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Ceramics
Volume 1
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Ceramics, Volume 1, Issue 2 (December 2018) – 13 articles

Cover Story (view full-size image): Without doubt, nanosized ceramic powders can have remarkable properties. However, transferring such structures into dense, bulk ceramic materials compatible with e.g. use in fuel cells is highly challenging. The problem, from a practical point of view, is that the subsequent annealing and sintering steps would destroy these nanostructures. Here we use a novel approach to overcome this obstacle. By using well-defined nanosized particles with a core–shell structure, the desired nanoscale microstructure is already defined in the reactant. In combination with spark plasma sintering, which enables very short sintering times at temperatures lower than in conventional sintering methods, it is possible to partly preserve the core–shell architecture of the primary powders after densification. View this paper.
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18 pages, 6864 KiB  
Article
Calcium Phosphate Powder Synthesized from Calcium Acetate and Ammonium Hydrophosphate for Bioceramics Application
by Tatiana Safronova, Valery Putlayev, Yaroslav Filippov, Tatiana Shatalova, Evgeny Karpushkin, Dmitrii Larionov, Gilyana Kazakova and Yury Shakhtarin
Ceramics 2018, 1(2), 375-392; https://doi.org/10.3390/ceramics1020030 - 15 Dec 2018
Cited by 8 | Viewed by 5559
Abstract
Calcium phosphate powder was synthesized at room temperature from aqueous solutions of ammonium hydrophosphate and calcium acetate without pH adjusting at constant Ca/P molar ratio 1.5. Phase composition of the as-synthesized powder depended on the precursors concentration: At 2.0 M of calcium acetate [...] Read more.
Calcium phosphate powder was synthesized at room temperature from aqueous solutions of ammonium hydrophosphate and calcium acetate without pH adjusting at constant Ca/P molar ratio 1.5. Phase composition of the as-synthesized powder depended on the precursors concentration: At 2.0 M of calcium acetate in the starting solution, poorly crystallized hydroxyapatite was formed, 0.125 M solution of calcium acetate afforded brushite, and the powders synthesized from 0.25–1.0 M calcium acetate solutions were mixtures of the mentioned phases. Firing at 1100 °C led to complete elimination of the reaction by-products, yet the phase composition of the annealed compacted samples was the following: When 2.0 M solution of calcium acetate was used, the obtained ceramics consisted of β-Ca3(PO4)2, whereas at 0.125 to 1.0 M of calcium acetate, the ceramics was a mixture of β-Ca3(PO4)2 and β-Ca2P2O7. Synthesized calcium phosphate powders can be used as the powdered precursors for biocompatible bioresorbable composite ceramics production. Full article
(This article belongs to the Special Issue Ceramics for Biomedical Applications)
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Figure 1
<p>Change in the reaction zone pH during addition of aqueous solution of (NH<sub>4</sub>)<sub>2</sub>HPO<sub>4</sub> to that of Ca(CH<sub>3</sub>COO)<sub>2</sub> to the final Ca/P ratio of 1.5. Concentrations of the Ca(CH<sub>3</sub>COO)<sub>2</sub> solutions are shown in the legend.</p> Full article ">Figure 2
<p>Calculated pH curves during addition of 0.083 and 0.667 M (NH<sub>4</sub>)<sub>2</sub>HPO<sub>4</sub> solutions to 100 mL of 0.125 and 1.0 M Ca(CH<sub>3</sub>COO)<sub>2</sub> solutions respectively (<b>a</b>) and concentrations of the crystalline phases corresponding to the “0.125 M_A,O,B” line (<b>b</b>). The allowed crystalline phases were hydroxyapatite Ca<sub>10</sub>(PO<sub>4</sub>)<sub>6</sub>(OH)<sub>2</sub> (H), amorphous calcium phosphates Ca<sub>3</sub>(PO<sub>4</sub>)<sub>2</sub> (A), octacalcium phosphate Ca<sub>8</sub>H<sub>2</sub>(PO<sub>4</sub>)<sub>6</sub>·5H<sub>2</sub>O (O), and brushite CaHPO<sub>4</sub>·2H<sub>2</sub>O (B).</p> Full article ">Figure 3
<p>XRD patterns of the powders synthesized from aqueous solutions of Ca(CH<sub>3</sub>COO)<sub>2</sub> of different concentrations (shown near the curves) and (NH<sub>4</sub>)<sub>2</sub>HPO<sub>4</sub> at Ca/P = 1.5. o—Ca<sub>10</sub>(PO<sub>4</sub>)<sub>6</sub>(OH)<sub>2</sub>, PDF card # 9-432; *—brushite CaHPO<sub>4</sub>·2H<sub>2</sub>O, PDF card # 9-77.</p> Full article ">Figure 4
<p>SEM images of the powders synthesized from aqueous solutions of Ca(CH<sub>3</sub>COO)<sub>2</sub> and (NH<sub>4</sub>)<sub>2</sub>HPO<sub>4</sub> at Ca/P = 1.5. Calcium acetate concentration: 0.125 (<b>a</b>), 0.25 (<b>b</b>), 0.5 (<b>c</b>), 1.0 (<b>d</b>), and 2.0 (<b>e</b>).</p> Full article ">Figure 5
<p>Integral (<b>a</b>) and differential (<b>b</b>) curves of the particles size distributions of the powders synthesized from aqueous solutions of Ca(CH<sub>3</sub>COO)<sub>2</sub> and (NH<sub>4</sub>)<sub>2</sub>HPO<sub>4</sub> at Ca/P = 1.5. Calcium acetate concentration: 2.0 M (<b>1</b>), 1.0 M (<b>2</b>), 0.5 M (<b>3</b>), 0.25 M (<b>4</b>), and 0.125 M (<b>5</b>).</p> Full article ">Figure 6
<p>Thermal analysis (TA) curves of the as-prepared powder samples synthesized from aqueous solutions of Ca(CH<sub>3</sub>COO)<sub>2</sub> and (NH<sub>4</sub>)<sub>2</sub>HPO<sub>4</sub> at Ca/P = 1.5. Calcium acetate concentration is shown near the corresponding curves.</p> Full article ">Figure 7
<p>Mass spectra MS profiles of the evolved gaseous decomposition products of the as-prepared powder samples synthesized from aqueous solutions of Ca(CH<sub>3</sub>COO)<sub>2</sub> and (NH<sub>4</sub>)<sub>2</sub>HPO<sub>4</sub> at Ca/P = 1.5. Calcium acetate concentration is shown in the legend; <span class="html-italic">m</span>/<span class="html-italic">z</span>: 18 (<b>a</b>), 44 (<b>b</b>), 30 (<b>c</b>).</p> Full article ">Figure 8
<p>The density of the green powder compacts pressed at 100 MPa. The labels at the horizontal axis show the concentration of Ca(CH<sub>3</sub>COO)<sub>2</sub> solution used for the synthesis.</p> Full article ">Figure 9
<p>XRD patterns of the ceramic samples prepared from the synthesized powders after firing at 1100 °C. Concentrations of Ca(CH<sub>3</sub>COO)<sub>2</sub> precursor are shown near the curves. o—β-Ca<sub>3</sub>(PO<sub>4</sub>)<sub>2</sub>, PDF card # 1-169; *—β-Ca<sub>2</sub>P<sub>2</sub>O<sub>7</sub>, PDF card # 9-346.</p> Full article ">Figure 10
<p>The content of the phases in the ceramic samples prepared from the synthesized powders after firing at 1100 °C; the labels at the horizontal axis show the concentration of Ca(CH<sub>3</sub>COO)<sub>2</sub> solution used for the synthesis.</p> Full article ">Figure 11
<p>Relative diameter (<b>a</b>) and mass (<b>b</b>) of ceramic samples after firing at 1100 °C: <span class="html-italic">D</span> and <span class="html-italic">m</span>—diameter and mass of a sample after firing at 1100 °C, respectively; <span class="html-italic">D</span><sub>0</sub> and <span class="html-italic">m</span><sub>0</sub>—diameter and mass of the corresponding green powder compacts before firing; the labels at the horizontal axis show the concentration of Ca(CH<sub>3</sub>COO)<sub>2</sub> solution used for the synthesis.</p> Full article ">Figure 12
<p>Density of the ceramic samples after firing at 1100 °C. The labels at the horizontal axis show the concentration of Ca(CH<sub>3</sub>COO)<sub>2</sub> solution used for the synthesis.</p> Full article ">Figure 13
<p>SEM images of the ceramic materials based on powders synthesized from aqueous solutions of Ca(CH<sub>3</sub>COO)<sub>2</sub> and (NH<sub>4</sub>)<sub>2</sub>HPO<sub>4</sub> at Ca/P = 1.5 after firing at 1100 °C. Calcium acetate concentration: 0.125 (<b>a</b>), 0.25 (<b>b</b>), 0.5 (<b>c</b>), 1.0 (<b>d</b>), and 2.0 M (<b>e</b>).</p> Full article ">Figure 13 Cont.
<p>SEM images of the ceramic materials based on powders synthesized from aqueous solutions of Ca(CH<sub>3</sub>COO)<sub>2</sub> and (NH<sub>4</sub>)<sub>2</sub>HPO<sub>4</sub> at Ca/P = 1.5 after firing at 1100 °C. Calcium acetate concentration: 0.125 (<b>a</b>), 0.25 (<b>b</b>), 0.5 (<b>c</b>), 1.0 (<b>d</b>), and 2.0 M (<b>e</b>).</p> Full article ">
11 pages, 13396 KiB  
Article
On the Enlargement of the Emission Spectra from the 4I13/2 Level of Er3+ in Silica-Based Optical Fibers through Lanthanum or Magnesium Co-Doping
by Manuel Vermillac, Jean-François Lupi, Stanislaw Trzesien, Michele Ude and Wilfried Blanc
Ceramics 2018, 1(2), 364-374; https://doi.org/10.3390/ceramics1020029 - 11 Dec 2018
Cited by 9 | Viewed by 3659
Abstract
Improving optical fiber amplifiers requires the elaboration and use of new materials and new compositions. In this sense, we prepared erbium-doped optical fiber samples that were co-doped with magnesium or lanthanum by gradual-time solution doping. Doping concentrations and thermal processes induce the formation [...] Read more.
Improving optical fiber amplifiers requires the elaboration and use of new materials and new compositions. In this sense, we prepared erbium-doped optical fiber samples that were co-doped with magnesium or lanthanum by gradual-time solution doping. Doping concentrations and thermal processes induce the formation of nanoparticles. The effect of lanthanum and magnesium contents on the width of the spontaneous emission of the 4 I 13 / 2 level of Er 3 + was characterized in the nanoparticle-rich fiber samples. For that purpose, the width was characterized by the effective linewidth and the full-width at half-maximum (FWHM). The results indicate the robustness of the effective linewidth to strong variations in the intensity profiles of the 4 I 13 / 2 spontaneous emission. Increasing the doping concentrations of both magnesium and lanthanum increases the FWHM and the effective linewidth, along with optical losses. Results show that the fabrication of nanoparticle-rich optical fibers through lanthanum or magnesium doping induces an FHWM broadening of 54% and 64%, respectively, or an effective linewidth broadening of 59% (for both elements) while maintaining a transparency that is compatible with fiber laser and amplifier applications. Full article
(This article belongs to the Special Issue Design, Fabrication and Assessment of Nanostructured Materials and Systems for Advanced Optical Applications)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>Pictures of the La-doped (<b>top</b>) and Mg-doped (<b>bottom</b>) optical preforms. For both preforms, the core whitens with the increase in La or Mg content along the length of the preform. Estimated positions in the preform of the extrema fiber samples are indicated by vertical lines.</p> Full article ">Figure 2
<p>Magnesium and lanthanum contents for analyzed gradual-doping fiber samples.</p> Full article ">Figure 3
<p>Secondary electron microscopy (SEM) images of the core of two optical fiber samples. The core sections correspond to the longitudinal view (along the drawing axis). The image on the left shows the core of the La-doped sample containing 1 atom % of La. The image on the right shows the highly doped part of the core of the Mg-doped sample containing 2.1 atom % of Mg.</p> Full article ">Figure 4
<p>Optical losses of lanthanum- and magnesium-doped fiber samples measured at 900 and 1100 nm.</p> Full article ">Figure 5
<p>Maximum-normalized (<b>on the left</b>) and area-normalized <b>(on the right</b>) Er<math display="inline"> <semantics> <msup> <mrow/> <mrow> <mn>3</mn> <mo>+</mo> </mrow> </msup> </semantics> </math> emission spectra of magnesium-doped samples. Samples are labeled according to their content in magnesium.</p> Full article ">Figure 6
<p>Maximum-normalized (<b>on the left</b>) and area-normalized (<b>on the right</b>) Er<math display="inline"> <semantics> <msup> <mrow/> <mrow> <mn>3</mn> <mo>+</mo> </mrow> </msup> </semantics> </math> emission spectra of lanthanum-doped samples. Samples are labeled according to their content in lanthanum; <math display="inline"> <semantics> <mi>ϵ</mi> </semantics> </math> is used for the sample where lanthanum content was too low to be measurable.</p> Full article ">Figure 7
<p>Comparison of the maximum normalization of Er<math display="inline"> <semantics> <msup> <mrow/> <mrow> <mn>3</mn> <mo>+</mo> </mrow> </msup> </semantics> </math> emission spectra of the <math display="inline"> <semantics> <msup> <mrow/> <mn>4</mn> </msup> </semantics> </math>I<math display="inline"> <semantics> <msub> <mrow/> <mrow> <mn>13</mn> <mo>/</mo> <mn>2</mn> </mrow> </msub> </semantics> </math> level in fiber samples. Emission spectra of highly doped Mg and La fiber samples are displayed with Ref-silica and low La content samples.</p> Full article ">Figure 8
<p>Emission peak wavelengths for fiber samples (FS) as a function of magnesium or lanthanum content.</p> Full article ">Figure 9
<p>Evolution of Er<math display="inline"> <semantics> <msup> <mrow/> <mrow> <mn>3</mn> <mo>+</mo> </mrow> </msup> </semantics> </math> emission bandwidth with lanthanum (black) and magnesium (red) content. Ref-silica values are shown in blue. Circle symbols are used for the effective linewidth (on the right scale) and crosses are used for FWHM (on the left scale).</p> Full article ">Figure 10
<p>Evolution of Er<math display="inline"> <semantics> <msup> <mrow/> <mrow> <mn>3</mn> <mo>+</mo> </mrow> </msup> </semantics> </math> emission width with optical losses measured at 900 nm for lanthanum and magnesium co-doping.</p> Full article ">
11 pages, 3268 KiB  
Article
Fabricating MOF/Polymer Composites via Freeze Casting for Water Remediation
by Coral Rogers, Daniel Pun, Qingshan Fu and Haifei Zhang
Ceramics 2018, 1(2), 353-363; https://doi.org/10.3390/ceramics1020028 - 28 Nov 2018
Cited by 15 | Viewed by 4346
Abstract
Various porous materials have been used as adsorbents for water remediation. Among them, metal-organic framework (MOF) particles have been explored intensively, due to their size-controlled micropores and high surface areas. MOF nanoparticles are often used because of high external surface area and easy [...] Read more.
Various porous materials have been used as adsorbents for water remediation. Among them, metal-organic framework (MOF) particles have been explored intensively, due to their size-controlled micropores and high surface areas. MOF nanoparticles are often used because of high external surface area and easy access to the micropores. However, recovering MOF nanoparticles, usually by filtration or centrifugation, is time-consuming and is difficult to scale up. We report here the preparation of porous MOF/polymer monoliths by freeze casting for water remediation. Chitosan and UiO-66 (Universitetet i Oslo) nanoparticles (including different surface functional groups) are used to prepare such monoliths. In order to improve the mechanical stability and the tendency of disintegrating in water, the freeze-dried UiO-66/chitosan monoliths are further treated by heating, washing with aqueous NaOH solution, or chemical crosslinking with glutaraldehyde. All these treated monoliths are used for adsorption of a herbicide methylchlorophenoxypropionic acid (MCPP) from aqueous solution. Particularly, the crosslinked chitosan/UiO-66 monolith achieves an adsorption capacity of 47.67 mg g−1, with a 60 ppm MCPP solution. It is superior to that presented by the sole UiO-66 nanoparticles, exhibiting over a 30% increase in the adsorption capacity. The monoliths can be easily removed using tweezers, providing facile recyclability, which is advantageous for upscaling. The recycled monolith upheld approximately 75% of the adsorption capacity compared to the original monolith after three reuse cycles. Full article
(This article belongs to the Special Issue Ice-Templated and Freeze-Cast Ceramics)
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Figure 1

Figure 1
<p>Pore structure by SEM imaging of (<b>A</b>) UiO-66-NO<sub>2</sub> nanoparticles, (<b>B</b>,<b>C</b>) CM/UiO-66-NO<sub>2</sub> (1:1) freeze dried composite, (<b>D</b>,<b>E</b>) CM/UiO-66-NO<sub>2</sub> composite after heat treatment (1:1), (<b>F</b>,<b>G</b>) CM/UiO-66-NO<sub>2</sub> (1:1) composite after base treatment, and (<b>H</b>,<b>I</b>) glutaraldehyde-crosslinked CM/UiO-66 (1:1).</p> Full article ">Figure 2
<p>Profiles of the adsorbed quantity of methylchlorophenoxypropionic acid (MCPP) versus soaking time by dispersing (10 mg) of nanoparticles in 10 mL of aqueous solution of MCPP (60 ppm). Each of the tests in this study were carried out three times.</p> Full article ">Figure 3
<p>Profiles of the adsorbed quantity of MCPP versus soaking time by immersing the composites (10 mg) in 10 mL of aqueous solution of MCPP (60 ppm).</p> Full article ">Figure 4
<p>Absorption profiles of the absorbent materials including UiO-66 nanoparticles, base-treated (BT) medium weight chitosan/UiO-66 (polymer: MOF = 1:1 in mass), and glutaraldehyde (GA)-crosslinked medium weight chitosan/UiO-66 (polymer: MOF = 1:1 in mass) against the soaking time within 60 ppm MCPP solution.</p> Full article ">Figure 5
<p>Reusability of glutaraldehyde-crosslinked CM/UiO-66 (polymer: MOF = 1:1) for the adsorptive removal of MCPP from 10 mL of 60 ppm MCPP aqueous solution.</p> Full article ">
10 pages, 2475 KiB  
Article
Equimolar Yttria-Stabilized Zirconia and Samaria-Doped Ceria Solid Solutions
by Reginaldo Muccillo, Daniel Z. De Florio and Eliana N. S. Muccillo
Ceramics 2018, 1(2), 343-352; https://doi.org/10.3390/ceramics1020027 - 22 Nov 2018
Cited by 2 | Viewed by 4047
Abstract
Compositions of (ZrO2)0.92(Y2O3)0.08 (zirconia: 8 mol % yttria—8YSZ) and (CeO2)0.8(Sm2O3)0.2 (ceria: 20 mol % samaria—SDC20) ceramic powders were prepared by attrition milling to form an [...] Read more.
Compositions of (ZrO2)0.92(Y2O3)0.08 (zirconia: 8 mol % yttria—8YSZ) and (CeO2)0.8(Sm2O3)0.2 (ceria: 20 mol % samaria—SDC20) ceramic powders were prepared by attrition milling to form an equimolar powder mixture, followed by uniaxial and isostatic pressing. The pellets were quenched to room temperature from 1200 °C, 1300 °C, 1400 °C and 1500 °C to freeze the defects configuration attained at those temperatures. X-ray diffraction analyses, performed in all quenched pellets, show the evolution of the two (8YSZ and SDC20) cubic fluorite structural phases to a single phase at 1500 °C, identified by Rietveld analysis as a tetragonal phase. Impedance spectroscopy analyses were carried out in pellets either quenched or slowly cooled from 1500 °C. Heating the quenched pellets to 1000 °C decreases the electrical resistivity while it increases in the slowly cooled pellets; the decrease is ascribed to annealing of defects created by lattice micro-tensions during quenching while the increase to partial destabilization of the tetragonal phase. Full article
(This article belongs to the Special Issue Novel Processing Routes of Ceramics for Functional Applications)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>X-ray diffraction patterns showing the phase evolution of equimolar ZrO<sub>2</sub>: 8 mol % Y<sub>2</sub>O<sub>3</sub> + CeO<sub>2</sub>: 20 mol % Sm<sub>2</sub>O<sub>3</sub> ceramic pellets quenched from 1200, 1300, 1400 and 1500 °C.</p> Full article ">Figure 2
<p>X-ray diffraction patterns, from bottom to top, of ZrO<sub>2</sub>: 8 mol % Y<sub>2</sub>O<sub>3</sub> (8YSZ), CeO<sub>2</sub>: 20 mol % Sm<sub>2</sub>O<sub>3</sub> (SDC20), and equimolar 8YSZ + SDC20 ceramic pellets quenched from 1500 °C. Numbers refer to 2θ peak diffraction angles (top) and peak diffraction angles with known Miller indices (middle and bottom).</p> Full article ">Figure 3
<p>X-ray diffraction patterns of equimolar ZrO<sub>2</sub>: 8 mol % Y<sub>2</sub>O<sub>3</sub> + CeO<sub>2</sub>: 20 mol % Sm<sub>2</sub>O<sub>3</sub> ceramic pellets quenched from 1500 °C and the calculated lines corresponding to the best fitted structure after the Rietveld analysis.</p> Full article ">Figure 4
<p>Scanning electron microscopy micrographs of polished and thermally etched surfaces of ZrO<sub>2</sub>: 8 mol % Y<sub>2</sub>O<sub>3</sub> (8YSZ), CeO<sub>2</sub>: 20 mol % Sm<sub>2</sub>O<sub>3</sub> (SDC20) and equimolar 8YSZ + SDC20 mixture, all quenched from 1500 °C.</p> Full article ">Figure 5
<p>Impedance spectroscopy diagrams, of ZrO<sub>2</sub>: 8 mol % Y<sub>2</sub>O<sub>3</sub>, CeO<sub>2</sub>: 20 mol % Sm<sub>2</sub>O<sub>3</sub> and equimolar ZrO<sub>2</sub>: 8 mol % Y<sub>2</sub>O<sub>3</sub> + CeO<sub>2</sub>: 20 mol % Sm<sub>2</sub>O<sub>3</sub> ceramic pellets quenched from 1500 °C. Numbers stand for log f (f: Hz). Temperature of measurement: 530 °C. Inset: Enlarged view of the low resistance region with the 8YSZ and SDC20 impedance diagrams.</p> Full article ">Figure 6
<p>(<b>a</b>): impedance diagrams of ceramic pellets composed of equimolar mixture of ZrO<sub>2</sub>: 8 mol % Y<sub>2</sub>O<sub>3</sub> + CeO<sub>2</sub>: 20 mol % Sm<sub>2</sub>O<sub>3</sub>, quenched from 1500 °C and annealed at 1000 °C for different times; numbers stand for log f (f: Hz). (<b>b</b>): corresponding Bode diagrams of the imaginary component of the impedance. Temperature of measurement: 530 °C.</p> Full article ">Figure 7
<p>Impedance diagrams of equimolar ZrO<sub>2</sub>: 8 mol % Y<sub>2</sub>O<sub>3</sub> + CeO<sub>2</sub>: 20 mol % Sm<sub>2</sub>O<sub>3</sub> sintered at 1500 °C/2 h with 0.17 °C/s heating and cooling rates, and after further annealing at 1000 °C/30 min. Temperature of measurement: (<b>a</b>) 310 °C and (<b>b</b>) 530 °C. Numbers stand for log f (f: Hz).</p> Full article ">Figure 8
<p>X-ray diffraction patterns of equimolar ZrO<sub>2</sub>: 8 mol % Y<sub>2</sub>O<sub>3</sub> + CeO<sub>2</sub>: 20 mol % Sm<sub>2</sub>O<sub>3</sub> sintered at 1500 °C/2 h with 0.17 °C/s heating and cooling rates (thkl), and after further annealing at 1000 °C/1 min (chkl); thkl and chkl stand for tetragonal and cubic phases with the Miller indices.</p> Full article ">
14 pages, 3504 KiB  
Article
Processing of Macroporous Alumina Ceramics Using Pre-Expanded Polymer Microspheres as Sacrificial Template
by Marina Ciurans Oset, Jan Nordin and Farid Akhtar
Ceramics 2018, 1(2), 329-342; https://doi.org/10.3390/ceramics1020026 - 6 Nov 2018
Cited by 13 | Viewed by 5825
Abstract
Shaped porous ceramics have proven to be the most adapted materials for several industrial applications, both at low and high temperatures. Recent research has been focused on developing shaping techniques, allowing for a better control over the total porosity and the pores characteristics. [...] Read more.
Shaped porous ceramics have proven to be the most adapted materials for several industrial applications, both at low and high temperatures. Recent research has been focused on developing shaping techniques, allowing for a better control over the total porosity and the pores characteristics. In this study, macroporous alumina foams were fabricated by gel-casting using pre-expanded polymeric microspheres with average sizes of 40 ?m, 20 ?m, and 12 ?m as sacrificial templates. The gel-casting method, as well as the drying, debinding, and presintering conditions were investigated and optimized to process mechanically strong and highly porous alumina scaffolds. Furthermore, a reliable model relating the amount of pre-expanded polymeric microspheres and the total porosity of the presintered foams was developed and validated by mercury intrusion porosimetry measurements. The electron microscopy investigation of the presintered foams revealed that the size distribution and the shape of the pores could be tailored by controlling the particle size distribution and the shape of the wet pre-expanded microspheres. Highly uniform and mechanically stable alumina foams with bimodal porosity ranging from 65.7 to 80.2 vol. % were processed, achieving compressive strengths from 3.3 MPa to 43.6 MPa. Given the relatively open pore structure, the pore size distribution, the presintered mechanical strength, and the high porosity achieved, the produced alumina foams could potentially be used as support structures for separation, catalytic, and filtration applications. Full article
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Figure 1

Figure 1
<p>Scanning electron microscopy (SEM) micrographs of the three types of wet expanded Expancel<sup>®</sup> microspheres used in this study: (<b>a</b>) WE 40 μm, (<b>b</b>) WE 20 μm, and (<b>c</b>) WE 12 μm.</p> Full article ">Figure 2
<p>Flowchart of the gel-casting and sacrificial templating procedure followed in the present study to prepare the macroporous alumina specimens.</p> Full article ">Figure 3
<p>Presintered alumina foams prepared using different dispersing agents: (<b>a</b>) sodium dodecyl sulfate (SDS), an anionic surfactant and (<b>b</b>) polyacrylic acid (PAA), a water-soluble anionic polyelectrolyte.</p> Full article ">Figure 4
<p>Fracture surface of presintered alumina foams prepared using different gelation methods: (<b>a</b>) thermal gelation and (<b>b</b>) catalytic gelation.</p> Full article ">Figure 5
<p>Effect of the drying conditions on (<b>a</b>) cumulative mass loss and (<b>b</b>) cumulative radial shrinkage. The composition and properties of batches 1 to 6 are summarized in <a href="#ceramics-01-00026-t002" class="html-table">Table 2</a>.</p> Full article ">Figure 6
<p>Thermal analysis of a high porosity (66.7 vol. %) alumina foam produced with WE 40 μm microspheres: (<b>a</b>) differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA) curves as a function of time and (<b>b</b>) TGA curve as a function of temperature.</p> Full article ">Figure 7
<p>Fracture surface of presintered alumina foams: low magnification SEM micrographs of (<b>a</b>) 75.1 vol. % porosity foam produced with WE 40 μm, (<b>b</b>) 76.0 vol. % porosity foam produced with WE 20 μm and (<b>c</b>) 80.2 vol. % porosity foam produced with WE 12 μm; (<b>d</b>) high magnification SEM micrograph of a 65.7 vol. % porosity foam produced with WE 12 μm.</p> Full article ">Figure 8
<p>Particle size distribution by number of the three types of wet pre-expanded microspheres used in this study: (<b>a</b>) WE 40 μm, (<b>b</b>) WE 20 μm, and (<b>c</b>) WE 12 μm. Primary macropore size distribution by number of presintered porous alumina foams: (<b>d</b>) 75.1 vol. % porosity foam produced with WE 40 μm, (<b>e</b>) 76.0 vol. % porosity foam produced with WE 20 μm and (<b>f</b>) 80.2 vol. % porosity foam produced with WE 12 μm.</p> Full article ">Figure 9
<p>Pore size distribution of presintered alumina foams prepared with WE 20 μm Expancel<sup>®</sup> microspheres, as determined by mercury intrusion porosimetry: (<b>a</b>) 76.0 vol. % porosity and (<b>b</b>) 61.6 vol. % porosity foams.</p> Full article ">
10 pages, 2026 KiB  
Article
Preparation of Electrically Conductive Calcium Phosphate Composite Foams by Particle-Stabilized Emulsion Route
by Wenjea J. Tseng and Wuei-Hung Kao
Ceramics 2018, 1(2), 319-328; https://doi.org/10.3390/ceramics1020025 - 28 Oct 2018
Cited by 4 | Viewed by 3523
Abstract
Macroporous composite foams consisting of β-tricalcium phosphate (β-TCP) and titanium nitride (TiN) have been prepared by a facile emulsion route involving sintering at elevated temperatures after shaping. Commercially available hydroxyapatite and titanium particles are used as the starting material; to which, the surface [...] Read more.
Macroporous composite foams consisting of β-tricalcium phosphate (β-TCP) and titanium nitride (TiN) have been prepared by a facile emulsion route involving sintering at elevated temperatures after shaping. Commercially available hydroxyapatite and titanium particles are used as the starting material; to which, the surface of the particles has been modified by preferential adsorption of hexadecylamine to change from hydrophilic to hydrophobic character in water. This renders stable air-in-water emulsions from the particle-filled suspensions by simple mechanical frothing. Sintered β-TCP/TiN foams with a porosity of 65–70%, pore size of 20–2000 nm, and three-point rupture strength of 25–43 kPa have been obtained. Electrical resistance has been found to reduce pronouncedly when the initial titanium loading exceeds 15 vol.% for the composite foams sintered at 1000 °C under reducing nitrogen-hydrogen atmosphere. Full article
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<p>Experimental flowchart.</p> Full article ">Figure 2
<p>SEM micrographs of the as-received (<b>a</b>) hydroxyapatite and (<b>b</b>) titanium particles.</p> Full article ">Figure 3
<p>XRD patterns of the as-received hydroxyapatite and titanium particles.</p> Full article ">Figure 4
<p>Effect of HDA concentration on the contact angle of hydroxyapatite and titanium particles, respectively.</p> Full article ">Figure 5
<p>SEM micrographs of the macroporous composite foams with volumetric titanium fraction of (<b>a</b>) 0.1, (<b>b</b>) 0.15, (<b>c</b>) 0.2, and (<b>d</b>) 0.25 after sintering at 1000 °C at 95 N<sub>2</sub>/5H<sub>2</sub> atmosphere for 2 h.</p> Full article ">Figure 6
<p>Porosity and rupture strength of the sintered composite foams with volumetric titanium fraction ranging from 0.1 to 0.5. The foams were sintered at 1000 °C at 95 N<sub>2</sub>/5H<sub>2</sub> atmosphere for 2 h.</p> Full article ">Figure 7
<p>Pore-size distribution of the sintered composite foams with volumetric titanium fraction from 0.1 to 0.5. The foams were sintered at 1000 °C at 95 N<sub>2</sub>/5H<sub>2</sub> atmosphere for 2 h.</p> Full article ">Figure 8
<p>XRD patterns of the sintered composite foams with volumetric titanium fraction of 0.25. The foams were sintered at 1000 and 1200 °C at 95 N<sub>2</sub>/5H<sub>2</sub> atmosphere for 2 h.</p> Full article ">Figure 9
<p>Effect of Ti fraction on the electrical resistance of the sintered composite foams with volumetric titanium fraction from 0.1 to 0.5. The foams were sintered at 1000 °C at 95 N<sub>2</sub>/5H<sub>2</sub> atmosphere for 2 h.</p> Full article ">
15 pages, 3712 KiB  
Article
Mechanochemically-Assisted Synthesis of Lead-Free Piezoelectric CaZrO3-Modified (K,Na,Li)(Nb,Ta)O3-Solid Solution
by Kristian Radan, Brigita Kmet, Silvo Drnovšek, Uroš Prah, Tadej Rojac and Barbara Malič
Ceramics 2018, 1(2), 304-318; https://doi.org/10.3390/ceramics1020024 - 17 Oct 2018
Cited by 7 | Viewed by 3680
Abstract
Lead-free piezoelectric 0.95(Na0.49K0.49Li0.02)(Nb0.8Ta0.2)O3–0.05CaZrO3 with 2 wt % MnO2 addition was prepared using mechanochemically-assisted solid-state synthesis. Upon mechanochemical activation of the mixture of reagents partial amorphization occurs which contributes to [...] Read more.
Lead-free piezoelectric 0.95(Na0.49K0.49Li0.02)(Nb0.8Ta0.2)O3–0.05CaZrO3 with 2 wt % MnO2 addition was prepared using mechanochemically-assisted solid-state synthesis. Upon mechanochemical activation of the mixture of reagents partial amorphization occurs which contributes to a significantly lower temperature of completion of the solid-state reaction, ~600 °C as opposed to ~700 °C for the conventional solid-state synthesis as determined by thermal analysis. The ceramic specimens prepared by the mechanochemically-assisted route exhibit improved compositional homogeneity and slightly enhanced piezoelectric properties, achieved in a considerably shorter processing time compared to the conventional solid-state synthesis route, which was studied as a reference. Full article
(This article belongs to the Special Issue Functional Ceramics for Energy Applications)
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<p>XRD patterns of nonactivated mixture (0 h) and after activation by high-energy milling for 1, 2, 4, and 8 h. Notations: ●, Nb<sub>2</sub>O<sub>5</sub>; *, Ta<sub>2</sub>O<sub>5</sub>; ▽, A<sub>2</sub>CO<sub>3</sub> (A = K and/or Na).</p> Full article ">Figure 2
<p>SEM micrographs of the powder mixture before (<b>a</b>) and after (<b>b</b>) mechanochemical activation induced by 8 h of high-energy milling.</p> Full article ">Figure 3
<p>TGA, DTA, and EGA (CO<sub>2</sub>, H<sub>2</sub>O) plots of the powder mixtures before (0 h, green) and after (8 h, red, dashed) mechanochemical activation.</p> Full article ">Figure 4
<p>Dimensional changes with temperature of the powder mixtures before (0 h, green) and after (8 h, red, dashed) mechanochemical activation, pressed into pellets.</p> Full article ">Figure 5
<p>XRD patterns of conventional solid-state synthesis (CSS) (above) and activated powder mixture (MCA) (below) after calcination. Diffraction peaks corresponding to the perovskite phase were indexed with a cubic unit cell (PDF 18-7023, (K<sub>0.47</sub>Na<sub>0.51</sub>Li<sub>0.02</sub>)(Ta<sub>0.1</sub>Nb<sub>0.9</sub>)O<sub>3</sub>, [<a href="#B41-ceramics-01-00024" class="html-bibr">41</a>]). The insets show the ZrO<sub>2</sub> (PDF 36-0420) and MnO<sub>2</sub> (PDF 81-2261) peaks in the 2-theta range from 23 ° to 65 °.</p> Full article ">Figure 6
<p>XRD patterns of CSS (<b>a</b>) and MCA (<b>b</b>) samples after sintering at 1150 °C. The insets show the detection of Mn<sub>3</sub>O<sub>4</sub> (PDF 24-0734) in the 2-theta range from 26° to 68°.</p> Full article ">Figure 7
<p>Measured and calculated XRD profiles and their difference plots for the Rietveld refinement of the CSS (<b>a</b>) and MCA (<b>b</b>) samples.</p> Full article ">Figure 8
<p>Cross-section SEM images of CSS ((<b>a</b>) polished surface with the inset showing a magnified Ta-rich phase and (<b>b</b>) thermally etched) and MCA ((<b>c</b>) polished surface and (<b>d</b>) thermally etched) ceramics. The arrows mark the secondary phases.</p> Full article ">Figure 9
<p>Polarization hysteresis loops of CSS (black) and MCA (blue, dashed) ceramic samples, measured at room temperature with a fixed frequency of 100 Hz.</p> Full article ">
17 pages, 3115 KiB  
Article
Influence of Aging on Mechanical Properties of Yttria-Doped Zirconia
by Yuta Kimura, Takuto Kushi, Atsushi Unemoto, Koji Amezawa and Tatsuya Kawada
Ceramics 2018, 1(2), 287-303; https://doi.org/10.3390/ceramics1020023 - 12 Oct 2018
Cited by 9 | Viewed by 3768
Abstract
We evaluated the influence of aging on mechanical properties of 8% yttria-doped zirconia (8YSZ) from room temperature to 1200 K. The temperature dependence of the dynamic Young’s and shear moduli of 8YSZ with and without the aging treatment was investigated by using a [...] Read more.
We evaluated the influence of aging on mechanical properties of 8% yttria-doped zirconia (8YSZ) from room temperature to 1200 K. The temperature dependence of the dynamic Young’s and shear moduli of 8YSZ with and without the aging treatment was investigated by using a resonance method. The dynamic Young’s and shear moduli of 8YSZ without the aging treatment decreased by 33% below 700 K and gradually increased at higher temperatures with increasing temperature. On the other hand, those with the aging treatments decreased by around 20% below 600 K while did not significantly change above 600 K with increasing temperature. These demonstrated the effect of aging on the dynamic Young’s and shear moduli of 8YSZ was most remarkable at intermediate temperatures (600~1000 K). Although it was suggested that the existence ratio of the metastable tetragonal phase was increased during the aging treatment, it is likely that the influence of this phase transition on the dynamic Young’s and shear moduli was not significant. It seemed that the difference in the dynamic Young’s and shear moduli of 8YSZ with and without the aging treatment at intermediate temperatures was due to the local ordering of the oxygen vacancies. Full article
(This article belongs to the Special Issue Functional Ceramics for Energy Applications)
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<p>(<b>a</b>) Electric furnace used for the aging treatment in this study. Several samples of 8YSZ were placed in the furnace and heat-treated at 1073 K in air for a maximum of seven months. Conducting wires were connected to one of the samples in the furnace for the electrical conductivity measurements. (<b>b</b>) Sample for the electrical conductivity measurements. The electrical conductivity was measured by the AC four-terminal method in a frequency range from 10<sup>5</sup> to 10 Hz with the potential amplitude of 10 mV.</p> Full article ">Figure 2
<p>(<b>a</b>) Temperature dependence of the dynamic Young’s and shear moduli and the Poisson’s ratio of 8YSZ without aging treatment under <span class="html-italic">P</span>(O<sub>2</sub>) of 0.01 bar. Black circle, red square and blue triangle represent the dynamic Young’s modulus, the dynamic shear modulus and the Poisson’s ratio, respectively. The black broken line expresses the fitting curve of the dynamic Young’s modulus of 8YSZ below 323 K by Wachtman’s equation [<a href="#B16-ceramics-01-00023" class="html-bibr">16</a>]. (<b>b</b>) Temperature dependence of the internal friction and the change in the dynamic Young’s modulus with temperature, d<span class="html-italic">E</span>/d<span class="html-italic">T</span>, of 8YSZ without aging treatment under <span class="html-italic">P</span>(O<sub>2</sub>) of 0.01 bar. Green symbol and open black circle show the internal friction and d<span class="html-italic">E</span>/d<span class="html-italic">T</span>, respectively.</p> Full article ">Figure 3
<p>(<b>a</b>) X-ray diffraction patterns at 2<span class="html-italic">θ</span> = 70~78° and (<b>b</b>) Raman spectrum of the 8YSZ sample without the aging treatment measured at room temperature.</p> Full article ">Figure 4
<p>Raman spectra of the 8YSZ sample without the aging treatment measured at room temperature, 373, 473, 573, 673, 773 and 873 K.</p> Full article ">Figure 5
<p>Peak height ratios of the peaks at around 610 and 470 cm<sup>−1</sup>, I<sub>4</sub>/I<sub>6</sub>, in the Raman spectrum of the 8YSZ sample with and without the aging treatment as a function of temperature. Red/black circles represent the result with the sample with/without the aging treatment, respectively.</p> Full article ">Figure 6
<p>Conductivity of the 8YSZ with the aging treatment at 1073 K in air as a function of aging time. Red and blue circles represent the conductivity obtained in this work and the one reported by Kondoh et al. [<a href="#B7-ceramics-01-00023" class="html-bibr">7</a>].</p> Full article ">Figure 7
<p>Temperature dependence of the dynamic Young’s modulus of the 8YSZ sample without the aging treatment and those with the aging treatment for 720, 1465, 2210 and 5110 h.</p> Full article ">Figure 8
<p>Temperature dependence of (<b>a</b>) the dynamic shear modulus and (<b>b</b>) the Poisson’s ratio of the 8YSZ sample without the aging treatment and those with the aging treatment for 720, 1465, 2210 and 5110 h.</p> Full article ">Figure 9
<p>Temperature dependence of (<b>a</b>) the internal friction and (<b>b</b>) the change in the dynamic Young’s modulus with temperature, d<span class="html-italic">E</span>/d<span class="html-italic">T</span>, of the 8YSZ sample without the aging treatment and those with the aging treatment for 720, 1465, 2210 and 5110 h.</p> Full article ">Figure 10
<p>Temperature dependence of the dynamic Young’s modulus of the 8YSZ sample with aging treatment for 720 h, that after re-annealing at 1673 K for 1 h and that without the aging treatment.</p> Full article ">Figure 11
<p>Change ratio of the dynamic Young’s modulus of 8YSZ sample as a function of aging time at low (room temperature), intermediate (773 K) and high (1073 K) temperatures. The broken lines are guides to the eye.</p> Full article ">Figure 12
<p>(<b>a</b>) X-ray diffraction patterns at 2<span class="html-italic">θ</span> = 70~78° and (<b>b</b>) Raman spectra of the 8YSZ samples with and without the aging treatment at 1073 K for 2930 h. Both the X-ray diffraction measurements and the Raman spectroscopy were conducted at room temperature.</p> Full article ">
13 pages, 5993 KiB  
Article
Devitrification Behavior of Sol-Gel Derived ZrO2-SiO2 Rare-Earth Doped Glasses: Correlation between Structural and Optical Properties
by Masato Isogai, Alexander Veber, Maria Rita Cicconi, Tomokatsu Hayakawa and Dominique De Ligny
Ceramics 2018, 1(2), 274-286; https://doi.org/10.3390/ceramics1020022 - 7 Oct 2018
Cited by 8 | Viewed by 3743
Abstract
Optical and structural properties of glasses and glass-ceramics (GC) obtained by different heat-treatment of Tb and Tb-Yb doped sol-gel derived 30ZrO2-70SiO2 materials were investigated. A glass was formed after treatment at 700 °C whereas devitrification of the media after the [...] Read more.
Optical and structural properties of glasses and glass-ceramics (GC) obtained by different heat-treatment of Tb and Tb-Yb doped sol-gel derived 30ZrO2-70SiO2 materials were investigated. A glass was formed after treatment at 700 °C whereas devitrification of the media after the treatment at 1000 and 1100 °C, led to the formation of GC containing up to three different crystalline phases, namely, tetragonal ZrO2, Yb-disilicate and cristobalite. The modification of the optical properties through the heat treatment was caused by redistribution of the rare earth elements (REE) among the different phases: both Tb and Yb entered the t-ZrO2 lattice, Yb can also be present in the form of a Yb2Si2O7 crystal. Devitrification led to an increase in Tb?Yb energy transfer efficiency as compared to the glass, though it was higher in the samples heat-treated at 1000 °C than in those treated at 1100 °C. The most intensive Yb3+ luminescence, induced by the energy transfer from the Tb3+ ion, was observed at the interface between t-ZrO2 and the glassy phases, due to the high concentration of REE in this area caused by the inability of ZrO2 to accept larger amounts of the REE. The mechanisms of the Tb?Yb energy transfer vary between different phases of the GC. The results obtained in this study are important for the development of spectral down-converters for potential solar energy applications based on Tb-Yb co-doped glass-ceramics. Full article
(This article belongs to the Special Issue Functional Ceramics for Energy Applications)
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<p>Emission spectra of Yb<sup>3+</sup> (<sup>5</sup>F<sub>5/2</sub>→<sup>2</sup>F<sub>7/2</sub>) obtained for the samples after different heat treatment when the ion was excited directly with λ<sub>ex</sub> = 920 nm (<b>a</b>); due to the Tb→Yb energy transfer with λ<sub>ex</sub> = 488 nm (<b>b</b>). In the latter case, the spectra were recorded with a lower spectral resolution. Temperature and time of the treatment are shown in the plots on the right side of each spectrum.</p> Full article ">Figure 2
<p>Emission spectra of Tb<sup>3+</sup> obtained for Tb-doped (<b>a</b>) and Tb-Yb-doped (<b>b</b>) samples after different heat treatment, λ<sub>ex</sub> = 488 nm. Temperature and time of the treatment are given in the plots on the right side of each spectrum.</p> Full article ">Figure 3
<p>Luminescence decay curves recorded under excitation at 488 nm and emission at 542 nm for (<b>a</b>) Tb-doped samples treated at 1000 °C; (<b>b</b>) Tb-doped samples treated at 1100 °C; (<b>c</b>) Tb-Yb-doped samples treated at 1000 °C, and <b>(d</b>) Tb-Yb-doped samples treated at 1100 °C. The calculated effective times are shown in the insets.</p> Full article ">Figure 4
<p>The Tb→Yb energy transfer efficiency for different heat treatment periods and temperatures.</p> Full article ">Figure 5
<p>Raman scattering spectra of (<b>a</b>) Tb-only and (<b>b</b>) Tb-Yb samples after different heat treatment. The spectra were recorded under 532 nm laser excitation. Temperature and time of the treatment are given in the plots on the right side of each spectrum.</p> Full article ">Figure 6
<p>X-ray diffraction pattern of the Tb-Yb doped samples heat treated at 1100 °C.</p> Full article ">Figure 7
<p>Selected spectra observed with the Raman spectroscopy (λ<sub>ex</sub> = 532 nm) in different points of <span class="html-italic">1100/30 Tb-Yb</span> sample. The observed phases in the spectra are (<b>a</b>) glass phase; (<b>b</b>) tetragonal ZrO<sub>2</sub>; (<b>c</b>) cristobalite; and (<b>d</b>) Yb<sub>2</sub>Si<sub>2</sub>O<sub>7</sub>. All the spectra represent a superposition of bands originating from all the phases. Characteristic vibrational lines correspondent to a specific phase are marked in each plot.</p> Full article ">Figure 8
<p>Results of detailed investigation of <span class="html-italic">1100/30 Tb-Yb</span> sample using Raman microspectrometer (λ<sub>ex</sub> = 532 nm): the microscopic image of the investigated area (<b>a</b>); distribution of tetragonal ZrO<sub>2</sub> (<b>b</b>) SiO<sub>2</sub> cristobalite (<b>c</b>) and Yb<sub>2</sub>Si<sub>2</sub>O<sub>7</sub> (<b>d</b>) phases. A rainbow palette was used for all the plots–violet and red colors correspond to the lowest and the highest value of the calculated ratios/phase amount respectively.</p> Full article ">Figure 9
<p>Results of detailed investigaion of <span class="html-italic">1100/30 Tb-Yb</span> sample using fluorescence microscopes: emission intensities of Tb<sup>3+</sup>/λ<sub>ex</sub> = 488 nm (<b>a</b>), Yb<sup>3+</sup>/λ<sub>ex</sub> = 780 nm (<b>c</b>) and Yb<sup>3+</sup>/λ<sub>ex</sub> = 488 nm (<b>f</b>); relative intensities of the emission peaks calculated as I(555)/I(544) for Tb<sup>3+</sup> (<b>b</b>), I(966)/I(973) and I(1036)/I(973) for Yb<sup>3+</sup> excited under 780 nm (<b>d</b> and <b>e,</b> respectively) and under 488 nm (<b>g</b> and <b>h,</b> respectively). A rainbow palette was used for all the plots–violet and red colors correspond to the lowest and the highest value of the plotted value, respectively.</p> Full article ">Figure 10
<p>Intensity of Yb<sup>3+</sup> luminescence versus incident laser power measured at different spots in 1100/30 Tb-Yb sample, λ<sub>ex</sub> = 488 nm, laser spot diameter during the experiment was ~1.5 μm.</p> Full article ">
13 pages, 6134 KiB  
Article
Evolution of Copper Electrodes Fabricated by Electroless Plating on BaZr0.7Ce0.2Y0.1O3-? Proton-Conducting Ceramic Membrane: From Deposition to Testing in Methane
by Steven P. Harvey, Sandrine Ricote, David R. Diercks, Chun-Sheng Jiang, Neil S. Patki, Anthony Manerbino, Brian Gorman and Mowafak Al-Jassim
Ceramics 2018, 1(2), 261-273; https://doi.org/10.3390/ceramics1020021 - 2 Oct 2018
Cited by 2 | Viewed by 3828
Abstract
We investigated copper electrodes deposited onto a BaZr0.7Ce0.2Y0.1O3-δ (BZCY72) proton-conducting membrane via a novel electroless plating method, which resulted in significantly improved performance when compared to a traditional painted copper electrode. The increased performance was examined [...] Read more.
We investigated copper electrodes deposited onto a BaZr0.7Ce0.2Y0.1O3-δ (BZCY72) proton-conducting membrane via a novel electroless plating method, which resulted in significantly improved performance when compared to a traditional painted copper electrode. The increased performance was examined with a multiscale multitechnique characterization method including time-of-flight secondary-ion mass spectroscopy (TOF-SIMS), transmission electron microscopy (TEM), scanning spreading-resistance microscopy (SSRM), and atom-probe tomography (APT). Through this method, we observed that a palladium catalyst layer alloys with the copper electrode. We also explored the nature of a non-coking-induced carbon-rich phase that may be involved with the improved performance of the electrode. Full article
(This article belongs to the Special Issue Functional Ceramics for Energy Applications)
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Full article ">Figure 1
<p>Stages at which the BZCY72-Pd/Cu electrode interface was investigated with advanced characterization including time-of-flight secondary-ion mass spectrometry (TOF-SIMS), transmission electron microscopy (TEM), scanning spreading-resistance microscopy (SSRM), and atom-probe tomography (APT).</p> Full article ">Figure 2
<p>Procedure for the catalytic membrane reactor test on the BZCY72/Ni//BZCY72//PdCu tubular sample.</p> Full article ">Figure 3
<p>SSRM results of a polished cross-section of the completed device before testing (Step 4 in <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>) and related TEM results from the same processing step. (<b>A</b>) SSRM map (10 µm × 10 µm) of the two-phase composition scaffold; the high-conductivity Ni phase appears bright. (<b>B</b>) SSRM map (10 µm × 10 µm) of the dense BZCY membrane; a high-conductivity phase is seen at the grain boundaries. (<b>C</b>) SSRM map of the BZCY/copper electrode interface. No evidence for a phase with different electronic conductivity is observed between the BZCY/and copper electrode. (<b>D</b>) STEM high-angle annular dark-field (HAADF) image of several grain boundaries in the dense BZCY/membrane. (<b>E</b>) corresponding STEM EDS map of the image in D, showing a nickel-rich phase at the grain boundaries of the BZCY electrolyte.</p> Full article ">Figure 4
<p>TOF-SIMS 3-D tomography results investigating a cross-section of the BaZr<sub>0.7</sub>Ce<sub>0.2</sub>Y<sub>0.1</sub>O<sub>3-δ</sub> (BZCY72)/Ni 2-phase composite support for the dense BZCY electrolyte, illustrating the strength of TOF-SIMS to elucidate the distribution of impurities in the materials. (<b>A</b>) Barium signal (from BZCY ceramic phase); (<b>B</b>) Nickel signal; (<b>C</b>) Impurity signal seen accumulating at the interface between BZCY and Ni in the composite (sodium in this case).</p> Full article ">Figure 5
<p>Backscattered electron images (<b>A</b>&amp;<b>C</b>), and secondary electron image (<b>B</b>) of the palladium catalyst particles during various steps in the electroless deposition process. (<b>A</b>) discrete palladium particles are observed after activation and calcination (Step 2 in <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>); (<b>B</b>) discrete palladium particles are observed under the deposited copper electrode (Step 3 in <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>); (<b>C</b>) no discrete palladium particles are observed after the annealing of the copper electrode (Step 4 in <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>).</p> Full article ">Figure 6
<p>TOF-SIMS 3-D tomography results showing the palladium signal at various processing steps. The same dataset is shown in the top and bottom images for (<b>A</b>–<b>C</b>), with the top image showing the top view of the dataset, and the bottom images showing the side view. The intensity data is on a blue-to-red colorscale, which is presented at the right. (<b>A</b>,<b>B</b>) After the palladium activation step, discrete palladium particles are observed on the surface (Step 2 in <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>). (<b>C</b>,<b>D</b>) After copper plating, discrete palladium particles are still observed under the deposited copper electrode (Step 3 in <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>). (<b>E</b>,<b>F</b>) After the annealing of the copper electrode (Step 4 in <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>), no discrete palladium particles are observed; the palladium appears to have alloyed with the copper electrode.</p> Full article ">Figure 7
<p>(<b>A</b>) Transmission electron microscopy image and (<b>B</b>) atom probe tomography reconstruction of the interface between the copper electrode (orange) and the dense BZCY electrolyte (green) for a sample after copper plating (Step 3 of <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>). A carbon-rich phase (red) is observed at the interface between the electrode and the electrolyte, consistent with the TOF-SIMS data.</p> Full article ">Figure 8
<p>TOF-SIMS 3-D tomography results showing the signal from a carbon-rich phase (C<sub>2</sub><sup>+</sup> secondary ion signal) at various processing steps. The same dataset is shown in the top and bottom images for (<b>A</b>–<b>D</b>), with the top image showing the top view of the dataset, and the bottom image showing the side view. The intensity data is on a blue-to-red colorscale, which is presented at the right. (<b>A</b>) After the palladium activation step, the carbon-rich phase is uniformly distributed near the surface palladium layer. (<b>B</b>) After the copper plating step, the carbon-rich phase persists through the electrode thickness. (<b>C</b>) After the anneal in reducing atmosphere, the carbon-rich phase persists as uniformly distributed through the electrode thickness. (<b>D</b>) After testing, the carbon-rich phase is more concentrated at the electrode/BZCY interface.</p> Full article ">Figure 9
<p>TOF-SIMS 3-D tomography results showing the signal from a carbon-rich phase (C<sub>2</sub><sup>+</sup> secondary ion signal) at various processing steps for a copper electrode deposited by the electroless method described in this work, as well as the conventional means of painted copper paste. The same dataset is shown in the top and bottom images for (<b>A</b>–<b>D</b>), with the top image showing the top view of the dataset, and the bottom image showing the side view. The intensity data is on a blue-to-red colorscale, which is presented at the right. (<b>A</b>) Electroless copper after hydrogen gas anneal (Step 4 of <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>); (<b>B</b>) Painted copper after a similar hydrogen gas anneal; the amount and distribution of the carbon-rich phase is vastly different compared to A; (<b>C</b>) Electroless copper after testing; (<b>D</b>) Painted copper after testing; the amount and distribution of the carbon-rich phase is vastly different compared to <a href="#ceramics-01-00021-f009" class="html-fig">Figure 9</a>C. These results suggest that the carbon-rich phase observed here is unique to the electroless copper deposition process and is not simply the result of coking during testing.</p> Full article ">Figure 10
<p>TEM and STEM EDS results of a carbon-rich phase located between the copper electrode and the BZCY electrolyte after testing (Step 5 of <a href="#ceramics-01-00021-f001" class="html-fig">Figure 1</a>). The carbon rich phase appears non-continuous due to preferential milling with the FIB during TEM specimen preparation. The HAADF STEM image is in the upper left, and the STEM EDS maps of the same area (each frame labeled with the signal) are shown in the subsequent frames. The apparent low concentration of copper observed in the BZCY layer is the result of a background signal in BZCY.</p> Full article ">
15 pages, 4017 KiB  
Article
Hydrothermal Synthesis, Characterization, and Sintering Behavior of Core-Shell Particles: A Principle Study on Lanthanum Strontium Cobaltite Coated with Nanosized Gadolinium Doped Ceria
by Yu Xu, Philipp Zielke, Ngo Van Nong, Stéven Pirou, Raquel Reolon, Xiaoqing Si, Søren Bredmose Simonsen, Poul Norby, Henning Lühmann, Wolfgang Bensch and Ragnar Kiebach
Ceramics 2018, 1(2), 246-260; https://doi.org/10.3390/ceramics1020020 - 2 Oct 2018
Cited by 3 | Viewed by 4603
Abstract
In this work, nanostructured (La0.6Sr0.4)0.99CoO3 (LSC)-Ce0.8Gd0.2O1.9 (CGO) core-shell particles were prepared by precipitating CGO nanoparticles on the surface of LSC particles under hydrothermal conditions. The as-prepared core-shell particles were sintered by [...] Read more.
In this work, nanostructured (La0.6Sr0.4)0.99CoO3 (LSC)-Ce0.8Gd0.2O1.9 (CGO) core-shell particles were prepared by precipitating CGO nanoparticles on the surface of LSC particles under hydrothermal conditions. The as-prepared core-shell particles were sintered by spark plasma sintering (SPS) and conventional sintering, and the microstructure evolution and densification behavior were studied. Dense microstructures were reached using both sintering methods at relatively low temperatures. In the case of SPS, the core-shell architecture was partially maintained and nano-structured CGO grains were formed, while conventional sintering led to the formation of larger CGO grains. This work covers a detailed characterization of (a) the individual LSC-CGO core-shell particles and (b) the composites after densification. Full article
(This article belongs to the Special Issue Functional Ceramics for Energy Applications)
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Graphical abstract

Graphical abstract
Full article ">Figure 1
<p>(<b>a</b>) Rietveld refinement of the powder X-ray diffraction of the commercial lanthanum strontium cobaltite (LSC) powder; the observed pattern, the calculated pattern, Bragg positions and the differential profile are presented by red circles, the black line, blue vertical bars and the green line, respectively; (<b>b</b>) an SE2 image of the LSC powders; (<b>c</b>) Rietveld refinement of the powder X-ray diffraction of the particle composed of LSC core and the shell of gadolinium-doped ceria (CGO) before sintering; red vertical bars corresponding to Bragg positions of cubic CGO are added; (<b>d</b>) an SE2 image of the LSC-CGO particles before sintering.</p> Full article ">Figure 2
<p>(<b>a</b>) A BF-TEM image of the LSC-CGO core-shell particles; the first inset is an FFT of the shell composed of nanosized CGO particles randomly oriented; the second inset is a high-resolution image of a CGO nanoparticle in the shell. (<b>b</b>) a DF-STEM image of LSC-CGO particles; (<b>c</b>) Mapping of all involved metallic elements by STEM-EDS; X-ray photons from La Lα, Sr Kα, Co Kα, Ce Lα and Gd Lα.</p> Full article ">Figure 3
<p>Powder X-ray diffraction patterns of (<b>a</b>) LSC-CGO composite densified by SPS of core-shell particles; (<b>b</b>) LSC-CGO composite densified by conventional sintering of core-shell particles; (<b>c</b>) LSC-CGO core-shell particles before sintering.</p> Full article ">Figure 4
<p>(<b>a</b>) A BSE image of the polished cross-section of LSC-CGO densified by spark plasma sintering; (<b>b</b>) segmentation of the CGO phase observed in (<b>a</b>); (<b>c</b>) a BSE image of the polished cross-section of LSC-CGO densified by conventional sintering; (<b>d</b>) segmentation of the CGO phase observed in (<b>c</b>).</p> Full article ">Figure 5
<p>(<b>a</b>) A BSE image of the cross-section of the LSC-CGO composite densified by SPS recorded with a primary energy of 5 kV; (<b>b</b>–<b>f</b>) SEM-EDS elemental maps of La L, Sr L, Co L, Ce L and Gd M.</p> Full article ">Figure 6
<p>A BSE image of the polished cross-section of the LSC-CGO sample by conventional sintering recorded with a primary energy of 15 kV and SEM-EDS element line distributions along two selected directions as indicated.</p> Full article ">Figure 7
<p>(<b>a</b>) XRD pattern of the SPS-densified LSC-CGO composite after the thermal cycle; (<b>b</b>) a BSE micrograph of the polished cross-section of the sample.</p> Full article ">
17 pages, 9675 KiB  
Article
Exploring the Processing of Tubular Chromite- and Zirconia-Based Oxygen Transport Membranes
by Astri Bjørnetun Haugen, Lev Martinez Aguilera, Kawai Kwok, Tesfaye Molla, Kjeld Bøhm Andersen, Stéven Pirou, Andreas Kaiser, Peter Vang Hendriksen and Ragnar Kiebach
Ceramics 2018, 1(2), 229-245; https://doi.org/10.3390/ceramics1020019 - 29 Sep 2018
Cited by 5 | Viewed by 4384
Abstract
Tubular oxygen transport membranes (OTMs) that can be directly integrated in high temperature processes have a large potential to reduce CO2 emissions. However, the challenging processing of these multilayered tubes, combined with strict material stability requirements, has so far hindered such a [...] Read more.
Tubular oxygen transport membranes (OTMs) that can be directly integrated in high temperature processes have a large potential to reduce CO2 emissions. However, the challenging processing of these multilayered tubes, combined with strict material stability requirements, has so far hindered such a direct integration. We have investigated if a porous support based on (Y2O3)0.03(ZrO2)0.97 (3YSZ) with a dense composite oxygen membrane consisting of (Y2O3)0.01(Sc2O3)0.10(ZrO2)0.89 (10Sc1YSZ) as an ionic conductor and LaCr0.85Cu0.10Ni0.05O3?? (LCCN) as an electronic conductor could be fabricated as a tubular component, since these materials would provide outstanding chemical and mechanical stability. Tubular components were made by extrusion, dip coating, and co-sintering, and their chemical and mechanical integrity was evaluated. Sufficient gas permeability (?10?14 m2) and mechanical strength (?50 MPa) were achieved with extruded 3YSZ porous support tubes. The high co-sintering temperature required to densify the 10ScYSZ/LCCN membrane on the porous support, however, causes challenges related to the evaporation of chromium from the membrane. This chemical degradation caused loss of the LCCN electronic conducting phase and the formation of secondary lanthanum zirconate compounds and fractures. LCCN is therefore not suitable as the electronic conductor in a tubular OTM, unless means to lower the sintering temperature and reduce the chromium evaporation are found that are applicable to the large-scale fabrication of tubular components. Full article
(This article belongs to the Special Issue Functional Ceramics for Energy Applications)
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Figure 1

Figure 1
<p>Flexural strength of tubular, porous 3YSZ supports plotted vs. (<b>a</b>) sintering temperature and (<b>b</b>) porosity. (<b>c</b>) Weibull plot of room temperature strength of specimens with 61.5 vol % pore former and sintered at 1300 °C.</p> Full article ">Figure 2
<p>Micrographs of the porous 3YSZ support tubes with 61.5 vol % pore formers after sintering for 2 h at different temperatures. (<b>a</b>,<b>b</b>) 1250 °C, (<b>c</b>) 1300 °C, and (<b>d</b>) 1400 °C.</p> Full article ">Figure 3
<p>Dilatometry of 3YSZ and 30/70 vol % LCCN/10Sc1YSZ as pellets and as the actual thermoplastic feedstock with 61.5 vol % pore formers for the support as a tube, and the 30/70 vol % LCCN/10Sc1YSZ as a casted tape rolled up to a cylinder.</p> Full article ">Figure 4
<p>Photographs of (<b>a</b>) as-extruded support tubes, (<b>b</b>) tubes coated with inner activation layer and membrane layer after calcination and pre-sintering (the white parts on the ends are uncoated support), and (<b>c</b>) tubes after application of all layers (the white parts on the ends are uncoated support, the white part in the middle is the outer 10Sc1YSZ activation layer. The ruler shows length in millimeters).</p> Full article ">Figure 5
<p>SEM micrographs of the tubular OTM layers, from top to bottom outer activation layer, composite membrane, inner activation layer, and porous support, after sintering at (<b>a</b>) 1400 °C, 2 h (leaks); (<b>b</b>) 1425 °C, 2 h (leaks); (<b>c</b>) 1425 °C, 6 h (tight); and (<b>d</b>) 1450 °C, 6 h (tight, with some leaky spots).</p> Full article ">Figure 6
<p>Effect of thickness on crack formation in the tubular OTMs. (<b>a</b>) Transversal crack in thick membranes. (<b>b</b>) Longitudinal cracks and (<b>c</b>) thin membrane without cracks, but with severe loss of electronic conducting phase.</p> Full article ">Figure 7
<p>Leakage dependence on thickness and density of the LCCN/10Sc1YSZ asymmetric membranes, with a qualitative estimate of regimes where the membrane leaks or is leak tight.</p> Full article ">Figure 8
<p>Micrograph of the top surface of a tubular OTM (without outer porous activation layer) showing decomposition of the LCCN phase into Cr and La rich regions.</p> Full article ">Figure 9
<p>SEM image (left) and EDS signals of La and Cr from a close-up of the tubular OTMs, showing evaporation of Cr and coarsening of the nominal LCCN grains (red circle) from the outer side from the crack (above the red line).</p> Full article ">Figure 10
<p>SEM image (left) and EDS signals from Cr, Zr, and La from the same area of a cross section of the tubular OTMs with outer porous activation layer containing 60/40 vol % LCCN/10Sc1YSZ. The yellow line is a guide for the eye.</p> Full article ">
18 pages, 637 KiB  
Article
Elastic and Dielectric Evaluation of the Piezoelectric Response of Ferroelectrics Using Unpoled Ceramics
by Francesco Cordero
Ceramics 2018, 1(2), 211-228; https://doi.org/10.3390/ceramics1020018 - 27 Sep 2018
Cited by 5 | Viewed by 4557
Abstract
The evaluation of the piezoelectric properties of ferroelectric ceramics generally has a high level of uncertainty, due to incomplete poling, porosity, domain wall clamping and other effects. In addition, the poling process is often difficult and dangerous, due to the risk of breaking [...] Read more.
The evaluation of the piezoelectric properties of ferroelectric ceramics generally has a high level of uncertainty, due to incomplete poling, porosity, domain wall clamping and other effects. In addition, the poling process is often difficult and dangerous, due to the risk of breaking or damaging the sample. A method is described for the evaluation of the potential intrinsic piezoelectric response that a ceramic would have after full poling, without poling it. The method relies on the fact that any material undergoes an elastic softening below the ferroelectric transition temperature, whose magnitude can be expressed in terms of the intrinsic piezoelectric and dielectric coefficients of the material. Such a softening is equivalent to an electromechanical coupling factor averaged over all the components, due to the unpoled state of the sample, and can be deduced from a single temperature scan of an elastic modulus of a ceramic sample, spanning the ferroelectric and paraelectric states. The strengths, limits and possible applications of the method are discussed. Full article
(This article belongs to the Special Issue Functional Ceramics for Energy Applications)
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Figure 1

Figure 1
<p>The two mechanisms causing the piezoelectric softening: direct and converse piezoelectric effects arising from the electrostrictive coupling with spontaneous polarization <math display="inline"><semantics> <mrow> <msub> <mi>P</mi> <mn>0</mn> </msub> <mo>.</mo> </mrow> </semantics></math></p> Full article ">Figure 2
<p>Various factors that determine the actual softening in the FE phase: (<b>a</b>) piezoelectric softening within the FE phase; (<b>b</b>) linear anharmonic stiffening of the background compliance; (<b>c</b>) fluctuations and thermoelastic effect; (<b>d</b>) additional terms in the FE free energy with respect to the simplest expansion (<a href="#FD11-ceramics-01-00018" class="html-disp-formula">11</a>).</p> Full article ">Figure 3
<p>Compliance and dielectric permittivity of PbZr<math display="inline"><semantics> <msub> <mrow/> <mrow> <mn>0.86</mn> </mrow> </msub> </semantics></math>Ti<math display="inline"><semantics> <msub> <mrow/> <mrow> <mn>0.14</mn> </mrow> </msub> </semantics></math>O<math display="inline"><semantics> <msub> <mrow/> <mn>3</mn> </msub> </semantics></math>, with anomalies at <math display="inline"><semantics> <msub> <mi>T</mi> <mi mathvariant="normal">C</mi> </msub> </semantics></math> and the octahedral tilting transition at <math display="inline"><semantics> <msub> <mi>T</mi> <mi mathvariant="normal">T</mi> </msub> </semantics></math> (data from Ref. [<a href="#B31-ceramics-01-00018" class="html-bibr">31</a>])</p> Full article ">Figure 4
<p>(<b>a</b>) Compliance of three samples of BaTiO<math display="inline"><semantics> <msub> <mrow/> <mn>3</mn> </msub> </semantics></math> from different laboratories and with different porosities; (<b>b</b>) After rescaling to the stiffer sample, with smaller porosity, in the PE phase; the dashed line is <math display="inline"><semantics> <mrow> <msup> <mi>s</mi> <mn>0</mn> </msup> <mfenced open="(" close=")"> <mi>T</mi> </mfenced> </mrow> </semantics></math> extrapolated from &gt;800 K; (<b>c</b>) After subtraction of the extrapolated <math display="inline"><semantics> <mrow> <msup> <mi>s</mi> <mn>0</mn> </msup> <mfenced open="(" close=")"> <mi>T</mi> </mfenced> <mo>;</mo> </mrow> </semantics></math> curves 4 and 5 are <math display="inline"><semantics> <mrow> <mo>Δ</mo> <msup> <mi>s</mi> <mi>piezo</mi> </msup> </mrow> </semantics></math> calculated from different sets of the <math display="inline"><semantics> <mi>ϵ</mi> </semantics></math> and <span class="html-italic">d</span> tensors from the literature; (<b>d</b>) Change of the magnitude of the polarization vector <math display="inline"><semantics> <mi mathvariant="bold">P</mi> </semantics></math> at the PE/FE transition, with the corresponding longitudinal fluctuations in red, and change of the direction of <math display="inline"><semantics> <mi mathvariant="bold">P</mi> </semantics></math> at the transition between tetragonal and orthorhombic FE phases, with transverse fluctuations.</p> Full article ">Figure 5
<p>Preliminary measurement of the compliance of a sample of (Pb,Ca)TiO<math display="inline"><semantics> <msub> <mrow/> <mn>3</mn> </msub> </semantics></math>.</p> Full article ">Figure 6
<p>(<b>a</b>) Hypothetical compliance curves obtained by varying a material parameter, for example doping; (<b>b</b>) The same curves after normalization to <math display="inline"><semantics> <mrow> <msub> <mi>s</mi> <mn>0</mn> </msub> <mfenced open="(" close=")"> <mi>T</mi> </mfenced> </mrow> </semantics></math> in the PE phase, in order to remove the dependence on porosity; (<b>c</b>) Dielectric permittivity measured of the same samples or compositions; (<b>d</b>) Effective piezoelectric coefficient, with the effect of different porosities removed.</p> Full article ">Figure 7
<p>Hypothetical compliance curves obtained by doping in a manner that changes the FE transition into an orbital/charge order transition.</p> Full article ">
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