT 1100 °C), and 1200 °C (HT 1200 °C).</p>
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<p>SEM micrographs of (<b>a</b>) heat-treated at 1000 °C, (<b>b</b>) heat-treated at 1100 °C, (<b>c</b>) heat-treated at 1200 °C coating surface, and (<b>d</b>) heat-treated at 1200 °C coating cross-sections ((<b>a</b>–<b>c</b>)—magnification—50k×, (<b>d</b>)—magnification 10k×).</p>
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<p>Nanoindentation result comparison of (<b>a</b>) indentation load of 30mN (coating cross-section) and (<b>b</b>) hardness and elastic modulus calculation results.</p>
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<p><span class="html-italic">H</span><sup>2</sup> = <span class="html-italic">f</span>(1/<span class="html-italic">h</span> − <span class="html-italic">h<sub>el</sub></span>) dependence for 30 mN indentation load for coatings as-sprayed and heat-treated at 1000 °C (HT_1000), 1100 °C (HT_1100), and 1200 °C (HT_1200).</p>
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<p>The scratch test results of the coatings (<b>a</b>) as-sprayed, (<b>b</b>) heat-treated at 1000 °C, (<b>c</b>) heat-treated at 1100 °C, and (<b>d</b>) heat-treated at 1200 °C. The bottom graph illustrates the optical micrographs of the scratch at the specified distances.</p>
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<p>Temperature at the back side of the samples (coated with 250 and 550 µm thick coating heat-treated at 1200 °C) during flame torch test.</p>
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<p>8YSZ-coated surface of samples heat-treated at 1200 °C after flame torch test (<b>a</b>) an overview of the entire sample surface, (<b>b</b>) a close-up view of region 1, which is marked in (<b>a</b>), and (<b>c</b>) a close-up view of region 2, which is marked in (<b>a</b>).</p>
Full article ">Figure 10
<p>8YSZ-coated 3YSZ-CNT shell-shaped composite heat-treated at 1200 °C: (<b>a</b>) face and (<b>b</b>) sides of the composite.</p>
Full article ">
16 pages, 2783 KiB  
Article
Development of Solid-State Lithium-Ion Batteries (LIBs) to Increase Ionic Conductivity through Interactions between Solid Electrolytes and Anode and Cathode Electrodes
by Majid Monajjemi and Fatemeh Mollaamin
Energies 2024, 17(18), 4530; https://doi.org/10.3390/en17184530 - 10 Sep 2024
Viewed by 486
Abstract
Although in general ions are not able to migrate in the solid-state position due to rigid skeletal structure, in some solid electrolytes with a low energy barrier and high ionic conductivities, these ion transition can occur. In this work, we considered several solid [...] Read more.
Although in general ions are not able to migrate in the solid-state position due to rigid skeletal structure, in some solid electrolytes with a low energy barrier and high ionic conductivities, these ion transition can occur. In this work, we considered several solid electrolytes including lithium phosphorus oxy-nitride (LIPON), a lithium super-ionic conductor (SILICON), and thio-LISICON. For the fabrication and characterization of the solid electrolyte’s fabrication, we used a single-step ball milling (SSBM) procedure. Through this research on all-solid-state rechargeable lithium-ion batteries, our target is to discuss solving several problems in solid LIBs that have recently escalated due to raised concerns relating to safety hazards such as solvent leakage and the flammability of the liquid electrolytes used for commercial LIBs. Through this research, we tested the conductivity amounts of various substrates containing amorphous glass, SSBM, and glass-ceramic samples. Obviously, the SSBM glass-ceramics increased the conductivity, and we also found that the values for conductivity attained by SSBM were higher than those values for glass-ceramics. Using an SSBM technique, silicon nanoparticles were used as an anode material and it was found that the charge and discharge curves in the battery cell cycled between 0.009 and 1.45 V versus Li+/Li at a current density of 210 mA g−1 at room temperature. Since high resistance causes degradation between the cathode material (LiCoO2) and the solid electrolyte, we added GeS2 and SiS2 to the Li2S-P2S5 system to obtain higher conductivities and better stability of the electrode–electrolyte interface. Full article
(This article belongs to the Section D2: Electrochem: Batteries, Fuel Cells, Capacitors)
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Figure 1

Figure 1
<p>XRD of crystals glass showing a general phase diagram of the 75% Li<sub>2</sub>S with %25P<sub>2</sub>S<sub>5</sub>.</p>
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<p>(<b>A</b>) Anode electrode containing acetylene black. (<b>B</b>) Multi-wall carbon nanotubes in anode. (<b>C</b>) Solid-state lithium battery including titanium.</p>
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<p>Titanium test with Li<sub>2</sub>S–GeS<sub>2</sub>–P<sub>2</sub>S<sub>5</sub> SSE, (<b>A</b>) solid electrolyte battery, and (<b>B</b>) conductivity tester.</p>
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<p>XRD patterns for all tested electrolytes, (a) Li<sub>2</sub>S, (b) P<sub>2</sub>S<sub>5</sub>, (c) x = GeS<sub>2</sub>, (d) x = 64, (f) x = 76.</p>
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<p>(a) Conductivity map for Li<sub>(4–x)</sub>Ge<sub>(1–x)</sub>P<sub>x</sub>S<sub>4</sub>; (b) glass, (c) glass-ceramic, (d) SSBM glass-ceramic.</p>
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<p>Comparison showing superior performance of MWCNT as a conductive additive for all solid–state lithium batteries over acetylene black in different voltages and three temperatures.</p>
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<p>Charge–discharge curves of all-solid-state cell fabricated with n-Si as anode material by different voltage ranges. (<b>A</b>–<b>F</b>) compare the cycling performance of solid state and liquid electrolyte systems using nano-Si. (<b>G</b>,<b>H</b>) exhibit the charge-discharge performance of ion cells, tested with lithium phosphorus oxy-nitride (LIPON) and lithium super-ionic conductor (SILICON) solid electrolyte, respectively.</p>
Full article ">Figure 7 Cont.
<p>Charge–discharge curves of all-solid-state cell fabricated with n-Si as anode material by different voltage ranges. (<b>A</b>–<b>F</b>) compare the cycling performance of solid state and liquid electrolyte systems using nano-Si. (<b>G</b>,<b>H</b>) exhibit the charge-discharge performance of ion cells, tested with lithium phosphorus oxy-nitride (LIPON) and lithium super-ionic conductor (SILICON) solid electrolyte, respectively.</p>
Full article ">Scheme 1
<p>Operating LIBs.</p>
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11 pages, 9819 KiB  
Article
Wear and Abrasion Resistance of Nitride Coatings on Ceramic Substrates Processed with Fast Argon Atoms
by Sergey N. Grigoriev, Alexander S. Metel, Marina A. Volosova, Enver S. Mustafaev and Yury A. Melnik
Surfaces 2024, 7(3), 714-724; https://doi.org/10.3390/surfaces7030046 - 4 Sep 2024
Viewed by 284
Abstract
The surfaces of ceramic products are replete with numerous defects, such as those that appear during the diamond grinding of sintered SiAlON ceramics. The defective surface layer is the reason for the low effectiveness of TiZrN coatings under abrasive and fretting wear. An [...] Read more.
The surfaces of ceramic products are replete with numerous defects, such as those that appear during the diamond grinding of sintered SiAlON ceramics. The defective surface layer is the reason for the low effectiveness of TiZrN coatings under abrasive and fretting wear. An obvious solution is the removal of an up to 4-µm-thick surface layer containing the defects. It was proposed in the present study to etch the layer with fast argon atoms. At the atom energy of 5 keV and a 0.5 mA/cm2 current density, the ions were converted into fast atoms and the sputtering rate for the SiAlON samples reached 20 μm/h. No defects were observed in the microstructures of coatings deposited after beam treatment for half an hour. The treatment reduced the volumetric abrasive wear by five times. The fretting wear was reduced by three to four times. Full article
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Figure 1
<p>Diagram of the experimental setup.</p>
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<p>Topside view of the experimental setup.</p>
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<p>Beam diameter <span class="html-italic">D</span> versus distance to the grid <span class="html-italic">Z</span>.</p>
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<p>Scheme of the abrasion tests (<b>a</b>) and equipment for the measurement of the abrasion resistance (<b>b</b>).</p>
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<p>Scheme of the fretting tests (<b>a</b>) and equipment for the measurement of the fretting wear resistance (<b>b</b>).</p>
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<p>The sputtering coefficient Y versus the energy <span class="html-italic">E</span> of argon atoms.</p>
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<p>Dependence of the removed layer thickness Δ on the treatment time <span class="html-italic">t</span> at the anode current 2 A (1) and 1 A (2).</p>
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<p>Profilograms of the SiAlON samples before treatment (<b>a</b>) and after treatment for 1 h with 5 keV argon atoms (<b>b</b>).</p>
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<p>SEM images of the SiAlON samples before treatment (<b>a</b>) and after treatment with fast argon atoms (<b>b</b>).</p>
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<p>SEM images of TiZrN coatings deposited on the samples before (<b>a</b>) and after (<b>b</b>) treatment with fast argon atoms.</p>
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<p>Profilograms of wear spots after 20 min of abrasive exposure on samples composed of SiAlON ceramics with TiZrN coatings deposited before (<b>a</b>) and after (<b>b</b>) treatment with fast argon atoms.</p>
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<p>Wear spots after 10<sup>5</sup> cycles of friction of counter body with SiAlON samples coated with TiZrN before (<b>a</b>) and after (<b>b</b>) treatment with fast argon atoms.</p>
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14 pages, 15274 KiB  
Article
Microstructure, Mechanical and Tribological Properties of Cu40Zn-Ti3AlC2 Composites by Powder Metallurgy
by Fangdian Peng, Shidong Zhou, Tao Yang, Liwei Wu, Jianbo Wu, Puyou Ying, Ping Zhang, Changhong Lin, Yabo Fu, Zhibiao Tu, Tianle Wang, Xin Zhang, Nikolai Myshkin and Vladimir Levchenko
Lubricants 2024, 12(9), 306; https://doi.org/10.3390/lubricants12090306 - 31 Aug 2024
Viewed by 545
Abstract
The exploration of unleaded free-cutting Cu40Zn brass with excellent mechanical and tribological properties has always drawn the attention of researchers. Due to its attractive properties combining metals and ceramics, Ti3AlC2 was added to Cu40Zn brass using high-energy milling and hot-pressing [...] Read more.
The exploration of unleaded free-cutting Cu40Zn brass with excellent mechanical and tribological properties has always drawn the attention of researchers. Due to its attractive properties combining metals and ceramics, Ti3AlC2 was added to Cu40Zn brass using high-energy milling and hot-pressing sintering. The effects of Ti3AlC2 on the microstructure, mechanical and tribological properties of Cu40Zn-Ti3AlC2 composites were studied. The results showed that Ti3AlC2 could suppress the formation of ZnO by adsorbing oxygen impurity and promote the formation of the β phase by releasing the β-forming element Al to the substrate. The hardness and wear resistance of Cu40Zn-Ti3AlC2 composites increased with increasing Ti3AlC2 content from 0 to 5 wt.%. The proper Ti3AlC2 additive was beneficial to both the strength and plasticity of the composites. The underlying mechanisms were discussed. Full article
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Figure 1

Figure 1
<p>XRD spectra and SEM images of raw materials: (<b>a</b>,<b>b</b>) Ti<sub>3</sub>AlC<sub>2</sub>; (<b>c</b>,<b>d</b>) water-atomized Cu40Zn.</p>
Full article ">Figure 1 Cont.
<p>XRD spectra and SEM images of raw materials: (<b>a</b>,<b>b</b>) Ti<sub>3</sub>AlC<sub>2</sub>; (<b>c</b>,<b>d</b>) water-atomized Cu40Zn.</p>
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<p>XRD spectra of Cu40Zn-Ti<sub>3</sub>AlC<sub>2</sub> powder mixtures after 10 h of milling.</p>
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<p>XRD spectra (<b>a</b>), OM images (<b>b</b>–<b>f</b>) and phase content (<b>g</b>) of the sintered Cu40Zn-Ti<sub>3</sub>AlC<sub>2</sub> composites: (<b>b</b>) M0; (<b>c</b>) M0.5; (<b>d</b>) M1; (<b>e</b>) M3; (<b>f</b>) M5.</p>
Full article ">Figure 3 Cont.
<p>XRD spectra (<b>a</b>), OM images (<b>b</b>–<b>f</b>) and phase content (<b>g</b>) of the sintered Cu40Zn-Ti<sub>3</sub>AlC<sub>2</sub> composites: (<b>b</b>) M0; (<b>c</b>) M0.5; (<b>d</b>) M1; (<b>e</b>) M3; (<b>f</b>) M5.</p>
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<p>SEM image and EDS spectra of sample M5.</p>
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<p>Hardness of Cu40Zn-Ti<sub>3</sub>AlC<sub>2</sub> composites as a function of Ti<sub>3</sub>AlC<sub>2</sub> content.</p>
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<p>Typical load–displacement curves (<b>a</b>) and flexural performances (<b>b</b>) of Cu40Zn-Ti<sub>3</sub>AlC<sub>2</sub> composites.</p>
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<p>SEM fractographs for Cu40Zn−Ti3AlC2 composites: (<b>a</b>,<b>a1</b>) M0; (<b>b</b>,<b>b1</b>) M0.5; (<b>c</b>,<b>c1</b>) M1; (<b>d</b>,<b>d1</b>) M5.</p>
Full article ">Figure 7 Cont.
<p>SEM fractographs for Cu40Zn−Ti3AlC2 composites: (<b>a</b>,<b>a1</b>) M0; (<b>b</b>,<b>b1</b>) M0.5; (<b>c</b>,<b>c1</b>) M1; (<b>d</b>,<b>d1</b>) M5.</p>
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<p>COF curves (<b>a</b>) and wear constant (<b>b</b>) of Cu40Zn-Ti<sub>3</sub>AlC<sub>2</sub> composites.</p>
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<p>Wear track morphologies (<b>a</b>–<b>e</b>) and sectional profiles of the wear tracks (<b>f</b>) of Cu40Zn-Ti<sub>3</sub>AlC<sub>2</sub> composites: (<b>a</b>) M0; (<b>b</b>) M0.5; (<b>c</b>) M1; (<b>d</b>) M3; (<b>e</b>) M5.</p>
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<p>SEM images of the worn surfaces of the Cu40Zn-Ti<sub>3</sub>AlC<sub>2</sub> composites: (<b>a</b>,<b>a1</b>) M0; (<b>b</b>–<b>b2</b>) M1; (<b>c</b>–<b>c2</b>) M5.</p>
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<p>SEM images of the worn surfaces of the counterparts for (<b>a</b>,<b>a1</b>) M0; (<b>b</b>,<b>b1</b>) M1; (<b>c</b>,<b>c1</b>) M5.</p>
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27 pages, 30719 KiB  
Article
The Effect of Multiple Solder Reflows on the Formation of Cu6Sn5 Intermetallics and the Decomposition of SnAg3.0Cu0.5 Solder Joints in the Framework of Rework and Reuse of MLCC Components
by Erik Wiss and Steffen Wiese
Metals 2024, 14(9), 986; https://doi.org/10.3390/met14090986 - 29 Aug 2024
Viewed by 350
Abstract
A rework of electronic assemblies and the reuse of electronic components are the most effective ways to reduce electronic waste. Since neither components nor substrates were developed with the intention of multiple usage, the question of how the integrity of lead-free solder joints [...] Read more.
A rework of electronic assemblies and the reuse of electronic components are the most effective ways to reduce electronic waste. Since neither components nor substrates were developed with the intention of multiple usage, the question of how the integrity of lead-free solder joints is affected by multiple reflow operations is crucial for the implementation of any reuse strategy. Therefore, various types of 1206 multilayer ceramic capacitors (MLCCs) differing in their capacitance value and dielectric type (X5R, X7R, Y5V, NP0) were soldered on test printed circuit boards (PCBs) having a pure Cu-metallization surface in order to investigate the intermetallic reactions during multiple reflows. The metallization system on the MLCC-component side consisted of a thick film of Ni covered by galvanic-deposited Sn. The reflow experiments were conducted using a hypoeutectic SnAgCu solder. The results show the formation of a Cu6Sn5 intermetallic phase on both metallizations, which grows homogeneously with the number of reflows. Moreover, an ongoing decomposition of the solder into Ag-enriched and depleted zones was observed. The effect of these microstructural changes on the functionality of the solder joint was investigated by mechanical shear experiments and electrical four-point capacitance measurements. Full article
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Figure 1

Figure 1
<p>Layout of the electrical test-board (length 160 mm and width 100 mm) having a capacity to carry eight samples (C1–C8, solder pad layout for MLCC with 1206 size). Each pair of the MLCC soldering pads is connected to four pads on the right edge of the PCB to enable precise four-wire measurements of the MLCC’s capacitances.</p>
Full article ">Figure 2
<p>Layout of a test-PCB for the samples that were subjected to a shear test. The pads are connected to two additional pads on the left side of the PCB to allow in-situ monitoring of the solder joint health state during the mechanical test. The three holes in the middle are used to align and attach the samples to the holder of the testing machine.</p>
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<p>Time-temperature curve of the reflow profile that was used for the experiments. The graph summarizes the recordings of the subsequent reflow process in a typical and idealized form (without noise). Every reflow process had an identical profile. The reflow soldering has been carried out in a batch reflow oven, ProtoFlow S (LPKF), using a nitrogen-protective atmosphere.</p>
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<p>Specially developed measurement adapter that is connected to the measurement device via four BNC connectors, which bring the shielded measurement connections to the adapter. These measurement connections are redistributed by the PCB to four pin heads, whose grid exactly matches one of the Cu pads. After the pin heads have been brought into contact with the pads, the measurement is initiated (manual trigger).</p>
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<p>Equivalent circuit diagram of a real capacitor, consisting of a series resistance R<sub>s</sub>, a series inductance L<sub>s</sub>, a parallel resistance R<sub>p,</sub> and the capacitance C, adapted from [<a href="#B27-metals-14-00986" class="html-bibr">27</a>].</p>
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<p>Overview of the mechanical test setup presented in [<a href="#B29-metals-14-00986" class="html-bibr">29</a>,<a href="#B30-metals-14-00986" class="html-bibr">30</a>]. The main components of the setup are labeled as follows: A: Actuator. B: Three-axis force sensor. C: Measurement amplifiers (for force sensors). D: Stereo zoom microscope with illumination units. E: Crack detection circuit.</p>
Full article ">Figure 7
<p>Detailed view of the experimental setup. The displacement of the actuator and thus a force are transferred to the resistor via a small PCB, which is electrically isolated to not influence the electrical measurement. A crack formation within one of the solder joints results in an interruption of the electrical pathway between the two wires on the left side, which are connected to the crack detection circuit.</p>
Full article ">Figure 8
<p>Light microscopy image of an X7R MLCC (Magnification: 5×). The red box between the MLCC and the Cu pad marks the location of the SEM and EDX analysis. A: PCB base material. B: Cu pad of the PCB. C: Solder joint. D: Cu metallization and Ni termination. E: MLCC ceramic body. F: Inner electrodes.</p>
Full article ">Figure 9
<p>SEM images in SE mode with an enlarged view of the red box area of <a href="#metals-14-00986-f008" class="html-fig">Figure 8</a>. Path along which the line scans were performed after the (<b>a</b>) first, (<b>b</b>) second, (<b>c</b>) fourth, and (<b>d</b>) eighth reflow cycles, starting within the Ni-thick film layer (termination) of the MLCC and ending in the Cu metallization of the PCB.</p>
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<p>EDX analysis of an X7R MLCC after the first reflow cycle with a magnification of 4000×. (<b>a</b>) SEM image in SE mode. (<b>b</b>) Combined EDX maps of Ni, Cu, Ag, and Sn. (<b>c</b>) EDX map of Ni. (<b>d</b>) EDX map of Cu. (<b>e</b>) EDX map of Ag. (<b>f</b>) EDX map of Sn.</p>
Full article ">Figure 11
<p>EDX analysis of an X7R MLCC after the second reflow cycle with a magnification of 4000×. (<b>a</b>) SEM image in SE mode. (<b>b</b>) Combined EDX maps of Ni, Cu, Ag, and Sn. (<b>c</b>) EDX map of Ni. (<b>d</b>) EDX map of Cu. (<b>e</b>) EDX map of Ag. (<b>f</b>) EDX map of Sn.</p>
Full article ">Figure 12
<p>EDX analysis of an X7R MLCC after the fourth reflow cycle with a magnification of 4000×. (<b>a</b>) SEM image in SE mode. (<b>b</b>) Combined EDX maps of Ni, Cu, Ag, and Sn. (<b>c</b>) EDX map of Ni. (<b>d</b>) EDX map of Cu. (<b>e</b>) EDX map of Ag. (<b>f</b>) EDX map of Sn.</p>
Full article ">Figure 13
<p>EDX analysis of an X7R MLCC after the eighth reflow cycle with a magnification of 4000×. (<b>a</b>) SEM image in SE mode. (<b>b</b>) Combined EDX maps of Ni, Cu, Ag, and Sn. (<b>c</b>) EDX map of Ni. (<b>d</b>) EDX map of Cu. (<b>e</b>) EDX map of Ag. (<b>f</b>) EDX map of Sn.</p>
Full article ">Figure 14
<p>Line scan overviews along the given path (compare <a href="#metals-14-00986-f009" class="html-fig">Figure 9</a>) of an X7R MLCC after the (<b>a</b>,<b>b</b>) first, (<b>c</b>,<b>d</b>) second, (<b>e</b>,<b>f</b>) fourth, and (<b>g</b>,<b>h</b>) eighth reflow cycle, showing the evolution of the Cu<sub>6</sub>Sn<sub>5</sub> IMC (left-hand side) and the distribution of Ag within the solder (right-hand side), standardized to a total concentration of 100% (<span class="html-italic">y</span>-axis). Green: Ni. Red: Cu. Gray: Sn. Blue: Ag. While a continuous IMC layer has already formed at the interface solder/PCB metallization after the first reflow cycle, it takes four reflow cycles to form a proper IMC layer at the interface solder/component termination. During several reflow cycles, the Ag depletion zones at the interfaces diminish, and the Ag seems to distribute more and more uniformly within the solder paste. The depicted diagrams were taken directly from the EDX software QUANTAX ESPRIT (version 2.0) without editing. The ordinate shows the concentration in %, and the abscise (‘Weg/µm’) shows the distance along the scanning path. The algorithm of the used EDX software seems not always to be able to standardize the total concentration to 100%, which causes the differences in the tin concentration in the center of the solder joints.</p>
Full article ">Figure 15
<p>Averaged height of the Cu6Sn5 IMC at the interface between (<b>a</b>) the solder paste and the PCB metallization (R<sup>2</sup> = 0.96), and (<b>b</b>) the solder paste and the MLCC component (R<sup>2</sup> = 0.54). As the values were only measured after one, two, four, and eight reflow cycles, the remaining values were interpolated using a linear fit.</p>
Full article ">Figure 16
<p>Cycle-Capacitance diagram of the measured values of KEMET C1206C106K4PAC7800+ (dielectric X5R, size 1206, nominal capacitance 10 µF ± 10%) for eight reflow cycles. The values behave in a similar way, including an increase after the first two reflow cycles, a decrease between the third and the fifth one, and further increases during the last three ones.</p>
Full article ">Figure 17
<p>Cycle-Capacitance diagram of the measured values of KEMET C1206C474K5RACTU (dielectric X7R, size 1206, nominal capacitance 470 nF ± 10%) for eight reflow cycles. Except for sample number 6, the values behave in a similar way, including an increase after the first two reflow cycles, a decrease between the third and the fifth one, and further increases during the last three ones.</p>
Full article ">Figure 18
<p>Cycle-Capacitance diagram of the measured values of Yageo CC1206ZPY5V7BB475 (dielectric Y5V, size 1206, nominal capacitance 4.7 µF ± 20%) for eight reflow cycles. The values behave in a similar way, including an increase after the first two reflow cycles, a decrease between the third and the fifth one, and further increases during the last three ones.</p>
Full article ">Figure 19
<p>Cycle-Capacitance diagram of the measured values of Yageo CC1206JRNPO9BN681 (dielectric NP0, size 1206, nominal capacitance 680 pF ± 5%) for eight reflow cycles. The values remain mainly constant, with very small deviations.</p>
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<p>Results of the shear test with respect to the number of reflow cycles conducted The green box marks the first and third quartiles, the black line within the box marks the median, and the red dot marks the mean.</p>
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<p>EDX analysis of a shear tested resistor after one reflow cycle. (<b>a</b>) SEM image in SE mode with a magnification of 300× and (<b>b</b>) associated combined EDX map of Cu, Ni, Ag, Sn and Al. (<b>c</b>) Detailed SEM image in SE mode of the location between the ragged resistor and the Cu pad of the PCB with a higher magnification of 2000× and (<b>d</b>) associated combined EDX map of Cu, Ag, Sn, and Al (here, the Ni barrier neither adhered to the ceramic body nor to the solder). For each pair of pictures, the same settings were used.</p>
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<p>EDX analysis of a shear tested resistor after two reflow cycles. (<b>a</b>) SEM image in SE mode with a magnification of 300× and (<b>b</b>) associated combined EDX map of Cu, Ni, Ag, Sn and Al. (<b>c</b>) Detailed SEM image in SE mode of the location between the ragged resistor and the Cu pad of the PCB with a higher magnification of 2000× and (<b>d</b>) associated combined EDX map of Cu, Ni, Ag, Sn, and Al. For each pair of pictures, the same settings were used.</p>
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<p>EDX analysis of a shear tested resistor after two reflow cycles. (<b>a</b>) SEM image in SE mode with a magnification of 300× and (<b>b</b>) associated combined EDX map of Cu, Ni, Ag, Sn and Al. (<b>c</b>) Detailed SEM image in SE mode of the location between the ragged resistor and the Cu pad of the PCB with a higher magnification of 2000× and (<b>d</b>) associated combined EDX map of Cu, Ni, Ag, Sn, and Al. For each pair of pictures, the same settings were used.</p>
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<p>EDX analysis of a shear tested resistor after four reflow cycles. (<b>a</b>) SEM image in SE mode with a magnification of 300× and (<b>b</b>) associated combined EDX map of Cu, Ni, Ag, Sn and Al. (<b>c</b>) Detailed SEM image in SE mode of the location between the ragged resistor and the Cu pad of the PCB with a higher magnification of 2000× and (<b>d</b>) associated combined EDX map of Cu, Ni, Ag, Sn, and Al. For each pair of pictures, the same settings were used.</p>
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<p>EDX analysis of a shear tested resistor after eight reflow cycles. (<b>a</b>) SEM image in SE mode with a magnification of 300× and (<b>b</b>) associated combined EDX map of Cu, Ni, Ag, Sn and Al. (<b>c</b>) Detailed SEM image in SE mode of the location between the ragged resistor and the Cu pad of the PCB with a higher magnification of 2000× and (<b>d</b>) associated combined EDX map of Cu, Ni, Ag, Sn, and Al. For each pair of pictures, the same settings were used.</p>
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<p>Light microscopy images with a magnification of 5× of the non-ragged side of resistors after the shear test after (<b>a</b>) one, (<b>b</b>) two, (<b>c</b>) four, and (<b>d</b>) eight reflow cycles.</p>
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15 pages, 3542 KiB  
Article
Effect of (Ba1/3Nb2/3)4+ Substitution on Microstructure, Bonding Properties and Microwave Dielectric Properties of Ce2Zr3(MoO4)9 Ceramics
by Huamin Gao, Xiangyu Xu, Xinwei Liu, Xiaoyu Zhang, Mingling Li, Jialun Du and Haitao Wu
Ceramics 2024, 7(3), 1172-1186; https://doi.org/10.3390/ceramics7030077 - 29 Aug 2024
Viewed by 304
Abstract
In this study, Ce2[Zr1−x(Ba1/3Nb2/3)x]3(MoO4)9 (0.02 ≤ x ≤ 0.1, CZ1−xNx) ceramics were sintered at 600 °C and 700 °C using the traditional [...] Read more.
In this study, Ce2[Zr1−x(Ba1/3Nb2/3)x]3(MoO4)9 (0.02 ≤ x ≤ 0.1, CZ1−xNx) ceramics were sintered at 600 °C and 700 °C using the traditional solid-state method. An analysis conducted through XRD and Rietveld refinement confirmed that all the CZ1−xNx ceramics displayed a single phase with a trigonal structure (space group R-3c). The observed increases in cell volume with increasing x values indicate the successful substitution of (Ba1/3Nb2/3)4+. The high densification of the synthesized phase was validated by the density and SEM results. Additionally, the P-V-L theory demonstrates a strong correlation between the Ce-O bond and εr, as well as τf, and between the Mo-O bond and Q×f. Notably, the CZ0.98N0.02 ceramics demonstrated superior performance at 675 °C, exhibiting εr = 10.41, Q×f = 53,296 GHz, and τf = −23.45 ppm/°C. Finally, leveraging CZ0.98N0.02 ceramics as substrate materials enabled the design of a patch antenna suitable for the 5G communication band, demonstrating its significant potential in this field. Full article
(This article belongs to the Special Issue Advances in Electronic Ceramics)
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<p>XRD patterns of CZ<sub>1−<span class="html-italic">x</span></sub>N<span class="html-italic"><sub>x</sub></span> ceramics sintered at the optimal temperature with different <span class="html-italic">x</span> values.</p>
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<p>Rietveld refinement of CZ<sub>1−<span class="html-italic">x</span></sub>N<span class="html-italic"><sub>x</sub></span> ceramics at the optimal sintering temperature with different <span class="html-italic">x</span> values.</p>
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<p>Lattice parameter changes (<b>a</b>) <span class="html-italic">a</span> and <span class="html-italic">b</span>, (<b>b</b>) <span class="html-italic">c</span>, and (<b>c</b>) <span class="html-italic">V<sub>m</sub></span> of CZ<sub>1−<span class="html-italic">x</span></sub>N<span class="html-italic"><sub>x</sub></span> ceramics as a function of <span class="html-italic">x</span> values.</p>
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<p>Crystal structure diagram of CZ<sub>1−<span class="html-italic">x</span></sub>N<span class="html-italic"><sub>x</sub></span> ceramics.</p>
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<p>(<b>a</b>) Diameter shrinkage, (<b>b</b>) apparent density (the relative density at the optimal sintering temperature as a function of <span class="html-italic">x</span> values are shown in the inset) of CZ<sub>1−<span class="html-italic">x</span></sub>N<span class="html-italic"><sub>x</sub></span> ceramics at 600 to 700 °C.</p>
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<p>Microstructure of CZ<sub>1−<span class="html-italic">x</span></sub>N<span class="html-italic"><sub>x</sub></span> ceramics at optimal sintering temperature with different <span class="html-italic">x</span> values.</p>
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<p>CZ<sub>1−<span class="html-italic">x</span></sub>N<span class="html-italic"><sub>x</sub></span> ceramics (<b>a</b>) <span class="html-italic">ε<sub>r</sub></span> at a sintering temperature of 600 ℃ to 700 ℃, (<b>b</b>) <span class="html-italic">ε<sub>r</sub></span>, (<b>c</b>) <span class="html-italic">ε<sub>corr.</sub></span>, (<b>d</b>) <span class="html-italic">α<sub>theo.</sub></span>, and (<b>e</b>) <span class="html-italic">f<sub>iave.</sub></span><sub>(Ce-O)</sub> at the optimal sintering temperature as a function of <span class="html-italic">x</span> values.</p>
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<p>CZ<sub>1−<span class="html-italic">x</span></sub>N<span class="html-italic"><sub>x</sub></span> ceramics (<b>a</b>) <span class="html-italic">Q</span>×<span class="html-italic">f</span> at a sintering temperature of 600 ℃ to 700 ℃, (<b>b</b>) <span class="html-italic">Q</span>×<span class="html-italic">f</span>, and (<b>c</b>) <span class="html-italic">U<sub>ave.</sub></span><sub>(Mo-O)</sub> at the optimal sintering temperature as a function of <span class="html-italic">x</span> values.</p>
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<p>CZ<sub>1−<span class="html-italic">x</span></sub>N<span class="html-italic"><sub>x</sub></span> ceramics (<b>a</b>) <span class="html-italic">τ<sub>f</sub></span>, (<b>b</b>) <span class="html-italic">α<sub>ave.</sub></span><sub>(Ce-O)</sub>, and (<b>c</b>) <span class="html-italic">E<sub>ave.</sub></span><sub>(Mo-O)</sub> at the optimal sintering temperature as a function of <span class="html-italic">x</span> values.</p>
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<p>(<b>a</b>) The design model and dimensions, (<b>b</b>) simulated S11 parameters, (<b>c</b>) 3D radiation pattern, (<b>d</b>) E-plane, and (<b>e</b>) H-plane of the antenna.</p>
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13 pages, 6650 KiB  
Article
Influence of Bond Coat Roughness on Adhesion of Thermal Barrier Coatings Deposited by the Electron Beam–Physical Vapour Deposition Process
by Grzegorz Maciaszek and Andrzej Nowotnik
Appl. Sci. 2024, 14(16), 7401; https://doi.org/10.3390/app14167401 - 22 Aug 2024
Viewed by 363
Abstract
Thermal barrier coatings (TBCs) are effective protective and insulative coatings on hot section components of turbine engines. The quality and subsequent performance of the TBCs are strongly dependent on the adhesion between the coating and the metal substrate. The adhesion strength of TBCs [...] Read more.
Thermal barrier coatings (TBCs) are effective protective and insulative coatings on hot section components of turbine engines. The quality and subsequent performance of the TBCs are strongly dependent on the adhesion between the coating and the metal substrate. The adhesion strength of TBCs varies depending on the substrate materials and coating, the coating technique used, the coating application parameters, the substrate surface treatments, and environmental conditions. Therefore, the roughness of the substrate surface has a significant effect on the performance of the TBC system. In this work, the roughness and microstructure of the 7YSZ (7 wt.% yttria-stabilised zirconia) top coat under different bond coat roughness treatments were studied. The purpose of this paper was to investigate the influence of the roughness of the bond coat on the adhesion of 7YSZ TBCs prepared by the electron beam–physical vapour deposition (EB-PVD) process. The VPA (vapour phase aluminium) bond coat was deposited on Inconel 718 nickel superalloy substrate using the above-the-pack technique. The ceramic top coat was applied to the bond coat using the EB-PVD process. The dependence between the TBC coating roughness and the bond coat roughness was determined. Adhesion strength measurements were performed according to the ASTM C 633 standard test method. The highest adhesion value observed in the tensile adhesion tests was 105 MPa. However, it was not determined whether the surface roughness of the bond coat affects the adhesion of the 7YSZ top coat. Full article
(This article belongs to the Section Surface Sciences and Technology)
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<p>Samples in the holder: (<b>a</b>) before the ED-PVD process; (<b>b</b>) mounted in a middle rake position; (<b>c</b>) after the EB-PVD process.</p>
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<p>The SMART Coater at Rzeszow University of Technology.</p>
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<p>The coated sample glued to counterpart (<b>a</b>) before the tensile adhesion test; and (<b>b</b>) after the tensile adhesion test.</p>
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<p>Influence of bond coat surface roughness on top coat roughness approximated by the equation <span class="html-italic">y</span> = 0.5328<span class="html-italic">x</span><sup>2</sup> − 0.0531<span class="html-italic">x</span> + 0.4277.</p>
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<p>Surface profiles of (<b>a</b>) sample I as-coated bond coat with Ra of 1.065 µm; (<b>b</b>) sample I top coat with Ra of 1.118 µm; (<b>c</b>) sample VI rough-grinded bond coat with Ra of 0.475 µm; (<b>d</b>) sample VI top coat with Ra of 0.521 µm; (<b>e</b>) sample VII grinded bond coat with Ra of 0.287 µm; (<b>f</b>) sample VII top coat with Ra of 0.437 µm; (<b>g</b>) sample IV polished bond coat with Ra of 0.018 µm; and (<b>h</b>) sample IV top coat with Ra of 0.385 µm.</p>
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<p>Surface profiles of (<b>a</b>) sample I as-coated bond coat with Ra of 1.065 µm; (<b>b</b>) sample I top coat with Ra of 1.118 µm; (<b>c</b>) sample VI rough-grinded bond coat with Ra of 0.475 µm; (<b>d</b>) sample VI top coat with Ra of 0.521 µm; (<b>e</b>) sample VII grinded bond coat with Ra of 0.287 µm; (<b>f</b>) sample VII top coat with Ra of 0.437 µm; (<b>g</b>) sample IV polished bond coat with Ra of 0.018 µm; and (<b>h</b>) sample IV top coat with Ra of 0.385 µm.</p>
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<p>SEM cross-sectional morphologies of the 7YSZ top coat deposited on (<b>a</b>) sample I as-coated bond coat; (<b>b</b>) sample II rough-grinded bond coat; (<b>c</b>) sample III grinded bond coat; (<b>d</b>) sample IV polished bond coat. Figures (<b>e</b>–<b>h</b>) show higher-magnification images of the coating microstructure presented in (<b>a</b>–<b>d</b>).</p>
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<p>The coated samples: (<b>a</b>) VI; (<b>b</b>) VII; and their counterparts after the tensile adhesion strength test.</p>
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<p>Tensile stress–strain curves obtained for adhesion strength of (<b>a</b>) 65.7 MPa; (<b>b</b>) 91.5 MPa; (<b>c</b>) 91.6 MPa; (<b>d</b>) 105.0 MPa; (<b>e</b>) 78.7 MPa; (<b>f</b>) 82.5 MPa; (<b>g</b>) 66.2 MPa; and (<b>h</b>) 74.6 MPa.</p>
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16 pages, 4424 KiB  
Article
Mechanism of Ag-SiO2-TiO2 Nanocomposite Coating Formation on NiTi Substrate for Enhanced Functionalization
by Karolina Dudek, Mateusz Dulski, Jacek Podwórny, Magdalena Kujawa and Patrycja Rawicka
Coatings 2024, 14(8), 1055; https://doi.org/10.3390/coatings14081055 - 18 Aug 2024
Viewed by 512
Abstract
The functionality of the NiTi shape memory alloy was improved through engineering Ag-SiO2-TiO2 nanocomposite coatings. For this purpose, an anaphoretic deposition process, conducted at a constant voltage of 40 V and deposition times ranging from 1 to 10 min, was [...] Read more.
The functionality of the NiTi shape memory alloy was improved through engineering Ag-SiO2-TiO2 nanocomposite coatings. For this purpose, an anaphoretic deposition process, conducted at a constant voltage of 40 V and deposition times ranging from 1 to 10 min, was used. Scanning electron microscopy (SEM) analysis demonstrated that the deposition parameters significantly impacted the morphology of the coatings. Complementary Raman Spectroscopy and X-ray diffraction (XRD) analyses confirmed the successful formation of distinct nanocomposite layers, and revealed the details of their crystalline structure and chemical composition. After that, the adhesion between the NiTi substrate and the electrophoretically deposited ceramic coatings was improved through a post-deposition heat treatment. To prevent excessive shrinkage and cracking of the coating, tests were carried out to characterize the behavior of the coating material at elevated temperatures. The nanocomposite coatings were exposed to a temperature of 800 °C for 2 h. The annealing induced significant structural and morphological transformations, resulting in layers that were distinctly different from both the original materials and those produced solely through electrophoretic deposition. The thermal treatment resulted in the formation of a new kind of nanocomposite structure with enhanced reactivity. Full article
(This article belongs to the Section Surface Characterization, Deposition and Modification)
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<p>The Zeta potential and particle size of the colloidal suspension at different pH.</p>
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<p>SEM images at different magnifications (<b>a</b>–<b>e</b>), and element distribution (<b>f</b>,<b>g</b>) in the Ag-SiO<sub>2</sub>-TiO<sub>2</sub> coatings deposited at different conditions.</p>
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<p>XRD patterns collected for the deposited coating.</p>
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<p>A chemical and structural differentiation image of the Ag-SiO<sub>2</sub>-TiO<sub>2</sub> coating visualized in (<b>a</b>) X-, Y-, and (<b>b</b>) X-, Z-, with (<b>c</b>) depth scan profiles of the four exemplary places among the AB cross line. Dash-lines on the depth profiles determine the boundary of the layer. (<b>d</b>) FTIR average spectrum gathered for the red dash-line square. (<b>e</b>) Averaged Raman spectra corresponded to individual color-highlighted phases of the Raman maps.</p>
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<p>Change in linear dimensions of the nanocomposite powder with temperature (<b>a</b>), dimensional changes where T<sub>s</sub> is a sintering start temperature and T<sub>A</sub> is the sintering finish temperature, (<b>b</b>) and photographs of the sample during the test in chosen temperatures in the microscope (<b>c</b>).</p>
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<p>SEM images at different magnifications, (<b>a</b>) and elements distribution (<b>b</b>) in the Ag-SiO<sub>2</sub>-TiO<sub>2</sub> after heat treatment.</p>
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<p>XRD patterns collected for the coating after heat treatment.</p>
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<p>A chemical and structural differentiation image of the coating after the heat treatment visualized in (<b>a</b>) X-, Y-, and (<b>b</b>) X-, Z-, with (<b>c</b>) depth scan profiles of the four exemplary places among the AB cross line. Dash-lines on the depth profiles determine the boundary of the Ag-SiO<sub>2</sub>-TiO<sub>2</sub> layer, while the blue-colored area corresponds to the thin interlayer. (<b>d</b>) FTIR average spectrum gathered of the red dash-line square. (<b>e</b>) Averaged Raman spectra corresponded to individual color-highlighted phases of the Raman maps.</p>
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15 pages, 6364 KiB  
Article
Microstructure and Wear Resistance of In Situ Synthesized Ti(C, N) Ceramic-Reinforced Nickel-Based Coatings by Laser Cladding
by Juncai Li, Ying Chen, Chuang Guan, Chao Zhang, Ji Zhao and Tianbiao Yu
Materials 2024, 17(15), 3878; https://doi.org/10.3390/ma17153878 - 5 Aug 2024
Cited by 1 | Viewed by 713
Abstract
In recent years, laser cladding technology has been widely used in surface modification of titanium alloys. To improve the wear resistance of titanium alloys, ceramic-reinforced nickel-based composite coatings were prepared on a TC4 alloy substrateusing coaxial powder feeding laser cladding technology. Ti (C, [...] Read more.
In recent years, laser cladding technology has been widely used in surface modification of titanium alloys. To improve the wear resistance of titanium alloys, ceramic-reinforced nickel-based composite coatings were prepared on a TC4 alloy substrateusing coaxial powder feeding laser cladding technology. Ti (C, N) ceramic was synthesized in situ by laser cladding by adding different contents (10%, 20%, 30%, and 40%) of TiN, pure Ti powder, graphite, and In625 powder. Thisestudy showed that small TiN particles were decomposed and directly formed the Ti (C, N) phase, while large TiN particles were not completely decomposed. The in situ synthetic TiCxN1−x phase was formed around the large TiN particles. With the increase in the proportion of powder addition, the wear volume of the coating shows a decreasing trend, and the wear resistance of the surface coating is improving. The friction coefficient of the sample with 40% TiN, pure Ti powder, and graphite powder is 0.829 times that of the substrate. The wear volume is 0.145 times that of the substrate. The reason for this is that with the increase in TiN, Ti, and graphite in the powder, there are more ceramic phases in the cladding layer, and the hard phases such as TiC, Ti(C, N) and Ti2Ni play the role in the structure of the “backbone”, inhibit the damage caused by micro-cutting, and impede the movement of the tearing point of incision, so that the coating has a higher abrasion resistance. Full article
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<p>Experimental setup diagram of laser cladding system.</p>
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<p>Preparation process of test samples.</p>
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<p>Temperature—dependent Gibbs free energy curves for various potential reactions during the solidification process.</p>
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<p>X-ray diffraction (XRD) patterns of different samples.</p>
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<p>EDS results of sample 1#: (<b>a</b>) electronic image, (<b>b</b>) EDS layered image, (<b>c</b>) Ti element, (<b>d</b>) N element, (<b>e</b>) C element, (<b>f</b>) Ni element, (<b>g</b>) Cr element, (<b>h</b>) Nb element.</p>
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<p>EDS surface scan image of sample 4#: (<b>a</b>) electronic image, (<b>b</b>) EDS layered image, (<b>c</b>) Ti element, (<b>d</b>) N element, (<b>e</b>) C element, (<b>f</b>) Ni element, (<b>g</b>) Cr element, (<b>h</b>) Nb element.</p>
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<p>Local line scanning results of sample 4.</p>
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<p>Friction coefficients of the matrix and various samples.</p>
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<p>Wear surface morphology of the substrate and various samples.</p>
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<p>Histograms of wear volume of the substrate and each sample.</p>
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<p>Microhardness and average microhardness on TC4 substrate and specimens 1#–4#.</p>
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<p>Optical surface morphology of abrasion marks on TC4 substrate and specimens 1#–4#.</p>
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17 pages, 2702 KiB  
Review
Microorganisms in Red Ceramic Building Materials—A Review
by Elżbieta Stanaszek-Tomal
Coatings 2024, 14(8), 985; https://doi.org/10.3390/coatings14080985 - 5 Aug 2024
Viewed by 828
Abstract
Ceramic materials have a very long tradition of use in construction. Their durability is related to the surface of the material and the action of the corrosive environment. One of the corrosive factors acting on ceramic materials is microorganisms. They can contribute to [...] Read more.
Ceramic materials have a very long tradition of use in construction. Their durability is related to the surface of the material and the action of the corrosive environment. One of the corrosive factors acting on ceramic materials is microorganisms. They can contribute to the deterioration of the technical and performance properties of building materials. Aesthetic, physical, and chemical deterioration are considered to be the main destructive processes in ceramic materials. This work shows how the different types of the most commonly used ceramic materials, i.e., brick and tiles, are damaged. Each of these types is susceptible to microbial growth. Most microorganisms that occur on ceramic materials produce staining substances and thus form coloured biofilms. The direct action of metabolic products secreted by organisms on inorganic substrates is the main cause of chemical biodeterioration. Therefore, this work presents the impact of microorganisms on ceramic building materials. Full article
(This article belongs to the Section Bioactive Coatings and Biointerfaces)
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<p>The classification of ceramic building materials.</p>
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<p>The most important causes behind damage to ceramic materials.</p>
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<p>Mechanisms of biodeterioration of natural rocks.</p>
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<p>Ordinary brick samples: first contaminated with MgSO<sub>4</sub> salt and then with <span class="html-italic">Cladosporium herbarum</span> fungi: (<b>a</b>) contaminated sample; (<b>b</b>) microstructure of bricks with visible mycelium hyphals (own research).</p>
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<p>Clay tiles with visible biodeterioration effects.</p>
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<p>Visible traces of microorganisms on bricks and mortar: (<b>a</b>) view of the wall, (<b>b</b>) contaminated bricks–enlargement.</p>
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<p>Processes occurring during biodeterioration under the influence of metabolic activity of microorganisms, based on [<a href="#B99-coatings-14-00985" class="html-bibr">99</a>], CC BY-NC-ND 4.0.</p>
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13 pages, 17643 KiB  
Article
Zirconia and Crofer Joint Made by Reactive Air Brazing Using the Silver Base Paste and Cu-Ti Coating Layer
by Shu-Wei Chang, Ren-Kae Shiue and Liang-Wei Huang
Materials 2024, 17(15), 3822; https://doi.org/10.3390/ma17153822 - 2 Aug 2024
Viewed by 435
Abstract
This study proposes a method to enhance the airtightness of the joint between the ZrO2 and Crofer alloy using coating technology. With the aid of vacuum sputtering technology, a titanium–copper alloy layer with a thickness between 1.5 μm and 6 μm was [...] Read more.
This study proposes a method to enhance the airtightness of the joint between the ZrO2 and Crofer alloy using coating technology. With the aid of vacuum sputtering technology, a titanium–copper alloy layer with a thickness between 1.5 μm and 6 μm was first deposited on the surface of ZrO2 and Crofer, respectively. The chemical composition of the deposited reaction layer was 70.2 Cu and 29.8 Ti in at%. Then, using silver as the base material in the reactive air brazing (RAB) process, we explore the use of this material design to improve the microstructure and reaction mechanism of the joint surface between ceramics and metal, compare the effects of different pretreatment thicknesses on the microstructure, and evaluate its effectiveness through air tightness tests. The results show that a coating of Cu-Ti alloy on the ZrO2 substrate can significantly improve bonding between the Ag filler and ZrO2. The Cu-Ti metallization layer on the ZrO2 substrate is beneficial to the RAB. After the brazing process, the coated Cu-Ti layers form suitable reaction interfaces between the filler, the metal, the filler, and the ceramic. In terms of coating layer thickness, the optimized 3 μm coated Cu-Ti alloy layer is achieved from the experiment. Melting and dissolving the Cu-Ti coated layer into the ZrO2 substrate results in a defect-free interface between the Ag-rich braze and the ZrO2. The air tightness test result shows no leakage under 2 psig at room temperature for 28 h. The pressure condition can still be maintained even under high-temperature conditions of 600 °C for 24 h. Full article
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<p>The schematic diagram of the pressure drop test: (<b>a</b>) the brazed specimen, (<b>b</b>) test infrastructure.</p>
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<p>(<b>a</b>) SEI, (<b>b</b>) BEI of the ZrO<sub>2</sub>/Ag/Crofer joint reactive air brazed at 960 °C for 1200 s; (<b>c</b>) higher magnification at the Ag/Crofer interface, (<b>d</b>) higher magnification at the ZrO<sub>2</sub>/Ag interface.</p>
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<p>(<b>a</b>) SEI, (<b>b</b>) BEI of the ZrO<sub>2</sub>/Ag/Crofer joint reactive air brazed with 3 μm coating layer at 960 °C for 1200 s; (<b>c</b>) higher magnification at the Ag/Crofer interface, (<b>d</b>) higher magnification at the ZrO<sub>2</sub>/Ag interface.</p>
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<p>EPMA element mappings of the Crofer side interface in <a href="#materials-17-03822-f003" class="html-fig">Figure 3</a>c: (<b>a</b>) Ag, (<b>b</b>) Cr, (<b>c</b>) Cu, (<b>d</b>) Fe, (<b>e</b>) O, (<b>f</b>) Ti, (<b>g</b>) Y, and (<b>h</b>) Zr.</p>
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<p>EPMA element mappings of the ZrO<sub>2</sub> side interface in <a href="#materials-17-03822-f003" class="html-fig">Figure 3</a>d: (<b>a</b>) Ag, (<b>b</b>) Cr, (<b>c</b>) Cu, (<b>d</b>) Fe, (<b>e</b>) O, (<b>f</b>) Ti, (<b>g</b>) Y, and (<b>h</b>) Zr.</p>
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<p>(<b>a</b>) SEI, (<b>b</b>) BEI of the ZrO<sub>2</sub>/Ag/Crofer joint brazed with 1.5 μm coated film at 960 °C for 1200 s, (<b>c</b>) higher magnification at the Ag/Crofer interface, (<b>d</b>) higher magnification at the ZrO<sub>2</sub>/Ag interface.</p>
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<p>EPMA element mappings of the Crofer side interface in <a href="#materials-17-03822-f006" class="html-fig">Figure 6</a>c: (<b>a</b>) Ag, (<b>b</b>) Cr, (<b>c</b>) Cu, (<b>d</b>) Fe, (<b>e</b>) O, (<b>f</b>) Ti, (<b>g</b>) Y, and (<b>h</b>) Zr.</p>
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<p>EPMA element mappings of the ZrO<sub>2</sub> side interface in <a href="#materials-17-03822-f006" class="html-fig">Figure 6</a>d: (<b>a</b>) Ag, (<b>b</b>) Cr, (<b>c</b>) Cu, (<b>d</b>) Fe, (<b>e</b>) O, (<b>f</b>) Ti, (<b>g</b>) Y, and (<b>h</b>) Zr.</p>
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<p>(<b>a</b>) SEI, (<b>b</b>) BEI of the ZrO<sub>2</sub>/Ag/Crofer joint brazed with 6 μm coated film at 960 °C for 1200 s; (<b>c</b>) higher magnification at the Ag/Crofer interface, (<b>d</b>) higher magnification at the ZrO<sub>2</sub>/Ag interface.</p>
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<p>EPMA element mappings of the Crofer side interface in <a href="#materials-17-03822-f009" class="html-fig">Figure 9</a>c: (<b>a</b>) Ag, (<b>b</b>) Cr, (<b>c</b>) Cu, (<b>d</b>) Fe, (<b>e</b>) O, (<b>f</b>) Ti, (<b>g</b>) Y, and (<b>h</b>) Zr.</p>
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<p>EPMA element mappings of the ZrO<sub>2</sub> side interface in <a href="#materials-17-03822-f009" class="html-fig">Figure 9</a>d: (<b>a</b>) Ag, (<b>b</b>) Cr, (<b>c</b>) Cu, (<b>d</b>) Fe, (<b>e</b>) O, (<b>f</b>) Ti, (<b>g</b>) Y, and (<b>h</b>) Zr.</p>
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<p>The pressure drop in the joint leak test was performed at room temperature for 28 h.</p>
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<p>The pressure drop in the joint leak test was at 600 °C for 24 h: (<b>a</b>) temperature and (<b>b</b>) pressure profiles.</p>
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12 pages, 5491 KiB  
Article
Direct Ink Writing of SiCN/RuO2/TiB2 Composite Ceramic Ink for High-Temperature Thin-Film Sensors
by Yusen Wang, Lida Xu, Xiong Zhou, Fuxin Zhao, Jun Liu, Siqi Wang, Daoheng Sun and Qinnan Chen
Materials 2024, 17(15), 3792; https://doi.org/10.3390/ma17153792 - 1 Aug 2024
Viewed by 464
Abstract
Direct ink writing (DIW) of high-temperature thin-film sensors holds significant potential for monitoring extreme environments. However, existing high-temperature inks face a trade-off between cost and performance. This study proposes a SiCN/RuO2/TiB2 composite ceramic ink. The added TiB2, after [...] Read more.
Direct ink writing (DIW) of high-temperature thin-film sensors holds significant potential for monitoring extreme environments. However, existing high-temperature inks face a trade-off between cost and performance. This study proposes a SiCN/RuO2/TiB2 composite ceramic ink. The added TiB2, after annealing in a high-temperature atmospheric environment, forms B2O3 glass, which synergizes with the SiO2 glass phase formed from the SiCN precursor to effectively encapsulate RuO2 particles. This enhances the film’s density and adhesion to the substrate, preventing RuO2 volatilization at high temperatures. Additionally, the high conductivity of TiB2 improves the film’s overall conductivity. Test results indicate that the SiCN/RuO2/TiB2 film exhibits high linearity from room temperature to 900 °C, high stability (resistance drift rate of 0.1%/h at 800 °C), and high conductivity (4410 S/m). As a proof of concept, temperature sensors and a heat flux sensor were successfully fabricated on a metallic hemispherical surface. Performance tests in extreme environments using high-power lasers and flame guns verified that the conformal thin-film sensor can accurately measure spherical temperature and heat flux, with a heat flux sensor response time of 53 ms. In conclusion, the SiCN/RuO2/TiB2 composite ceramic ink developed in this study offers a high-performance and cost-effective solution for high-temperature conformal thin-film sensors in extreme environments. Full article
(This article belongs to the Special Issue Surface Technology and Coatings Materials)
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<p>Preparation of SiCN/RuO<sub>2</sub>/TiB<sub>2</sub> composite ink and sensor manufacturing. (<b>a</b>) Ink preparation; (<b>b</b>) 3D-printing patterning.</p>
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<p>Surface characteristics of the SiCN/RuO<sub>2</sub>/TiB<sub>2</sub> composite film. (<b>a</b>) Surface morphology of the composite film at different sintering temperatures. (<b>b</b>) Cross-sectional changes of the film after annealing at different sintering temperatures. (<b>c</b>) Graph of conductivity versus sintering temperature. (<b>d</b>) XRD pattern.</p>
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<p>(<b>a</b>)Surface morphology and EDS elemental mapping of the printed SiCN/RuO<sub>2</sub>/TiB<sub>2</sub> composite film. (<b>b</b>) Cross-sectional surface morphology and EDS elemental mapping of the film sintered at 800 °C for 1 h. (<b>c</b>) Sintering process schematic diagram.</p>
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<p>Thermistor performance testing of the SiCN/RuO<sub>2</sub>/TiB<sub>2</sub> composite film thermistor. (<b>a</b>) Test object image. (<b>b</b>) Cold and hot cycle test. (<b>c</b>) Second-order fitting of experimental data. (<b>d</b>) Test results after 30 h of holding at 800 °C. (<b>e</b>) Durability test under high-temperature cyclic conditions.</p>
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<p>Harsh environment testing of the SiCN/RuO<sub>2</sub>/TiB<sub>2</sub> composite material thermistor array. (<b>a</b>) Schematic of laser irradiation on the metallic hemispherical surface thermistor array. (<b>b</b>) Thermistor array design scheme. (<b>c</b>) Test results of gradually increasing laser power on a single thermistor on the metallic hemispherical surface. (<b>d</b>) Infrared thermal imaging at 60s under a certain power of laser irradiation. (<b>e</b>) Comparison of conformal film thermistor and K-type thermocouple under the flame gun impact. (<b>f</b>) Infrared thermal imaging under flame gun impact.</p>
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<p>(<b>a</b>) Schematic of laser irradiation on the metallic hemispherical surface heat flux sensor. (<b>b</b>) Structural principle schematic of the heat flux sensor. (<b>c</b>) Test results under different heat fluxes. (<b>d</b>) Test results of 21 cycles under the same heat flux. (<b>e</b>) Response time test results. (<b>f</b>) Optical image during flame gun impact test. (<b>g</b>) Test results of the heat flux sensor under flame gun impact.</p>
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18 pages, 6190 KiB  
Article
Physicochemical Properties of (La,Sr)CoO3 Thick Films on Fe-25Cr Steel under Exposure to SOFC Cathode Operating Conditions
by Janusz Prażuch, Michał Pyzalski, Daniel Fernández González and Tomasz Brylewski
Materials 2024, 17(15), 3791; https://doi.org/10.3390/ma17153791 - 1 Aug 2024
Viewed by 442
Abstract
La0.6Sr0.4CoO3 (LSC) coatings with a thickness of 50–100 µm were deposited on Fe-25Cr ferritic stainless steel (DIN 50049) via screen printing. The required suspension had been prepared using fine LSC powders synthesised using EDTA gel processes. In its [...] Read more.
La0.6Sr0.4CoO3 (LSC) coatings with a thickness of 50–100 µm were deposited on Fe-25Cr ferritic stainless steel (DIN 50049) via screen printing. The required suspension had been prepared using fine LSC powders synthesised using EDTA gel processes. In its bulk form, the LSC consisted entirely of the rhombohedral phase with space group R-3c, and it exhibited high electrical conductivity (~144 S·cm−1). LSC-coated steel was oxidised in air at 1073 K, i.e., under conditions corresponding to SOFC cathode operation, for times of up to 144 h. The in situ electrical resistance of the steel/La0.6Sr0.4CoO3 layered system during oxidation was measured. The products formed on the samples after the oxidation reaction resulting from exposure to the corrosive medium were investigated using XRD, SEM-EDS, and TEM-SAD. The microstructural, nanostructural, phase, and chemical analysis of films was performed with a focus on the film/substrate interface. It was determined that the LSC coating interacts with the oxidised steel in the applied conditions, and a multi-layer interfacial zone is formed. Detailed TEM-SAD observations indicated the formation of a main layer consisting of SrCrO4, which was the reaction product of (La,Sr)CoO3, and the Cr2O3 scale formed on the metal surface. The formation of the SrCrO4 phase resulted in improved electrical conductivity of the investigated metal/ceramics layered composite material, as demonstrated by the low area-specific resistance values of 5 mΩ·cm2, thus making it potentially useful as a SOFC interconnect material operating at the tested temperature. In addition, the evaporation rate of chromium measured for the uncoated steel and the steel/La0.6Sr0.4CoO3 layered system likewise indicates that the coating is capable of acting as an effective barrier against the formation of volatile compounds of chromium. Full article
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<p>Graphical summary of the research concept related to the development of a layered system consisting of ferritic steel Fe-25Cr and an (La,Sr)CoO<sub>3</sub> coating.</p>
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<p>Particle size distribution and cumulated distribution of the (La,Sr)CoO<sub>3</sub> powder after calcination at 1273 K for 25 h in air.</p>
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<p>SEM micrograph of (La,Sr)CoO<sub>3</sub> powder after calcination at 1273 K for 25 h in air.</p>
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<p>SEM micrograph of the fractured cross-section of the La<sub>0.6</sub>Sr<sub>0.4</sub>CoO<sub>3</sub> compact sintered at 1473 K in air for 12 h.</p>
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<p>Electrical conductivity vs. temperature in the Arrhenius layout for the La<sub>0.6</sub>Sr<sub>0.4</sub>CoO<sub>3</sub> compact sintered at 1473 K in air for 12 h.</p>
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<p>SEM images of the oxide scale formed on Fe-25Cr steel after oxidation for 144 h in air at 1073 K: (<b>a</b>) image from the polished taper cross-section and (<b>b</b>) image of the scale surface with higher magnification on the insert.</p>
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<p>(<b>a</b>) SEM images of Fe-25Cr steel covered with La<sub>0.6</sub>Sr<sub>0.4</sub>CoO<sub>3</sub> oxidised for 144 h in air at 1073 K; (<b>b</b>) images from the polished taper cross-section; (<b>c</b>) the EDS line scan runs along the black line in (<b>b</b>) across the metal/oxide interphase.</p>
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<p>XRD patterns of (La,Sr)CoO<sub>3</sub> film on Fe-25Cr steel after oxidation in air at 1073 K for 144 h for different depths from the gas/film to the film/metal interface.</p>
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<p>TEM cross-section micrograph of the multilayer metal/oxide interface between the (La,Sr)CoO<sub>3</sub> coating and the Fe-25Cr substrate.</p>
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<p>EDS spectrum of the chromia layer formed at the Fe-25Cr/(La,Sr)CoO<sub>3</sub> interface.</p>
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<p>(<b>a</b>) TEM cross-section micrograph, and (<b>b</b>) SAD pattern with the [211] zone axis of the SrCrO4 layer.</p>
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<p>Quantitative EDS line scan analysis across the Fe-25Cr substrate/(La,Sr)CoO<sub>3</sub> film interface shown in <a href="#materials-17-03791-f009" class="html-fig">Figure 9</a>.</p>
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<p>Area-specific resistance vs. time for (La,Sr)CoO<sub>3</sub>-coated and uncoated Fe-25Cr steels and for the ceramic (La,Sr)CrO<sub>3</sub> at 1073 K.</p>
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<p>Chromium transport rate determined for Fe-25Cr alloy, Fe-25Cr steel, and Fe-25Cr steel coated with (La,Sr)CoO<sub>3</sub>. Test temperature: 1073 K. Test atmosphere: humidified air.</p>
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15 pages, 5169 KiB  
Article
Aluminium Nitride Surface Characterization by Grinding with Laser–Ultrasonic Coupling
by He Zhang, Cong Sun, Yuan Hong, Yansheng Deng and Liang Ma
Materials 2024, 17(15), 3772; https://doi.org/10.3390/ma17153772 - 1 Aug 2024
Viewed by 441
Abstract
Aluminium nitride (AlN) materials are widely used in heat-dissipation substrates and electronic device packages. However, the application of aluminium nitride ceramics is hindered by the obvious anisotropy and high brittleness of its crystals, leading to poor material surface integrity and high grinding force. [...] Read more.
Aluminium nitride (AlN) materials are widely used in heat-dissipation substrates and electronic device packages. However, the application of aluminium nitride ceramics is hindered by the obvious anisotropy and high brittleness of its crystals, leading to poor material surface integrity and high grinding force. With the rapid development of microelectronics, the requirements for the material’s dimensional accuracy, machining efficiency, and surface accuracy are increasing. Therefore, a new machining process is proposed, combining laser and ultrasonic vibration with grinding. The laser–ultrasonic-assisted grinding (LUAG) of aluminium nitride is simulated by molecular dynamics (MD). Meanwhile, the effects of different processing techniques on grinding force, stress distribution, matrix damage mechanism, and subsurface damage depth are systematically investigated and verified by experiments. The results show that laser–ultrasonic-assisted grinding produces 50% lower grinding forces compared to traditional grinding (TG). The microhardness of AlN can reach more than 1200 HV, and the coefficient of friction and wear is reduced by 42.6%. The dislocation lines of the AlN substrate under this process are short but interlaced, making the material prone to phase transformation. Moreover, the subsurface damage depth is low, realising the substrate’s material hardening and wear resistance. These studies not only enhance the comprehension of material build-up and stress damage under the synergistic impact of laser, ultrasonic, and abrasive processing but also indicate that the proposed method can facilitate and realise high-performance machining of aluminium nitride substrate surfaces. Full article
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<p>Three-dimensional MD simulation model.</p>
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<p>The RDF of Al–N, Al–Al, and N–N bonds of AlN via the four processing technologies.</p>
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<p>The grinding force and the amount of growth via the four processing technologies. (<b>a</b>) The grinding force (<b>b</b>) The grinding force.</p>
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<p>The atom flow field via the four processing technologies. (<b>a</b>) TG (<b>b</b>) LAG (<b>c</b>) UVAG (<b>d</b>) LUAG.</p>
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<p>The von Mises shear stress via the four processing technologies. (<b>a1</b>–<b>a2</b>) TG (<b>b1</b>–<b>b2</b>) LAG (<b>c1</b>–<b>c2</b>) UVAG (<b>d1</b>–<b>d2</b>) LUAG.</p>
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<p>The subsurface damage depth via the four processing technologies. (<b>a</b>) TG (<b>b</b>) LAG (<b>c</b>) UVAG (<b>d</b>) LUAG.</p>
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<p>The surface mechanical properties via the four processing technologies. (<b>a</b>) TG (<b>b</b>) LAG (<b>c</b>) UVAG (<b>d</b>) LUAG.</p>
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<p>LUAG experimental platform.</p>
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<p>Microhardness and surface roughness of the machining area.</p>
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<p>Measurement results of dynamic grinding force.</p>
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<p>Statistics of friction and wear test results.</p>
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<p>X-ray diffraction (XRD) spectrum of AlN ceramics. (<b>a</b>) Raw surface (<b>b</b>) TG surface (<b>c</b>) LAG surface (<b>d</b>) UVAG surface (<b>e</b>) LUAG surface.</p>
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13 pages, 17421 KiB  
Communication
The Direct Cold Sintering of α-Al2O3 Ceramics in a Pure Water Medium
by Anastasia A. Kholodkova, Maxim V. Kornyushin, Arseniy N. Khrustalev, Levko A. Arbanas, Andrey V. Smirnov and Yurii D. Ivakin
Ceramics 2024, 7(3), 1030-1042; https://doi.org/10.3390/ceramics7030067 - 31 Jul 2024
Viewed by 592
Abstract
Porous α-Al2O3 ceramics are a highly sought-after material with a multitude of applications; for example, they are used as filters, substrates, biomedicine materials, etc. Despite the availability of raw materials, a challenge associated with this technology is the high energy [...] Read more.
Porous α-Al2O3 ceramics are a highly sought-after material with a multitude of applications; for example, they are used as filters, substrates, biomedicine materials, etc. Despite the availability of raw materials, a challenge associated with this technology is the high energy budget caused by sintering above 1500 °C. For the cold sintering processing (CSP) of ceramics, lowering the α-Al2O3 sintering temperature is one of the most urgent challenges in the background of its rapid development. This paper is the first to demonstrate a solution to this problem using the CSP of α-alumina ceramics in the presence of pure water as a transient liquid. The manufactured materials were examined using XRD analysis; the evolution of their microstructures during CSP was revealed by SEM; and the porosity was evaluated using the Archimedes method. Ceramics with an open porosity up to 36% were produced at 380–450 °C and 220 MPa in 30 min. An increase in the pressure was found to impede α-Al2O3 formation from γ-AlOOH. The development of the microstructure was discussed within the framework of the dissolution–precipitation model and homogenous nucleation. The results of the SEM study pointed to the coalescence of γ-AlOOH grains during CSP. Full article
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<p>Alumina ceramics manufactured by CSP with an addition of 20 wt.% of distilled water (processing parameters: 450 °C, 220 MPa).</p>
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<p>XRD patterns (<b>a</b>) and phase fractions (<b>b</b>) of the ceramic samples manufactured by CSP at 450 °C and the indicated mechanical pressure (90–350 MPa). Miller indices correspond to the following phases: γ-AlOOH (PDF2 #000-83-2384), α-Al<sub>2</sub>O<sub>3</sub> (PDF2 #000-71-1683), and χ-Al<sub>2</sub>O<sub>3</sub> (PDF2 #000-04-0880). Phase fractions are presented excluding foreign impurities.</p>
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<p>XRD patterns (<b>a</b>) and phase fractions (<b>b</b>) of the ceramic samples manufactured by CSP at a pressure of 220 MPa and indicated temperature (380–450 °C). Miller indices correspond to the following phases: γ-AlOOH (PDF2 #000-83-2384) and α-Al<sub>2</sub>O<sub>3</sub> (PDF2 #000-71-1683). Phase fractions are presented excluding foreign impurities.</p>
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<p>XRD patterns (<b>a</b>) and phase fractions (<b>b</b>) of the ceramic samples manufactured by CSP at a pressure of 350 MPa and the indicated temperature (380–450 °C). Miller indices correspond to the following phases: γ-AlOOH (PDF2 #000-83-2384) and α-Al<sub>2</sub>O<sub>3</sub> (PDF2 #000-71-1683).</p>
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<p>Fractured surfaces SEM images of the ceramics manufactured by CSP at a mechanical pressure of 220 MPa and a temperature of 380 °C (<b>a</b>) and 450 °C (<b>b</b>). Corresponding grain size distributions were calculated on the basis of 400 measurements of the grain size in each of the samples.</p>
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<p>SEM images of fractured surfaces of ceramics manufactured by CSP at a mechanical pressure of 350 MPa and a temperature of 380 °C (<b>a</b>), 410 °C (<b>b</b>), and 450 °C (<b>c</b>,<b>d</b>).</p>
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<p>SEM images of fractured surfaces of ceramics manufactured by CSP at a mechanical pressure of 350 MPa and a temperature of 380 °C (<b>a</b>), 410 °C (<b>b</b>), and 450 °C (<b>c</b>,<b>d</b>).</p>
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<p>A schematic representation of the microstructural transformations during the CSP: (<b>a</b>) formation of fine-grained boehmite structure; (<b>b</b>) formation of plate-like boehmite grains by coalescence; (<b>c</b>) α-Al<sub>2</sub>O<sub>3</sub> formation.</p>
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<p>XRD pattern (<b>a</b>) and SEM image (<b>b</b>) of the initial gibbsite (γ-Al(OH)<sub>3</sub>) powder. The observed peaks correspond to PDF2 #000-076-1782. The size distribution is based on the measurements of 300 particles.</p>
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<p>XRD pattern (<b>a</b>) and SEM image (<b>b</b>) of a commercial α-Al<sub>2</sub>O<sub>3</sub> additive with the initial γ-Al(OH)<sub>3</sub> powder. Miller indices correspond to the α-Al<sub>2</sub>O<sub>3</sub> phase (PDF2 #000-075-0782), asterisks indicate the Al<sub>22</sub>O<sub>34</sub>(H<sub>2</sub>O)<sub>2</sub> phase (PDF2 #000-070-1204). The calculated major phase content is 90.5 wt.%. The size distribution is based on the measurements of 1300 particles.</p>
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<p>Schematic representation of the mold prepared for CSP.</p>
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