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Materials, Volume 14, Issue 16 (August-2 2021) – 423 articles

Cover Story (view full-size image): The strain field offers crucial information regarding lattice distortion and phase transformation in the native state or under external perturbation. It is well defined in the continuum limit to measure the deformation, which can be alternatively calculated from the arrangement of atoms in discrete lattices through methods such as geometrical phase analysis from transmission electron imaging, bond distortion or virial stress from atomic structures obtained from molecular simulations. In this paper, these methods are evaluated for the representative problems including structural distortion near defects, dislocations, and the process of fracture where inhomogeneous strain is built up locally. View this paper
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15 pages, 7004 KiB  
Article
Strain Rate Effect upon Mechanical Behaviour of Hydrogen-Charged Cycled NiTi Shape Memory Alloy
by Fehmi Gamaoun
Materials 2021, 14(16), 4772; https://doi.org/10.3390/ma14164772 - 23 Aug 2021
Cited by 3 | Viewed by 2339
Abstract
The rate dependence of thermo-mechanical responses of superelastic NiTi with different imposed strain rates after cycling from 1 to 50 cycles under applied 10−5s−1, 10−4s−1 and 10−3s−1 strain rates, immersion for 3 h [...] Read more.
The rate dependence of thermo-mechanical responses of superelastic NiTi with different imposed strain rates after cycling from 1 to 50 cycles under applied 10−5s−1, 10−4s−1 and 10−3s−1 strain rates, immersion for 3 h and ageing has been investigated. The loaded and unloaded as-received NiTi alloy under an imposed strain of 7.1% have shown an increase in the residual deformation at zero stress with an increase in strain rates. It has been found that after 13 cycles and hydrogen charging, the amount of absorbed hydrogen (291 mass ppm) was sufficient to cause the embrittlement of the tensile loaded NiTi alloy with 10−5s−1. However, no premature fracture has been detected for the imposed strain rates of 10−4s−1 and 10−3s−1. Nevertheless, after 18 cycles and immersion for 3 h, the fracture has occurred in the plateau of the austenite martensite transformation during loading with 10−4s−1. Despite the higher quantity of absorbed hydrogen, the loaded specimen with a higher imposed strain rate of 10−3s−1 has kept its superelasticity behaviour, even after 20 cycles. We attribute such a behaviour to the interaction between the travelling distance during the growth of the martensitic domains while introducing the martensite phase and the amount of diffused hydrogen. Full article
(This article belongs to the Special Issue Mechanical Behavior of Shape Memory Alloys: 2022)
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<p>Loading protocol of the monotonic and cyclic loading: (<b>a</b>) Schematic representation of the Instron tensile machine, (<b>b</b>) used part of the as-received alloy, and (<b>c</b>) clamping method of the NiTi archwire.</p> Full article ">Figure 2
<p>(<b>a</b>) Scheme of hydrogen charging (<b>b</b>) setup by electrolysis.</p> Full article ">Figure 3
<p>Typical engineering stress–strain curves of superelastic Ni-Ti alloy immersed for 3 h and aged for 24 h at different strain rates of 10<sup>−5</sup>s<sup>−1</sup>, 10<sup>−4</sup>s<sup>−1</sup> and 10<sup>−5</sup>s<sup>−1</sup> and as-received at 10<sup>−5</sup>s<sup>−1</sup>.</p> Full article ">Figure 4
<p>Simplified NiTi phase diagram showing the increase in the martensite starting stress from σ<sub>1</sub> to σ<sub>2</sub> when the temperature goes up from T<sub>1</sub> to T<sub>2</sub><sub>,</sub> respectively.</p> Full article ">Figure 5
<p>Typical strain cycling curve after 50 cycles for deformed specimen until 7.1%, at imposed strain rate of (<b>a</b>) 10<sup>−</sup><sup>5</sup>s<sup>−1</sup>, (<b>b</b>) 10<sup>−</sup><sup>4</sup>s<sup>−1</sup> and (<b>c</b>) 10<sup>−</sup><sup>3</sup>s<sup>−1</sup>, showing reduction in phase-transformation yield stress and increase in residual strain.</p> Full article ">Figure 6
<p>Evolution of residual strain at zero stress as a function of number of cycles for cyclic deformed specimen with 7.1% strain after different imposed strain rates of 10<sup>−</sup><sup>5</sup>s<sup>−1</sup>, 10<sup>−</sup><sup>4</sup>s<sup>−1</sup> and 10<sup>−</sup><sup>3</sup>s<sup>−1</sup> (represents the fitting curve).</p> Full article ">Figure 7
<p>Hydrogen thermal desorption curves for specimens immersed for 3 h and aged for 24 h at room temperature after 13 cycles with imposed deformation of (<b>a</b>) 10<sup>−5</sup>s<sup>−1</sup>, (<b>b</b>) 10<sup>−4</sup>s<sup>−1</sup> and (<b>c</b>) 10<sup>−3</sup>s<sup>−1</sup>, showing an increase in the amount of absorbed hydrogen with stain rates.</p> Full article ">Figure 8
<p>Amount of absorbed hydrogen vs. number of cycles of loaded and unloaded specimens at imposed 10<sup>−5</sup>s<sup>−1</sup>, 10<sup>−4</sup>s<sup>−1</sup>, and 10<sup>−3</sup>s<sup>−1</sup> strain rates and hydrogen charging for 3 h (represents the fitting curve).</p> Full article ">Figure 9
<p>(<b>a</b>) Tensile curve at 10<sup>−5</sup>s<sup>−1</sup> showing fracture after 13 cycles with imposed strain rate of 10<sup>−5</sup>s<sup>−1</sup>, and (<b>b</b>) amount of absorbed hydrogen after the same number of cycles at 10<sup>−5</sup>s<sup>−1</sup>, 10<sup>−4</sup>s<sup>−1</sup>, and 10<sup>−3</sup>s<sup>−1</sup>.</p> Full article ">Figure 10
<p>(<b>a</b>) Tensile curves obtained after 18 cycles and immersion for 3 h showing the embrittlement of loaded specimen with imposed strain rates of 10<sup>−5</sup>s<sup>−1</sup> and 10<sup>−4</sup>s<sup>−1</sup> and superelastic behaviour of loaded specimens with 10<sup>−3</sup>s<sup>−1</sup>, and (<b>b</b>) comparison between critical amount of absorbed hydrogen causing fracture at 10<sup>−5</sup>s<sup>−1</sup> and quantity of absorbed hydrogen after 18 cycles with 10<sup>−3</sup>s<sup>−1</sup> and immersion.</p> Full article ">
12 pages, 4987 KiB  
Article
Analysis of the Displacement of Thin-Walled Workpiece Using a High-Speed Camera during Peripheral Milling of Aluminum Alloys
by Jakub Czyżycki, Paweł Twardowski and Natalia Znojkiewicz
Materials 2021, 14(16), 4771; https://doi.org/10.3390/ma14164771 - 23 Aug 2021
Cited by 12 | Viewed by 2487
Abstract
The paper presents the possibilities of a high-speed camera in recording displacements of thin-walled workpiece during milling made of aluminum alloys, which allowed for an analysis in which it was compared to other methods of testing the deflection of such elements. The tests [...] Read more.
The paper presents the possibilities of a high-speed camera in recording displacements of thin-walled workpiece during milling made of aluminum alloys, which allowed for an analysis in which it was compared to other methods of testing the deflection of such elements. The tests were carried out during peripheral milling with constant cutting parameters. Deflection of thin-walled workpiece due to cutting forces was measured using a high-speed camera and a laser displacement sensor. Additionally, the experimental results were compared with the theoretical results obtained with the use of the finite element method. The research proved the effectiveness of the use of high-speed camera in diagnostics of thin-walled workpieces during milling with an accuracy of up to 11% compared to measurements made with a displacement laser sensor. Full article
(This article belongs to the Special Issue Innovative and Modern Technologies of Material Machining in Cutting and Abrasive Processes)
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<p>Research sample containing thin walls.</p> Full article ">Figure 2
<p>Experimental setup.</p> Full article ">Figure 3
<p>Photo of test stand.</p> Full article ">Figure 4
<p>A frame of film recorded with a high-speed camera while milling a thin wall.</p> Full article ">Figure 5
<p>The plot of the deflection of the thin-walled workpiece obtained on the basis of the recording analysis.</p> Full article ">Figure 6
<p>The plot of the deflection of the thin-walled workpiece obtained by measuring the deformation with a displacement laser sensor.</p> Full article ">Figure 7
<p>Comparison of the results of the measurement of deflection during thin-wall milling measured with a laser displacement sensor and a high-speed camera.</p> Full article ">Figure 8
<p>An example of the amplitude-frequency analysis of the displacement of the milled wall.</p> Full article ">Figure 9
<p>Thin wall deflection simulation made in Fusion 360: (<b>a</b>) constant force 100 N; (<b>b</b>) pressure of the end mill on the thin wall.</p> Full article ">Figure 10
<p>Finite element analysis—wall deflection under the influence of a constant force, test a.</p> Full article ">Figure 11
<p>Finite element analysis—wall deflection under influence of the cutter, test b.</p> Full article ">Figure 12
<p>Measurement uncertainty when measuring thin walled workpiece deflection with a high-speed camera and displacement laser sensor.</p> Full article ">
13 pages, 5878 KiB  
Article
Characteristics of 3D Printable Bronze PLA-Based Filament Composites for Gaskets
by Marcela Sava, Ramona Nagy and Karoly Menyhardt
Materials 2021, 14(16), 4770; https://doi.org/10.3390/ma14164770 - 23 Aug 2021
Cited by 8 | Viewed by 2723
Abstract
Composite materials can be tailored for various properties, but the manufacturing process can be quite lengthy depending on the complexity of the final product. Instead, we focused our attention on the relatively new technology of additive manufacturing (3D printing) that can produce complex [...] Read more.
Composite materials can be tailored for various properties, but the manufacturing process can be quite lengthy depending on the complexity of the final product. Instead, we focused our attention on the relatively new technology of additive manufacturing (3D printing) that can produce complex geometries for a limited number of samples. Due to the weak bond between successive printed layers, these objects will have weaker mechanical properties in relation to cast or sintered materials. Thus, the orientation of the printed layers can make a huge difference in the behavior of the products. In this paper, a 3D printed composite made from bronze-filled PLA is mechanically characterized in order to be used as a substitute for sintered compacted bronze products for compression loads. Thus, cylindrical samples grown with the base horizontally and vertically were subjected to compression loads to determine their stress-strain curves at room temperature as well as in the glass transition region. Due to a lack of published research in this area, this study offers an insight into the usability of bronze-filled PLA for gaskets or other objects subjected to compression loads. Full article
(This article belongs to the Topic Structure and Properties Modification in Additive Manufacturing)
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Figure 1
<p>Selection of gasket geometry, from simple to complex.</p> Full article ">Figure 2
<p>Bronze powder samples: (<b>a</b>) green compact, (<b>b</b>) sintered compact.</p> Full article ">Figure 3
<p>PLA bronze-filled samples: (<b>a</b>) vertically layered printed sample; (<b>b</b>) schematic of vertically layered printed sample; (<b>c</b>) top view of the vertically layered printed sample; (<b>d</b>) horizontally layered printed sample; (<b>e</b>) schematic of horizontally layered printed sample, (<b>f</b>) top view of the horizontally layered printed sample.</p> Full article ">Figure 4
<p>Compressive modulus of the sintered powder bronze samples versus sintering time and temperature.</p> Full article ">Figure 5
<p>Vertically layered printed samples after compression load.</p> Full article ">Figure 6
<p>Horizontally layered printed samples after compression load.</p> Full article ">Figure 7
<p>Experimental results for vertically layered printed samples: engineering compressive stress versus compressive strain.</p> Full article ">Figure 8
<p>Details from <a href="#materials-14-04770-f007" class="html-fig">Figure 7</a>: (<b>a</b>) for sample 2; (<b>b</b>) for sample 3; (<b>c</b>) for sample 4; (<b>d</b>) for sample 5.</p> Full article ">Figure 9
<p>Experimental results for horizontally layered printed samples: engineering compressive stress versus compressive strain.</p> Full article ">Figure 10
<p>Proof of compression elastic range up to 3000 N for vertically layered samples.</p> Full article ">Figure 11
<p>Proof of compression elastic range up to 3000 N for horizontally layered samples.</p> Full article ">Figure 12
<p>Experimental results for vertically layered printed samples: engineering compressive stress versus compressive strain at different temperatures.</p> Full article ">
19 pages, 4269 KiB  
Article
Simulations of the Ultra-Fast Kinetics in Ni-Si-C Ternary Systems under Laser Irradiation
by Salvatore Sanzaro, Corrado Bongiorno, Paolo Badalà, Anna Bassi, Ioannis Deretzis, Marius Enachescu, Giovanni Franco, Giuseppe Fisicaro, Patrizia Vasquez, Alessandra Alberti and Antonino La Magna
Materials 2021, 14(16), 4769; https://doi.org/10.3390/ma14164769 - 23 Aug 2021
Cited by 7 | Viewed by 2405
Abstract
We present a method for the simulation of the kinetic evolution in the sub µs timescale for composite materials containing regions occupied by alloys, compounds, and mixtures belonging to the Ni-Si-C ternary system. Pulsed laser irradiation (pulses of the order of 100 ns) [...] Read more.
We present a method for the simulation of the kinetic evolution in the sub µs timescale for composite materials containing regions occupied by alloys, compounds, and mixtures belonging to the Ni-Si-C ternary system. Pulsed laser irradiation (pulses of the order of 100 ns) promotes this evolution. The simulation approach is formulated in the framework of the phase-field theory and it consists of a system of coupled non-linear partial differential equations (PDEs), which considers as variables the following fields: the laser electro-magnetic field, the temperature, the phase-field and the material (Ni, Si, C, C clusters and Ni-silicides) densities. The model integrates a large set of materials and reaction parameters which could also self-consistently depend on the model variables. A parameter calibration is also proposed, specifically suited for the wavelength of a widely used class of excimer lasers (λ = 308 nm). The model is implemented on a proprietary laser annealing technology computer-aided design (TCAD) tool based on the finite element method (FEM). This integration allows, in principle, numerical solutions in systems of any dimension. Here we discuss the complex simulation trend in the one-dimensional case, considering as a starting state, thin films on 4H-SiC substrates, i.e., a configuration reproducing a technologically relevant case study. Simulations as a function of the laser energy density show an articulated scenario, also induced by the variables’ dependency of the materials’ parameters, for the non-melting, partial-melting and full-melting process conditions. The simulation results are validated by post-process experimental analyses of the microstructure and composition of the irradiated samples. Full article
(This article belongs to the Special Issue SiC Materials and Applications)
►▼ Show Figures

Figure 1

Figure 1
<p>Global maximum temperature (purple line) as a function of the time obtained for <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi mathvariant="italic">dens</mi> </mrow> </msub> <mo>=</mo> <mn>2.2</mn> <mo> </mo> <mi mathvariant="normal">J</mi> <mo>/</mo> <msup> <mrow> <mi>cm</mi> </mrow> <mn>2</mn> </msup> </mrow> </semantics></math> energy density laser process of a Ni (100 nm) + 4H-SiC stack. The power density released by the laser pulse in the Ni layer is shown as green line.</p> Full article ">Figure 2
<p>Total atomic fraction as a function of the position of the three elements of the Ni-Si-C ternary system (Ni blue line, Si red line, C green line) after 100 ns (panel (<b>a</b>)) and at the end (panel (<b>b</b>)) of the simulated irradiation at <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> <mo>=</mo> <mn>2.2</mn> <mo> </mo> <msup> <mrow> <mrow> <mi mathvariant="normal">J</mi> <mo>/</mo> <mi>cm</mi> </mrow> </mrow> <mn>2</mn> </msup> </mrow> </semantics></math> energy density.</p> Full article ">Figure 3
<p>Simulated local density (left axis scale) at the end of the process of the Ni<sub>3</sub>Si (dark yellow lines), Ni<sub>5</sub>Si<sub>2</sub> class (dark green lines) and Ni<sub>2</sub>Si (blue lines) for Laser Annealing processes with fluences <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> <mo>=</mo> <mn>2.2</mn> <mo> </mo> <msup> <mrow> <mrow> <mi mathvariant="normal">J</mi> <mo>/</mo> <mi>cm</mi> </mrow> </mrow> <mn>2</mn> </msup> </mrow> </semantics></math>. Simulated local maximum temperature <math display="inline"><semantics> <mrow> <msup> <mi>T</mi> <mrow> <mi>m</mi> <mi>a</mi> <mi>x</mi> </mrow> </msup> <mfenced> <mi>x</mi> </mfenced> </mrow> </semantics></math> (right axis scale), achieved in the different positions of the irradiated structure.</p> Full article ">Figure 4
<p>Phase-field (black line), temperature (dark red line) and Ni total atomic fraction (blue line) as a function of the position after 220 ns of the simulated irradiation at <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> <mo>=</mo> <mn>2.5</mn> <mo> </mo> <msup> <mrow> <mrow> <mi mathvariant="normal">J</mi> <mo>/</mo> <mi>cm</mi> </mrow> </mrow> <mn>2</mn> </msup> </mrow> </semantics></math> energy density. We note the phase and atomic fraction have the same range of variation.</p> Full article ">Figure 5
<p>Phase-field (black line and left axis) and simulated local density (right axis scale) of the Ni<sub>3</sub>Si (dark yellow lines), Ni<sub>5</sub>Si<sub>2</sub> class (dark green lines) and Ni<sub>2</sub>Si (blue lines) for a laser annealing process with fluence <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> <mo>=</mo> <mn>2.5</mn> <mo> </mo> <msup> <mrow> <mrow> <mi mathvariant="normal">J</mi> <mo>/</mo> <mi>cm</mi> </mrow> </mrow> <mn>2</mn> </msup> </mrow> </semantics></math>. Snapshots (<b>a</b>–<b>d</b>) are taken at <span class="html-italic">t</span> = 120, 160, 200, 350 ns.</p> Full article ">Figure 6
<p>Phase-field (black line and left axis) and simulated local density (right axis scale) of the Ni<sub>3</sub>Si (dark yellow lines), Ni<sub>5</sub>Si<sub>2</sub> class (dark green lines) and Ni<sub>2</sub>Si (blue lines) at the end of a laser annealing process with fluence <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> <mo>=</mo> <mn>2.5</mn> <mo> </mo> <msup> <mrow> <mrow> <mi mathvariant="normal">J</mi> <mo>/</mo> <mi>cm</mi> </mrow> </mrow> <mn>2</mn> </msup> </mrow> </semantics></math>.</p> Full article ">Figure 7
<p>Maximum melting extension <math display="inline"><semantics> <mrow> <msub> <mi>D</mi> <mrow> <mi>m</mi> <mi>a</mi> <mi>x</mi> </mrow> </msub> <mfenced> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> </mrow> </mfenced> </mrow> </semantics></math> (melt depth for fluence <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> <mo>&gt;</mo> <mn>2.8</mn> <mo> </mo> <msup> <mrow> <mrow> <mi mathvariant="normal">J</mi> <mo>/</mo> <mi>cm</mi> </mrow> </mrow> <mn>2</mn> </msup> </mrow> </semantics></math>) as a function of the fluence for an irradiated Ni-4HSiC stack. The different regimes (non-melting, partial melting at the Ni-SiC interface, partial melting at the surface, full melting) are indicated by means of the colored areas.</p> Full article ">Figure 8
<p>Phase-field (black line and left axis) and simulated local density (right axis scale) of the Ni<sub>3</sub>Si (dark yellow lines), Ni<sub>5</sub>Si<sub>2</sub> class (dark green lines) and Ni<sub>2</sub>Si (blue lines) for a laser annealing process with fluence <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> <mo>=</mo> <mn>3.2</mn> <mo> </mo> <mi mathvariant="normal">J</mi> <mo>/</mo> <msup> <mrow> <mi>cm</mi> </mrow> <mn>2</mn> </msup> </mrow> </semantics></math>. Snapshots (<b>a</b>–<b>d</b>) are taken at <span class="html-italic">t</span> = 120, 160, 200, 350 ns.</p> Full article ">Figure 9
<p>Phase-field (black line and left axis) and simulated local density (right axis scale) of the Ni<sub>3</sub>Si (dark yellow lines), Ni<sub>5</sub>Si<sub>2</sub> class (dark green lines) and Ni<sub>2</sub>Si (blue lines) at the end of a laser annealing process with fluence <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> <mo>=</mo> <mn>3.2</mn> <mo> </mo> <mi mathvariant="normal">J</mi> <mo>/</mo> <mi mathvariant="normal">c</mi> <msup> <mi mathvariant="normal">m</mi> <mn mathvariant="normal">2</mn> </msup> </mrow> </semantics></math>.</p> Full article ">Figure 10
<p>Extension of the Ni-rich layer M<sub>Ni-rich</sub>(<math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> </mrow> </semantics></math>) (green line and square) and the extension of the silicide layer M<sub>Silicide</sub>(<math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> </mrow> </semantics></math>) (purple line and circles) as a function of <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> </mrow> </semantics></math>. Experimental values from [<a href="#B9-materials-14-04769" class="html-bibr">9</a>], of the silicide layers 18, 36 and 62 nm are also reported as black crosses for the 2.4, 3.2 and 3.8 J/cm<sup>2</sup> cases, respectively.</p> Full article ">Figure 11
<p>Carbon cluster density ratio as a function of the fluence of the <math display="inline"><semantics> <mrow> <msub> <mi>E</mi> <mrow> <mi>d</mi> <mi>e</mi> <mi>n</mi> <mi>s</mi> </mrow> </msub> </mrow> </semantics></math>.</p> Full article ">
11 pages, 10549 KiB  
Article
Direct One-Step Growth of Bimetallic Ni2Mo3N on Ni Foam as an Efficient Oxygen Evolution Electrocatalyst
by Sang Heon Park, Soon Hyung Kang and Duck Hyun Youn
Materials 2021, 14(16), 4768; https://doi.org/10.3390/ma14164768 - 23 Aug 2021
Cited by 16 | Viewed by 3425
Abstract
A simple and economical synthetic route for direct one-step growth of bimetallic Ni2Mo3N nanoparticles on Ni foam substrate (Ni2Mo3N/NF) and its catalytic performance during an oxygen evolution reaction (OER) are reported. The Ni2Mo [...] Read more.
A simple and economical synthetic route for direct one-step growth of bimetallic Ni2Mo3N nanoparticles on Ni foam substrate (Ni2Mo3N/NF) and its catalytic performance during an oxygen evolution reaction (OER) are reported. The Ni2Mo3N/NF catalyst was obtained by annealing a mixture of a Mo precursor, Ni foam, and urea at 600 °C under N2 flow using one-pot synthesis. Moreover, the Ni2Mo3N/NF exhibited high OER activity with low overpotential values (336.38 mV at 50 mA cm−2 and 392.49 mV at 100 mA cm−2) and good stability for 5 h in Fe-purified alkaline electrolyte. The Ni2Mo3N nanoparticle surfaces converted into amorphous surface oxide species during the OER, which might be attributed to the OER activity. Full article
(This article belongs to the Special Issue Advances in Nanostructured Catalysts)
►▼ Show Figures

Figure 1

Figure 1
<p>(<b>a</b>) SEM image of Ni<sub>2</sub>Mo<sub>3</sub>N/NF and SEM-EDS elemental mapping images (scale bar = 300 μm). (<b>b</b>,<b>c</b>) TEM images of Ni<sub>2</sub>Mo<sub>3</sub>N/NF.</p> Full article ">Figure 2
<p>(<b>a</b>) XRD patterns of Ni<sub>2</sub>Mo<sub>3</sub>N/NF. XPS spectra of Ni<sub>2</sub>Mo<sub>3</sub>N/NF for (<b>b</b>) Ni 2p, (<b>c</b>) Mo 3d, and (<b>d</b>) N 1 s.</p> Full article ">Figure 3
<p>Electrochemical characterization of the prepared catalysts. (<b>a</b>) Polarization curves (1.0 M KOH solution), (<b>b</b>) bar graphs showing overpotentials at 50 mA cm<sup>−2</sup> and current densities at an overpotential of 400 mV, (<b>c</b>) Nyquist plots, and (<b>d</b>) durability measurement.</p> Full article ">Figure 4
<p>(<b>a</b>) XRD patterns of Ni<sub>2</sub>Mo<sub>3</sub>N/NF (fresh and after the durability test). XPS spectra of Ni<sub>2</sub>Mo<sub>3</sub>N/NF after the durability test. (<b>b</b>) Ni 2p, (<b>c</b>) Mo 3d, and (<b>d</b>) N 1s.</p> Full article ">Figure 5
<p>Ni<sub>2</sub>Mo<sub>3</sub>N/NF characterization results after the durability tests: (<b>a</b>) SEM images; (<b>b</b>) SEM-EDS elemental mapping images (scale bar = 300 μm); (<b>c</b>,<b>d</b>) TEM images.</p> Full article ">
18 pages, 2422 KiB  
Review
Lateral Formwork Pressure for Self-Compacting Concrete—A Review of Prediction Models and Monitoring Technologies
by Yaser Gamil, Jonny Nilimaa, Mats Emborg and Andrzej Cwirzen
Materials 2021, 14(16), 4767; https://doi.org/10.3390/ma14164767 - 23 Aug 2021
Cited by 13 | Viewed by 3863
Abstract
The maximum amount of lateral formwork pressure exerted by self-compacting concrete is essential to design a technically correct, cost-effective, safe, and robust formwork. A common practice of designing formwork is primarily based on using the hydrostatic pressure. However, several studies have proven that [...] Read more.
The maximum amount of lateral formwork pressure exerted by self-compacting concrete is essential to design a technically correct, cost-effective, safe, and robust formwork. A common practice of designing formwork is primarily based on using the hydrostatic pressure. However, several studies have proven that the maximum pressure is lower, thus potentially enabling a reduction in the cost of formwork by, for example, optimizing the casting rate. This article reviews the current knowledge regarding formwork pressure, parameters affecting the maximum pressure, prediction models, monitoring technologies and test setups. The currently used pressure predicting models require further improvement to consider several pressures influencing parameters, including parameters related to fresh and mature material properties, mix design and casting methods. This study found that the maximum pressure is significantly affected by the concretes’ structural build-up at rest, which depends on concrete rheology, temperature, hydration rate and setting time. The review indicates a need for more in-depth studies. Full article
(This article belongs to the Special Issue Concrete and Construction Materials)
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<p>Correlation between slump flow and formwork pressure for different mix designs.</p> Full article ">Figure 2
<p>Lateral pressure calculation method [<a href="#B46-materials-14-04767" class="html-bibr">46</a>] (reprinted with permission from © ACI).</p> Full article ">Figure 3
<p>Maximum lateral pressure for concretes with a final setting time of 10 h [<a href="#B75-materials-14-04767" class="html-bibr">75</a>] (reprinted with permission from © ACI).</p> Full article ">Figure 4
<p>Normalized maximum pressure for high flowable concrete [<a href="#B42-materials-14-04767" class="html-bibr">42</a>] (reprinted with permission from © John Wiley and Sons).</p> Full article ">Figure 5
<p>Laboratory setups for lateral pressure measurement (reprinted from [<a href="#B74-materials-14-04767" class="html-bibr">74</a>], with permission © John Wiley and Sons).</p> Full article ">Figure 6
<p>Full-scale experimental setup for pressure measurement: (<b>a</b>) location of sensors; (<b>b</b>) mounted pressure sensor; (<b>c</b>) strain gauge; (<b>d</b>) load cells [<a href="#B80-materials-14-04767" class="html-bibr">80</a>] (reprinted with permission from © ACI).</p> Full article ">Figure 7
<p>Laboratory setup with steel formwork (<b>a</b>) pressure sensor; (<b>b</b>) welded pressure diaphragm (<b>c</b>) formwork setup [<a href="#B40-materials-14-04767" class="html-bibr">40</a>] (reprinted with permission from © 2021 Elsevier Ltd.).</p> Full article ">Figure 8
<p>Laboratory setup. (<b>a</b>) Flush diaphragm sensor. (<b>b</b>) Laboratory setup for pressure monitoring [<a href="#B11-materials-14-04767" class="html-bibr">11</a>] (reprinted with permission from © 2021 Elsevier Ltd.).</p> Full article ">
21 pages, 8763 KiB  
Article
Modeling the Average and Instantaneous Friction Coefficient of a Disc Brake on the Basis of Bench Tests
by Wojciech Sawczuk, Armando Miguel Rilo Cañás, Dariusz Ulbrich and Jakub Kowalczyk
Materials 2021, 14(16), 4766; https://doi.org/10.3390/ma14164766 - 23 Aug 2021
Cited by 11 | Viewed by 3143
Abstract
This article presents the results of tests conducted on the average and instantaneous friction coefficients of railway vehicle disc brakes. The tests were carried out independently of various states of wear on the friction linings and the brake disc. The requirements of the [...] Read more.
This article presents the results of tests conducted on the average and instantaneous friction coefficients of railway vehicle disc brakes. The tests were carried out independently of various states of wear on the friction linings and the brake disc. The requirements of the International Union of Railways (UIC) regarding the approval of brake linings for use were taken into account. Based on many years of research using a brake bench to test railway disc brakes, the authors developed multiple regression models for the average friction coefficient and fluctuations (tolerances) in the instantaneous friction coefficient and achieved 870 results. The models proposed three types of variables: the input braking parameters (speed, pressure, and mass to be braked), operational parameters (the wear on the friction linings and the brake disc), and design parameters (perforations in the form of holes on the disc surface). The above two models were validated on the basis of 384 brakes, and in subsequent stages a further evaluation was performed. The coefficients were determined to be, respectively, 0.99 for the model of the average friction coefficient and 0.71 for the model of tolerance (fluctuations) of the instantaneous friction coefficient. Full article
(This article belongs to the Special Issue Friction and Wear of Materials Surfaces)
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<p>Brake bench for testing railway disc brakes: (<b>a</b>) drive part of the brake stand with rotating masses, (<b>b</b>) brake disc type 610 × 110 mounted on a brake bench.</p> Full article ">Figure 2
<p>View of the brake disc: (<b>a</b>) diagram of the Archimedes spiral with one turn limited by the functions of a circle, (<b>b</b>) disc on the brake bench.</p> Full article ">Figure 3
<p>View of the friction linings used during bench tests: (<b>a</b>) view with visible expansion grooves, (<b>b</b>) side view with visible lining thickness.</p> Full article ">Figure 4
<p>Characteristics of the instantaneous coefficient of friction μ<sub>a</sub> on the braking initiation speed during braking with a pressure of 44 kN and mass per disc of 7.5 t: (<b>a</b>) for a new disc, (<b>b</b>) for a worn disc.</p> Full article ">Figure 5
<p>Characteristics of the average friction coefficient μ<sub>m</sub> on the braking start speed with a pressure of 44 kN and mass per disc of 7.5 t: (<b>a</b>) for a new disc, (<b>b</b>) for a worn disc.</p> Full article ">Figure 6
<p>Characteristics of the instantaneous coefficient of friction μ<sub>a</sub> on the braking initiation speed with a pressure of 36 kN, and mass per disc of 4.7 t: (<b>a</b>) for a classic disc, (<b>b</b>) for a disc perforated on the friction surface.</p> Full article ">Figure 7
<p>Characteristics of the average coefficient of friction μ<sub>m</sub> on the braking initiation speed with a pressure of 36 kN and mass per disc of 4.7 t: (<b>a</b>) for a classic disc, (<b>b</b>) for a disc perforated on the friction surface.</p> Full article ">Figure 8
<p>Characteristics of the friction coefficient (average and its spread) on the braking initiation speed with a pressure of 36 kN and mass per disc of 4.7 t: (<b>a</b>) for a classic disc, (<b>b</b>) for a disc perforated on the friction surface.</p> Full article ">Figure 9
<p>Temperature distribution on a disc at the moment braking ended from a speed of 200 km/h: (<b>a</b>) classic (smooth), (<b>b</b>) perforated on the friction surface.</p> Full article ">Figure 10
<p>Distribution of the absolute value of the Pearson correlation coefficient for the variables of the regression model; (<b>a</b>) coefficient of friction and (<b>b</b>) tolerance (changes in the instantaneous coefficient of friction); red indicates an increase in the variable that caused a decrease in the value of µ<sub>m</sub> or T<sub>μa</sub>; green indicates an increase in µ<sub>m</sub> or T<sub>μa</sub> with an increase in the value of the variable.</p> Full article ">Figure 11
<p>The μ<sub>m</sub> changes from the tests with the multiple regression model during braking with N = 44 kN, M = 7.5 t for: (<b>a</b>) new disc, (<b>b</b>) worn disc.</p> Full article ">Figure 12
<p>The changes of μ<sub>m</sub> from the tests with multiple regression model during braking on the G<sub>1</sub> lining with N = 25 kN, M = 5.7 t for: (<b>a</b>) smooth disc, (<b>b</b>) perforated disc.</p> Full article ">Figure 13
<p>The changes in μ<sub>m</sub> from the tests with multiple regression model during braking on the G<sub>2</sub> lining with N = 25 kN, M = 5.7 t for (<b>a</b>) smooth disc and (<b>b</b>) perforated disc.</p> Full article ">Figure 14
<p>The changes of μ<sub>m</sub> from the tests with multiple regression model during braking on the G<sub>3</sub> lining with N = 25 kN, M = 5.7 t for: (<b>a</b>) smooth disc, (<b>b</b>) perforated disc.</p> Full article ">Figure 15
<p>Thermal image of the brake disc, type: (<b>a</b>) 590 × 110 (new disc), (<b>b</b>) 640 × 110 (worn disc).</p> Full article ">Figure 16
<p>The changes of μ<sub>m</sub> from the tests with multiple regression model during braking on a new disc type 590 × 110 z: (<b>a</b>) N = 25 kN and M = 5.7 t, (<b>b</b>) N = 36 kN and M = 5.7 t.</p> Full article ">Figure 17
<p>The changes of μ<sub>m</sub> from the tests with multiple regression model during braking on a worn disc type 640 × 110 z: (<b>a</b>) N = 28 kN and M = 6.7 t, (<b>b</b>) N = 40 kN and M = 6.7 t.</p> Full article ">
16 pages, 3550 KiB  
Article
A Green, Simple and Facile Way to Synthesize Silver Nanoparticles Using Soluble Starch. pH Studies and Antimicrobial Applications
by Bogdan Pascu, Adina Negrea, Mihaela Ciopec, Narcis Duteanu, Petru Negrea, Nicoleta Sorina Nemeş, Corina Seiman, Eleonora Marian and Otilia Micle
Materials 2021, 14(16), 4765; https://doi.org/10.3390/ma14164765 - 23 Aug 2021
Cited by 11 | Viewed by 3541
Abstract
Along with the progress of nanoscience and nanotechnology came the means to synthesize nanometric scale materials. While changing their physical and chemical properties, they implicitly changed their application area. The aim of this paper was the synthesis of colloidal silver nanoparticles (Ag-NPs by [...] Read more.
Along with the progress of nanoscience and nanotechnology came the means to synthesize nanometric scale materials. While changing their physical and chemical properties, they implicitly changed their application area. The aim of this paper was the synthesis of colloidal silver nanoparticles (Ag-NPs by ultrasonic disruption), using soluble starch as a reducing agent and further as a stabilizing agent for produced Ag-NPs. In this context, an important parameter for Ag-NPs preparation is the pH, which can determine the particle size and stability. The physical-chemical behavior of the synthesized Ag-NPs (shape, size, dispersion, electric charge) is strongly influenced by the pH value (experiment being conducted for pH values in the range between 8 and 13). The presence of a peak located at 412 nm into the UV-VIS spectra demonstrates the presence of silver nano-spheres into the produced material. In UV/VIS spectra, we observed a specific peak for yellow silver nano-spheres located at 412 nm. Samples characterization was performed by scanning electron microscopy, SEM, energy-dispersive X-ray spectroscopy, EDX, Fourier-transform infrared spectroscopy, and FT-IR. For all Ag-NP samples, we determined the zeta and observed that the Ag-NP particles obtained at higher pH and have better stability. Due to the intrinsic therapeutic properties and broad antimicrobial spectrum, silver nanoparticles have opened new horizons and new approaches for the control of different types of infections and wound healing abilities. In this context, the present study also aims to confirm the antimicrobial effect of prepared Ag-NPs against several bacterial strains (indicator and clinically isolated strains). In this way, it was confirmed that the antimicrobial activity of synthesized Ag-NPs was good against Staphylococcus aureus (ATCC 25923 and S. aureus MSSA) and Escherichia coli (ATTC 25922 and clinically isolated strain). Based on this observation, we conclude that the prepared Ag-NPs can represent an alternative or auxiliary material used for controlling important nosocomial pathogens. The fungal reference strain Candida albicans was more sensitive at Ag-NPs actions (zone of inhibition = 20 mm) compared with the clinically isolated strain (zone of inhibition = 10 mm), which emphasizes the greater resistance of fungal strains at antimicrobial agent’s action. Full article
(This article belongs to the Special Issue Adsorption and Desorption Behavior for Rare Earth Metal Ions)
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<p>UV-VIS spectra recorded for prepared Ag-NPs.</p> Full article ">Figure 2
<p>Particle size distribution of synthesized Ag-NPs vs. pH.</p> Full article ">Figure 3
<p>Scanning electron micrographs recorded for samples prepared at pH values between 8 and 13.</p> Full article ">Figure 3 Cont.
<p>Scanning electron micrographs recorded for samples prepared at pH values between 8 and 13.</p> Full article ">Figure 3 Cont.
<p>Scanning electron micrographs recorded for samples prepared at pH values between 8 and 13.</p> Full article ">Figure 4
<p>Energy-dispersive X-ray spectra recorded for samples prepared at pH values between 8 and 13.</p> Full article ">Figure 5
<p>FT-IR spectra were recorded for Ag-NPs and Starch-NaOH.</p> Full article ">
22 pages, 11835 KiB  
Article
Using the Stochastic Gradient Descent Optimization Algorithm on Estimating of Reactivity Ratios
by Iosif Sorin Fazakas-Anca, Arina Modrea and Sorin Vlase
Materials 2021, 14(16), 4764; https://doi.org/10.3390/ma14164764 - 23 Aug 2021
Cited by 3 | Viewed by 2588
Abstract
This paper describes an improved method of calculating reactivity ratios by applying the neuronal networks optimization algorithm, named gradient descent. The presented method is integral and has been compared to the following existing methods: Fineman–Ross, Tidwell–Mortimer, Kelen–Tüdös, extended Kelen–Tüdös and Error in Variable [...] Read more.
This paper describes an improved method of calculating reactivity ratios by applying the neuronal networks optimization algorithm, named gradient descent. The presented method is integral and has been compared to the following existing methods: Fineman–Ross, Tidwell–Mortimer, Kelen–Tüdös, extended Kelen–Tüdös and Error in Variable Methods. A comparison of the reactivity ratios that obtained different levels of conversions was made based on the Fisher criterion. The new calculation method for reactivity ratios shows better results than these other methods. Full article
(This article belongs to the Special Issue Collection of Papers in Material Science from Romania)
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<p>Evolution of reactivity ratio searching (<b>A</b>) and error for last 60 searching steps (<b>B</b>) with GD method for <span class="html-italic">r</span><sub>1</sub> = 0.05, <span class="html-italic">r</span><sub>2</sub> = 0.5 and <span class="html-italic">P<sub>n</sub></span> ∊ (1–10%).</p> Full article ">Figure 2
<p>Evolution of reactivity ratio searching (<b>A</b>) and error for last 60 searching steps (<b>B</b>) with GD method for <span class="html-italic">r</span><sub>1</sub> = 0.8, <span class="html-italic">r</span><sub>2</sub> = 1.8 and <span class="html-italic">P<sub>n</sub></span> ∊ (1–10%).</p> Full article ">Figure 3
<p>Evolution of reactivity ratio searching (<b>A</b>) and error for last 60 searching steps (<b>B</b>) with GD method for <span class="html-italic">r</span><sub>1</sub> = 0.4, <span class="html-italic">r</span><sub>2</sub> = 0.8 and <span class="html-italic">P<sub>n</sub></span> ∊ (1–10%).</p> Full article ">Figure 4
<p>Evolution of reactivity ratio searching (<b>A</b>) and error for last 60 searching steps (<b>B</b>) with GD method for <span class="html-italic">r</span><sub>1</sub> = 0.05, <span class="html-italic">r</span><sub>2</sub> = 0.5 and <span class="html-italic">P<sub>n</sub></span> ∊ (10–25%).</p> Full article ">Figure 5
<p>Evolution of reactivity ratio searching (<b>A</b>) and error for last 60 searching steps (<b>B</b>) with GD method for <span class="html-italic">r</span><sub>1</sub> = 0.8, <span class="html-italic">r</span><sub>2</sub> = 1.8 and <span class="html-italic">P<sub>n</sub></span> ∊ (10–25%).</p> Full article ">Figure 6
<p>Evolution of reactivity ratio searching (<b>A</b>) and error for last 60 searching steps (<b>B</b>) with GD method for <span class="html-italic">r</span><sub>1</sub> = 0.4, <span class="html-italic">r</span><sub>2</sub> = 0.8 and <span class="html-italic">P<sub>n</sub></span> ∊ (10–25%).</p> Full article ">Figure 7
<p>Evolution of reactivity ratio searching (<b>A</b>) and error for last 60 searching steps (<b>B</b>) with GD method for <span class="html-italic">r</span><sub>1</sub> = 0.05, <span class="html-italic">r</span><sub>2</sub> = 0.5 and <span class="html-italic">P<sub>n</sub></span> ∊ (30–55%).</p> Full article ">Figure 8
<p>Evolution of reactivity ratio searching (<b>A</b>) and error for last 60 searching steps (<b>B</b>) with GD method for <span class="html-italic">r</span><sub>1</sub> = 0.8, <span class="html-italic">r</span><sub>2</sub> = 1.8 and <span class="html-italic">P<sub>n</sub></span> ∊ (30–55%).</p> Full article ">Figure 9
<p>Evolution of reactivity ratio searching (<b>A</b>) and error for last 60 searching steps (<b>B</b>) with GD method for <span class="html-italic">r</span><sub>1</sub> = 0.4, <span class="html-italic">r</span><sub>2</sub> = 0.8 and <span class="html-italic">P<sub>n</sub></span> ∊ (30–55%).</p> Full article ">Figure 10
<p>The JCR for following initial conditions <span class="html-italic">r</span><sub>1</sub> = 0.05, <span class="html-italic">r</span><sub>2</sub> = 0.50 <span class="html-italic">P<sub>n</sub></span> ∊ (1–10%).</p> Full article ">Figure 11
<p>The JCR for following initial conditions <span class="html-italic">r</span><sub>1</sub> = 0.8, <span class="html-italic">r</span><sub>2</sub> = 1.8 <span class="html-italic">P<sub>n</sub></span> ∊ (1–10%).</p> Full article ">Figure 12
<p>The JCR for following initial conditions <span class="html-italic">r</span><sub>1</sub> = 0.4, <span class="html-italic">r</span><sub>2</sub> = 0.8 <span class="html-italic">P<sub>n</sub></span> ∊ (1–10%).</p> Full article ">Figure 13
<p>The JCR for initial conditions <span class="html-italic">r</span><sub>1</sub> = 0.05, <span class="html-italic">r</span><sub>2</sub> = 0.5, <span class="html-italic">P<sub>n</sub></span> ∊ (10–25%).</p> Full article ">Figure 14
<p>Details of JCR for GD given in <a href="#materials-14-04764-f013" class="html-fig">Figure 13</a>.</p> Full article ">Figure 15
<p>The JCR for initial conditions <span class="html-italic">r</span><sub>1</sub> = 0.80, <span class="html-italic">r</span><sub>2</sub> = 1.80, <span class="html-italic">P<sub>n</sub></span> ∊ (10–25%).</p> Full article ">Figure 16
<p>Details of JCR for GD, EVM and e-KT given in <a href="#materials-14-04764-f015" class="html-fig">Figure 15</a>.</p> Full article ">Figure 17
<p>The JCR for initial conditions <span class="html-italic">r</span><sub>1</sub> = 0.4, <span class="html-italic">r</span><sub>2</sub> = 0.8, <span class="html-italic">P<sub>n</sub></span> ∊ (10–25%).</p> Full article ">Figure 18
<p>Details of JCR for GD given in <a href="#materials-14-04764-f017" class="html-fig">Figure 17</a>.</p> Full article ">Figure 19
<p>The JCR domain for initial data <span class="html-italic">r</span><sub>1</sub> = 0.05, <span class="html-italic">r</span><sub>2</sub> = 0.5 <span class="html-italic">P<sub>n</sub></span> ∊ (30–55%).</p> Full article ">Figure 20
<p>The JCR domain for initial data <span class="html-italic">r</span><sub>1</sub> = 0.8, <span class="html-italic">r</span><sub>2</sub> = 1.8 <span class="html-italic">P<sub>n</sub></span> ∊ (30–55%).</p> Full article ">Figure 21
<p>Details of JCR for GD given in <a href="#materials-14-04764-f020" class="html-fig">Figure 20</a>.</p> Full article ">Figure 22
<p>The JCR domain for initial data <span class="html-italic">r</span><sub>1</sub> = 0.40, <span class="html-italic">r</span><sub>2</sub> = 0.80 <span class="html-italic">P<sub>n</sub></span> ∊ (30–55%).</p> Full article ">Figure 23
<p>The JCR domain for copolymerization of n-butyl methacrylate with n-butyl acrylate [<a href="#B23-materials-14-04764" class="html-bibr">23</a>].</p> Full article ">Figure 24
<p>The JCR domain for copolymerization of 2-isopropenyl-2-oxazoline with methyl methacrylate [<a href="#B24-materials-14-04764" class="html-bibr">24</a>].</p> Full article ">Figure 25
<p>The JCR domain for copolymerization of N-(4-carboxyphenyl) maleimide with N-vinyl-2-pyrrolidone [<a href="#B25-materials-14-04764" class="html-bibr">25</a>].</p> Full article ">
16 pages, 4660 KiB  
Article
Treatment of Rhenium-Containing Effluents Using Environmentally Friendly Sorbent, Saccharomyces cerevisiae Biomass
by Inga Zinicovscaia, Nikita Yushin, Dmitrii Grozdov, Konstantin Vergel, Pavel Nekhoroshkov and Elena Rodlovskaya
Materials 2021, 14(16), 4763; https://doi.org/10.3390/ma14164763 - 23 Aug 2021
Cited by 10 | Viewed by 2157
Abstract
Yeast Saccharomyces cerevisiae biomass was applied for rhenium and accompanying elements (copper and molybdenum) removal from single- and multi-component systems (Re, Re-Mo, Re-Cu, and Re-Mo-Cu). Yeast biomass was characterized using X-ray diffraction, scanning electron microscopy, and Fourier transform infrared spectroscopy. The effects of [...] Read more.
Yeast Saccharomyces cerevisiae biomass was applied for rhenium and accompanying elements (copper and molybdenum) removal from single- and multi-component systems (Re, Re-Mo, Re-Cu, and Re-Mo-Cu). Yeast biomass was characterized using X-ray diffraction, scanning electron microscopy, and Fourier transform infrared spectroscopy. The effects of biosorption experimental parameters such as solution pH (2.0–6.0), rhenium concentration (10–100 mg/L), time of interaction (5–120 min), and temperature (20–50 °C) have been discussed in detail. Maximum removal of rhenium (75–84%) and molybdenum (85%) was attained at pH 2.0, while pH 3.0–5.0 was more favorable for copper ions removal (53–68%). The Langmuir, Freundlich, and Temkin isotherm models were used to describe the equilibrium sorption of rhenium on yeast biomass. Langmuir isotherm shows the maximum yeast adsorption capacities toward rhenium ions ranged between 7.7 and 33 mg/g. Several kinetic models (pseudo-first-order, pseudo-second-order, and Elovich) were applied to define the best correlation for each metal. Biosorption of metal ions was well-fitted by Elovich and pseudo-first-order models. The negative free energy reflected the feasibility and spontaneous nature of the biosorption process. Saccharomyces cerevisiae biomass can be considered as a perspective biosorbent for metal removal. Full article
(This article belongs to the Special Issue Low-Cost Water Treatment - New Materials and New Approaches)
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<p>SEM image (<b>left</b>) and XRD pattern (<b>right</b>) of S. cerevisiae yeast biomass before biosorption experiments.</p> Full article ">Figure 2
<p>Effect of initial solution pH on the adsorption of metal ions present in analyzed systems: (<b>a</b>) Re, (<b>b</b>) Re-Cu, (<b>c</b>) Re-Mo, and (<b>d</b>) Re-Mo-Cu.</p> Full article ">Figure 3
<p>Effect of temperature on metal ions removal from analyzed systems: (<b>a</b>) Re, (<b>b</b>) Re-Cu, (<b>c</b>) Re-Mo, and (<b>d</b>) Re-Mo-Cu.</p> Full article ">Figure 4
<p>Kinetics of rhenium biosorption on yeast biomass.</p> Full article ">Figure 5
<p>Kinetics of rhenium and copper biosorption on yeast biomass.</p> Full article ">Figure 6
<p>Kinetics of rhenium and copper biosorption on yeast biomass.</p> Full article ">Figure 7
<p>Kinetics of rhenium and copper biosorption on yeast biomass.</p> Full article ">Figure 8
<p>The adsorption isotherms and experimental data for rhenium ion sorption on yeast biomass. (<b>a</b>) Re, (<b>b</b>) Re-Cu, (<b>c</b>) Re-Mo, and (<b>d</b>) Re-Mo-Cu.</p> Full article ">Figure 9
<p>FTIR spectra of control and metal-loaded yeast biomass.</p> Full article ">
22 pages, 3938 KiB  
Article
The Design and Development of Recycled Concretes in a Circular Economy Using Mixed Construction and Demolition Waste
by Marcos Díaz González, Pablo Plaza Caballero, David Blanco Fernández, Manuel Miguel Jordán Vidal, Isabel Fuencisla Sáez del Bosque and César Medina Martínez
Materials 2021, 14(16), 4762; https://doi.org/10.3390/ma14164762 - 23 Aug 2021
Cited by 17 | Viewed by 3452
Abstract
This research study analysed the effect of adding fine—fMRA (0.25% and 50%)—and coarse—cMRA (0%, 25% and 50%)—mixed recycled aggregate both individually and simultaneously in the development of sustainable recycled concretes that require a lower consumption of natural resources. For this purpose, we first [...] Read more.
This research study analysed the effect of adding fine—fMRA (0.25% and 50%)—and coarse—cMRA (0%, 25% and 50%)—mixed recycled aggregate both individually and simultaneously in the development of sustainable recycled concretes that require a lower consumption of natural resources. For this purpose, we first conducted a physical and mechanical characterisation of the new recycled raw materials and then analysed the effect of its addition on fresh and hardened new concretes. The results highlight that the addition of fMRA and/or cMRA does not cause a loss of workability in the new concrete but does increase the amount of entrained air. Regarding compressive strength, we observed that fMRA and/or cMRA cause a maximum increase of +12.4% compared with conventional concrete. Tensile strength increases with the addition of fMRA (between 8.7% and 5.5%) and decreases with the use of either cMRA or fMRA + cMRA (between 4.6% and 7%). The addition of fMRA mitigates the adverse effect that using cMRA has on tensile strength. Regarding watertightness, all designed concretes have a structure that is impermeable to water. Lastly, the results show the feasibility of using these concretes to design elements with a characteristic strength of 25 MPa and that the optimal percentage of fMRA replacement is 25%. Full article
(This article belongs to the Topic Sustainability in Buildings: New Trends in the Management of Construction and Demolition Waste)
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<p>View of the aggregates used to manufacture concretes. Legend: (<b>a</b>) View of natural aggregate; (<b>b</b>) View of natural aggregate after grinding; (<b>c</b>) View of recycled aggregates and (<b>d</b>) View of recycled aggregates after grinding.</p> Full article ">Figure 2
<p>Granulometric distribution of aggregates (NA and MRA).</p> Full article ">Figure 3
<p>Flow diagram corresponding to the manufacturing process and tests carried out on the studied mixtures.</p> Full article ">Figure 4
<p>Concrete slump test: (<b>a</b>) Concrete with 100% NA (M1) and (<b>b</b>) Concrete with 50% fMRA and cMRA (M9).</p> Full article ">Figure 5
<p>Connection between the percentage of mixed recycled sand and entrained air.</p> Full article ">Figure 6
<p>Compressive strength variation of recycled concretes (M2–M9) compared with conventional concrete (M1).</p> Full article ">Figure 7
<p>Type of failure of the concretes: (<b>a</b>) M1; (<b>b</b>) M3; (<b>c</b>) M6; and (<b>d</b>) M9.</p> Full article ">Figure 8
<p>Appearance of the concretes when performing the tensile test 28 days after producing the mortar: (<b>a</b>) conventional concrete (M1) and (<b>b</b>) concrete with 50% cMRA + 50% fMRA (M9).</p> Full article ">Figure 9
<p>Water penetration of the concretes under pressure. Limits established in EHE-08 and EC-2 (Note. IIIa: marine class—subclass: aerial; IIIb: marine class—subclass: submerged; IIIc: marine class—subclass: tidal and splash zones; Qa: aggressive chemical class—subclass: weak; Qb: aggressive chemical class—subclass: average; Qc: aggressive chemical class—subclass: strong; H: with frost class—subclass: without deicing salts; and F: with frost class—subclass: without deicing salts).</p> Full article ">Figure 10
<p>Penetration fronts of the concretes: (<b>a</b>) M1; (<b>b</b>) M3; (<b>c</b>) M6; and (<b>d</b>) M9.</p> Full article ">
15 pages, 31217 KiB  
Article
Electrochemically Synthesized Poly(3-hexylthiophene) Nanowires as Photosensitive Neuronal Interfaces
by Szilveszter Gáspár, Tiziana Ravasenga, Raluca-Elena Munteanu, Sorin David, Fabio Benfenati and Elisabetta Colombo
Materials 2021, 14(16), 4761; https://doi.org/10.3390/ma14164761 - 23 Aug 2021
Cited by 1 | Viewed by 2523
Abstract
Poly(3-hexylthiophene) (P3HT) is a hole-conducting polymer that has been intensively used to develop organic optoelectronic devices (e.g., organic solar cells). Recently, P3HT films and nanoparticles have also been used to restore the photosensitivity of retinal neurons. The template-assisted electrochemical synthesis of polymer nanowires [...] Read more.
Poly(3-hexylthiophene) (P3HT) is a hole-conducting polymer that has been intensively used to develop organic optoelectronic devices (e.g., organic solar cells). Recently, P3HT films and nanoparticles have also been used to restore the photosensitivity of retinal neurons. The template-assisted electrochemical synthesis of polymer nanowires advantageously combines polymerization and polymer nanostructuring into one, relatively simple, procedure. However, obtaining P3HT nanowires through this procedure was rarely investigated. Therefore, this study aimed to investigate the template-assisted electrochemical synthesis of P3HT nanowires doped with tetrabutylammonium hexafluorophosphate (TBAHFP) and their biocompatibility with primary neurons. We show that template-assisted electrochemical synthesis can relatively easily turn 3-hexylthiophene (3HT) into longer (e.g., 17 ± 3 µm) or shorter (e.g., 1.5 ± 0.4 µm) P3HT nanowires with an average diameter of 196 ± 55 nm (determined by the used template). The nanowires produce measurable photocurrents following illumination. Finally, we show that primary cortical neurons can be grown onto P3HT nanowires drop-casted on a glass substrate without relevant changes in their viability and electrophysiological properties, indicating that P3HT nanowires obtained by template-assisted electrochemical synthesis represent a promising neuronal interface for photostimulation. Full article
(This article belongs to the Special Issue Advanced Designs of Materials, Devices and Techniques for Biosensing)
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<p>Schematic representation of the template-assisted, electrochemical synthesis of P3HT nanowires. A nanoporous alumina membrane is used as template. After covering the “branched” side of the membrane with Cu by physical vapor deposition, three electrochemical steps are carried out, namely, electrodeposition of Cu, electrodeposition of Au, and electropolymerization of P3HT. Following these three electrochemical steps, the template is disassembled by using two metal etching solutions (one to remove Cu and one to remove Au) and a NaOH solution (to dissolve the alumina). A suspension of P3HT nanowires is obtained. The nanowires are then washed until a neutral pH suspension is obtained (not shown).</p> Full article ">Figure 2
<p>Electrochemical signals recorded during: deposition of Cu into the “branched” pores of the alumina template (<b>A</b>), deposition of Au segments into the cylindrical pores of the alumina template (<b>B</b>), and electropolymerization of P3HT into the cylindrical pores of the alumina template (<b>C</b>). Noise affecting current signals is due to occasionally “stirring” the solution during the electrochemical deposition. See the Materials and Methods section for additional experimental details.</p> Full article ">Figure 3
<p>AFM images of “long” P3HT nanowires (<b>A</b>) and of “short” P3HT nanowires (<b>B</b>).</p> Full article ">Figure 4
<p>Distribution of P3HT nanowire diameters and lengths. The distribution of the nanowire diameters (<b>A</b>) was obtained by analyzing the AFM images of 140 nanowires from eight different batches. The distribution of the lengths of “short” P3HT nanowires (<b>B</b>) was obtained by analyzing the AFM images of 240 nanowires from five different batches. The distribution of the lengths of “long” P3HT nanowires (<b>C</b>) was obtained by analyzing either AFM or optical microscopy images of 170 nanowires from three different batches.</p> Full article ">Figure 5
<p>SEM, EDS and Raman spectroscopy characterization of the nanowires. (<b>A</b>,<b>B</b>) show representative SEM images of “long” and “short” nanowires, respectively. (<b>C</b>) EDS spectra showing the evident K<sub>α</sub> peak of sulfur (S) from P3HT and no other elements apart from the contribution of the glass coverslip substrate, whose spectrum is also presented. (<b>D</b>) Raman spectra acquired at 633 nm (1 s) showing peaks at ~1380 and 1440 cm<sup>−1</sup>, due to C–C intra-ring stretching and C=C symmetric stretching, respectively, both of them characteristic for P3HT.</p> Full article ">Figure 6
<p>Typical absorbance spectra of suspensions containing P3HT nanowires (<b>A</b>) and photocurrents observed with a P3HT nanowires-modified glassy carbon electrode (<b>B</b>). Normalized absorbance spectra are shown for suspensions with either “short” or “long” P3HT nanowires. Photocurrents are shown for the bare glassy carbon electrode (black signal), for the glassy carbon electrode modified by drop-casting with 18 µL of “short” P3HT nanowire suspension (red signal), and for the glassy carbon electrode modified by drop-casting with 36 µL of “short” P3HT nanowire suspension (blue signal). See the Materials and Methods section for additional experimental details.</p> Full article ">Figure 7
<p>Neuronal viability. Representative bright field (#) and epifluorescence images (∆) of neurons grown in the presence of P3HT nanowires (2 or under control conditions (1). P3HT nanowires are black and red in the transmitted and epifluorescence images, respectively, thanks to their optical properties. Cultures were stained with Hoechst-33342 for nuclear visualization (blue) and propidium iodide (PI, red) for cell death quantification (scale bar: 100 µm) (<b>A</b>). Cell viability was evaluated by fluorescence microscopy at DIV7 and DIV14. The number of PI-positive cells was subtracted to the total number of Hoechst-positive cells to obtain the number of living cells. The percentages of live neurons with respect to the total number of Hoechst-positive cells was assessed for each experimental group. Viability values were normalized to the respective average value of the control sample and plotted as Means ± SEM with superimposed individual experimental points. No significant changes in cell viability were observed under the various experimental conditions (unpaired Student’s <span class="html-italic">t</span>-test, <span class="html-italic">n</span> = 18 fields per experimental condition, from two independent neuronal preparations) (<b>B</b>). Representative confocal microscopy image showing primary mouse cortical neurons (grown onto poly(<span class="html-small-caps">l</span>-lysine)-coated glass further modified with endogenously fluorescent P3HT nanowires (red)) immunolabeled for β3 tubulin (green) and with Hoechst-33342 (white) (scale bar: 10 µm) (<b>C</b>).</p> Full article ">Figure 8
<p>Electrophysiological properties of primary neurons grown in the presence of P3HT nanowires. From left to right: resting membrane potential, membrane resistance, membrane capacitance (<b>A</b>). Representative traces of AP firing evoked by 350 pA current injection under both experimental conditions, and representative phase-plane plot analysis (dV/dt vs. V) of AP waveform obtained under control conditions (left) and in the presence of P3HT nanowires (right) (<b>B</b>). Minimal current injection intensity necessary to evoke an AP (rheobase) (<b>C</b>). From left to right: peak, threshold potential (V<sub>th</sub>), and width of AP (<b>D</b>). Instantaneous firing frequency (IFF) at each inter-spike interval (ISI) evoked by the injection of 350 pA depolarizing current (left) and mean firing frequency (MFF) plotted as a function of the injected current (right) (<b>E</b>). All data are shown as Means ± SEM with superimposed individual experimental points. All the analyzed parameters were not significantly different between the two experimental groups; Student’s <span class="html-italic">t</span>-test (<b>A</b>–<b>D</b>) and two-way ANOVA with Bonferroni’s multiple comparison tests (<b>E</b>); <span class="html-italic">n</span> = 9–13 cells, from two independent neuronal preparations).</p> Full article ">Figure 8 Cont.
<p>Electrophysiological properties of primary neurons grown in the presence of P3HT nanowires. From left to right: resting membrane potential, membrane resistance, membrane capacitance (<b>A</b>). Representative traces of AP firing evoked by 350 pA current injection under both experimental conditions, and representative phase-plane plot analysis (dV/dt vs. V) of AP waveform obtained under control conditions (left) and in the presence of P3HT nanowires (right) (<b>B</b>). Minimal current injection intensity necessary to evoke an AP (rheobase) (<b>C</b>). From left to right: peak, threshold potential (V<sub>th</sub>), and width of AP (<b>D</b>). Instantaneous firing frequency (IFF) at each inter-spike interval (ISI) evoked by the injection of 350 pA depolarizing current (left) and mean firing frequency (MFF) plotted as a function of the injected current (right) (<b>E</b>). All data are shown as Means ± SEM with superimposed individual experimental points. All the analyzed parameters were not significantly different between the two experimental groups; Student’s <span class="html-italic">t</span>-test (<b>A</b>–<b>D</b>) and two-way ANOVA with Bonferroni’s multiple comparison tests (<b>E</b>); <span class="html-italic">n</span> = 9–13 cells, from two independent neuronal preparations).</p> Full article ">
24 pages, 12790 KiB  
Article
Effect of Initial Microstructure on the Toughness of Coarse-Grained Heat-Affected Zone in a Microalloyed Steel
by Minghao Shi, Man Di, Jian Zhang, Rangasayee Kannan, Jing Li, Xiaoguang Yuan and Leijun Li
Materials 2021, 14(16), 4760; https://doi.org/10.3390/ma14164760 - 23 Aug 2021
Cited by 3 | Viewed by 2256
Abstract
Toughness of the coarse-grained-heat-affected-zone (CGHAZ) strongly depends on the prior austenite grain size. The prior austenite grain size is affected not only by chemical composition, thermal cycle, and dissolution of second-phase particles, but also by the initial microstructure. The effect of base metal [...] Read more.
Toughness of the coarse-grained-heat-affected-zone (CGHAZ) strongly depends on the prior austenite grain size. The prior austenite grain size is affected not only by chemical composition, thermal cycle, and dissolution of second-phase particles, but also by the initial microstructure. The effect of base metal microstructure (ferrite/pearlite obtained by air cooling and martensite obtained by water-quenching) on Charpy impact toughness of the CGHAZ has been investigated for different heat inputs for high-heat input welding of a microalloyed steel. A welding thermal cycle with a heat input of 100 kJ/cm and 400 kJ/cm were simulated on the MMS-300 system. Despite a similar microstructure in the CGHAZ of both the base metals, the average Charpy impact energy for the air-cooled base metal was found to be higher than the water-quenched base metal. Through thermo-kinetic simulations, it was found that a higher enrichment of Mn/C at the ferrite/austenite transformation interface of the CGHAZ of water-quenched base metal resulted in stabilizing austenite at a lower A1 temperature, which resulted in a coarser austenite grain size and eventually lowering the toughness of the CGHAZ. Full article
(This article belongs to the Special Issue Advances in Novel Composites and Their Mechanical Properties)
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<p>Recorded welding thermal cycles for simulation of the coarse-grained heat-affected zone.</p> Full article ">Figure 2
<p>Typical microstructure of air-cooled base metal (<b>a</b>), CGHAZ with 100 kJ/cm in air-cooled initial microstructure (<b>b</b>), and CGHAZ with 400 kJ/cm in air-cooled initial microstructure (<b>c</b>), water-quenched base metal (<b>d</b>), CGHAZ with 100 kJ/cm in water-quenched initial microstructure (<b>e</b>), and the CGHAZ with 400 kJ/cm in water-quenched initial microstructure (<b>f</b>). P indicates pearlite, which is the back-etching grains; PF indicates polygonal ferrite, which mostly are equiaxed grains; M indicates martensite, which are thin laths; AF indicates acicular ferrite, which is ferrite that formed by combination of shear and displacive mechanism as sharp needles; GBF indicates grain boundary ferrite; WF indicates Widmanstatten ferrite; and IN indicates inclusions.</p> Full article ">Figure 2 Cont.
<p>Typical microstructure of air-cooled base metal (<b>a</b>), CGHAZ with 100 kJ/cm in air-cooled initial microstructure (<b>b</b>), and CGHAZ with 400 kJ/cm in air-cooled initial microstructure (<b>c</b>), water-quenched base metal (<b>d</b>), CGHAZ with 100 kJ/cm in water-quenched initial microstructure (<b>e</b>), and the CGHAZ with 400 kJ/cm in water-quenched initial microstructure (<b>f</b>). P indicates pearlite, which is the back-etching grains; PF indicates polygonal ferrite, which mostly are equiaxed grains; M indicates martensite, which are thin laths; AF indicates acicular ferrite, which is ferrite that formed by combination of shear and displacive mechanism as sharp needles; GBF indicates grain boundary ferrite; WF indicates Widmanstatten ferrite; and IN indicates inclusions.</p> Full article ">Figure 2 Cont.
<p>Typical microstructure of air-cooled base metal (<b>a</b>), CGHAZ with 100 kJ/cm in air-cooled initial microstructure (<b>b</b>), and CGHAZ with 400 kJ/cm in air-cooled initial microstructure (<b>c</b>), water-quenched base metal (<b>d</b>), CGHAZ with 100 kJ/cm in water-quenched initial microstructure (<b>e</b>), and the CGHAZ with 400 kJ/cm in water-quenched initial microstructure (<b>f</b>). P indicates pearlite, which is the back-etching grains; PF indicates polygonal ferrite, which mostly are equiaxed grains; M indicates martensite, which are thin laths; AF indicates acicular ferrite, which is ferrite that formed by combination of shear and displacive mechanism as sharp needles; GBF indicates grain boundary ferrite; WF indicates Widmanstatten ferrite; and IN indicates inclusions.</p> Full article ">Figure 3
<p>Prior austenite grain diameters in the CGHAZ with heat input of 100 kJ/cm and 400 kJ/cm. AC-100 is CGHAZ with 100 kJ/cm in air-cooled initial microstructure; AC-400 is CGHAZ with 400 kJ/cm in air-cooled initial microstructure; WQ-100 is CGHAZ with 100 kJ/cm in water-quenched initial microstructure; WQ-400 is CGHAZ with 400 kJ/cm in water-quenched initial microstructure.</p> Full article ">Figure 4
<p>Stress–strain curves for the specimens with air-cooled base metal (AC-BM) and water-quenched base metal (WQ-BM).</p> Full article ">Figure 5
<p>Typical load and absorbed energy versus displacement plots of base metal and CGHAZ by instrumented impact test at −20 °C. (<b>a</b>) base metal, (<b>b</b>) CGHAZ with 100 kJ/cm, (<b>c</b>) CGHAZ with 400 kJ/cm. AC indicates the air-cooled initial microstructure; WQ indicates the water-quenched initial microstructure.</p> Full article ">Figure 5 Cont.
<p>Typical load and absorbed energy versus displacement plots of base metal and CGHAZ by instrumented impact test at −20 °C. (<b>a</b>) base metal, (<b>b</b>) CGHAZ with 100 kJ/cm, (<b>c</b>) CGHAZ with 400 kJ/cm. AC indicates the air-cooled initial microstructure; WQ indicates the water-quenched initial microstructure.</p> Full article ">Figure 6
<p>SEM fractographs of Charpy impact tested at −20 °C of the base metal (<b>a</b>,<b>d</b>), CGHAZ with 100 kJ/cm (<b>b</b>,<b>e</b>), CGHAZ with 400 kJ/cm (<b>c</b>,<b>f</b>) of the air-cooled initial microstructure.</p> Full article ">Figure 7
<p>SEM fractographs of specimens Charpy impact tested at −20 °C of the base metal (<b>a</b>,<b>d</b>), CGHAZ with 100 kJ/cm (<b>b</b>,<b>e</b>), CGHAZ with 400 kJ/cm (<b>c</b>,<b>f</b>) of the water-quenched initial microstructure.</p> Full article ">Figure 8
<p>Prior austenite grains following a peak temperature of 1300 °C quenching by water for the air-cooled initial microstructure (<b>a</b>), for the water-quenched initial microstructure (<b>b</b>).</p> Full article ">Figure 8 Cont.
<p>Prior austenite grains following a peak temperature of 1300 °C quenching by water for the air-cooled initial microstructure (<b>a</b>), for the water-quenched initial microstructure (<b>b</b>).</p> Full article ">Figure 9
<p>On heating dilatation and the corresponding derivative for the air-cooled initial microstructure (<b>a</b>), and for the water-quenched initial microstructure (<b>b</b>), and (<b>c</b>) austenite phase fraction and transformation (<b>d</b>) for air-cooled and water-quenched initial microstructure.</p> Full article ">Figure 9 Cont.
<p>On heating dilatation and the corresponding derivative for the air-cooled initial microstructure (<b>a</b>), and for the water-quenched initial microstructure (<b>b</b>), and (<b>c</b>) austenite phase fraction and transformation (<b>d</b>) for air-cooled and water-quenched initial microstructure.</p> Full article ">Figure 9 Cont.
<p>On heating dilatation and the corresponding derivative for the air-cooled initial microstructure (<b>a</b>), and for the water-quenched initial microstructure (<b>b</b>), and (<b>c</b>) austenite phase fraction and transformation (<b>d</b>) for air-cooled and water-quenched initial microstructure.</p> Full article ">Figure 10
<p>Driving force (molar Gibbs free energy change) for on-heating ferrite to austenite transformation calculated using ThermoCalc. AC is for air-cooled initial microstructure; WQ is for water-quenched initial microstructure.</p> Full article ">Figure 11
<p>DICTRA and PRISMA calculations to understand the kinetics of austenite formation during CGHAZ simulation with different starting microstructures. (<b>a</b>) Evolution of phases as a function of time for the air-cooled and water-quenched starting microstructure, (<b>b</b>) normalized driving force for austenite formation as a function of temperature, (<b>c</b>) C concentration in ferrite as a function of transformation time, (<b>d</b>) Mn concentration in ferrite as a function of transformation time, (<b>e</b>) C concentration in austenite as a function of transformation time, and (<b>f</b>) Mn concentration in austenite as a function of transformation time.</p> Full article ">Figure 11 Cont.
<p>DICTRA and PRISMA calculations to understand the kinetics of austenite formation during CGHAZ simulation with different starting microstructures. (<b>a</b>) Evolution of phases as a function of time for the air-cooled and water-quenched starting microstructure, (<b>b</b>) normalized driving force for austenite formation as a function of temperature, (<b>c</b>) C concentration in ferrite as a function of transformation time, (<b>d</b>) Mn concentration in ferrite as a function of transformation time, (<b>e</b>) C concentration in austenite as a function of transformation time, and (<b>f</b>) Mn concentration in austenite as a function of transformation time.</p> Full article ">Figure 11 Cont.
<p>DICTRA and PRISMA calculations to understand the kinetics of austenite formation during CGHAZ simulation with different starting microstructures. (<b>a</b>) Evolution of phases as a function of time for the air-cooled and water-quenched starting microstructure, (<b>b</b>) normalized driving force for austenite formation as a function of temperature, (<b>c</b>) C concentration in ferrite as a function of transformation time, (<b>d</b>) Mn concentration in ferrite as a function of transformation time, (<b>e</b>) C concentration in austenite as a function of transformation time, and (<b>f</b>) Mn concentration in austenite as a function of transformation time.</p> Full article ">Figure 12
<p>Comparison between experimental austenite formation and model predicted austenite formation.</p> Full article ">Figure 13
<p>On cooling dilatation recordings and the corresponding derivative for the air-cooled initial microstructure (<b>a</b>), and for the water-quenched initial microstructure (<b>b</b>).</p> Full article ">
20 pages, 39966 KiB  
Article
Printing Parameter Requirements for 3D Printable Geopolymer Materials Prepared from Industrial Side Streams
by Qaisar Munir, Riku Peltonen and Timo Kärki
Materials 2021, 14(16), 4758; https://doi.org/10.3390/ma14164758 - 23 Aug 2021
Cited by 15 | Viewed by 2938
Abstract
The objective of this investigation is to study the printing parameter requirements for sustainable 3D printable geopolymer materials. Side streams of the paper, mining, and construction industries were applied as geopolymer raw materials. The effect of printing parameters in terms of buildability, mixability, [...] Read more.
The objective of this investigation is to study the printing parameter requirements for sustainable 3D printable geopolymer materials. Side streams of the paper, mining, and construction industries were applied as geopolymer raw materials. The effect of printing parameters in terms of buildability, mixability, extrudability, curing, Al-to-Si ratio, and waste materials utilisation on the fresh and hardened state of the materials was studied. The material performance of a fresh geopolymer was measured using setting time and shape stability tests. Standardised test techniques were applied in the testing of the hardened material properties of compressive and flexural strength. The majority of developed suitable 3D printable geopolymers comprised 56–58% recycled material. Heating was used to improve the buildability and setting of the material significantly. A reactive recyclable material content of greater than 20% caused the strength and material workability to decrease. A curing time of 7–28 days increased the compressive strength but decreased the flexural strength. The layers in the test samples exhibited decreased and increased strength, respectively, in compressive and flexural strength tests. Geopolymer development was found to be a compromise between different strength values and recyclable material contents. By focusing on specialised and complex-shape products, 3D printing of geopolymers can compete with traditional manufacturing in limited markets. Full article
(This article belongs to the Topic Geopolymers: Synthesis, Characterization and Applications)
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<p>Process flow chart of laboratory experiment for material testing of geopolymer for 3D printing.</p> Full article ">Figure 2
<p>Test specimens for testing: (<b>1</b>) Flexural test, (<b>2</b>) Setting (measurement points illustrated) and shape stability test, (<b>3</b>) Compressive test.</p> Full article ">Figure 3
<p>(<b>a</b>) Nozzles used in extruding with orifice dimensions, (<b>b</b>) Setup for measuring shape stability. Two digital Vernier callipers are fixed to the platform and the change in height is measured automatically.</p> Full article ">Figure 4
<p>(<b>a</b>) Machine used for compression testing (<b>b</b>) sample after compression testing result.</p> Full article ">Figure 5
<p>Machine used for flexural testing.</p> Full article ">Figure 6
<p>Compression values from initial compressive strength tests.</p> Full article ">Figure 7
<p>Initial setting time of geopolymer A at various temperatures.</p> Full article ">Figure 8
<p>Initial setting time of geopolymer B at various temperatures.</p> Full article ">Figure 9
<p>Shape stability of geopolymers A and B at 20 and 100 °C [<a href="#B47-materials-14-04758" class="html-bibr">47</a>].</p> Full article ">Figure 10
<p>Deformation of geopolymers A and B at 20 and 100 °C [<a href="#B47-materials-14-04758" class="html-bibr">47</a>].</p> Full article ">Figure 11
<p>Compression strength of geopolymers A and B after 7 and 28 days of curing.</p> Full article ">Figure 12
<p>Flexural strength of specimens A and B after 7 and 28 days of curing and their variations [<a href="#B47-materials-14-04758" class="html-bibr">47</a>].</p> Full article ">
34 pages, 12277 KiB  
Review
Multi-Scale Structure–Mechanical Property Relations of Graphene-Based Layer Materials
by Jingran Liu, Huasong Qin and Yilun Liu
Materials 2021, 14(16), 4757; https://doi.org/10.3390/ma14164757 - 23 Aug 2021
Cited by 13 | Viewed by 4864
Abstract
Pristine graphene is one of the strongest materials known in the world, and may play important roles in structural and functional materials. In order to utilize the extraordinary mechanical properties in practical engineering structures, graphene should be assembled into macroscopic structures such as [...] Read more.
Pristine graphene is one of the strongest materials known in the world, and may play important roles in structural and functional materials. In order to utilize the extraordinary mechanical properties in practical engineering structures, graphene should be assembled into macroscopic structures such as graphene-based papers, fibers, foams, etc. However, the mechanical properties of graphene-based materials such as Young’s modulus and strength are 1–2 orders lower than those of pristine monolayer graphene. Many efforts have been made to unveil the multi-scale structure–property relations of graphene-based materials with hierarchical structures spanning the nanoscale to macroscale, and significant achievements have been obtained to improve the mechanical performance of graphene-based materials through composition and structure optimization across multi-scale. This review aims at summarizing the currently theoretical, simulation, and experimental efforts devoted to the multi-scale structure–property relation of graphene-based layer materials including defective monolayer graphene, nacre-like and laminar nanostructures of multilayer graphene, graphene-based papers, fibers, aerogels, and graphene/polymer composites. The mechanisms of mechanical property degradation across the multi-scale are discussed, based on which some multi-scale optimization strategies are presented to further improve the mechanical properties of graphene-based layer materials. We expect that this review can provide useful insights into the continuous improvement of mechanical properties of graphene-based layer materials. Full article
(This article belongs to the Special Issue Structure, Function and Mechanics of Low-Dimensional Materials and Their Assemblies)
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<p>Multi-scale structures of graphene-based materials and the key loading bearing structures in each scale. Reprinted with permission from [<a href="#B8-materials-14-04757" class="html-bibr">8</a>,<a href="#B9-materials-14-04757" class="html-bibr">9</a>,<a href="#B10-materials-14-04757" class="html-bibr">10</a>,<a href="#B11-materials-14-04757" class="html-bibr">11</a>,<a href="#B12-materials-14-04757" class="html-bibr">12</a>,<a href="#B13-materials-14-04757" class="html-bibr">13</a>,<a href="#B14-materials-14-04757" class="html-bibr">14</a>].</p> Full article ">Figure 2
<p>In-plane mechanical behaviors of monolayer graphene. (<b>a</b>) Anisotropic toughness of pristine graphene. Reprinted with permission from [<a href="#B66-materials-14-04757" class="html-bibr">66</a>]. (<b>b</b>) Initial stress field induced by grain boundaries with different tilt angles along armchair (ac) and zigzag (zz) directions, and the strength-tilt angle relation. Adapted and reprinted with permission from [<a href="#B67-materials-14-04757" class="html-bibr">67</a>]. (<b>c</b>) Schematic plot of two grains showing GB tilt angle (<span class="html-italic">θ</span>) and tensile loading angle (<span class="html-italic">φ</span>). Reprinted with permission from [<a href="#B68-materials-14-04757" class="html-bibr">68</a>]. (<b>d</b>) Top and perspective views of graphene polycrystalline, and the survival probability predicted by the statistic model. Reprinted with permission from [<a href="#B69-materials-14-04757" class="html-bibr">69</a>].</p> Full article ">Figure 3
<p>Tension–shear models for the nacre-like structure. (<b>a</b>) Tension–shear chain (TSC) model. The mineral tablets were assumed to be rigid, and the shear deformation of protein was the same everywhere. Reprinted with permission from [<a href="#B74-materials-14-04757" class="html-bibr">74</a>]. (<b>b</b>) Deformable tension-shear (DTS) model. Two failure modes were distinguished, namely fracture of sheets (mode G) and failure of crosslinks (mode I). Adapted and reprinted with permission from [<a href="#B22-materials-14-04757" class="html-bibr">22</a>]. (<b>c</b>) Nonlinear tension–shear model. The interlayer interactions are elastic perfectly plastic. Reprinted with permission from [<a href="#B23-materials-14-04757" class="html-bibr">23</a>].</p> Full article ">Figure 4
<p>Bending models for laminar structures. (<b>a</b>) Multi-beam shear model, the in-plane extension of each sheet is neglected. Reprinted with permission from [<a href="#B78-materials-14-04757" class="html-bibr">78</a>]. (<b>b</b>) Modified Timoshenko beam model (MTBM). Reprinted with permission from [<a href="#B31-materials-14-04757" class="html-bibr">31</a>]. (<b>c</b>) Failure behaviors of GLNs under bending deformation, including interlayer shearing, rippling, and kink/delamination. Adapted and reprinted with permission from [<a href="#B32-materials-14-04757" class="html-bibr">32</a>].</p> Full article ">Figure 5
<p>Out-of-plane deformation of GLMs. (<b>a</b>) Intrinsic rippling. Reprinted with permission from [<a href="#B14-materials-14-04757" class="html-bibr">14</a>]. (<b>b</b>) Buckling induced by grain boundaries. Reprinted with permission from [<a href="#B69-materials-14-04757" class="html-bibr">69</a>]. (<b>c</b>) Intrinsic buckling of GLNs induced by external load parallel to the sheet direction. Reprinted with permission from [<a href="#B33-materials-14-04757" class="html-bibr">33</a>]. (<b>d</b>) Wrinkle propagation through flower-like GBs. Reprinted with permission from. [<a href="#B86-materials-14-04757" class="html-bibr">86</a>].</p> Full article ">Figure 6
<p>In-plane mechanical properties of defective graphene monolayer. (<b>a</b>) 2D elastic modulus of monolayer graphene as a function of vacancy density observed in the experiments. Adapted and reprinted with permission from [<a href="#B20-materials-14-04757" class="html-bibr">20</a>]. (<b>b</b>) 2D elastic modulus as a function of GB density. Adapted and reprinted with permission from [<a href="#B98-materials-14-04757" class="html-bibr">98</a>]. (<b>c</b>) 2D elastic modulus and strength as a function of GB density. Adapted and reprinted with permission from [<a href="#B100-materials-14-04757" class="html-bibr">100</a>] (<b>d</b>) Strength, failure strain, and toughness of vacancy defected monolayer graphene as a function of vacancy density. Reprinted with permission from [<a href="#B97-materials-14-04757" class="html-bibr">97</a>]. (<b>e</b>) Flower-like grain boundary. Reprinted with permission from [<a href="#B101-materials-14-04757" class="html-bibr">101</a>]. (<b>f</b>) Atomic-scale crack bridging. Reprinted with permission from [<a href="#B102-materials-14-04757" class="html-bibr">102</a>]. (<b>g</b>) Shielding effect of dislocation to crack propagation. Reprinted with permission from [<a href="#B73-materials-14-04757" class="html-bibr">73</a>].</p> Full article ">Figure 7
<p>Bending behaviors of multilayer graphene. (<b>a</b>) Direct measurements of bending rigidity of multilayer graphene, molybdenum disulfide (MoS2), and hexagonal boron nitride (hBN) based on pressurized bubbles. Reprinted with permission from [<a href="#B111-materials-14-04757" class="html-bibr">111</a>,<a href="#B112-materials-14-04757" class="html-bibr">112</a>]. (<b>b</b>) Measurements and results of bending rigidity of multilayer graphene as a function of bending angle. Reprinted with permission from [<a href="#B30-materials-14-04757" class="html-bibr">30</a>].</p> Full article ">Figure 8
<p>(<b>a</b>) Interlayer crosslinks including covalent (VCB1/2) and non-covalent (CB and HB) bonds. Reprinted with permission from [<a href="#B119-materials-14-04757" class="html-bibr">119</a>]. Interlayer crosslinks (<b>b</b>) and geometrical locking (<b>c</b>) can improve the in-plane shear properties of multilayer graphene. Reprinted with permission from [<a href="#B22-materials-14-04757" class="html-bibr">22</a>,<a href="#B120-materials-14-04757" class="html-bibr">120</a>].</p> Full article ">Figure 9
<p>Conformation phase map of GO in solution accompanying the potential energy landscape. Reprinted with permission from [<a href="#B126-materials-14-04757" class="html-bibr">126</a>].</p> Full article ">Figure 10
<p>Graphene fibers (GFs). (<b>a</b>) Improved tensile strength for GFs with introduced polymers among adjacent sheets. Reprinted with permission from [<a href="#B132-materials-14-04757" class="html-bibr">132</a>]. (<b>b</b>) Tensile stress curves for GFs with giant graphene sheets and different crosslinks. Reprinted with permission from [<a href="#B9-materials-14-04757" class="html-bibr">9</a>]. (<b>c</b>) Schematic plot of full-scale synergetic defect engineering to manage possible defects of GFs from the atomic scale to macroscale and tensile stress curves of resultant GFs. Reprinted with permission from [<a href="#B133-materials-14-04757" class="html-bibr">133</a>]. (<b>d</b>) Surface morphologies of shrunk GFs fabricated by dry-spinning method and corresponding strength and break elongation. Reprinted with permission from [<a href="#B129-materials-14-04757" class="html-bibr">129</a>].</p> Full article ">Figure 11
<p>Stiffening and strengthening strategies for graphene papers (GPs). (<b>a</b>) Schematic plots of crosslinks introduced to adjacent sheets including (top) divalent bonds, (left to right in the bottom row) borate, polymers, and π–π crosslinking. Reprinted with permission from [<a href="#B142-materials-14-04757" class="html-bibr">142</a>,<a href="#B147-materials-14-04757" class="html-bibr">147</a>,<a href="#B148-materials-14-04757" class="html-bibr">148</a>,<a href="#B149-materials-14-04757" class="html-bibr">149</a>]. (<b>b</b>) Self-stiffening approach by cyclic stretching with small strain amplitude. Reprinted with permission from [<a href="#B151-materials-14-04757" class="html-bibr">151</a>]. (<b>c</b>) Plasticizer-assistant stretching method to straighten the wrinkles of direct-cast GO papers, and the resultant stress–strain curves. Reprinted with permission from [<a href="#B143-materials-14-04757" class="html-bibr">143</a>].</p> Full article ">Figure 12
<p>Improving the flexibility of graphene papers (GPs). (<b>a</b>) Surface wrinkling method and top view of the resultant wrinkle morphology. Reprinted with permission from [<a href="#B153-materials-14-04757" class="html-bibr">153</a>]. (<b>b</b>) Rubber-like GPs with hierarchical crumples obtained by sheet collapsing approach. Reprinted with permission from [<a href="#B125-materials-14-04757" class="html-bibr">125</a>]. (<b>c</b>) Thermal annealing approach and resultant highly flexible GPs with micro folds. Reprinted with permission from [<a href="#B141-materials-14-04757" class="html-bibr">141</a>].</p> Full article ">Figure 13
<p>Graphene aerogels (GAs). (<b>a</b>) Regular microstructures of GAs resultant from template directing method. Reprinted with permission from [<a href="#B163-materials-14-04757" class="html-bibr">163</a>]. (<b>b</b>) Bio-mimicking lamellar GAs with exceptional strength and mechanical resilience. Reprinted with permission from [<a href="#B165-materials-14-04757" class="html-bibr">165</a>]. (<b>c</b>) 3D printed stretchable GAs. Reprinted with permission from [<a href="#B164-materials-14-04757" class="html-bibr">164</a>].</p> Full article ">Figure 14
<p>Graphene/polymer nanocomposites with graphene sheets dispersed in polymers. (<b>a</b>) Graphene sheets dispersed near randomly in polymers, enhanced strength and decreased break elongation. Reprinted with permission from [<a href="#B180-materials-14-04757" class="html-bibr">180</a>]. (<b>b</b>) Introducing interfacial crosslinks (covalent bonds here) between graphene sheets and polymers brings out overall enhanced mechanical properties. Reprinted with permission from [<a href="#B43-materials-14-04757" class="html-bibr">43</a>]. (<b>c</b>) Fractured surface of nanocomposites with covalent (<b>top</b>) and non-covalent (<b>bottom</b>) functionalized graphene sheets. Reprinted with permission from [<a href="#B43-materials-14-04757" class="html-bibr">43</a>,<a href="#B183-materials-14-04757" class="html-bibr">183</a>].</p> Full article ">Figure 15
<p>Graphene/polymer nanocomposites with polymer infiltrated into graphene scaffold and bi-continuous microstructure. (<b>a</b>) RGO/epoxy composites, in which the graphene scaffold was grown on a Ni foam template, showing enhanced flexural modulus, strength, and fracture toughness compared to pure epoxy. Reprinted with permission from [<a href="#B190-materials-14-04757" class="html-bibr">190</a>]. (<b>b</b>) RGO/epoxy composites with well-aligned rGO-CMC (carboxymethyl cellulose sodium) scaffold by ice templating, showing enhanced fracture strength and toughness. Reprinted with permission from [<a href="#B45-materials-14-04757" class="html-bibr">45</a>]. (<b>c</b>) Bi-continuous GO/PVA composites, showing overall enhancement of mechanical properties including Young’s modulus, strength, break elongation, and toughness compared to pure PVA. Reprinted with permission from [<a href="#B46-materials-14-04757" class="html-bibr">46</a>].</p> Full article ">Figure 16
<p>Nacre-like graphene/polymer nanocomposites. (<b>a</b>) PVA/GO composite films with enhanced tensile strength and fracture strain. Adapted and reprinted with permission from [<a href="#B191-materials-14-04757" class="html-bibr">191</a>]. (<b>b</b>) GO (rGO)/CMC composites with improved mechanical properties. Adapted and reprinted with permission from [<a href="#B179-materials-14-04757" class="html-bibr">179</a>].</p> Full article ">Figure 17
<p>Multi-scale optimization of GLMs. Reprinted with permission from [<a href="#B20-materials-14-04757" class="html-bibr">20</a>,<a href="#B23-materials-14-04757" class="html-bibr">23</a>,<a href="#B26-materials-14-04757" class="html-bibr">26</a>,<a href="#B43-materials-14-04757" class="html-bibr">43</a>,<a href="#B46-materials-14-04757" class="html-bibr">46</a>,<a href="#B67-materials-14-04757" class="html-bibr">67</a>,<a href="#B97-materials-14-04757" class="html-bibr">97</a>,<a href="#B102-materials-14-04757" class="html-bibr">102</a>,<a href="#B113-materials-14-04757" class="html-bibr">113</a>,<a href="#B125-materials-14-04757" class="html-bibr">125</a>,<a href="#B143-materials-14-04757" class="html-bibr">143</a>].</p> Full article ">
17 pages, 3359 KiB  
Article
Physicochemical and Optical Characterization of Citrus aurantium Derived Biochar for Solar Absorber Applications
by Nancy G. Gonzalez-Canche, Jose G. Carrillo, Beatriz Escobar-Morales, Iván Salgado-Tránsito, Neith Pacheco, Soledad Cecilia Pech-Cohuo and Manuel I. Peña-Cruz
Materials 2021, 14(16), 4756; https://doi.org/10.3390/ma14164756 - 23 Aug 2021
Cited by 18 | Viewed by 3690
Abstract
Agro-industrial waste valorization is an attractive approach that offers new alternatives to deal with shrinkage and residue problems. One of these approaches is the synthesis of advanced carbon materials. Current research has shown that citrus waste, mainly orange peel, can be a precursor [...] Read more.
Agro-industrial waste valorization is an attractive approach that offers new alternatives to deal with shrinkage and residue problems. One of these approaches is the synthesis of advanced carbon materials. Current research has shown that citrus waste, mainly orange peel, can be a precursor for the synthesis of high-quality carbon materials for chemical adsorption and energy storage applications. A recent approach to the utilization of advanced carbon materials based on lignocellulosic biomass is their use in solar absorber coatings for solar-thermal applications. This study focused on the production of biochar from Citrus aurantium orange peel by a pyrolysis process at different temperatures. Biochars were characterized by SEM, elemental analysis, TGA-DSC, FTIR, DRX, Raman, and XPS spectroscopies. Optical properties such as diffuse reflectance in the UV−VIS−NIR region was also determined. Physical-chemical characterization revealed that the pyrolysis temperature had a negative effect in yield of biochars, whereas biochars with a higher carbon content, aromaticity, thermal stability, and structural order were produced as the temperature increased. Diffuse reflectance measurements revealed that it is possible to reduce the reflectance of the material by controlling its pyrolysis temperature, producing a material with physicochemical and optical properties that could be attractive for use as a pigment in solar absorber coatings. Full article
(This article belongs to the Special Issue Carbon Materials Applied for Biomass Conversion)
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<p>SEM images of SOPRAW, SOP400, SOP600, and SOP800.</p> Full article ">Figure 2
<p>FTIR spectra of SOPRAW and SOP400, SOP600, and SOP800.</p> Full article ">Figure 3
<p>(<b>a</b>) TG-DTG curves; (<b>b</b>) DSC curves of SOPRAW.</p> Full article ">Figure 4
<p>TG curves of the SOPRAW and biochars.</p> Full article ">Figure 5
<p>XRD patterns of SOPRAW and biochars SOP400, SOP600, and SOP800.</p> Full article ">Figure 6
<p>Raman spectra SOPRAW and biochars SOP400, SOP600, and SOP800.</p> Full article ">Figure 7
<p>Raman evolution of D band width vs disorder parameter.</p> Full article ">Figure 8
<p>XPS survey spectra of SOPRAW and biochars.</p> Full article ">Figure 9
<p>UV−VIS−NIR diffuse reflectance of SOPRAW and biochars.</p> Full article ">
24 pages, 5593 KiB  
Article
Comparative Gravimetric Studies on Carbon Steel Corrosion in Selected Fruit Juices and Acidic Chloride Media (HCl) at Different pH
by Stanley Udochukwu Ofoegbu
Materials 2021, 14(16), 4755; https://doi.org/10.3390/ma14164755 - 23 Aug 2021
Cited by 6 | Viewed by 3677
Abstract
Food contamination due to metal corrosion and the consequent leakage of metals into foods is a problem. Understanding the mechanism(s) of metal corrosion in food media is vital to evaluating, mitigating, and predicting contamination levels. Fruit juices have been employed as model corrosive [...] Read more.
Food contamination due to metal corrosion and the consequent leakage of metals into foods is a problem. Understanding the mechanism(s) of metal corrosion in food media is vital to evaluating, mitigating, and predicting contamination levels. Fruit juices have been employed as model corrosive media to study the corrosion behaviour of metallic material in food media. Carbon steel corrosion in fresh juices of tomato, orange, pineapple, and lemon, as well as dilute hydrochloric acid solutions at varied pH, was studied using scanning electron microscopy, gravimetric and spectrophotometric techniques, and comparisons made between the corrosivity of these juices and mineral acids of comparable pH. The corrosion of carbon steel in fruit juices and HCl solutions manifests as a combination of uniform and pitting corrosion. Gravimetric data acquired after one hour of immersion at ambient temperature (22 °C) indicated corrosion rates of 0.86 mm yr−1 in tomato juice (pH ≈ 4.24), 1.81 mm yr−1 in pineapple juice (pH ≈ 3.94), 1.52 mm yr−1 in orange juice (pH ≈ 3.58), and 2.89 mm yr−1 in lemon juice (pH ≈ 2.22), compared to 2.19 mm yr−1 in 10−2 M HCl (pH ≈ 2.04), 0.38 mm yr−1 in 10−3 M HCl (pH ≈ 2.95), 0.17 mm yr−1 in 10−4 M HCl (pH ≈ 3.95), and 0.04 mm yr−1 in 10−5 M HCl (pH ≈ 4.98). The correlation of gravimetrically acquired corrosion data with post-exposure spectrophotometric analysis of fruit juices enabled de-convolution of iron contamination rates from carbon steel corrosion rates in fruit juices. Elemental iron contamination after 50 h of exposure to steel samples was much less than the values predicted from corrosion data (≈40%, 4.02%, 8.37%, and 9.55% for tomato, pineapple, orange, and lemon juices, respectively, relative to expected values from corrosion (weight loss) data). Tomato juice (pH ≈ 4.24) was the least corrosive to carbon steel compared to orange juice (pH ≈ 3.58) and pineapple juice (pH ≈ 3.94). The results confirm that though the fruit juices are acidic, they are generally much less corrosive to carbon steel compared to hydrochloric acid solutions of comparable pH. Differences in the corrosion behaviour of carbon steel in the juices and in the different mineral acid solutions are attributed to differences in the compositions and pH of the test media, the nature of the corrosion products formed, and their dissolution kinetics in the respective media. The observation of corrosion products (iron oxide/hydroxide) in some of the fruit juices (tomato, pineapple, and lemon juices) in the form of apparently hollow microspheres indicates the feasibility of using fruit juices and related wastes as “green solutions” for the room-temperature and hydrothermal synthesis of metal oxide/hydroxide particles. Full article
(This article belongs to the Section Corrosion)
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<p>Illustration of the three major effects of corrosion and relevance to different industries.</p> Full article ">Figure 2
<p>(<b>a</b>) SEM image of as-received carbon steel sheets; (<b>b</b>) microstructural (optical) image of CR1 Carbon Steel Sheet after etching in 2% Nital.</p> Full article ">Figure 3
<p>Plot of specific weight loss as a function of immersion time, respectively, for carbon steel in (<b>a</b>) different fruit juices and (<b>b</b>) HCl solutions of comparable pH.</p> Full article ">Figure 4
<p>Plot of corrosion rate as a function of immersion time, respectively, for carbon steel in different fruit juices and HCl solutions of comparable pH from gravimetric data.</p> Full article ">Figure 5
<p>SEM images of carbon steel sheet (CR1 ≈ EN 1.0338) surfaces after immersion in (<b>a</b>) tomato juice, (<b>b</b>) 10<sup>−5</sup> M HCl, (<b>c</b>) pineapple juice, (<b>d</b>) 10<sup>−4</sup> M HCl, (<b>e</b>) orange juice, (<b>f</b>) 10<sup>−3</sup> M HCl, (<b>g</b>) lemon juice, and (<b>h</b>) 10<sup>−2</sup> M HCl, respectively.</p> Full article ">Figure 6
<p>SEM/EDS mapping images of carbon steel sheet (CR1 ≈ EN 1.0338) surfaces after immersion in tomato juice (pH = 4.24).</p> Full article ">Figure 7
<p>SEM/EDS mapping images of carbon steel sheet (CR1 ≈ EN 1.0338) surfaces after immersion in pineapple juice (pH = 3.94).</p> Full article ">Figure 8
<p>SEM/EDS mapping images of carbon steel sheet (CR1 ≈ EN 1.0338) surfaces after immersion in orange juice (pH = 3.58).</p> Full article ">Figure 9
<p>SEM/EDS mapping images of carbon steel sheet (CR1 ≈ EN 1.0338) surfaces after immersion in lemon juice (pH = 2.22).</p> Full article ">Figure 10
<p>Cross-sectional SEM image with elemental maps for carbon steel sheet surfaces after 12 h of immersion in 10<sup>−3</sup> M HCl solution (pH = 2.95).</p> Full article ">Figure 11
<p>Cross-sectional SEM image with elemental maps for carbon steel sheet surfaces after 12 h of immersion in tomato juice (pH = 4.24).</p> Full article ">Figure 12
<p>X-ray diffraction patterns acquired from carbon steel (CR1 ≈ EN 1.0338) surfaces after 12 h of immersion in (<b>a</b>) lemon juice, (<b>b</b>) orange juice, (<b>c</b>) pineapple juice, and (<b>d</b>) tomato juice, (Ak is Akaganéite, Ma is Maghemite, and Mt is Magnetite).</p> Full article ">Figure 13
<p>Plot of corrosion rate vs. test media pH for carbon steels after 5 h immersion in fruit juices, and acidic chloride solutions of varying pH evaluated from gravimetric data.</p> Full article ">
17 pages, 5357 KiB  
Article
Effects of FSW Tool Plunge Depth on Properties of an Al-Mg-Si Alloy T-Joint: Thermomechanical Modeling and Experimental Evaluation
by Shabbir Memon, Dariusz Fydrych, Aintzane Conde Fernandez, Hamed Aghajani Derazkola and Hesamoddin Aghajani Derazkola
Materials 2021, 14(16), 4754; https://doi.org/10.3390/ma14164754 - 23 Aug 2021
Cited by 39 | Viewed by 4157
Abstract
One of the main challenging issues in friction stir welding (FSW) of stiffened structures is maximizing skin and flange mixing. Among the various parameters in FSW that can affect the quality of mixing between skin and flange is tool plunge depth (TPD). In [...] Read more.
One of the main challenging issues in friction stir welding (FSW) of stiffened structures is maximizing skin and flange mixing. Among the various parameters in FSW that can affect the quality of mixing between skin and flange is tool plunge depth (TPD). In this research, the effects of TPD during FSW of an Al-Mg-Si alloy T-joint are investigated. The computational fluid dynamics (CFD) method can help understand TPD effects on FSW of the T-joint structure. For this reason, the CFD method is employed in the simulation of heat generation, heat distribution, material flow, and defect formation during welding processes at various TPD. CFD is a powerful method that can simulate phenomena during the mixing of flange and skin that are hard to assess experimentally. For the evaluation of FSW joints, macrostructure visualization is carried out. Simulation results showed that at higher TPD, more frictional heat is generated and causes the formation of a bigger stir zone. The temperature distribution is antisymmetric to the welding line, and the concentration of heat on the advancing side (AS) is more than the retreating side (RS). Simulation results from viscosity changes and material velocity study on the stir zone indicated that the possibility of the formation of a tunnel defect on the skin–flange interface at the RS is very high. Material flow and defect formation are very sensitive to TPD. Low TPD creates internal defects with incomplete mixing of skin and flange, and high TPD forms surface flash. Higher TPD increases frictional heat and axial force that diminish the mixing of skin and flange in this joint. The optimum TPD was selected due to the best materials flow and final mechanical properties of joints. Full article
(This article belongs to the Special Issue Effect of Hot Manufacturing Methods on Material Processing by Finite Element Modelling)
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<p>Schematic view of thermocouple position on (<b>a</b>) FSW tool and (<b>b</b>) skin and flange side.</p> Full article ">Figure 2
<p>(<b>a</b>) Stress–strain graph of AA6068 aluminum alloy at various temperatures. (<b>b</b>) Specific heat and (<b>c</b>) thermal conductivity of AA6068 aluminum alloy at various temperatures. (<b>d</b>) Meshed domain.</p> Full article ">Figure 3
<p>(<b>a</b>) Recorded temperature inside the FSW tool. (<b>b</b>) Recorded surface temperature of the joint. (<b>c</b>) Recorded temperature inside the FSW tool at various TPDs. (<b>d</b>) Cross-sectional view of the simulation results of internal heat flow. (<b>e</b>) Simulation results of generated heat at various TPDs.</p> Full article ">Figure 4
<p>(<b>a</b>) Simulation results of internal heat flow at 0.2 mm TPD. (<b>b</b>) Simulation results of internal heat flow at 0.2 mm TPD. (<b>c</b>) Experimental and (<b>d</b>) maximum simulated temperature and heat diffusion on skin. (<b>e</b>) Recorded and (<b>f</b>) simulated results of maximum temperature and heat diffusion on flange.</p> Full article ">Figure 5
<p>(<b>a</b>) Simulation results of internal viscosity changes at 0.2 mm and 0.4 mm TPD. (<b>b</b>) Force analysis of the FSW tool at 0.2 mm and 0.4 mm.</p> Full article ">Figure 6
<p>(<b>a</b>) Surface materials flow at 0.2 mm and 0.4 mm TPD. (<b>b</b>) Three-dimensional (3D) view of material velocity from simulation result. (<b>c</b>) Side view from velocity simulation result of various TPD.</p> Full article ">Figure 7
<p>(<b>a</b>) Simulation results of strain rate in 0.2 mm and 0.4 mm TPD joints. (<b>b</b>) Microstructure of SZ in 0.2 mm and 0.4 mm TPD joints. Statistic results of (<b>c</b>) joint size and (<b>d</b>) grain size in skin and flange at different cases.</p> Full article ">
13 pages, 3691 KiB  
Article
Effects of Sc and Zr Addition on Microstructure and Mechanical Properties of AA5182
by Jingxiao Li, Xiaofang Yang, Shihua Xiang, Yongfa Zhang, Jie Shi, Youcai Qiu and Robert Edward Sanders
Materials 2021, 14(16), 4753; https://doi.org/10.3390/ma14164753 - 23 Aug 2021
Cited by 7 | Viewed by 2273
Abstract
The effects of 0.1 wt.% Sc and 0.1 wt.% Zr addition in AA5182 on microstructure and mechanical properties were investigated. Results show that Al3(ScxZr1−x) dispersoids formed in AA5182. Observation of ingots microstructures showed that the grain size [...] Read more.
The effects of 0.1 wt.% Sc and 0.1 wt.% Zr addition in AA5182 on microstructure and mechanical properties were investigated. Results show that Al3(ScxZr1−x) dispersoids formed in AA5182. Observation of ingots microstructures showed that the grain size of 5182-Sc-Zr alloy was 56% lower than that of based AA5182. Isothermal annealing between 230 °C and 500 °C for 2 h was performed to study the recrystallization, tensile properties and dispersoid coarsening. The recrystallization was inhibited by the dispersoids, and the alloy microstructure remained deformed after annealing. Al3(ScxZr1−x) in AA5182 was stable when annealing below 400 °C, while parts of dispersoids coarsened significantly when heating at 500 °C. The addition of Sc and Zr allowed YS of 5182 alloy to achieve 247.8 MPa, which is 100 MPa higher than the corresponding AA5182. The contributions of Orowan strengthening and grain boundary strengthening were obtained by calculation. Full article
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<p>Microstructure of as-cast samples: (<b>a</b>) 5182 alloy; (<b>b</b>) 5185-Sc-Zr alloy.</p> Full article ">Figure 2
<p>Microstructures after isothermal annealing for 2 h: (<b>a</b>) 5182 alloy annealed at 230 °C; (<b>b</b>) 5182-Sc-Zr alloy annealed at 230 °C; (<b>c</b>) 5182 alloy annealed at 250 °C; (<b>d</b>) 5182-Sc-Zr alloy annealed at 250 °C; (<b>e</b>) 5182 alloy annealed at 300 °C; (<b>f</b>) 5182-Sc-Zr alloy annealed at 300 °C; (<b>g</b>) 5182 alloy annealed at 400 °C; (<b>h</b>) 5182-Sc-Zr alloy annealed at 400 °C; (<b>i</b>) 5182 alloy annealed at 500 °C; (<b>j</b>) 5182-Sc-Zr alloy annealed at 500 °C.</p> Full article ">Figure 3
<p>TEM of 5182 alloy: (<b>a</b>) annealed at 250 °C; (<b>b</b>) annealed at 300 °C.</p> Full article ">Figure 4
<p>TEM of annealed 5182-Sc-Zr alloy: (<b>a</b>–c) annealed at 250 °C; (<b>d</b>–<b>f</b>) annealed at 300 °C; (<b>g</b>–<b>i</b>) annealed at 400 °C; (<b>j</b>–<b>l</b>) annealed at 500 °C.</p> Full article ">Figure 5
<p>Average diameter of the Al<sub>3</sub>(Sc<sub>x</sub>Zr<sub>1−x</sub>) dispersoids after annealing.</p> Full article ">Figure 6
<p>Stress–strain curves, samples with cold rolling reduction rate of 70% (<b>a</b>) cold rolled; (<b>b</b>) cold rolled and annealed at 230 °C; (<b>c</b>) cold rolled and annealed at 250 °C; (<b>d</b>) cold rolled and annealed at 300 °C; (<b>e</b>) cold rolled and annealed at 400 °C; (<b>f</b>) cold rolled and annealed at 500 °C.</p> Full article ">
19 pages, 7990 KiB  
Article
New Equivalent Thermal Conductivity Model for Size-Dependent Convection-Driven Melting of Spherically Encapsulated Phase Change Material
by Feng Hou, Shihao Cao and Hui Wang
Materials 2021, 14(16), 4752; https://doi.org/10.3390/ma14164752 - 23 Aug 2021
Cited by 6 | Viewed by 2536
Abstract
Spherically encapsulated phase change materials (PCMs) are extensively incorporated into matrix material to form composite latent heat storage system for the purposes of saving energy, reducing PCM cost and decreasing space occupation. Although the melting of PCM sphere has been studied comprehensively by [...] Read more.
Spherically encapsulated phase change materials (PCMs) are extensively incorporated into matrix material to form composite latent heat storage system for the purposes of saving energy, reducing PCM cost and decreasing space occupation. Although the melting of PCM sphere has been studied comprehensively by experimental and numerical methods, it is still challenging to quantitatively depict the contribution of complex natural convection (NC) to the melting process in a practically simple and acceptable way. To tackle this, a new effective thermal conductivity model is proposed in this work by focusing on the total melting time (TMT) of PCM, instead of tracking the complex evolution of solid–liquid interface. Firstly, the experiment and finite element simulation of the constrained and unconstrained meltings of paraffin sphere are conducted to provide a deep understanding of the NC-driven melting mechanism and exhibit the difference of melting process. Then the dependence of NC on the particle size and heating temperature is numerically investigated for the unconstrained melting which is closer to the real-life physics than the constrained melting. Subsequently, the contribution of NC to the TMT is approximately represented by a simple effective thermal conductivity correlation, through which the melting process of PCM is simplified to involve heat conduction only. The effectiveness of the equivalent thermal conductivity model is demonstrated by rigorous numerical analysis involving NC-driven melting. By addressing the TMT, the present correlation thoroughly avoids tracking the complex evolution of melting front and would bring great convenience to engineering applications. Full article
(This article belongs to the Special Issue Recent Developments in Phase Change Materials for Energy Storage Applications)
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<p>Schematic diagrams of: (<b>a</b>) equivalent thermal conductivity model, (<b>b</b>) constrained melting, and (<b>c</b>) unconstrained melting.</p> Full article ">Figure 2
<p>(<b>a</b>) The spherical enclosure filled with paraffin wax and (<b>b</b>) the schematic diagram of the constrained computational domain embedding with a thermocouple located at the center of the sphere.</p> Full article ">Figure 3
<p>(<b>a</b>) Schematic diagram and (<b>b</b>) actual picture of the experimental setup for the constrained sample.</p> Full article ">Figure 4
<p>The mesh grid used in the two-dimensional axisymmetric finite element model.</p> Full article ">Figure 5
<p>Transient shapes of the solid paraffin for the constrained melting inside the spherical container with <math display="inline"><semantics> <mrow> <mi>R</mi> <mo>=</mo> <mn>31.02</mn> <mrow> <mo> </mo> <mi>mm</mi> <mo> </mo> <mi>and</mi> <mo> </mo> </mrow> <mo> </mo> <msub> <mi>T</mi> <mi>w</mi> </msub> <mo>=</mo> <mn>318.15</mn> <mo> </mo> <mi mathvariant="normal">K</mi> </mrow> </semantics></math>.</p> Full article ">Figure 6
<p>Bottom waviness profile of the sold PCM at 60 min for the constrained melting inside the spherical container with <math display="inline"><semantics> <mrow> <mi>R</mi> <mo>=</mo> <mn>31.02</mn> <mrow> <mo> </mo> <mi>mm</mi> <mo> </mo> <mi>and</mi> <mo> </mo> </mrow> <mo> </mo> <msub> <mi>T</mi> <mi>w</mi> </msub> <mo>=</mo> <mn>318.15</mn> <mo> </mo> <mi mathvariant="normal">K</mi> </mrow> </semantics></math>.</p> Full article ">Figure 7
<p>Comparison of the computational and measured temperatures at the central point of the spherical container during the constrained melting process with <math display="inline"><semantics> <mrow> <mi>R</mi> <mo>=</mo> <mn>31.02</mn> <mrow> <mo> </mo> <mi>mm</mi> <mo> </mo> <mi>and</mi> <mo> </mo> </mrow> <mo> </mo> <msub> <mi>T</mi> <mi>w</mi> </msub> <mo>=</mo> <mn>318.15</mn> <mo> </mo> <mi mathvariant="normal">K</mi> </mrow> </semantics></math>.</p> Full article ">Figure 8
<p>Thermally stable and unstable regions for the constrained melting at 70 min inside the spherical container with <math display="inline"><semantics> <mrow> <mi>R</mi> <mo>=</mo> <mn>31.02</mn> <mrow> <mo> </mo> <mi>mm</mi> <mo> </mo> <mi>and</mi> <mo> </mo> </mrow> <mo> </mo> <msub> <mi>T</mi> <mi>w</mi> </msub> <mo>=</mo> <mn>318.15</mn> <mo> </mo> <mi mathvariant="normal">K</mi> </mrow> </semantics></math>.</p> Full article ">Figure 9
<p>The PCM melting process without natural convection for the constrained melting inside the spherical container with <math display="inline"><semantics> <mrow> <mi>R</mi> <mo>=</mo> <mn>31.02</mn> <mrow> <mo> </mo> <mi>mm</mi> <mo> </mo> <mi>and</mi> <mo> </mo> </mrow> <mo> </mo> <msub> <mi>T</mi> <mi>w</mi> </msub> <mo>=</mo> <mn>318.15</mn> <mo> </mo> <mi mathvariant="normal">K</mi> </mrow> </semantics></math>.</p> Full article ">Figure 10
<p>Transient shapes of the solid paraffin for the unconstrained melting process inside the spherical container with <math display="inline"><semantics> <mrow> <mi>R</mi> <mo>=</mo> <mn>31.02</mn> <mrow> <mo> </mo> <mi>mm</mi> <mo> </mo> <mi>and</mi> </mrow> <mo> </mo> <msub> <mi>T</mi> <mi>w</mi> </msub> <mo>=</mo> <mn>318.15</mn> <mo> </mo> <mi mathvariant="normal">K</mi> </mrow> </semantics></math>.</p> Full article ">Figure 11
<p>Comparison of liquid fraction calculated by the PC and NC models for the case of <math display="inline"><semantics> <mrow> <mo>Δ</mo> <mi>T</mi> <mo>=</mo> <mn>10</mn> <mrow> <mo> </mo> <mi mathvariant="normal">K</mi> </mrow> </mrow> </semantics></math>.</p> Full article ">Figure 12
<p>Variations of TMT ratio for different sphere sizes and heating temperatures.</p> Full article ">
13 pages, 4980 KiB  
Article
Electrochemically Inert Li2MnO3: The Key to Improving the Cycling Stability of Li-Rich Manganese Oxide Used in Lithium-Ion Batteries
by Lian-Bang Wang, He-Shan Hu, Wei Lin, Qing-Hong Xu, Jia-Dong Gong, Wen-Kui Chai and Chao-Qi Shen
Materials 2021, 14(16), 4751; https://doi.org/10.3390/ma14164751 - 23 Aug 2021
Cited by 1 | Viewed by 3461
Abstract
Lithium-rich manganese oxide is a promising candidate for the next-generation cathode material of lithium-ion batteries because of its low cost and high specific capacity. Herein, a series of xLi2MnO3·(1 − x)LiMnO2 nanocomposites were designed via an ingenious one-step [...] Read more.
Lithium-rich manganese oxide is a promising candidate for the next-generation cathode material of lithium-ion batteries because of its low cost and high specific capacity. Herein, a series of xLi2MnO3·(1 − x)LiMnO2 nanocomposites were designed via an ingenious one-step dynamic hydrothermal route. A high concentration of alkaline solution, intense hydrothermal conditions, and stirring were used to obtain nanoparticles with a large surface area and uniform dispersity. The experimental results demonstrate that 0.072Li2MnO3·0.928LiMnO2 nanoparticles exhibit a desirable electrochemical performance and deliver a high capacity of 196.4 mAh g−1 at 0.1 C. This capacity was maintained at 190.5 mAh g−1 with a retention rate of 97.0% by the 50th cycle, which demonstrates the excellent cycling stability. Furthermore, XRD characterization of the cycled electrode indicates that the Li2MnO3 phase of the composite is inert, even under a high potential (4.8 V), which is in contrast with most previous reports of lithium-rich materials. The inertness of Li2MnO3 is attributed to its high crystallinity and few structural defects, which make it difficult to activate. Hence, the final products demonstrate a favorable electrochemical performance with appropriate proportions of two phases in the composite, as high contents of inert Li2MnO3 lower the capacity, while a sufficient structural stability cannot be achieved with low contents. The findings indicate that controlling the composition through a dynamic hydrothermal route is an effective strategy for developing a Mn-based cathode material for lithium-ion batteries. Full article
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<p>The XRD patterns (<b>a</b>,<b>b</b>) and ratio of the two phases (<b>c</b>) of the four synthesized samples.</p> Full article ">Figure 2
<p>The XPS spectra of the various samples (<b>a</b>–<b>d</b>).</p> Full article ">Figure 3
<p>The SEM images of the various samples (<b>a</b>–<b>d</b>), TEM image of LMO-2 (<b>e</b>), and particle size distribution of the samples (<b>f</b>–<b>i</b>).</p> Full article ">Figure 4
<p>Cyclic voltammograms (CVs) of the various samples in the potential range of 2.0–4.8 V (<b>a</b>–<b>d</b>).</p> Full article ">Figure 5
<p>The charge/discharge profiles of the various samples in the first cycle (<b>a</b>) and the corresponding dQ/dV curves (<b>b</b>).</p> Full article ">Figure 6
<p>Cycling performance at 0.1 C with the 2.0–4.5 V range (<b>a</b>), 0.1 C with the 2.0–4.8 V range (<b>b</b>), and 1 C with the 2.0–4.8 V range (<b>c</b>).</p> Full article ">Figure 7
<p>The charging–discharging curves of various samples in the 50th cycle at potential range of 2.0−4.8 V (<b>a</b>–<b>d</b>).</p> Full article ">Figure 8
<p>The XRD pattern of the discharged LMO-2 cathode disk after 15 cycles.</p> Full article ">Figure 9
<p>Electrochemical impedance spectra (EIS) of the various samples before (<b>a</b>) and after 15 cycles (<b>b</b>,<b>c</b>).</p> Full article ">
9 pages, 1452 KiB  
Article
Debris Removal Using a Hydroxyapatite Nanoparticle-Containing Solution (Vector Polish) with Sonic or Ultrasonic Agitation
by Michael Hülsmann, Christoph Beckmann and Steffi Baxter
Materials 2021, 14(16), 4750; https://doi.org/10.3390/ma14164750 - 23 Aug 2021
Cited by 1 | Viewed by 2244
Abstract
Chemomechanical preparation of the root canal system is considered to be the most important part of root canal treatment, including both mechanical removal of tissue remnants and dentine chips, and chemical elimination of biofilm and microorganisms. A number of different solutions and agitation [...] Read more.
Chemomechanical preparation of the root canal system is considered to be the most important part of root canal treatment, including both mechanical removal of tissue remnants and dentine chips, and chemical elimination of biofilm and microorganisms. A number of different solutions and agitation techniques have been proposed for that purpose. It was the aim of the present study to investigate whether root canal cleanliness can be improved by using a hydroxyapatite nanoparticle-containing solution with and without sonic or ultrasonic agitation. Seventy-four single-rooted teeth were divided into four experimental groups (n = 15) and two control groups (n = 7). All teeth were split longitudinally and a groove and three holes were cut into the root canal wall and filled with dentinal debris. Final irrigation was performed using sodium hypochlorite or a hydroxyapatite nanoparticle-containing solution (Vector polish) activated with a sonically or an ultrasonically driven endodontic file. Two calibrated investigators rated the remaining debris using a four-score scale. The results were analyzed using a non-parametric test with α < 0.05. Sonic and ultrasonic irrigation with sodium hypochlorite cleaned the grooves and holes well from debris. The hydroxyapatite nanoparticles activated by a sonic file cleaned grooves and holes equally well. Ultrasonically activated nanoparticles performance was clearly inferior. The syringe control-group left large amounts of debris in grooves and holes. The use of the hydroxyapatite nanoparticles used in this study did not improve removal of debris. Full article
(This article belongs to the Special Issue Endodontics)
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<p>(<b>a</b>) Groove prepared into one half of the split root; (<b>b</b>) Three holes prepared into the other half of the split root.</p> Full article ">Figure 2
<p>(<b>a</b>) Score 0: 0–25% of the holes filled with debris; (<b>b</b>) Score 1: 26–50% of the holes filled with debris; (<b>c</b>) Score 2: 51–75% of the holes filled with debris; (<b>d</b>) Score 3: 76–100% of the holes filled with debris.</p> Full article ">Figure 3
<p>(<b>a</b>): Score 0: 0–25% of the groove filled with debris; (<b>b</b>) Score 1: 26–50% of the groove filled with debris; (<b>c</b>): Score 2: 51–75% of the groove filled with debris; (<b>d</b>) Score 3: 76–100 of the groove filled with debris.</p> Full article ">Figure 3 Cont.
<p>(<b>a</b>): Score 0: 0–25% of the groove filled with debris; (<b>b</b>) Score 1: 26–50% of the groove filled with debris; (<b>c</b>): Score 2: 51–75% of the groove filled with debris; (<b>d</b>) Score 3: 76–100 of the groove filled with debris.</p> Full article ">
13 pages, 2975 KiB  
Article
Monitoring Electrical Biasing of Pb(Zr0.2Ti0.8)O3 Ferroelectric Thin Films In Situ by DPC-STEM Imaging
by Alexander Vogel, Martin F. Sarott, Marco Campanini, Morgan Trassin and Marta D. Rossell
Materials 2021, 14(16), 4749; https://doi.org/10.3390/ma14164749 - 23 Aug 2021
Cited by 5 | Viewed by 3975
Abstract
Increased data storage densities are required for the next generation of nonvolatile random access memories and data storage devices based on ferroelectric materials. Yet, with intensified miniaturization, these devices face a loss of their ferroelectric properties. Therefore, a full microscopic understanding of the [...] Read more.
Increased data storage densities are required for the next generation of nonvolatile random access memories and data storage devices based on ferroelectric materials. Yet, with intensified miniaturization, these devices face a loss of their ferroelectric properties. Therefore, a full microscopic understanding of the impact of the nanoscale defects on the ferroelectric switching dynamics is crucial. However, collecting real-time data at the atomic and nanoscale remains very challenging. In this work, we explore the ferroelectric response of a Pb(Zr0.2Ti0.8)O3 thin film ferroelectric capacitor to electrical biasing in situ in the transmission electron microscope. Using a combination of high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) and differential phase contrast (DPC)-STEM imaging we unveil the structural and polarization state of the ferroelectric thin film, integrated into a capacitor architecture, before and during biasing. Thus, we can correlate real-time changes in the DPC signal with the presence of misfit dislocations and ferroelastic domains. A reduction in the domain wall velocity of 24% is measured in defective regions of the film when compared to predominantly defect-free regions. Full article
(This article belongs to the Special Issue Materials Characterizations Using In-Situ Techniques)
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<p>(<b>a</b>) False colored scanning electron microscopy (SEM) image and (<b>b</b>) false colored high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) image of a lamella for in situ biasing prepared by focused ion beam (FIB). Yellow: Pt electrodes of the microelectromechanical systems (MEMS) chip. Orange: FIB-deposited Pt. Green: Si-MEMS chip, Blue: Nb-doped SrTiO<sub>3</sub> (Nb:STO) substrate. Red: Pb(Zr<sub>0.2</sub>Ti<sub>0.8</sub>)O<sub>3</sub> (PZT) thin film. The scale bars are 5 µm and 2 µm, respectively. (<b>c</b>) Applied voltage and measured current profiles during the in situ biasing experiments.</p> Full article ">Figure 2
<p>(<b>a</b>) Schematic representation of the basic principle of differential phase contrast (DPC)-STEM imaging. (<b>Left</b>) Electrons passing through vacuum. (<b>Right</b>) The electron beam is deflected by an electric field represented by the red ar-row. (<b>b</b>) Differential signals between two opposing detector segments (A − C, B − D). (<b>c</b>) Vector map of the DPC signal (<b>left</b>) and magnitude of the DPC signal (<b>right</b>). The scale bars are 50 nm.</p> Full article ">Figure 3
<p>(<b>a</b>) HAADF-STEM image of the PZT thin film grown on Nb:STO and Pt electrode. (<b>b</b>) Relative strain maps along x- and y-directions calculated by geometric phase analysis (GPA) of the HAADF-STEM image shown in (<b>a</b>). The substrate was used as zero strain reference. Area 1 highlights a row of misfit dislocations at the PZT/Nb:STO interface. (<b>c</b>) Atomic resolution HAADF-STEM image showing a row of dislocations at the interface and a section of a needle-like domain at the top left corner. Magnified views of the dislocations with the Burgers circuits and resulting Burgers vector are shown at the bottom. (<b>d</b>) Polarization map superimposed to the ferroelastic needle-like domain showing the polarization rotation, calculated from the atomic displacements of the Ti ions with respect to the 4 neighboring Pb ions. The color and size of the arrows correspond to the direction and magnitude of the polarization, respectively. (<b>e</b>) Map of the polarization direction of the region comprising the two dislocations on the left side (blue rectangle). A region near the bottom interface with reversed (downwards) polarization is clearly observed in blue color. Scale bars in (<b>a</b>,<b>b</b>) are 50 nm and in (<b>c</b>) 5 nm, respectively.</p> Full article ">Figure 4
<p>(<b>a</b>) DPC vector color plots of the electric field of the PZT film at 5 V (used as a reference), 0.8 V and −0.4 V, where the hue gives the field direction and the saturation is proportional to the vector modulus. Note the difficulty of detecting the changes in the polarization state of the film in the DPC vector plots. (<b>b</b>) Phase difference map obtained by subtracting the 5 V and 0.8 V DPC vector plots. (<b>c</b>) Phase difference map obtained by subtracting the 5 V and −0.4 V vector plots. The line profiles of the difference maps are extracted over the full image widths. The scale bars are 50 nm.</p> Full article ">Figure 5
<p>(<b>a</b>) In-plane ε<sub>xx</sub> strain map with an overlay of the estimated position of the switching front throughout the experiment. The red and blue rectangles correspond to the regions shown in (<b>b</b>,<b>d</b>), respectively. The scale bar is 50 nm. (<b>b</b>,<b>d</b>) Slices of the phase change plots at different voltages, showing the progression of the switching front in a predominantly defect free region (<b>b</b>) and a region enclosing a ferroelastic needle-like domain (<b>d</b>). (<b>c</b>,<b>e</b>) Line profiles of the average phase change extracted along horizontal direction in the regions shown in (<b>b</b>,<b>d</b>), respectively.</p> Full article ">
14 pages, 2129 KiB  
Article
Evolution of Structural and Magnetic Properties of Fe-Co Wire-like Nanochains Caused by Annealing Atmosphere
by Marcin Krajewski, Mateusz Tokarczyk, Sabina Lewińska, Katarzyna Brzózka, Kamil Bochenek and Anna Ślawska-Waniewska
Materials 2021, 14(16), 4748; https://doi.org/10.3390/ma14164748 - 23 Aug 2021
Cited by 1 | Viewed by 2172
Abstract
Thermal treatment is a post-synthesis treatment that aims to improve the crystallinity and interrelated physical properties of as-prepared materials. This process may also cause some unwanted changes in materials like their oxidation or contamination. In this work, we present the post-synthesis annealing treatments [...] Read more.
Thermal treatment is a post-synthesis treatment that aims to improve the crystallinity and interrelated physical properties of as-prepared materials. This process may also cause some unwanted changes in materials like their oxidation or contamination. In this work, we present the post-synthesis annealing treatments of the amorphous Fe1−xCox (x = 0.25; 0.50; 0.75) Wire-like nanochains performed at 400 °C in two different atmospheres, i.e., a mixture of 80% nitrogen and 20% hydrogen and argon. These processes caused significantly different changes of structural and magnetic properties of the initially-formed Fe-Co nanostructures. All of them crystallized and their cores were composed of body-centered cubic Fe-Co phase, whereas their oxide shells comprised of a mixture of CoFe2O4 and Fe3O4 phases. However, the annealing carried out in hydrogen-containing atmosphere caused a decomposition of the initial oxide shell layer, whereas a similar process in argon led to its slight thickening. Moreover, it was found that the cores of thermally-treated Fe0.25Co0.75 nanochains contained the hexagonal closest packed (hcp) Co phase and were covered by the nanosheet-like shell layer in the case of annealing performed in argon. Considering the evolution of magnetic properties induced by structural changes, it was observed that the coercivities of annealed Fe-Co nanochains increased in comparison with their non-annealed counterparts. The saturation magnetization (MS) of the Fe0.25Co0.75 nanomaterial annealed in both atmospheres was higher than that for the non-annealed sample. In turn, the MS of the Fe0.75Co0.25 and Fe0.50Co0.50 nanochains annealed in argon were lower than those recorded for non-annealed samples due to their partial oxidation during thermal processing. Full article
(This article belongs to the Special Issue Research in Structural and Magnetic Properties of Ferromagnetic Materials)
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<p>SEM images of as-prepared (<b>a</b>) Fe<sub>0.75</sub>Co<sub>0.25</sub>, (<b>b</b>) Fe<sub>0.50</sub>Co<sub>0.50</sub>, and (<b>c</b>) Fe<sub>0.25</sub>Co<sub>0.75</sub> Wire-like nanochains; SEM images of (<b>d</b>) Fe<sub>0.75</sub>Co<sub>0.25</sub>, (<b>e</b>) Fe<sub>0.50</sub>Co<sub>0.50</sub>, and (<b>f</b>) Fe<sub>0.25</sub>Co<sub>0.75</sub> Wire-like nanochains annealed in hydrogen-containing atmosphere; SEM images of (<b>g</b>) Fe<sub>0.75</sub>Co<sub>0.25</sub>, (<b>h</b>) Fe<sub>0.50</sub>Co<sub>0.50</sub>, and (<b>i</b>) Fe<sub>0.25</sub>Co<sub>0.75</sub> Wire-like nanochains annealed in argon.</p> Full article ">Figure 2
<p>TEM images of as-prepared (<b>a</b>) Fe<sub>0.75</sub>Co<sub>0.25</sub>, (<b>b</b>) Fe<sub>0.50</sub>Co<sub>0.50</sub>, and (<b>c</b>) Fe<sub>0.25</sub>Co<sub>0.75</sub> Wire-like nanochains; TEM images of (<b>d</b>) Fe<sub>0.75</sub>Co<sub>0.25</sub>, (<b>e</b>) Fe<sub>0.50</sub>Co<sub>0.50</sub>, and (<b>f</b>) Fe<sub>0.25</sub>Co<sub>0.75</sub> Wire-like nanochains annealed in hydrogen-containing atmosphere; TEM images of (<b>g</b>) Fe<sub>0.75</sub>Co<sub>0.25</sub>, (<b>h</b>) Fe<sub>0.50</sub>Co<sub>0.50</sub>, and (<b>i</b>) Fe<sub>0.25</sub>Co<sub>0.75</sub> Wire-like nanochains annealed in argon.</p> Full article ">Figure 3
<p>XRD patterns of as-prepared and annealed (<b>a</b>) Fe<sub>0.75</sub>Co<sub>0.25</sub>, (<b>b</b>) Fe<sub>0.50</sub>Co<sub>0.50</sub>, and (<b>c</b>) Fe<sub>0.25</sub>Co<sub>0.75</sub> Wire-like nanochains; the bottom panels indicate the reference positions of the Bragg peaks based on the JCPDS database for the crystalline bcc Fe-Co, Fe<sub>3</sub>O<sub>4</sub>, CoFe<sub>2</sub>O<sub>4</sub>, and hcp Co phases; the insets present the magnified view of XRD patterns associated with oxides in the range between 30 and 80°.</p> Full article ">Figure 4
<p>(<b>a</b>) Room temperature Mössbauer spectra collected for non-annealed Fe-Co nanochains, and (<b>b</b>) their corresponding hyperfine magnetic field distributions; component 1—green line, component 2—red line, component 3—violet and blue lines.</p> Full article ">Figure 5
<p>Room temperature Mössbauer spectra collected for the Fe-Co nanochains annealed in (<b>a</b>) hydrogen-containing atmosphere, and (<b>b</b>) argon atmosphere; component I—blue lines, component II—red line, component III—orange line, and component IV—olive and grey lines.</p> Full article ">Figure 6
<p>Room temperature magnetic hysteresis loops of as-prepared and annealed (<b>a</b>) Fe<sub>0.75</sub>Co<sub>0.25,</sub> (<b>b</b>) Fe<sub>0.50</sub>Co<sub>0.50</sub>, and (<b>c</b>) Fe<sub>0.25</sub>Co<sub>0.75</sub> Wire-like nanochains.</p> Full article ">
15 pages, 6512 KiB  
Article
Zinc Oxide Nanoparticles for Water Purification
by Angela Spoială, Cornelia-Ioana Ilie, Roxana-Doina Trușcă, Ovidiu-Cristian Oprea, Vasile-Adrian Surdu, Bogdan Ștefan Vasile, Anton Ficai, Denisa Ficai, Ecaterina Andronescu and Lia-Mara Dițu
Materials 2021, 14(16), 4747; https://doi.org/10.3390/ma14164747 - 23 Aug 2021
Cited by 59 | Viewed by 6114
Abstract
In this study, zinc oxide nanoparticles were synthesized through a simple co-precipitation method starting from zinc acetate dihydrate and sodium hydroxide as reactants. The as-obtained ZnO nanoparticles were morphologically and structurally characterized by Fourier transform infrared spectroscopy (FTIR), X-ray diffraction (XRD), scanning electron [...] Read more.
In this study, zinc oxide nanoparticles were synthesized through a simple co-precipitation method starting from zinc acetate dihydrate and sodium hydroxide as reactants. The as-obtained ZnO nanoparticles were morphologically and structurally characterized by Fourier transform infrared spectroscopy (FTIR), X-ray diffraction (XRD), scanning electron microscopy (SEM), transmission electron microscopy (TEM), photocatalytic activity, and by determining the antimicrobial activity against Gram-negative and Gram-positive bacteria. The XRD pattern of the zinc oxide nanoparticles showed the wurtzite hexagonal structure, and its purity highlighted that the crystallinity correlated with the presence of a single product, zinc oxide. The ZnO nanoparticles have an average crystallite size of 19 ± 11 nm, which is in accordance with the microscopic data. ZnO nanoparticles were tested against methyl orange, used as a model pollutant, and it was found that they exhibit strong photocatalytic activity against this dye. The antibacterial activity of ZnO nanoparticles was tested against Gram-negative and Gram-positive strains (Escherichia coli, Staphylococcus aureus, and Candida albicans). The strongest activity was found against Gram-positive bacteria (S. aureus). Full article
(This article belongs to the Special Issue Research of Photocatalytic, Antibacterial and Biocompatible Properties of Nanoparticles)
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<p>Thermal analysis of ZnO powder.</p> Full article ">Figure 2
<p>FTIR 3D chromatogram of evolved gases.</p> Full article ">Figure 3
<p>FTIR spectra of ZnO nanoparticles.</p> Full article ">Figure 4
<p>XRD pattern for ZnO nanoparticles.</p> Full article ">Figure 5
<p>SEM images for ZnO powder, (<b>a</b>) 5000× magnification, (<b>b</b>) 10,000× magnification, (<b>c</b>) 100,000× magnification, (<b>d</b>) 200,000× magnification.</p> Full article ">Figure 6
<p>TEM images for ZnO nanoparticles, (<b>a</b>) scale bar 50 nm, (<b>b</b>) scale bar 100 nm, (<b>c</b>) scale bar 200 nm, (<b>d</b>) scale bar 500 nm.</p> Full article ">Figure 7
<p>Photocatalytic activity against methyl orange (MO) for ZnO nanoparticles.</p> Full article ">Figure 8
<p>The pseudo-first-order rate constant <span class="html-italic">k</span><sub>app</sub> (min<sup>−1</sup>) was calculated from the slope of ln(C0/C) versus irradiation time <span class="html-italic">t</span>.</p> Full article ">Figure 9
<p>Absorbance values for evaluating anti-adhesion activity for ZnO nanoparticles against <span class="html-italic">S. aureus</span>. * indicates the MBEC values and differences between groups are considered statistically significant (<span class="html-italic">p</span> &lt; 0.001).</p> Full article ">Figure 10
<p>Absorbance values for evaluating anti-adhesion activity for ZnO nanoparticles against <span class="html-italic">E. coli</span>. * indicates the MBEC values and differences between groups are considered statistically significant (<span class="html-italic">p</span> &lt; 0.001).</p> Full article ">Figure 11
<p>Absorbance values for evaluation of anti-adhesion activity for ZnO nanoparticles against <span class="html-italic">C. albicans</span>. * indicates the MBEC values and differences between groups are considered statistically significant (<span class="html-italic">p</span> &lt; 0.001).</p> Full article ">
16 pages, 5941 KiB  
Article
Bonding Performance of Universal Adhesives Applied to Nano-Hydroxyapatite Desensitized Dentin Using Etch-and-Rinse or Self-Etch Mode
by Yuchen Meng, Fan Huang, Silin Wang, Meiwen Li, Yi Lu, Dandan Pei and Ang Li
Materials 2021, 14(16), 4746; https://doi.org/10.3390/ma14164746 - 22 Aug 2021
Cited by 9 | Viewed by 3758
Abstract
The study assessed the bonding performance of three universal adhesives on desensitized dentin with etch-and-rinse mode or self-etch mode after nano-hydroxyapatite (nHAp)-based desensitizers application. Simulated sensitive dentin specimens were prepared and separated into four groups: no treatment as the negative control, groups desensitized [...] Read more.
The study assessed the bonding performance of three universal adhesives on desensitized dentin with etch-and-rinse mode or self-etch mode after nano-hydroxyapatite (nHAp)-based desensitizers application. Simulated sensitive dentin specimens were prepared and separated into four groups: no treatment as the negative control, groups desensitized by Biorepair toothpaste, Dontodent toothpaste, or nHAp paste. Three universal adhesives of All-Bond Universal, Single Bond Universal, and Clearfil Universal Bond with etch-and-rinse or self-etch mode were bonded to the desensitized dentin specimens separately, followed by resin composite build-ups. Micro-tensile bond strength was measured using a micro-tensile tester. The wettability of desensitized dentin was evaluated by the contact angle of the adhesives. Resin infiltration was observed by confocal laser scanning microscopy. Dentin tubular occlusion and nanoleakage were observed by scanning electron microscope. The results showed that the etch-and-rinse or self-etch mode of each adhesive showed similar bond strength when bonding to nHAp-based desensitized dentin. The dentin surface was partially covered by desensitizers after desensitization. Compared with the self-etch mode, stronger demineralization and more reopened dentin tubules were observed in the etch-and-rinse mode after acid etching; longer resin tags and more nanoleakage in the resin–dentin interface were observed when using the etch-and-rinse mode. When bonding to nHAp-based desensitized dentin with universal adhesives, no significant difference in bond strength was found between self-etch mode or etch-and-rinse mode; while the latter produced more nanoleakage in the resin–dentin interfaces. Full article
(This article belongs to the Special Issue Adhesives and Composite Resins in Contemporary Restorative Dentistry: From the Lab to the Clinic)
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Figure 1
<p>Schematic illustration of the study design. Mid-coronal and disc specimens were prepared for sensitive models and applied by desensitizing treatments. Mid-coronal specimens were bonded with universal adhesives, restored by resin composite, then sectioned into beams and slabs. μTBS were tested on bonded beams. Adhesive interfaces were observed on bonded slabs by CLSM and SEM. Dentin tubular occlusion before and after adhesive application were analyzed by SEM. Wettability of universal adhesives was measured by the contact angle.</p> Full article ">Figure 2
<p>Micro-tensile bond strength and failure mode. (<b>a</b>) Effects of desensitizer types and bonding modes on bond strength of universal adhesives. Different superscript uppercase letters indicate significant differences with etch-and-rinse mode (<span class="html-italic">p</span> &lt; 0.05). Different superscript lowercase letters indicate significant differences with self-etch mode (<span class="html-italic">p</span> &lt; 0.05). * denotes a significant difference between the etch-and-rinse mode and self-etch mode (<span class="html-italic">p</span> &lt; 0.05). (<b>b</b>) Failure mode distribution of fracture specimens. M: mixed failure; A: adhesive failure; CD: cohesive failure within the dentin; and CC: cohesive failure within the composite resin.</p> Full article ">Figure 3
<p>General condition graphs (×90) and micrographs (×2000) of adhesive and mixed failure on dentin interfaces. (<b>a</b>,<b>b</b>) Mixed failure from control group bonded with etch-and-rinse mode; (<b>c</b>,<b>d</b>) adhesive failure from desensitizing groups bonded with self-etch mode; (<b>e</b>,<b>f</b>) typical adhesive failure; and (<b>g</b>,<b>h</b>) typical mixed failure. Circle: open dentin tubules; triangle: sealed dentin tubules; and solid arrowhead: resin tags.</p> Full article ">Figure 4
<p>SEM image observation and EDX spectrum of dentin surface after desensitization treatment (×2000 and ×5000).</p> Full article ">Figure 5
<p>Morphological SEM images of desensitized dentin surfaces treated by universal adhesives with etch-and-rinse mode or self-etch mode. (<b>a</b>) All-Bond Universal (ABU). (<b>b</b>) Single Bond Universal (SBU). (<b>c</b>) Clearfil Universal Bond (CUB). ER: etch-and-rinse mode; SE: self-etch mode.</p> Full article ">Figure 6
<p>Contact angle values and anterior view of universal adhesives on dentin substrates after desensitization. (<b>a</b>) All-Bond Universal. (<b>b</b>) Single Bond Universal. (<b>c</b>) Clearfil Universal Bond. * denotes a significant difference between the etch-and-rinse mode and self-etch mode (<span class="html-italic">p</span> &lt; 0.05).</p> Full article ">Figure 7
<p>Adhesive resin infiltration observation in etch-and-rinse mode or self-etch mode. (<b>a</b>) All-Bond Universal (ABU). (<b>b</b>) Single Bond Universal (SBU). (<b>c</b>) Clearfil Universal Bond (CUB). Long homogeneous resin tags formed in etch-and-rinse mode and short resin tags formed in self-etch mode. ER: etch-and-rinse mode; SE: self-etch mode.</p> Full article ">Figure 8
<p>Resin tags’ length of universal adhesives in etch-and-rinse mode or self-etch mode. (<b>a</b>) All-Bond Universal. (<b>b</b>) Single Bond Universal. (<b>c</b>) Clearfil Universal Bond. Different superscript uppercase letters indicate significant differences in etch-and-rinse mode (<span class="html-italic">p</span> &lt; 0.05). Different superscript lowercase letters indicate significant differences in self-etch mode (<span class="html-italic">p</span> &lt; 0.05). * denotes a significant difference between the etch-and-rinse mode and self-etch mode (<span class="html-italic">p</span> &lt; 0.05).</p> Full article ">Figure 9
<p>Nanoleakage within the resin–dentin bonding interfaces in etch-and-rinse mode or self-etch mode. (<b>a</b>) All-Bond Universal (ABU). (<b>b</b>) Single Bond Universal (SBU). (<b>c</b>) Clearfil Universal Bond (CUB). A silver-free zone distributed along adhesive interfaces in self-etch mode. Spotted silver deposits were restricted to the hybrid layer of etch-and-rinse mode. Solid arrowhead: nanoleakage indicated by silver deposits. ER: etch-and-rinse mode; SE: self-etch mode.</p> Full article ">
26 pages, 2627 KiB  
Review
Polymer Geogrids: A Review of Material, Design and Structure Relationships
by Mohammad Al-Barqawi, Rawan Aqel, Mark Wayne, Hani Titi and Rani Elhajjar
Materials 2021, 14(16), 4745; https://doi.org/10.3390/ma14164745 - 22 Aug 2021
Cited by 26 | Viewed by 8753
Abstract
Geogrids are a class of geosynthetic materials made of polymer materials with widespread transportation, infrastructure, and structural applications. Geogrids are now routinely used in soil stabilization applications ranging from reinforcing walls to soil reinforcement below grade or embankments with increased potential for remote-sensing [...] Read more.
Geogrids are a class of geosynthetic materials made of polymer materials with widespread transportation, infrastructure, and structural applications. Geogrids are now routinely used in soil stabilization applications ranging from reinforcing walls to soil reinforcement below grade or embankments with increased potential for remote-sensing applications. Developments in manufacturing procedures have allowed new geogrid designs to be fabricated in various forms of uniaxial, biaxial, and triaxial configurations. The design flexibility allows deployments based on the load-carrying capacity desired, where biaxial geogrids may be incorporated when loads are applied in both the principal directions. On the other hand, uniaxial geogrids provide higher strength in one direction and are used for mechanically stabilized earth walls. More recently, triaxial geogrids that offer a more quasi-isotropic load capacity in multiple directions have been proposed for base course reinforcement. The variety of structures, polymers, and the geometry of the geogrid materials provide engineers and designers many options for new applications. Still, they also create complexity in terms of selection, characterization, and long-term durability. In this review, advances and current understanding of geogrid materials and their applications to date are presented. A critical analysis of the various geogrid systems, their physical and chemical characteristics are presented with an eye on how these properties impact the short- and long-term properties. The review investigates the approaches to mechanical behavior characterization and how computational methods have been more recently applied to advance our understanding of how these materials perform in the field. Finally, recent applications are presented for remote sensing sub-grade conditions and incorporation of geogrids in composite materials. Full article
(This article belongs to the Special Issue Feature Collection in Advanced Composites Section)
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<p>Variety of geogrid materials being produced.</p> Full article ">Figure 2
<p>Ultimate tensile strength (kN/m) of tested geogrid specimens around 360 degrees loading directions, ref. [<a href="#B31-materials-14-04745" class="html-bibr">31</a>].</p> Full article ">Figure 3
<p>Creep strain rate vs. applied load for high-density polyethylene (HDPE) and polyethylene terephthalate (PET) geogrids [<a href="#B43-materials-14-04745" class="html-bibr">43</a>].</p> Full article ">Figure 4
<p>Manufacturing defects in geogrids: (<b>a</b>) Variability in aperture sizes, (<b>b</b>) bent ribs or variation in aperture shape, (<b>c</b>) splitting in geogrid ribs.</p> Full article ">Figure 5
<p>(<b>a</b>) Effect of normal stress on the: (<b>a</b>) tensile strength of geogrid (<b>b</b>) secant tensile stiffness of geogrid [<a href="#B85-materials-14-04745" class="html-bibr">85</a>].</p> Full article ">Figure 6
<p>Geogrid components showing stress results from digital image correlation from ref. [<a href="#B101-materials-14-04745" class="html-bibr">101</a>].</p> Full article ">Figure 7
<p>Strain evolution at various levels with the vertical bold dashed line representing the end of construction (<b>a</b>–<b>c</b>). In (<b>d</b>) is the accumulated strain inside the reinforced structure at different levels (Creative Commons CC BY, ref. [<a href="#B116-materials-14-04745" class="html-bibr">116</a>]).</p> Full article ">
18 pages, 3112 KiB  
Article
Operational Variables on the Processing of Porous Titanium Bodies by Gelation of Slurries with an Expansive Porogen
by Antonio Javier Sanchez-Herencia, Zoilo Gonzalez, Alejandro Rodriguez, Esther Molero and Begoña Ferrari
Materials 2021, 14(16), 4744; https://doi.org/10.3390/ma14164744 - 22 Aug 2021
Viewed by 2565
Abstract
Colloidal processing techniques, based on the suspension of powders in a liquid, are very versatile techniques to fabricate porous structures. They can provide customized pores, shapes and surfaces through the control of operational parameters, being the base of the alternative additive manufacture processes. [...] Read more.
Colloidal processing techniques, based on the suspension of powders in a liquid, are very versatile techniques to fabricate porous structures. They can provide customized pores, shapes and surfaces through the control of operational parameters, being the base of the alternative additive manufacture processes. In this work disperse and stable titanium aqueous slurries has been formulated in order to process porous materials by the incorporation of methylcellulose (MC) as a gelation agent and ammonium bicarbonate as an expansive porogen. After casting the slurries and heating at mild temperatures (60–80 °C) the methylcellulose gels and traps the gas bubbles generated by the ammonium bicarbonate decomposition to finally obtain stiff porous green structures. Using an experimental design method, the influence of the temperature as well as the concentration of gelation agent and porogen on the viscosity, apparent density and pore size distribution is analyzed by a second-order polynomial function in order to identifying the influence of the operating variables in the green titanium porous compact. After sintering at 1100 °C under high vacuum, titanium sponges with 39% of open porosity and almost no close porosity were obtained. Full article
(This article belongs to the Section Porous Materials)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>Scheme of Powder Metallurgy techniques for porous metal processing.</p> Full article ">Figure 2
<p>(<b>a</b>) Flow curves of Ti slurries with 50 vol. % of solid contents formulated without MC and with 8, 10 and 12 g/L of MC. (<b>b</b>) Viscosity values vs. Temperature for 50 vol.% slurries with 8, 10 and 12 g/L of MC.</p> Full article ">Figure 3
<p>Porous Ti green sample, (<b>left</b>) dome shaped by gelation and drying and (<b>right</b>) a detail of the internal porous structure observed after breaking apart the sample.</p> Full article ">Figure 4
<p>Evolution of the viscosity with the shear rate of 50 vol.% Ti slurries with 15, 20 and 25 wt.% of ammonium bicarbonate and 12 g/L of MC.</p> Full article ">Figure 5
<p>Pore distributions by Mercury Intrusion Porosimetry (MIP), for samples in <a href="#app1-materials-14-04744" class="html-app">Table S1</a>. Grouping operational variables by fixing X<sub>G</sub> and X<sub>P</sub> and variable X<sub>T</sub>: (<b>a</b>) sample 9 (8X<sub>G</sub>−15X<sub>P</sub>−60X<sub>T</sub>) and sample 7 (8X<sub>G</sub>−15X<sub>P</sub>−80X<sub>T</sub>); (<b>b</b>) sample 5 (8X<sub>G</sub>−25X<sub>P</sub>−60X<sub>T</sub>) and sample 3 (8X<sub>G</sub>−25X<sub>P</sub>−80X<sub>T</sub>); (<b>c</b>) sample 8 (12X<sub>G</sub>−15X<sub>P</sub>−60X<sub>T</sub>) and sample 6 (12X<sub>G</sub>−15X<sub>P</sub>−80X<sub>T</sub>); (<b>d</b>) sample 4 (12X<sub>G</sub>−25X<sub>P</sub>−60X<sub>T</sub>) and sample 2 (12X<sub>G</sub>−25X<sub>P</sub>−80X<sub>T</sub>).</p> Full article ">Figure 6
<p>Force-strain curves recorded in the compression tests of green samples: (<b>a</b>) 7 (8X<sub>G</sub>−15X<sub>P</sub>−80X<sub>T</sub>) and 3 (8X<sub>G</sub>−25X<sub>P</sub>−80X<sub>T</sub>); (<b>b</b>) samples 6 (12X<sub>G</sub>−15X<sub>P</sub>−80X<sub>T</sub>) and 2 (12X<sub>G</sub>−25X<sub>P</sub>−80X<sub>T</sub>) in <a href="#app1-materials-14-04744" class="html-app">Table S1</a>.</p> Full article ">Figure 7
<p>Pictures and micrographs of the porous titanium sintered bodies representing the different processing categories (<b>A</b>–<b>C</b>).</p> Full article ">
16 pages, 4069 KiB  
Article
Good Choice of Electrode Material as the Key to Creating Electrochemical Sensors—Characteristics of Carbon Materials and Transparent Conductive Oxides (TCO)
by Anna Cirocka, Dorota Zarzeczańska and Anna Wcisło
Materials 2021, 14(16), 4743; https://doi.org/10.3390/ma14164743 - 22 Aug 2021
Cited by 19 | Viewed by 4224
Abstract
The search for new electrode materials has become one of the goals of modern electrochemistry. Obtaining electrodes with optimal properties gives a product with a wide application potential, both in analytics and various industries. The aim of this study was to select, from [...] Read more.
The search for new electrode materials has become one of the goals of modern electrochemistry. Obtaining electrodes with optimal properties gives a product with a wide application potential, both in analytics and various industries. The aim of this study was to select, from among the presented electrode materials (carbon and oxide), the one whose parameters will be optimal in the context of using them to create sensors. Electrochemical impedance spectroscopy and cyclic voltammetry techniques were used to determine the electrochemical properties of the materials. On the other hand, properties such as hydrophilicity/hydrophobicity and their topological structure were determined using contact angle measurements and confocal microscopy, respectively. Based on the research carried out on a wide group of electrode materials, it was found that transparent conductive oxides of the FTO (fluorine doped tin oxide) type exhibit optimal electrochemical parameters and offer great modification possibilities. These electrodes are characterized by a wide range of work and high chemical stability. In addition, the presence of a transparent oxide layer allows for the preservation of valuable optoelectronic properties. An important feature is also the high sensitivity of these electrodes compared to other tested materials. The combination of these properties made FTO electrodes selected for further research. Full article
(This article belongs to the Special Issue Advanced Electrode Materials Dedicated for Electroanalysis)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>Schematic representation of the range of electrochemical stability of the tested group of electrodes in 0.5 M Na<sub>2</sub>SO<sub>4</sub> together with the designated area in which water electrolysis takes place (gray field).</p> Full article ">Figure 2
<p>Comparison of cyclic voltammograms of the reference redox systems (Fe(CN)<sub>6</sub>)<sup>3−/4−</sup> and H<sub>2</sub>Q/Q in an aqueous solution of Na<sub>2</sub>SO<sub>4</sub> (0.5 M) registered on the following electrodes: glassy carbon (GC), boron doped diamond on silicone (Si/BDD), carbon nanowalls on silicone (Si/CNW) and fluorine doped tin oxide (FTO). Scan rate: 0.1 Vs<sup>−1</sup>.</p> Full article ">Figure 3
<p>Comparison of cyclic voltammograms of the model redox system (Fe(CN)<sub>6</sub>)<sup>3</sup><sup>−/4</sup><sup>−</sup> (5 mM) in an aqueous solution of Na<sub>2</sub>SO<sub>4</sub> (0.5 M) registered on B-NCD electrodes on silicon and glass substrate with different (B)/(C) ratios.</p> Full article ">Figure 4
<p>Pictures of contact angle measurement (WCA) of electrode materials (<b>a</b>) GC, Si/BDD, Si/CNW, and FTO; (<b>b</b>) boron-doped nanodiamonds (B-NCD) on glass and silicon substrates with different (B)/(C) ratios.</p> Full article ">Figure 5
<p>Free surface energy (SFE) graph and contact angles of tested conductive materials: (<b>a</b>) GC, Si/BDD, Si/CNW, Si/B-NCD-10k, Glass/B-NCD-10k, and FTO; (<b>b</b>) boron doped nanodiamonds (B-NCD) on glass and silicon substrates with different (B)/(C) ratios; γ<sup>S</sup>-SFE, γ<sup>D</sup>: disperse part, γ<sup>P</sup>: polar part, and WCA: water contact angle.</p> Full article ">Figure 6
<p>Graph of the water contact angle (WCA) and reversibility of the redox process for (Fe(CN)<sub>6</sub>)<sup>3</sup><sup>−/4</sup><sup>−</sup> couple (ΔE) on the type of electrode: (<b>a</b>) GC, Si/BDD, Si/CNW, Si/B- NCD-10k, Glass/B-NCD-10k, and FTO; (<b>b</b>) from different levels of doping B-NCD electrodes on a glass and silicon substrate.</p> Full article ">Figure 7
<p>Keyence VK-X1000 confocal laser microscope photos of the surface of electrode materials: GC, Si/CNW, and FTO.</p> Full article ">
11 pages, 2832 KiB  
Article
One-Pot Synthesis of Alumina-Titanium Diboride Composite Powder at Low Temperature
by Xueyin Liu, Ke Bao, Junfeng Chen, Quanli Jia and Shaowei Zhang
Materials 2021, 14(16), 4742; https://doi.org/10.3390/ma14164742 - 22 Aug 2021
Cited by 2 | Viewed by 2086
Abstract
Alumina-titanium diboride (Al2O3-TiB2) composite powders were synthesised via aluminothermic reduction of TiO2 and B2O3, mediated by a molten chloride salt (NaCl, KCl, or MgCl2). The effects of salt type, initial [...] Read more.
Alumina-titanium diboride (Al2O3-TiB2) composite powders were synthesised via aluminothermic reduction of TiO2 and B2O3, mediated by a molten chloride salt (NaCl, KCl, or MgCl2). The effects of salt type, initial batch composition, and firing temperature/time on the phase formation and overall reaction extent were examined. Based on the results and equilibrium thermodynamic calculations, the mechanisms underpinning the reaction/synthesis processes were clarified. Given their evaporation losses at test temperatures, appropriately excessive amounts of Al and B2O3 are needed to complete the synthesis reaction. Following this, phase-pure Al2O3-TiB2 composite powders composed of 0.3–0.6 μm Al2O3 and 30–60 nm TiB2 particles were successfully fabricated in NaCl after 5 h at 1050 °C. By increasing the firing temperature to 1150 °C, the time required to complete the synthesis reaction could be reduced to 4 h, although the sizes of Al2O3 and TiB2 particles in the resultant phase pure composite powder increased slightly to 1–2 μm and 100–200 nm, respectively. Full article
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<p>XRD patterns of stoichiometric samples after 4 h firing at 850 °C in: (<b>a</b>) KCl, (<b>b</b>) NaCl, and (<b>c</b>) MgCl<sub>2</sub>.</p> Full article ">Figure 2
<p>XRD patterns of stoichiometric samples after 4 h firing in NaCl at: (<b>a</b>) 850, (<b>b</b>) 950, (<b>c</b>) 1050, and (<b>d</b>) 1150 °C.</p> Full article ">Figure 3
<p>XRD patterns of samples using (<b>a</b>) 0, (<b>b</b>) 20, (<b>c</b>) 25, and (<b>d</b>) 30 wt% excess Al, after 4 h firing in NaCl at 1150 °C.</p> Full article ">Figure 4
<p>Influence of excess B<sub>2</sub>O<sub>3</sub> on phase formation in samples using 30 wt% excess Al, after 4 h firing in NaCl at 1150 °C: (<b>a</b>) 0 (stoichiometric), (<b>b</b>) 10, and (<b>c</b>) 20 wt% excess B<sub>2</sub>O<sub>3</sub>.</p> Full article ">Figure 5
<p>XRD patterns of samples using 30 wt% excess Al and 20 wt% excess B<sub>2</sub>O<sub>3</sub> after firing in NaCl at 1050 °C for: (<b>a</b>) 4, (<b>b</b>) 5, and (<b>c</b>) 6 h.</p> Full article ">Figure 6
<p>XRD patterns of samples after firing in NaCl at 1050 °C for 5 h using 30 wt% excess Al, and (<b>a</b>) 20, (<b>b</b>) 25, and (<b>c</b>) 30 wt% excess B<sub>2</sub>O<sub>3</sub>.</p> Full article ">Figure 7
<p>TEM images (<b>a</b>,<b>b</b>) and corresponding EDS (<b>c</b>,<b>d</b>) of the Al<sub>2</sub>O<sub>3</sub>-TiB<sub>2</sub> composite powder resultant from 5 h firing in NaCl at 1050 °C (the small C and Cu peaks arose from the carbon coating and the Cu grid).</p> Full article ">Figure 8
<p>SEM images (<b>a</b>,<b>b</b>) and corresponding EDS (<b>c</b>,<b>d</b>) of the Al<sub>2</sub>O<sub>3</sub>-TiB<sub>2</sub> composite powder resultant from 4 h firing in NaCl at 1150 °C.</p> Full article ">Figure 9
<p>The standard Gibbs energy values corresponding to Reactions (1) and (3)–(10), as a function of temperature.</p> Full article ">
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