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SDC bi-layer construction
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ScienceDirect Journal of the European Ceramic Society 34 (2014) 347–354
Microstructure evaluation of Ni–SDC synthesized with an innovative method and Ni–SDC/SDC bi-layer construction Tina Skalar a,b , Klementina Zupan a,b , Marjan Marinˇsek a,b,∗ , Barbara Novosel a,b , Jadran Maˇcek a,b a
University of Ljubljana, Faculty of Chemistry and Chemical Technology, Aˇskerˇceva 5, 1000 Ljubljana, Slovenia b Center of Excellence Low-Carbon Technologies (CO-NOT), Hajdrihova 19, 1000 Ljubljana, Slovenia Received 26 July 2013; received in revised form 13 August 2013; accepted 15 August 2013 Available online 7 September 2013
Abstract Nickel oxide and samaria-doped ceria based anode material (NiO–SDC) for high temperature ceramic fuel cells was synthesized via an innovative, simplified method using the reaction of metal acetates and ethylene glycol. Ni content in the final product was 55 vol% while the Ce:Sm molar ratio in the SDC (45 vol%) was 80:20. The influences of preparation conditions including the calcination temperature and morphology characteristics of the prepared powders as well as the influence of the sintering temperature on the final microstructure of sintered bodies were investigated with EGA, BET, XRD, a heating microscope and FE-SEM. After calcination at 500 ◦ C, the one-step synthesis gave powdered products with the specific surface area of 72.5 m2 g−1 . The optimal sintering temperature that ensured relatively high sintered densities and excellent contact between grains was found at 1250 ◦ C. After sintering at 1250 ◦ C, the prepared bi-layer system electrolyte/anode showed the highest contiguity, excellent contact between phases and proper porosity (31.2%) confirmed from SEM micrographs. © 2013 Elsevier Ltd. All rights reserved. Keywords: Ni–SDC synthesis; Microstructure analysis; Sintering; Bi-layer system
1. Introduction Solid oxide fuel cells (SOFCs) are a very promising energy converter because they have high energy conversion efficiency and are environmentally friendly with respect to pollution.1–4 Conventional high temperature SOFCs operate at temperatures between 900 ◦ C and 1000 ◦ C. However, current research efforts are aimed at intermediate temperature SOFCs (IT-SOFCs) where the operating temperature of the stack ranges from 400 ◦ C to 700 ◦ C. The materials for IT-SOFCs are more efficient, reduce the thermal stresses in a cell and increase the system stability.1,5–7 Samaria-doped ceria (SDC) is a well-known advanced ceramic material due to its high oxygen ion conductivity; it can be used as an electrolyte in IT-SOFC.5,8–10,12 SDC has superb ionic conductivity compared to the conventional electrolyte YSZ. It was experimentally determined that the optimal
∗ Corresponding author at: University of Ljubljana, Faculty of Chemistry and Chemical Technology, Aˇskerˇceva 5, 1000 Ljubljana, Slovenia. Tel.: +386 1 2419 204; fax: +386 1 2419 220. E-mail address: [email protected] (M. Marinˇsek).
0955-2219/$ – see front matter © 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.jeurceramsoc.2013.08.020
ionic conductivity in the SDC solid solution is achieved when the molar ratio of Ce:Sm is 80:20.10–12 Additionally, a properly dense microstructure of the SDC electrolyte can be achieved at relatively low sintering temperatures.4,13 Ni–SDC is most commonly used for the anode side of IT-SOFC. The main reasons for Ni–SDC selection are in its relatively good ionic and electronic conductivity, acceptable catalytic activity for hydrogen or hydro-carbon electro-oxidation, and in its comparable thermal expansion coefficient with the SDC electrolyte.14–17 The anode microstructure should be optimized to ensure as many triple point boundaries (TPB) as possible, i.e. where the catalyst, ceramic phase and pores are in contact with each other.4,12 Previous works related to SDC and Ni–SDC synthesis employed various preparation paths. Mechanical mixing processes,4,12,18 spray pyrolysis techniques,15 co-precipitation reactions,4,15 gel-casting techniques,16 hydrothermal preparations,19 polymeric-complexing processes,19 and citrate/nitrate combustion synthesis4 have been mostly used to prepare small particle-sized powders. In general, it was found that material homogenization by mechanical mixing of the oxides leads to non-uniform particle distribution in the final
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material, high costs of fine particle milling, and the possibility of product contamination with impurities.19 However, the milling process is relatively easy to perform and enables the single-step preparation of the product. Due to higher product homogeneity, the wet-chemical techniques were recognized as more suitable for mixed oxide synthesis.20 Starting with a solution of metal ions, it is possible to manipulate the synthesis parameters at the ionic or molecular level, achieving final distribution of metal oxides at the nanoscale.21,22 Furthermore, it was found that the presence of fine SDC ceramic particles inhibits the excessive grain growth of NiO phase during sintering and eventually results in better homogeneity compared to mixtures prepared by mechanical mixing.20 The Pechini-type polymeric gel method seems to be the easiest sol-gel synthesis. Appropriate metal salts are dissolved in a solution of citric acid and lead into ethylene glycol to form a metal–citrate complex in which metal ions are immobilized in a polymer resin during the early stages of synthesis. This synthesis does not require special equipment, atmosphere or careful control of time and conditions. It is often used for the preparation of ferroelectric, superconducting, ferromagnetic, photocatalytic, fuel cell and other complex oxide materials.23 In this paper, we report on original, simple synthesis of SDC and NiO–SDC preparations. The main objective of the present work is to demonstrate that nano-sized, uniformly distributed and sinterable NiO–SDC dispersions can be prepared by the thermal treatment of Ni, Ce and Sm metal-organic precursors. The in-depth microstructure development from synthesized NiO–SDC nano-powder to Ni–SDC anode cermet is also described, which finally lead to successful preparation of the bi-layer electrolyte–anode (SDC/Ni–SDC) system. 2. Experimental procedure 2.1. Anode material 2.1.1. Synthesis Ce(C2 H3 O2 )3 ·xH2 O (Sigma Aldrich, 99.9% pure, metals basis, x = 1.7) and Sm(C2 H3 O2 )3 ·xH2 O (Sigma Aldrich, 99.9% pure, metals basis, x = 4.2) were dissolved in 150 mL of deionized water with an 80:20 Ce:Sm molar ratio. Water content (x) was determined by thermal analysis. After the dissolution, 100 mL of ethylene glycol (C2 H6 O2 , Sigma Aldrich, 99.5%, puriss. p.a.) was added. The mixture was heated up to 80 ◦ C and held at this temperature for 0.5 h. Subsequently, Ni(CH3 COO)2 ·4H2 O (Riedel-de Haën, min 98% purity) was added to this mixture in an amount which ensures 55 vol% of Ni in the final cermet Ni–SDC. Such a mixture was then gradually reheated to 80 ◦ C and held at 80 ◦ C for 2.5 h in a vacuum bath (Büchi Heating bath B-490 Rotavapor R-200, p = 20 mbar) to produce a bright green gel. Afterwards, the gel was air-dried at 80 ◦ C for three days to form the final crystals, which were used as an intermediate precursor for further NiO–SDC preparation. This precursor was calcined at various temperatures (500 ◦ C, 600 ◦ C and 900 ◦ C) for 1 h in an electric furnace and subsequently milled in attritor mill. The milled powders were compacted into cylindrical pellets (φ = 6 mm) with uniaxial
pressing at a pressure of 400 MPa; each pellet reached around 50% of relative green density. The pellets were sintered in air for 1 h at various temperatures (1200 ◦ C, 1250 ◦ C, 1300 ◦ C, 1350 ◦ C and 1400 ◦ C). To determine the final composite microstructure, all sintered tablets were polished, thermally etched and reduced in Ar–H2 (5 vol%) at 900 ◦ C for 3 h.
2.1.2. Characterization Thermal decomposition of the intermediate precursor was monitored with EGA analysis from ambient temperatures up to 800 ◦ C with a heating rate of 10 K min−1 in argon atmosphere using Netzsch STA 449 F3 Jupiter apparatus coupled with QMS 403 C Aëolos mass spectrometer. XRD spectra of the calcined intermediate precursor were recorded using a PANalytical X’Pert PRO MPD diffractometer in the range 2Θ from 10◦ to 70◦ in steps of 0.034◦ for 1 s/step. The crystallite sizes of powders calcined at different temperatures were determined using the Scherrer method. The specific surface areas of the calcined powders and tablets sintered at various temperatures were determined with the Brunauer–Emmett–Teller (BET) method, using an ASAP 2020Micromeritics instrument. The shrinkage curve of the pellets during sintering up to 1500 ◦ C with a constant heating rate of 10 K min−1 was obtained with a heating microscope (Leitz Wetzlar). Geometric densities in green, sintered and reduced states were determined and subsequently expressed as relative densities by comparing their values with the theoretical densities of NiO–SDC and Ni–SDC (6.913 g cm−3 and 8.154 g cm−3 , respectively). Microstructures of the sintered and reduced pellets were determined with the FE-SEM microscope (Zeiss Ultra Plus). The quantitative microstructure analysis of the sintered and reduced samples was performed on digital images (microstructure images were digitized into pixels with 255 different grey values) using Zeiss KS300 3.0 image-analysis software. The contiguities that represent a degree of contacts between grains of one or various phases in a three-phase mixture were determined using the Gurland principle,24,25 which was modified by Lee et al.26
2.2. Electrolyte material 2.2.1. Synthesis Ce(C2 H3 O2 )3 ·xH2 O (Sigma Aldrich, 99.9% pure, metals basis) and Sm(C2 H3 O2 )3 ·xH2 O (Sigma Aldrich, 99.9% pure, metals basis) were dissolved in deionized water. The molar ratio in solution was 80:20 for Ce:Sm. After the dissolution, ethylene glycol (C2 H6 O2 , Sigma Aldrich, 99.5%, puriss. p.a.) was added. The mixture was vacuum dried for 3 h (p = 20 mbar) at 80 ◦ C. This precursor was then transferred to a conventional drier and left at 80 ◦ C for three days to form tiny yellow needle-like crystals. These crystals were calcined at 900 ◦ C and milled in a Fritsch Pulverisette 5 planetary ball mill with 3 mm zirconia balls.
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2.2.2. Characterization The sintering curve of the pellet heated up to 1400 ◦ C with a constant heating rate of 10 ◦ C min−1 was obtained with a heating microscope (Leitz Wetzlar). 2.3. Bi-layer system 2.3.1. Preparation of bi-layer The prepared SDC powder was pressed into cylindrical pellets (φ = 12 mm) with 400 MPa. The tablets were pre-sintered at 1000 ◦ C in air for 1 h to achieve sufficient mechanical strength of the substrate on which the anode layer was screen-printed. The anode paste for screen-printing was composed of attritormilled NiO–SDC powder (calcined at 500 ◦ C) and organic component (a mixture of ethyl cellulose ␣-terpine-ol and (2(2-butoxyethoxy)ethyl)acetate). The weight ratio between the oxide powder and the organic component (carrier) was 60:40. The prepared bi-layer system was first co-sintered at 1250 ◦ C in air for 1 h and later reduced in Ar–H2 (5 vol%) at 900 ◦ C for 3 h. 3. Results and discussion 3.1. Intermediate anode precursor preparation and its characterization TG-EGA analysis was used to determine the thermal properties of the intermediate precursor material. The precursor decomposes to final NiO–SDC mixture in several consecutive steps through which water, acetone, acetic acid, carbon dioxide, carbon monoxide and ethylene glycol were detected as volatile products and are described by characteristic signals at m/e 18, m/e 58, m/e 60, m/e 44, m/e 28 and m/e 31, respectively (Fig. 1). To follow CO evolution path CO2 contribution was extracted from m/e signal 28. Since m/e 18 and m/e 31 were the only detected signals during the early stages of thermal analysis, any weight losses below 140 ◦ C were ascribed to the sample drying. The thermal decomposition starts at around 140 ◦ C and ends at ∼420 ◦ C. The measured overall mass loss from 30 ◦ C to 420 ◦ C is 62.7%. During the entire decomposition interval, water, acetone, CO, CO2 and ethylene glycol overlapped release may be followed. In contrast, acetic acid leaves the system only during the main mass loss interval (between 300 ◦ C and 350 ◦ C). In this temperature region, the sample loses ∼45% of its mass. The above-described decomposition pattern indicates that the transformation of the intermediate precursor reactive gel into oxide mixture NiO–SDC may not be described by simple chemical equations but rather follows a more complicated reaction path. The EGA results served also as a groundwork according to which the three different calcination temperatures (500 ◦ C, 600 ◦ C and 900 ◦ C) were determined. The role of the calcination process is to transform the intermediate precursor into a final mixture of metal oxides. Since higher temperatures of calcination also trigger prompt particle growth, the calcination temperature should be kept as low as possible.18 According to the obtained XRD patterns (Fig. 2), all calcined powders are crystalline solids composed of only two phases; the cubic NiO (PDF card number 000-75-0197) and the fluorite
Fig. 1. Thermal decomposition of intermediate precursor in Ar atmosphere (A) with EGA analysis (B and C).
Fig. 2. XRD patterns of NiO–SDC powders calcined at 500 ◦ C, 600 ◦ C and 900 ◦ C.
structure of SDC (PDF card number 000-75-0158). Metal oxides are already formed at 500 ◦ C. However, with increasing calcination temperature the degree of crystallization is increased. At 500 ◦ C, the XRD peaks are broad and the background in the spectra is still relatively high. In contrast, at 900 ◦ C the XRD peaks are narrow and the background in the spectra is negligible. From the X-ray broadening, the average crystallite sizes in calcined samples are calculated as 5.4 nm, 11.1 nm and 45.0 nm [plane SDC 1 1 1] and 15.5 nm, 44.6 nm and 49.1 nm [plane NiO 1 1 1] for samples calcined at 500 ◦ C, 600 ◦ C and 900 ◦ C, respectively. It is obvious that with
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increasing temperature the NiO phase grows faster than SDC phase. Nevertheless, taking into consideration that the process starts in a solution and rather similar values for crystallite sizes at 900 ◦ C suggests that both phases NiO and SDC should be uniformly distributed one into another. XRD analysis results are consistent with the results of thermal decomposition, because the final products (NiO and SDC) are already presented after calcination at minimum temperature (500 ◦ C). The BET specific surface measurements of NiO–SDC mixtures calcined at 500 ◦ C, 600 ◦ C and 900 ◦ C are 72.5 m2 g−1 , 20.0 m2 g−1 and 6.2 m2 g−1 , respectively. Since BET analysis enables the calculation of an average particle with a specific geometry, the diameters of spherical particles calculated from BET measurements are 11.0 nm, 43.3 nm and 139.0 nm for powders calcined at 500 ◦ C, 600 ◦ C and 900 ◦ C, respectively. By comparing these calculated values with crystallite sizes determined with the Scherrer calculation, we assume that the degree of agglomeration of individual particles is rather low at 500 ◦ C. With increasing calcination temperature, nano-particles tend to aggregate. 3.2. Sintering of anode material An impact of the calcination temperature and consequently the degree of particles’ aggregation is also reflected on the sintering behaviour of the powders. Powders calcined at 900 ◦ C milled in an attritor mill and pressed into tablets show relatively poor sinterability. The overall shrinkage during dynamic sintering up to 1500 ◦ C is only 14%. It is caused by specific surface area loss during the calcination process. In contrast, material calcined at 500 ◦ C exhibits much higher shrinkage (around 25% during dynamic sintering up to 1500 ◦ C). Due to low sinterability, samples calcined at 900 ◦ C were omitted in the further investigation. Similar findings regarding NiO–SDC sinterability were reported by Fang et al.15 Detailed sintering behaviour of NiO–SDC pellets prepared from powders calcined at 500 ◦ C were determined via an indepth microstructure analysis on sintered and reduced samples (Fig. 3) where A1, B1, C1, D1 and E1 denote samples sintered at 1200 ◦ C, 1250 ◦ C, 1300 ◦ C, 1350 ◦ C and 1400 ◦ C, respectively and notations from A2 to E2 of the same samples reduced at 900 ◦ C. The dark grey particles in Fig. 3 are NiO or Ni, the light grey phase is SDC, and the pores are black. A very important characteristic of the prepared composites, which originates from the preparation procedure, is the homogeneous distribution of
all phases. It can be seen that NiO or Ni grains are homogeneously distributed in the SDC phase. The distribution of pores in reduced samples is also uniform. Since all parameters important for exact cermet analysis are sometimes difficult to obtain simply from SEM micrographs, a detailed quantitative microstructure analysis on sintered and reduced samples was performed. For statistically reliable data at each chosen temperature, five to ten different regions on the surface of the tablets were analyzed. The results of the quantitative microstructure analysis are summarized in Tables 1 and 2. ¯ dx and dy are represented as the diameter of the Parameters d, area-analogue circle – DCIRCLE and intercept lengths x and y direction – FERETX, Y, respectively. Porosity ε was determined as geometrical porosity (from tablet dimensions). According to the results of quantitative microstructural analysis, the porosity of the sintered samples decreases with the increasing sintering temperature (Table 1). The higher sintering temperature also results in pronounced grain growth of both the NiO and SDC phases. The NiO phase grows somewhat faster and also reaches larger final dimensions.27 However, both phases remain well within submicrometer range after sintering up to 1400 ◦ C not exceeding ∼0.44 m and ∼0.38 m for NiO and SDC, respectively. Relatively small NiO and SDC grains after sintering as well as the minimum one phase dominance in sintered bodies is a consequence of the powder treatment after the synthesis which modified its morphology. The ball-milling process reduced the size of the particles and weakened the aggregation of the nanocrystallites.28 Another important parameter to follow the sintering progress is sintered density. From relative sintered density, the geometrical porosity of a sintered body can be easily calculated (Tables 1 and 2). The target value for the porosity of a sintered and reduced element is approximately 30% to ensure that pores remain a continuous phase throughout the final composite but, at the same time, that solid phases are also in good contact.29–32 It is evident from Tables 1 and 2 that sintered density increases with temperature while porosity decreases with sintering temperature. A relative sintered density 71.8% reached at 1200 ◦ C reveals that material sinters relatively poorly at that temperature. Consequently, NiO and SDC phases are not in good contact, making such material inappropriate for SOFC technology (Fig. 3). Enhanced grain growth and a relatively quick sintered density increase are observed at 1250 ◦ C. At the same time, at this temperature, the target value for porosity in reduced samples is also reached (31.2%). The shape factor does not change significantly with increasing the sintering temperatures. The most apparent differences
Table 1 Quantitative microstructure analysis of the sintered samples. Sample
A1 B1 C1 D1 E1
Geometrical porosity, ε (%)
28.2 15.2 9.0 6.1 3.7
d¯ (m)
Shape factor, ψ
dx (m)
BET surface area (m2 g−1 )
dy (m)
NiO
SDC
NiO
SDC
NiO
SDC
NiO
SDC
0.18 0.24 0.27 0.37 0.44
0.14 0.16 0.24 0.30 0.38
0.80 0.79 0.77 0.78 0.79
0.76 0.76 0.76 0.77 0.75
0.20 0.27 0.27 0.42 0.50
0.13 0.18 0.28 0.34 0.42
0.20 0.28 0.27 0.41 0.49
0.13 0.18 0.29 0.33 0.43
1.25 1.10 0.99 0.80 0.52
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Fig. 3. SEM micrographs of samples after sintering (left column) as NiO–SDC, and after reduction (right column) as Ni–SDC.
between the sintered and the reduced samples are the increased porosity of the samples and increased Ni particles relative to NiO particle size. It is a consequence of mass change when NiO is transformed to Ni while the overall volume of the composite
does not change due to the rigid framework of SDC. After the reduction, the formed Ni tends to sinter, which is the main cause that Ni grains grow relative to NiO. Ni sintering is a consequence of different surface energies of both phases, which causes the
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Table 2 Quantitative microstructure analysis of the reduced samples. Sample
A2 B2 C2 D2 E2
Geometrical porosity, ε (%)
44.7 31.2 27.4 24.6 21.8
d¯ (m)
Shape factor, ψ
dx (m)
BET surface area (m2 g−1 )
dy (m)
Ni
SDC
Ni
SDC
Ni
SDC
Ni
SDC
0.38 0.33 0.40 0.49 0.52
0.15 0.17 0.23 0.28 0.38
0.90 0.90 0.89 0.93 0.90
0.84 0.82 0.81 0.79 0.78
0.40 0.35 0.42 0.51 0.55
0.16 0.19 0.25 0.32 0.42
0.38 0.34 0.41 0.49 0.54
0.15 0.18 0.24 0.31 0.43
hetero-grains at the interface to loose adhesion of Ni to the SDC at elevated temperatures. However, the rigid SDC framework apparently suppresses excessive Ni sintering, keeping Ni grains within the submicrometer range. A more in-depth image analysis is performed using the line intercept method. Using this method, one can obtain information about the number N of contact points between grains of various phases in a unit length and calculate the corresponding contiguity values C (Fig. 4). As expected from the microstructure analysis, absolute values of NNi–SDC , NSDC–SDC and NNi–Ni reach their maximum at 1250 ◦ C. At this temperature, both phases already sinter; however, no extensive grain growth or severe phase separation is noticed. Below 1250 ◦ C, Ni and SDC phases sinter poorly, meaning that there are fewer contacts between phases. In contrast, with increasing sintering temperature above 1250 ◦ C, the number of contacts in a unit length decreases since grains of both phases grow. Relatively small one-phase dominance in the final Ni–SDC cermet at 1250 ◦ C results in the highest contiguity value CNi–SDC . However, at this temperature the contiguities CSDC–SDC , CNi–Ni are relatively low. Lower CSDC–SDC and CNi–Ni values at 1250 ◦ C are mostly due to significant increases of NNi–SDC . With higher NNi–SDC numbers, the denominator of the fraction to calculate the degree of contact of adequate phases (contiguity) is increased. Consequently, if CNi–SDC increases CSDC–SDC and CNi–Ni must decrease.
Fig. 4. Contiguity and absolute values of contacts between Ni–Ni, Ni–SDC and SDC–SDC grains in cermets submitted to various sintering temeperatures.
0.91 1.25 1.05 0.85 0.58
3.3. Construction of bi-layer system In an SOFC unit, anode material is sintered to an electrolyte where both layers have to exhibit fairly different microstructural properties. One of the main important factors in such a bilayer system is good adhesion at the layer boundary. Since 1250 ◦ C is the most appropriate sintering temperature for the prepared anode material with respect to its final microstructure, we attempted to use the same sintering temperature also for bi-layer system co-sintering. However, pure SDC and NiO–SDC sinter at somewhat different temperatures and also with different shrinkages (Fig. 5). Such dissimilar densification processes where NiO–SDC already shrinks ∼3% before SDC starts to sinter (∼900 ◦ C) also means that within the temperature interval where both materials sinter SDC undergoes ∼21% of dimensional change while NiO–SDC shrinks only an additional ∼13%. A ∼8% shrinkage mismatch from 900 ◦ C to 1250 ◦ C often led to cracks formation at the layer boundary. To bring the probability of crack formation at the layer boundary to a minimum SDC was pre-sintered. The pre-sintering temperature for SDC is chosen at 1000 ◦ C. In this manner, shrinkages of both materials are adjusted and at the same time the SDC substrate is strengthened for subsequent screen-printing of the NiO–SDC. After co-sintering at 1250 ◦ C for 1 h, the bi-layer system was reduced in Ar–H2 (5 vol%) at 900 ◦ C for 3 h. The final thickness of anode film after reduction is ∼20 m. The carefully adjusted shrinkages of both layers result in good adhesion, as confirmed
Fig. 5. Sintering curves of SDC and NiO–SDC materials calcined at 900 ◦ C.
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Fig. 6. SEM micrograph of anode–electrolyte contact in the bi-layer system.
by SEM micrograph (Fig. 6). The electrolyte at 1250 ◦ C reaches the final sintering stage (ρrelat. = 94%) and forms gas-tight electrolyte layer. It is also evident that the SDC phase in anode and electrolyte layers are continuously connected across the layer boundary, expanding the reaction zone for fuel oxidation into the anode layer. A precondition for the applicability of the anode material is successful co-sintering with the electrolyte material. Fig. 5 presents the sintering curve for SDC electrolyte material. The shrinkage of NiO–SDC after sintering at 1250 ◦ C was calculated from the relative green and sintered densities. SDC electrolyte was pre-sintered at 1000 ◦ C where 5% of shrinkage was reached; later, it was co-sintered with NiO–SDC phase at 1250 ◦ C where both materials shrank by a further 16%. The bi-layer system was later reduced in Ar–H2 (5 vol%) at 900 ◦ C for 3 h. The final thickness of the anode film after reduction was 20 m. Good contact between both components of the bi-layer was confirmed by SEM micrograph (Fig. 6). From Fig. 6, it is also evident that SDC in electrode and electrolyte layers are continuously connected, which is precondition for good ionic conductivity of anode material. According to the shrinkage curve, at 1250 ◦ C the electrolyte reaches the final sintering stage and forms a gas-tight electrolyte layer. 4. Conclusions Nano-sized NiO–SDC composites are prepared with an innovative, simple synthesis, easy to manipulate, via thermal decomposition of intermediate anode precursor. The EGA analysis indicates that the intermediate decomposes in several consecutive steps, during which water, acetone, acetic acid, carbon dioxide, carbon monoxide and ethylene glycol are detected as volatile products. After the thermal treatment of the metal acetate precursor at 500 ◦ C fine mixture of NiO–SDC with specific surface area 72.5 m2 g−1 is formed. The crystallite sizes are 5.4 nm and 15.5 nm for SDC and NiO, respectively. During sintering, the average grain sizes of both phases SDC and NiO are enlarged with the increasing sintering temperature, however not exceeding 500 nm when sintered up to 1400 ◦ C. NiO grains
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grow faster than SDC with the increasing sintering temperature. The dimensions of the final Ni grains formed during the subsequent reduction are influenced by the microstructure stability of the SDC framework established during the sintering. The smallest Ni grains, well in sub-micrometre range (d¯ = 0.33 m), as well as the highest contiguity value between Ni and SDC phases are obtained in the sample sintered at 1250 ◦ C. Smaller CNi–SDC values determined for samples sintered at temperatures lower than 1250 ◦ C or higher than 1250 ◦ C are ascribed to relatively larger Ni grains due to the non-rigid SDC phase or the increased one-phase dominance region of both phases, respectively. During the preparation of the bi-layer system SDC/NiO–SDC, the SDC substrate is strengthened before the screen-printing by pre-sintering at 1000 ◦ C. At the same time, the shrinkages of both components SDC/NiO–SDC are adjusted for the final sintering at 1250 ◦ C. After the reduction, good contact on the anode–electrolyte phase boundary with no cracks is prepared. The bi-layer system consists of a highly dense electrolyte (relative density 93%) and an anode layer with a porosity of 31.2%. It can be concluded that the present preparation procedure is one of more promising approaches for achieving homogeneous distribution, integrity and connectivity of all present phases. Acknowledgment This work was supported by the Ministry of Higher Education, Science and Technology of the Republic of Slovenia through grant P1-0175-103 and Centre of Excellence LowCarbon Technologies (CO NOT). References 1. Haile SM. Fuel cell materials and components. Acta Mater 2003;51(19):5981–6000. 2. Tietz F, Arul Raj I, Jungen W, Stöver D. High-temperature superconductor materials for contact layers in solid oxide fuel cells: I. Sintering behaviour and physical properties at operating temperatures. Acta Mater 2001;49:803–10. 3. Suzuki T, Hasan Z, Funahashi Y, Yamaguchi T, Fujishiro Y, Awano M. Impact of anode microstructure on solid oxide fuel cells. Sci Mag 2009;325:852–5. 4. Yin Y, Li S, Xia C, Meng G. Electrochemical performance of gel-cast NiO–SDC composite anodes in low-temperature SOFCs. Electrochim Acta 2006;51:2594–8. 5. Li J, Ikegami T, Mori T. Low temperature processing of dense samariumdoped CeO2 ceramics: sintering and grain growth behaviours. Acta Mater 2004;52:2221–8. 6. Misono T, Murata K, Fukui T, Chaichanawong J, Sato Ab H, Naito M. Ni–SDC cermet anode fabricated from NiO–SDC composite powder for intermediate temperature SOFC. J Power Sources 2006;157:754–7. 7. Hui S, Roller J, Yick S, Zhang X, Deces-Petit C, Xie Y, et al. A brief of the ionic conductivity enhancement for selected oxide electrolytes. J Power Sources 2007;172:493–502. 8. Wright J, Virkar A. Conductivity of porous Sm2 O3 -doped CeO2 as a function of temperature and oxygen partial pressure. J Power Sources 2011;196:6118–24. 9. Cheng J, Deng L, Zhang B, Shi P, Meng G. Properties and microstructure of NiO/SDC materials for SOFC anode applications. Rare Metals 2007;26:110–7.
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10. Chen M, Kim BH, Xu Q, Ahn BK, Kang WJ, Huang DP. Synthesis and electrical properties of Ce0.8 Sm0.2 O1.9 ceramics for IT-SOFC electrolytes by urea-combustion technique. Ceram Int 2009;35:1335–43. 11. Jung W, Park H, Kang Y, Yoon D. Lowering the sintering temperature of Gddoped ceria by mechanochemical activation, short communication. Ceram Int 2010;36:371–4. 12. Liu Q, Dong X, Yang C, Ma S, Chen F. Self-rising synthesis of Ni–SDC cermets as anodes for solid oxide fuel cells. J Power Sources 2010;195:1543–50. 13. Shih S, Li G, Cockayne D, Borisenko K. Mechanism of dopant distribution: an example of nickel-doped ceria nanoparticles. Scripta Mater 2009;61:832–5. 14. Chen M, Kim BH, Xu Q, Nam OJ, Ko JH. Synthesis and performances of Ni–SDC cermets for IT-SOFC anode. J Eur Ceram Soc 2008;28: 2947–53. 15. Fang X, Zhu G, Xia C, Liu X, Meng G. Synthesis and properties of Ni–SDC cermets for IT-SOFC anode by co-precipitation. Solid State Ionics 2004;168:31–6. 16. Yin Y, Zhu W, Xia C, Meng G. Gel-cast NiO–SDC composites as anodes for solid oxide fuel cells. J Power Sources 2004;132:36–41. 17. Grgicak CM, Green RG, Du WF, Giorgi JB. Synthesis and characterization of NiO–YSZ anode materials: precipitation, calcination and the effects on sintering. J Am Ceram Soc 2005;88:3081–7. 18. Hong JE, Inagaki T, Ida S, Ishihara T. Improved power generation performance of solid oxide fuel cells using doped LaGaO3 electrolyte films prepared by screen printing method II. Optimization of Ni-Ce0.8 Sm0.2 O1.9 cermet anode support. Int J Hydrogen Energ 2011;36:14632–42. 19. Shao Z, Zhou W, Zhu Z. Advanced synthesis of materials for intermediatetemperature solid oxide fuel cells. Prog Mater Sci 2012;57:804–74. 20. Han RK, Jeong Y, Lee H, Kim CS. Fabrication of NiO/YSZ anode material for SOFC via mixed NiO precursors. Mater Lett 2007;61:1242–5.
21. Grzebielucka EC, Chinelatto ASA, Tebcherani SM, Chinelatto AL. Synthesis and sintering of Y2 O3 -doped ZrO2 powders using two Pechini-type gel routes. Short communication. Ceram Int 2010;36:1737–42. 22. Cheng JG, Deng LP, Xiong ET, Shi P. NiO–SDC powder of solid oxide fuel cell anode applications by buffer solution method. Key Eng Mat 2007;336–338:440–3. 23. Sakka S. Handbook of sol–gel science and technology: processing characterization and applications. Kluwer Academic Publishers Group; 2004. 24. Gurland J. The measurements of grain contiguity in two phase alloys. Trans Metall Soc 1958;212:452–5. 25. Gurland J. An estimate of contact and contiguity of dispersions in opaque sample. Trans Metall Soc 1966;236:642–6. 26. Lee JH, Moon H, Lee HW, Kim J, Kim JD, Yoon KH. Quantitative analysis of microstructure and its related electrical property of SOFC anode, Ni–YSZ cermet. Solid State Ionics 2002;148:15–26. 27. Marinˇsek M, Zupan K, Maˇcek J. Ni–YSZ cermet anodes prepared by citrate/nitrate combustion synthesis. J Power Sources 2002;106:178–88. 28. Xu Q, Huang DP, Zhao K, Chen W, Chen M, Kim BH. Powder morphology and sinterability improvement of Ce0.8 Sm0.2 O1.9 derived from solution combustion process. Ceram Int 2011;37:913–20. 29. Hu W, Guan H, Sun X, Li S, Fukumoto M, Okane I. Electrical and thermal conductivities of nickel–zirconia cermets. J Am Ceram Soc 1998;81:2209–12. 30. Dees DW, Claar TD, Easler TE, Fee DC, Mrazek FC. Conductivity of porous Ni/ZrO2 –Y2 O3 cermets. J Electrochem Soc 1987;134:2141–6. 31. Kawashima T, Hishinuma M. Analysis of electrical conduction paths in Ni/YSZ participate composites using percolation theory. Mater Trans 1996;37:1397–403. 32. Anselmi-Tamburini U, Chiodelli G, Arimondi M, Maglia F, Spinolo G, Munir ZA. Electrical properties of Ni/YSZ cermets obtained by combustion synthesis. Solid State Ionics 1998;110:35–43.
ScienceDirect Journal of the European Ceramic Society 34 (2014) 347–354
Microstructure evaluation of Ni–SDC synthesized with an innovative method and Ni–SDC/SDC bi-layer construction Tina Skalar a,b , Klementina Zupan a,b , Marjan Marinˇsek a,b,∗ , Barbara Novosel a,b , Jadran Maˇcek a,b a
University of Ljubljana, Faculty of Chemistry and Chemical Technology, Aˇskerˇceva 5, 1000 Ljubljana, Slovenia b Center of Excellence Low-Carbon Technologies (CO-NOT), Hajdrihova 19, 1000 Ljubljana, Slovenia Received 26 July 2013; received in revised form 13 August 2013; accepted 15 August 2013 Available online 7 September 2013
Abstract Nickel oxide and samaria-doped ceria based anode material (NiO–SDC) for high temperature ceramic fuel cells was synthesized via an innovative, simplified method using the reaction of metal acetates and ethylene glycol. Ni content in the final product was 55 vol% while the Ce:Sm molar ratio in the SDC (45 vol%) was 80:20. The influences of preparation conditions including the calcination temperature and morphology characteristics of the prepared powders as well as the influence of the sintering temperature on the final microstructure of sintered bodies were investigated with EGA, BET, XRD, a heating microscope and FE-SEM. After calcination at 500 ◦ C, the one-step synthesis gave powdered products with the specific surface area of 72.5 m2 g−1 . The optimal sintering temperature that ensured relatively high sintered densities and excellent contact between grains was found at 1250 ◦ C. After sintering at 1250 ◦ C, the prepared bi-layer system electrolyte/anode showed the highest contiguity, excellent contact between phases and proper porosity (31.2%) confirmed from SEM micrographs. © 2013 Elsevier Ltd. All rights reserved. Keywords: Ni–SDC synthesis; Microstructure analysis; Sintering; Bi-layer system
1. Introduction Solid oxide fuel cells (SOFCs) are a very promising energy converter because they have high energy conversion efficiency and are environmentally friendly with respect to pollution.1–4 Conventional high temperature SOFCs operate at temperatures between 900 ◦ C and 1000 ◦ C. However, current research efforts are aimed at intermediate temperature SOFCs (IT-SOFCs) where the operating temperature of the stack ranges from 400 ◦ C to 700 ◦ C. The materials for IT-SOFCs are more efficient, reduce the thermal stresses in a cell and increase the system stability.1,5–7 Samaria-doped ceria (SDC) is a well-known advanced ceramic material due to its high oxygen ion conductivity; it can be used as an electrolyte in IT-SOFC.5,8–10,12 SDC has superb ionic conductivity compared to the conventional electrolyte YSZ. It was experimentally determined that the optimal
∗ Corresponding author at: University of Ljubljana, Faculty of Chemistry and Chemical Technology, Aˇskerˇceva 5, 1000 Ljubljana, Slovenia. Tel.: +386 1 2419 204; fax: +386 1 2419 220. E-mail address: [email protected] (M. Marinˇsek).
0955-2219/$ – see front matter © 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.jeurceramsoc.2013.08.020
ionic conductivity in the SDC solid solution is achieved when the molar ratio of Ce:Sm is 80:20.10–12 Additionally, a properly dense microstructure of the SDC electrolyte can be achieved at relatively low sintering temperatures.4,13 Ni–SDC is most commonly used for the anode side of IT-SOFC. The main reasons for Ni–SDC selection are in its relatively good ionic and electronic conductivity, acceptable catalytic activity for hydrogen or hydro-carbon electro-oxidation, and in its comparable thermal expansion coefficient with the SDC electrolyte.14–17 The anode microstructure should be optimized to ensure as many triple point boundaries (TPB) as possible, i.e. where the catalyst, ceramic phase and pores are in contact with each other.4,12 Previous works related to SDC and Ni–SDC synthesis employed various preparation paths. Mechanical mixing processes,4,12,18 spray pyrolysis techniques,15 co-precipitation reactions,4,15 gel-casting techniques,16 hydrothermal preparations,19 polymeric-complexing processes,19 and citrate/nitrate combustion synthesis4 have been mostly used to prepare small particle-sized powders. In general, it was found that material homogenization by mechanical mixing of the oxides leads to non-uniform particle distribution in the final
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material, high costs of fine particle milling, and the possibility of product contamination with impurities.19 However, the milling process is relatively easy to perform and enables the single-step preparation of the product. Due to higher product homogeneity, the wet-chemical techniques were recognized as more suitable for mixed oxide synthesis.20 Starting with a solution of metal ions, it is possible to manipulate the synthesis parameters at the ionic or molecular level, achieving final distribution of metal oxides at the nanoscale.21,22 Furthermore, it was found that the presence of fine SDC ceramic particles inhibits the excessive grain growth of NiO phase during sintering and eventually results in better homogeneity compared to mixtures prepared by mechanical mixing.20 The Pechini-type polymeric gel method seems to be the easiest sol-gel synthesis. Appropriate metal salts are dissolved in a solution of citric acid and lead into ethylene glycol to form a metal–citrate complex in which metal ions are immobilized in a polymer resin during the early stages of synthesis. This synthesis does not require special equipment, atmosphere or careful control of time and conditions. It is often used for the preparation of ferroelectric, superconducting, ferromagnetic, photocatalytic, fuel cell and other complex oxide materials.23 In this paper, we report on original, simple synthesis of SDC and NiO–SDC preparations. The main objective of the present work is to demonstrate that nano-sized, uniformly distributed and sinterable NiO–SDC dispersions can be prepared by the thermal treatment of Ni, Ce and Sm metal-organic precursors. The in-depth microstructure development from synthesized NiO–SDC nano-powder to Ni–SDC anode cermet is also described, which finally lead to successful preparation of the bi-layer electrolyte–anode (SDC/Ni–SDC) system. 2. Experimental procedure 2.1. Anode material 2.1.1. Synthesis Ce(C2 H3 O2 )3 ·xH2 O (Sigma Aldrich, 99.9% pure, metals basis, x = 1.7) and Sm(C2 H3 O2 )3 ·xH2 O (Sigma Aldrich, 99.9% pure, metals basis, x = 4.2) were dissolved in 150 mL of deionized water with an 80:20 Ce:Sm molar ratio. Water content (x) was determined by thermal analysis. After the dissolution, 100 mL of ethylene glycol (C2 H6 O2 , Sigma Aldrich, 99.5%, puriss. p.a.) was added. The mixture was heated up to 80 ◦ C and held at this temperature for 0.5 h. Subsequently, Ni(CH3 COO)2 ·4H2 O (Riedel-de Haën, min 98% purity) was added to this mixture in an amount which ensures 55 vol% of Ni in the final cermet Ni–SDC. Such a mixture was then gradually reheated to 80 ◦ C and held at 80 ◦ C for 2.5 h in a vacuum bath (Büchi Heating bath B-490 Rotavapor R-200, p = 20 mbar) to produce a bright green gel. Afterwards, the gel was air-dried at 80 ◦ C for three days to form the final crystals, which were used as an intermediate precursor for further NiO–SDC preparation. This precursor was calcined at various temperatures (500 ◦ C, 600 ◦ C and 900 ◦ C) for 1 h in an electric furnace and subsequently milled in attritor mill. The milled powders were compacted into cylindrical pellets (φ = 6 mm) with uniaxial
pressing at a pressure of 400 MPa; each pellet reached around 50% of relative green density. The pellets were sintered in air for 1 h at various temperatures (1200 ◦ C, 1250 ◦ C, 1300 ◦ C, 1350 ◦ C and 1400 ◦ C). To determine the final composite microstructure, all sintered tablets were polished, thermally etched and reduced in Ar–H2 (5 vol%) at 900 ◦ C for 3 h.
2.1.2. Characterization Thermal decomposition of the intermediate precursor was monitored with EGA analysis from ambient temperatures up to 800 ◦ C with a heating rate of 10 K min−1 in argon atmosphere using Netzsch STA 449 F3 Jupiter apparatus coupled with QMS 403 C Aëolos mass spectrometer. XRD spectra of the calcined intermediate precursor were recorded using a PANalytical X’Pert PRO MPD diffractometer in the range 2Θ from 10◦ to 70◦ in steps of 0.034◦ for 1 s/step. The crystallite sizes of powders calcined at different temperatures were determined using the Scherrer method. The specific surface areas of the calcined powders and tablets sintered at various temperatures were determined with the Brunauer–Emmett–Teller (BET) method, using an ASAP 2020Micromeritics instrument. The shrinkage curve of the pellets during sintering up to 1500 ◦ C with a constant heating rate of 10 K min−1 was obtained with a heating microscope (Leitz Wetzlar). Geometric densities in green, sintered and reduced states were determined and subsequently expressed as relative densities by comparing their values with the theoretical densities of NiO–SDC and Ni–SDC (6.913 g cm−3 and 8.154 g cm−3 , respectively). Microstructures of the sintered and reduced pellets were determined with the FE-SEM microscope (Zeiss Ultra Plus). The quantitative microstructure analysis of the sintered and reduced samples was performed on digital images (microstructure images were digitized into pixels with 255 different grey values) using Zeiss KS300 3.0 image-analysis software. The contiguities that represent a degree of contacts between grains of one or various phases in a three-phase mixture were determined using the Gurland principle,24,25 which was modified by Lee et al.26
2.2. Electrolyte material 2.2.1. Synthesis Ce(C2 H3 O2 )3 ·xH2 O (Sigma Aldrich, 99.9% pure, metals basis) and Sm(C2 H3 O2 )3 ·xH2 O (Sigma Aldrich, 99.9% pure, metals basis) were dissolved in deionized water. The molar ratio in solution was 80:20 for Ce:Sm. After the dissolution, ethylene glycol (C2 H6 O2 , Sigma Aldrich, 99.5%, puriss. p.a.) was added. The mixture was vacuum dried for 3 h (p = 20 mbar) at 80 ◦ C. This precursor was then transferred to a conventional drier and left at 80 ◦ C for three days to form tiny yellow needle-like crystals. These crystals were calcined at 900 ◦ C and milled in a Fritsch Pulverisette 5 planetary ball mill with 3 mm zirconia balls.
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2.2.2. Characterization The sintering curve of the pellet heated up to 1400 ◦ C with a constant heating rate of 10 ◦ C min−1 was obtained with a heating microscope (Leitz Wetzlar). 2.3. Bi-layer system 2.3.1. Preparation of bi-layer The prepared SDC powder was pressed into cylindrical pellets (φ = 12 mm) with 400 MPa. The tablets were pre-sintered at 1000 ◦ C in air for 1 h to achieve sufficient mechanical strength of the substrate on which the anode layer was screen-printed. The anode paste for screen-printing was composed of attritormilled NiO–SDC powder (calcined at 500 ◦ C) and organic component (a mixture of ethyl cellulose ␣-terpine-ol and (2(2-butoxyethoxy)ethyl)acetate). The weight ratio between the oxide powder and the organic component (carrier) was 60:40. The prepared bi-layer system was first co-sintered at 1250 ◦ C in air for 1 h and later reduced in Ar–H2 (5 vol%) at 900 ◦ C for 3 h. 3. Results and discussion 3.1. Intermediate anode precursor preparation and its characterization TG-EGA analysis was used to determine the thermal properties of the intermediate precursor material. The precursor decomposes to final NiO–SDC mixture in several consecutive steps through which water, acetone, acetic acid, carbon dioxide, carbon monoxide and ethylene glycol were detected as volatile products and are described by characteristic signals at m/e 18, m/e 58, m/e 60, m/e 44, m/e 28 and m/e 31, respectively (Fig. 1). To follow CO evolution path CO2 contribution was extracted from m/e signal 28. Since m/e 18 and m/e 31 were the only detected signals during the early stages of thermal analysis, any weight losses below 140 ◦ C were ascribed to the sample drying. The thermal decomposition starts at around 140 ◦ C and ends at ∼420 ◦ C. The measured overall mass loss from 30 ◦ C to 420 ◦ C is 62.7%. During the entire decomposition interval, water, acetone, CO, CO2 and ethylene glycol overlapped release may be followed. In contrast, acetic acid leaves the system only during the main mass loss interval (between 300 ◦ C and 350 ◦ C). In this temperature region, the sample loses ∼45% of its mass. The above-described decomposition pattern indicates that the transformation of the intermediate precursor reactive gel into oxide mixture NiO–SDC may not be described by simple chemical equations but rather follows a more complicated reaction path. The EGA results served also as a groundwork according to which the three different calcination temperatures (500 ◦ C, 600 ◦ C and 900 ◦ C) were determined. The role of the calcination process is to transform the intermediate precursor into a final mixture of metal oxides. Since higher temperatures of calcination also trigger prompt particle growth, the calcination temperature should be kept as low as possible.18 According to the obtained XRD patterns (Fig. 2), all calcined powders are crystalline solids composed of only two phases; the cubic NiO (PDF card number 000-75-0197) and the fluorite
Fig. 1. Thermal decomposition of intermediate precursor in Ar atmosphere (A) with EGA analysis (B and C).
Fig. 2. XRD patterns of NiO–SDC powders calcined at 500 ◦ C, 600 ◦ C and 900 ◦ C.
structure of SDC (PDF card number 000-75-0158). Metal oxides are already formed at 500 ◦ C. However, with increasing calcination temperature the degree of crystallization is increased. At 500 ◦ C, the XRD peaks are broad and the background in the spectra is still relatively high. In contrast, at 900 ◦ C the XRD peaks are narrow and the background in the spectra is negligible. From the X-ray broadening, the average crystallite sizes in calcined samples are calculated as 5.4 nm, 11.1 nm and 45.0 nm [plane SDC 1 1 1] and 15.5 nm, 44.6 nm and 49.1 nm [plane NiO 1 1 1] for samples calcined at 500 ◦ C, 600 ◦ C and 900 ◦ C, respectively. It is obvious that with
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increasing temperature the NiO phase grows faster than SDC phase. Nevertheless, taking into consideration that the process starts in a solution and rather similar values for crystallite sizes at 900 ◦ C suggests that both phases NiO and SDC should be uniformly distributed one into another. XRD analysis results are consistent with the results of thermal decomposition, because the final products (NiO and SDC) are already presented after calcination at minimum temperature (500 ◦ C). The BET specific surface measurements of NiO–SDC mixtures calcined at 500 ◦ C, 600 ◦ C and 900 ◦ C are 72.5 m2 g−1 , 20.0 m2 g−1 and 6.2 m2 g−1 , respectively. Since BET analysis enables the calculation of an average particle with a specific geometry, the diameters of spherical particles calculated from BET measurements are 11.0 nm, 43.3 nm and 139.0 nm for powders calcined at 500 ◦ C, 600 ◦ C and 900 ◦ C, respectively. By comparing these calculated values with crystallite sizes determined with the Scherrer calculation, we assume that the degree of agglomeration of individual particles is rather low at 500 ◦ C. With increasing calcination temperature, nano-particles tend to aggregate. 3.2. Sintering of anode material An impact of the calcination temperature and consequently the degree of particles’ aggregation is also reflected on the sintering behaviour of the powders. Powders calcined at 900 ◦ C milled in an attritor mill and pressed into tablets show relatively poor sinterability. The overall shrinkage during dynamic sintering up to 1500 ◦ C is only 14%. It is caused by specific surface area loss during the calcination process. In contrast, material calcined at 500 ◦ C exhibits much higher shrinkage (around 25% during dynamic sintering up to 1500 ◦ C). Due to low sinterability, samples calcined at 900 ◦ C were omitted in the further investigation. Similar findings regarding NiO–SDC sinterability were reported by Fang et al.15 Detailed sintering behaviour of NiO–SDC pellets prepared from powders calcined at 500 ◦ C were determined via an indepth microstructure analysis on sintered and reduced samples (Fig. 3) where A1, B1, C1, D1 and E1 denote samples sintered at 1200 ◦ C, 1250 ◦ C, 1300 ◦ C, 1350 ◦ C and 1400 ◦ C, respectively and notations from A2 to E2 of the same samples reduced at 900 ◦ C. The dark grey particles in Fig. 3 are NiO or Ni, the light grey phase is SDC, and the pores are black. A very important characteristic of the prepared composites, which originates from the preparation procedure, is the homogeneous distribution of
all phases. It can be seen that NiO or Ni grains are homogeneously distributed in the SDC phase. The distribution of pores in reduced samples is also uniform. Since all parameters important for exact cermet analysis are sometimes difficult to obtain simply from SEM micrographs, a detailed quantitative microstructure analysis on sintered and reduced samples was performed. For statistically reliable data at each chosen temperature, five to ten different regions on the surface of the tablets were analyzed. The results of the quantitative microstructure analysis are summarized in Tables 1 and 2. ¯ dx and dy are represented as the diameter of the Parameters d, area-analogue circle – DCIRCLE and intercept lengths x and y direction – FERETX, Y, respectively. Porosity ε was determined as geometrical porosity (from tablet dimensions). According to the results of quantitative microstructural analysis, the porosity of the sintered samples decreases with the increasing sintering temperature (Table 1). The higher sintering temperature also results in pronounced grain growth of both the NiO and SDC phases. The NiO phase grows somewhat faster and also reaches larger final dimensions.27 However, both phases remain well within submicrometer range after sintering up to 1400 ◦ C not exceeding ∼0.44 m and ∼0.38 m for NiO and SDC, respectively. Relatively small NiO and SDC grains after sintering as well as the minimum one phase dominance in sintered bodies is a consequence of the powder treatment after the synthesis which modified its morphology. The ball-milling process reduced the size of the particles and weakened the aggregation of the nanocrystallites.28 Another important parameter to follow the sintering progress is sintered density. From relative sintered density, the geometrical porosity of a sintered body can be easily calculated (Tables 1 and 2). The target value for the porosity of a sintered and reduced element is approximately 30% to ensure that pores remain a continuous phase throughout the final composite but, at the same time, that solid phases are also in good contact.29–32 It is evident from Tables 1 and 2 that sintered density increases with temperature while porosity decreases with sintering temperature. A relative sintered density 71.8% reached at 1200 ◦ C reveals that material sinters relatively poorly at that temperature. Consequently, NiO and SDC phases are not in good contact, making such material inappropriate for SOFC technology (Fig. 3). Enhanced grain growth and a relatively quick sintered density increase are observed at 1250 ◦ C. At the same time, at this temperature, the target value for porosity in reduced samples is also reached (31.2%). The shape factor does not change significantly with increasing the sintering temperatures. The most apparent differences
Table 1 Quantitative microstructure analysis of the sintered samples. Sample
A1 B1 C1 D1 E1
Geometrical porosity, ε (%)
28.2 15.2 9.0 6.1 3.7
d¯ (m)
Shape factor, ψ
dx (m)
BET surface area (m2 g−1 )
dy (m)
NiO
SDC
NiO
SDC
NiO
SDC
NiO
SDC
0.18 0.24 0.27 0.37 0.44
0.14 0.16 0.24 0.30 0.38
0.80 0.79 0.77 0.78 0.79
0.76 0.76 0.76 0.77 0.75
0.20 0.27 0.27 0.42 0.50
0.13 0.18 0.28 0.34 0.42
0.20 0.28 0.27 0.41 0.49
0.13 0.18 0.29 0.33 0.43
1.25 1.10 0.99 0.80 0.52
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Fig. 3. SEM micrographs of samples after sintering (left column) as NiO–SDC, and after reduction (right column) as Ni–SDC.
between the sintered and the reduced samples are the increased porosity of the samples and increased Ni particles relative to NiO particle size. It is a consequence of mass change when NiO is transformed to Ni while the overall volume of the composite
does not change due to the rigid framework of SDC. After the reduction, the formed Ni tends to sinter, which is the main cause that Ni grains grow relative to NiO. Ni sintering is a consequence of different surface energies of both phases, which causes the
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Table 2 Quantitative microstructure analysis of the reduced samples. Sample
A2 B2 C2 D2 E2
Geometrical porosity, ε (%)
44.7 31.2 27.4 24.6 21.8
d¯ (m)
Shape factor, ψ
dx (m)
BET surface area (m2 g−1 )
dy (m)
Ni
SDC
Ni
SDC
Ni
SDC
Ni
SDC
0.38 0.33 0.40 0.49 0.52
0.15 0.17 0.23 0.28 0.38
0.90 0.90 0.89 0.93 0.90
0.84 0.82 0.81 0.79 0.78
0.40 0.35 0.42 0.51 0.55
0.16 0.19 0.25 0.32 0.42
0.38 0.34 0.41 0.49 0.54
0.15 0.18 0.24 0.31 0.43
hetero-grains at the interface to loose adhesion of Ni to the SDC at elevated temperatures. However, the rigid SDC framework apparently suppresses excessive Ni sintering, keeping Ni grains within the submicrometer range. A more in-depth image analysis is performed using the line intercept method. Using this method, one can obtain information about the number N of contact points between grains of various phases in a unit length and calculate the corresponding contiguity values C (Fig. 4). As expected from the microstructure analysis, absolute values of NNi–SDC , NSDC–SDC and NNi–Ni reach their maximum at 1250 ◦ C. At this temperature, both phases already sinter; however, no extensive grain growth or severe phase separation is noticed. Below 1250 ◦ C, Ni and SDC phases sinter poorly, meaning that there are fewer contacts between phases. In contrast, with increasing sintering temperature above 1250 ◦ C, the number of contacts in a unit length decreases since grains of both phases grow. Relatively small one-phase dominance in the final Ni–SDC cermet at 1250 ◦ C results in the highest contiguity value CNi–SDC . However, at this temperature the contiguities CSDC–SDC , CNi–Ni are relatively low. Lower CSDC–SDC and CNi–Ni values at 1250 ◦ C are mostly due to significant increases of NNi–SDC . With higher NNi–SDC numbers, the denominator of the fraction to calculate the degree of contact of adequate phases (contiguity) is increased. Consequently, if CNi–SDC increases CSDC–SDC and CNi–Ni must decrease.
Fig. 4. Contiguity and absolute values of contacts between Ni–Ni, Ni–SDC and SDC–SDC grains in cermets submitted to various sintering temeperatures.
0.91 1.25 1.05 0.85 0.58
3.3. Construction of bi-layer system In an SOFC unit, anode material is sintered to an electrolyte where both layers have to exhibit fairly different microstructural properties. One of the main important factors in such a bilayer system is good adhesion at the layer boundary. Since 1250 ◦ C is the most appropriate sintering temperature for the prepared anode material with respect to its final microstructure, we attempted to use the same sintering temperature also for bi-layer system co-sintering. However, pure SDC and NiO–SDC sinter at somewhat different temperatures and also with different shrinkages (Fig. 5). Such dissimilar densification processes where NiO–SDC already shrinks ∼3% before SDC starts to sinter (∼900 ◦ C) also means that within the temperature interval where both materials sinter SDC undergoes ∼21% of dimensional change while NiO–SDC shrinks only an additional ∼13%. A ∼8% shrinkage mismatch from 900 ◦ C to 1250 ◦ C often led to cracks formation at the layer boundary. To bring the probability of crack formation at the layer boundary to a minimum SDC was pre-sintered. The pre-sintering temperature for SDC is chosen at 1000 ◦ C. In this manner, shrinkages of both materials are adjusted and at the same time the SDC substrate is strengthened for subsequent screen-printing of the NiO–SDC. After co-sintering at 1250 ◦ C for 1 h, the bi-layer system was reduced in Ar–H2 (5 vol%) at 900 ◦ C for 3 h. The final thickness of anode film after reduction is ∼20 m. The carefully adjusted shrinkages of both layers result in good adhesion, as confirmed
Fig. 5. Sintering curves of SDC and NiO–SDC materials calcined at 900 ◦ C.
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Fig. 6. SEM micrograph of anode–electrolyte contact in the bi-layer system.
by SEM micrograph (Fig. 6). The electrolyte at 1250 ◦ C reaches the final sintering stage (ρrelat. = 94%) and forms gas-tight electrolyte layer. It is also evident that the SDC phase in anode and electrolyte layers are continuously connected across the layer boundary, expanding the reaction zone for fuel oxidation into the anode layer. A precondition for the applicability of the anode material is successful co-sintering with the electrolyte material. Fig. 5 presents the sintering curve for SDC electrolyte material. The shrinkage of NiO–SDC after sintering at 1250 ◦ C was calculated from the relative green and sintered densities. SDC electrolyte was pre-sintered at 1000 ◦ C where 5% of shrinkage was reached; later, it was co-sintered with NiO–SDC phase at 1250 ◦ C where both materials shrank by a further 16%. The bi-layer system was later reduced in Ar–H2 (5 vol%) at 900 ◦ C for 3 h. The final thickness of the anode film after reduction was 20 m. Good contact between both components of the bi-layer was confirmed by SEM micrograph (Fig. 6). From Fig. 6, it is also evident that SDC in electrode and electrolyte layers are continuously connected, which is precondition for good ionic conductivity of anode material. According to the shrinkage curve, at 1250 ◦ C the electrolyte reaches the final sintering stage and forms a gas-tight electrolyte layer. 4. Conclusions Nano-sized NiO–SDC composites are prepared with an innovative, simple synthesis, easy to manipulate, via thermal decomposition of intermediate anode precursor. The EGA analysis indicates that the intermediate decomposes in several consecutive steps, during which water, acetone, acetic acid, carbon dioxide, carbon monoxide and ethylene glycol are detected as volatile products. After the thermal treatment of the metal acetate precursor at 500 ◦ C fine mixture of NiO–SDC with specific surface area 72.5 m2 g−1 is formed. The crystallite sizes are 5.4 nm and 15.5 nm for SDC and NiO, respectively. During sintering, the average grain sizes of both phases SDC and NiO are enlarged with the increasing sintering temperature, however not exceeding 500 nm when sintered up to 1400 ◦ C. NiO grains
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grow faster than SDC with the increasing sintering temperature. The dimensions of the final Ni grains formed during the subsequent reduction are influenced by the microstructure stability of the SDC framework established during the sintering. The smallest Ni grains, well in sub-micrometre range (d¯ = 0.33 m), as well as the highest contiguity value between Ni and SDC phases are obtained in the sample sintered at 1250 ◦ C. Smaller CNi–SDC values determined for samples sintered at temperatures lower than 1250 ◦ C or higher than 1250 ◦ C are ascribed to relatively larger Ni grains due to the non-rigid SDC phase or the increased one-phase dominance region of both phases, respectively. During the preparation of the bi-layer system SDC/NiO–SDC, the SDC substrate is strengthened before the screen-printing by pre-sintering at 1000 ◦ C. At the same time, the shrinkages of both components SDC/NiO–SDC are adjusted for the final sintering at 1250 ◦ C. After the reduction, good contact on the anode–electrolyte phase boundary with no cracks is prepared. The bi-layer system consists of a highly dense electrolyte (relative density 93%) and an anode layer with a porosity of 31.2%. It can be concluded that the present preparation procedure is one of more promising approaches for achieving homogeneous distribution, integrity and connectivity of all present phases. Acknowledgment This work was supported by the Ministry of Higher Education, Science and Technology of the Republic of Slovenia through grant P1-0175-103 and Centre of Excellence LowCarbon Technologies (CO NOT). References 1. Haile SM. Fuel cell materials and components. Acta Mater 2003;51(19):5981–6000. 2. Tietz F, Arul Raj I, Jungen W, Stöver D. High-temperature superconductor materials for contact layers in solid oxide fuel cells: I. Sintering behaviour and physical properties at operating temperatures. Acta Mater 2001;49:803–10. 3. Suzuki T, Hasan Z, Funahashi Y, Yamaguchi T, Fujishiro Y, Awano M. Impact of anode microstructure on solid oxide fuel cells. Sci Mag 2009;325:852–5. 4. Yin Y, Li S, Xia C, Meng G. Electrochemical performance of gel-cast NiO–SDC composite anodes in low-temperature SOFCs. Electrochim Acta 2006;51:2594–8. 5. Li J, Ikegami T, Mori T. Low temperature processing of dense samariumdoped CeO2 ceramics: sintering and grain growth behaviours. Acta Mater 2004;52:2221–8. 6. Misono T, Murata K, Fukui T, Chaichanawong J, Sato Ab H, Naito M. Ni–SDC cermet anode fabricated from NiO–SDC composite powder for intermediate temperature SOFC. J Power Sources 2006;157:754–7. 7. Hui S, Roller J, Yick S, Zhang X, Deces-Petit C, Xie Y, et al. A brief of the ionic conductivity enhancement for selected oxide electrolytes. J Power Sources 2007;172:493–502. 8. Wright J, Virkar A. Conductivity of porous Sm2 O3 -doped CeO2 as a function of temperature and oxygen partial pressure. J Power Sources 2011;196:6118–24. 9. Cheng J, Deng L, Zhang B, Shi P, Meng G. Properties and microstructure of NiO/SDC materials for SOFC anode applications. Rare Metals 2007;26:110–7.
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