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Highly durable nano-oxide dispersed ferritic stainless steel interconnects for intermediate temperature solid oxide fuel cells
Journal of Power Sources 439 (2019) 227109
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Highly durable nano-oxide dispersed ferritic stainless steel interconnects for intermediate temperature solid oxide fuel cells Muhammad Taqi Mehran a, b, Tae-Hun Kim a, c, Muhammad Zubair Khan a, Seung-Bok Lee a, Tak-Hyoung Lim a, Rak-Hyun Song a, c, * a b c
Fuel Cell Research Center, Korea Institute of Energy Research (KIER), 152 Gajeong-ro, Yuseong-gu, Daejeon, 34129, South Korea School of Chemical and Materials Engineering (SCME), National University of Sciences and Technology (NUST), Main Campus, H-12, Islamabad, 44000, Pakistan Korea University of Science and Technology, 217 Gajeong-ro, Yuseong-gu, Daejeon, 34113, South Korea
H I G H L I G H T S
G R A P H I C A L A B S T R A C T
� Nano-CeO2 dispersed SUS430 shows better oxidation properties for long-term testing. � Nano oxide dispersion in steel improves oxide-scale adhesion and thickness. � 3 wt% CeO2 dispersed steel alloy results 15 mΩcm2 ASR increase in 1000 h at 800 � C. � Presence of dispersed oxides at grain boundary influences the oxidation kinetics.
A R T I C L E I N F O
A B S T R A C T
Keywords: Metallic interconnects Reactive element Solid oxide fuel cell Area-specific resistance
Herein, ferritic stainless steel alloys are developed with the dispersion of nano-oxides of reactive elements to improve their oxidation characteristics for solid oxide fuel cell (SOFC) interconnects. Nano-oxides of yttrium, cerium, and lanthanum chromite are homogeneously dispersed in a ferritic stainless steel base powder and specimen pellets are prepared by pressing the powder and sintering it in hydrogen at 1400 � C. The nano-oxide dispersed alloys are tested for oxidation characteristics and electrical properties at 800 � C for 1000 h in air. The results show that the ferritic stainless steel with 3 wt% nano-ceria exhibits significantly improved areaspecific resistance (ASR) characteristics. The increase in the ASR for this alloy is found to be only 15 mΩ cm2/kh and is ascribed to the thin dense oxide layer (~0.4 μm), the presence of spinel (MnCrO4) at the outer side of the oxide layer, and the improved oxide layer adherence to the metal substrate. The addition of reactive element oxides altered the oxidation kinetics of the Fe-Cr alloys, thus making nano-oxide dispersed ferritic stainless steel highly durable and most suitable for SOFC interconnects. The oxidation kinetics of the nano-oxide dispersed steel was also discussed as related to cation diffusion.
* Corresponding author. Fuel Cell Research Center, Korea Institute of Energy Research (KIER), 152 Gajeong-ro, Yuseong-gu, Daejeon, 34129, South Korea. E-mail address: [email protected] (R.-H. Song). https://doi.org/10.1016/j.jpowsour.2019.227109 Received 11 July 2019; Received in revised form 26 August 2019; Accepted 3 September 2019 Available online 5 September 2019 0378-7753/© 2019 Elsevier B.V. All rights reserved.
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~10 mΩcm2. To control the scale composition and conductivity, various commercial alloys named as Crofer 22 APU, ZMG232, etc. have been structured by modifying the Fe-Cr alloy [24,34]. During recent years, the development of the alloy with the oxide dispersion has made considerable progress [36,37]. The better adher ence of oxide layer to base metal and deceleration of cations to the interface were caused by these oxide particles as a result of improved oxidation resistance of the base alloys [11,38]. When RE’s are added, in one of the following two forms of metallic or oxide dispersion, proved to be quite beneficial in enhancing the scale adhesion to the substrate [39–42]. Y, La, Ce, and Zr are mostly used as REs. Quadakkers et al. and Whittle et al. mentioned that, besides developing the adhesion proper ties, the REs also play a vital role in giving a considerable decrement in the scale growth of oxides and promoting the selective oxidization of Cr in low-Cr alloys [6,43]. The scale growth can be suppressed by the introduction of REs which reduces outward Cr transport. The oxidation resistance of steel alloys along with the better mechanical properties can be achieved by adding RE-oxides in dispersed phase [44,45]. Linderoth et al. reported that the oxide dispersed chromium alloys is having a very lower rate of oxidation than that of chromium alloys with availability of no reactive element oxides [46]. Quadakker et al. presented his work on improving the oxidation resistance of alloy by adding rare-earth ele ments/reactive elements (REs) that form RE oxides at the boundaries of grain [47]. The coatings of reactive element exceptionally improved oxidation properties of Iron–Chromium steels for long run applications of SOFC [15–17]. Fontana et al. [17] investigated the effect of coatings (thickness: 102– 2 � 102nm) of La2O3, Nd2O3, and Y2O3 on the proper ties of oxidation and conductivity for three different ferritic stainless steels, Crofer 22 APU, AL 453, and Haynes 230. The coated samples showed lower growth rates for oxide and improved scale adhesion to the substrate. Furthermore, the best result for conductivity was showed by La2O3 coated Crofer 22 APU. Fontana et al. in another study [48] investigated long-term exposure for more than two years at 800 � C in air of La2O3 and Y2O3 coated Crofer 22 APU. They reported enhanced oxidation resistance, slower growth rates of oxides, and reduced elec trical resistance within the coated samples compared to the uncoated material [49]. Commercial SUS430 has proved to be a cost-effective and cheap interconnect material in comparison to Crofer 22 APU, and when nanooxides of REs are alloyed in the SUS430 base material, gave the results for possible improvement in oxidation protection, electronic conduc tion, and chromium retention for SOFC systems [50,51]. In our previous study [52], we reported that the addition of nano-oxides significantly improves the properties of SUS430 as an interconnect. However, the oxide layer development and the influence of the addition of nano-oxides on the oxidation properties during the long-term test should be further investigated. The main focus of this work was to configure different ferritic steel-based nano-oxide dispersed alloys (nano-CeO2, -Y2O3, and -LaCrO3 dispersed SUS430 steel) by using an improved powder mixing method. The dispersion of nano-oxides of the reactive elements influenced the oxide scale thickness, adhesion, and chemistry, and we have therefore conducted a detailed study of the oxide scale and the electrical properties of RE-oxide dispersed SUS430 alloys during a 1000 h ASR test at 800 � C in air. The thermally grown oxide scale was accordingly characterized based on the composition, microstructure, and electrical properties. The evaluation of grown oxide was done to determine the suitability of the oxide dispersed alloys as interconnect materials for IT-SOFCs.
1. Introduction The solid oxide fuel cell (SOFC) stack is formed as a lamellar struc ture of planar cells in order to reduce the internal resistance and to in crease the effective electrode area per unit volume. To separate two adjacent cells in a stack, a separator or interconnector is employed [1,2]. To apply such a device as a SOFC interconnect, the material of the separator should have a thermal expansion coefficient (TEC) closer to that of the other cell components, as well as high conductivity. Because the SOFC’s operating temperature is high, the interconnect material must possess good corrosion resistance in fuel and air electrode envi ronments [3–5]. Typical interconnects are formed of electrically conductive ceramic materials such as LaCrO3. However, these ceramic materials are relatively expensive and difficult to fabricate. Also, due to recent improvements that have reduced the operating temperature of the SOFCs below 800 � C, many research groups are extensively search ing for metallic interconnects materials. Metallic materials exhibit higher thermal and electrical conductivities as compared to ceramic interconnects and are less costly and can be fabricated easily [1,6–8]. Ferritic stainless steel (Fe-Cr) based alloys are considered to be a suitable metallic interconnect material where preferential oxidation of Chromium to Chromium oxide would configure a continuous protective oxide layer. However, during long-term exposure to the high tempera ture oxidizing and reducing atmosphere in the SOFC, the Cr-containing steel experiences continuous growth of an oxide layer on the steel sub strate [9]. This oxide layer has lesser conductivity and a different TEC than those of the base material, thus causing a substantial increase in the area-specific resistance (ASR) and spallation of the oxide layer. The ASR values of the currently developed ferritic stainless steel alloys are pro jected to be 150 mΩ-cm2 after 40,000 h operation at 850 � C; the oxide thickness after long-term exposure is expected to be more than 30 μm, which can cause delamination and cracking during the thermal cycling operation of the SOFC [10,11]. In addition to the oxide scale growth, Cr poisoning is also a significant issue related to the use of Cr2O3-forming alloys. Cr-containing gaseous species are formed from the oxide scales, which results in poisoning of the electrodes and performance degrada tion of the SOFCs [8,12–14]. In order to improve the properties of Cr2O3-forming alloys such that they will be suitable for application in high-temperature SOFC service conditions, surface modification and coatings are applied. Reactive element (RE) [15–17], spinel [18,19], and perovskite [20–23] oxide coatings significantly mitigate the Cr-poisoning and the conductivity of the oxide scale is also reported to be improved. In particular, Cr-free spinel coatings such as (Mn, Co)3O4 [24,25], (Cu,Mn)3O4 [19], Cu-Fe spinel [26], and NiCo2O3 [27] have shown promising results for SOFC interconnects. Similarly, many reports suggest that the application of a protective coating of reactive element oxides (such as La [28], Ce [29], and/or Y [30]) on the metallic interconnects enhances the electrical properties [31]. However, while applying surface coatings can effec tively inhibit Cr poisoning, practical improvement of the electrical conductivity and oxidation resistance of the metallic interconnect cannot be achieved for long-term operation. Moreover, the coating process increases the manufacturing cost of the metallic interconnects. In order to meet the requirements of an appropriate TEC and suitable oxide scale properties for the metallic interconnects, modifi cation of Fe-Cr alloys is considered to be a viable approach [9,24,32,33]. Horita et al. [34] added Nb (0.35 wt%) and Mo (0.5 wt%) to Fe-20Cr alloy and investigated the oxidation properties in an H2/H2O atmo sphere at 800 � C for more than 1000 h. They reported that, in addition to frequently observed Cr2O3 and Cr-Mn spinel in the scale and at the grain boundary, an Nb lava phase is also formed. During a long-term test, the grain boundary lava-phase was found which resulted in an increment in the oxidation resistance and a decrease in the ASR. Similarly, Hua et al. [35] developed a promising ferritic steel alloy Fe-17Cr-1Mn-0.5Ti with the addition of La, Y, and Zr; After running 1000 h isothermal oxidation test at 750 � C in air, this alloy showed an ASR value as low as
2. Experimental 2.1. Preparation of powders and pellets Commercial nano-CeO2 and -Y2O3 powders (Sigma-Aldrich, USA, particle size >25 nm) were used for blending with the steel powder for commercial use such as SUS430 (Metal Player, Korea) to develop the 2
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steel with dispersed oxide; nano-LaCrO3 (Kanto Chemical Co., Japan) had an average particle size of < 50 nm. Crofer 22APU Fe-Cr alloy (ThyssenKrupp VDM, Germany) was used as a reference material. The chemical compositions of the starting alloys (SUS430 and Crofer 22 APU) are presented in Table 1. The composition of SUS430 was deter mined by using an Inductively Coupled Plasma Mass Spectrometer (ICPMS, model: NexION 300D, PerkinElmer). SUS430, SUS430þ3 wt% CeO2, SUS430þ3 wt% Y2O3, and SUS430þ3 wt% LaCrO3 were homogenized by mixing in a high energy planetary ball mill for several hours (12–48 h) using polyvinyl butyral (PVB) which binds a mixture of isopropyl alcohol and toluene in a stainless steel jar using stainless steel balls. The high-energy ball mill ensured homogenous mixing of the SUS430 and nano-oxide powders. The powder mixtures were then compacted at 600 MPa and 200 MPa via uniaxial and cold isostatic pressing, respectively, to yield specimen pellets of 5 mm � 8 mm � 25 mm dimensions. The pellets were sintered in hydrogen for 10 h at 1400 � C to obtain dense pellet samples. The details of sample processing and sintering are given in Fig. S1 in the supplementary material. 2.2. Material characterization The sintered pellets were first ground using 1000 grit SiC paper and dried at 100 � C for 2 h before density measurement by the Archimedes method. For each measurement, at least three pellets were used and the average relative density was determined. The sintered pellets were characterized by SEM and XRD analyses. The surface and cross-section of the selected specimens were analyzed by a scanning electron micro scope (SEM, Hitachi S4800) equipped with an energy dispersive spec trometer (EDS), which was used to assess the thickness, structural morphology, and element distribution of the oxide scales. For observa tion of the grain structure, the polished specimens were etched with aqua regia þ glycerol for 30 s before the SEM analysis. The crystalline phases of the oxidation products on the surfaces of these specimens were identified by X-Ray diffraction (XRD). Additionally, cross-sections of the specimens were also examined using SEM/EDS. The thermal expansion coefficient (TEC) of the alloys was evaluated from room temperature to 800 � C at a rate of 5 � C per minute. Air flowed into the chamber at 100 cc.min 1. To measure the TEC, the device was first calibrated with the standard Al2O3 specimen and the calibration was used to determine the thermal expansion behavior of the SUS430 alloys up to 800 � C.
Fig. 1. (a) Detailed schematics showing the positions of Pt wires and mesh attached to the pellet for ASR measurement test; (b–d) Jig and application of load on the sample using steel slabs to ensure appropriate contact. (e,f) Actual image of final ASR test samples.
% CeO2, and SUS430þ3 wt% Y2O3 to the SOFC cathode environment (800 � C, air) during 1000 h, and the ASR was measured every 20 h by the four-probe method [53]. To measure the ASR of the samples, 50–300 mA current was applied by a current source (Toyotech, DP 30-03 TP) at an interval of 50 mA and the respective values of voltage were measured using an Agilent Mul timeter (model: 34450A). The ASR was calculated for each sample using Equation (1) for each current applied and the average ASR value was determined. 2.4. Cr poisoning experiment
2.3. Long-term ASR testing set-up
To study relative Cr poisoning of a cathode in the developed alloy based interconnects, a half-cell was fabricated as reported elsewhere [19]. Gd2O3-doped CeO2 (GDC) pellet (thickness: 1 mm) was used as the electrolyte and a slurry of La0.6Sr0.4Co0.2Fe0.8O3 (LSCF) cathode mate rial was coated on one side of the sintered GDC coin cell using a screen-printing and sintered in air for 3 h at 1150 � C. The area of the coated LSCF working electrode was 1.0 cm2. The reference and counter electrodes were prepared by applying a conductive Pt paste on the other side and at the edge of GDC coin cell, respectively, and fired in air at for 2 h 1000 � C. Pt meshes and wires were used as current collectors and voltage probes. Three metallic interconnect specimen SUS430, SUS430þ3 wt%CeO2 and Crofer 22 APU were prepared with area of 1.5 cm2, thus entirely covering the cathode active area. The alloy specimens were placed onto the Pt mesh current collector for the working electrode. Humidified air (3% H2O) was supplied to the cell at 800 � C. Half-cell voltage was measured at a constant current density of 50 mA cm 2 and AC-impedance spectra were periodically measured at an OCV. Electrochemical experiments were carried out using Biologic Potentiostat (SP-240).
To prepare the samples for the long-term ASR test, the surfaces of the specimens were polished mechanically with SiC paper up to 2000 grade. The samples were then cleaned ultrasonically in acetone for 30 min. Preoxidation was conducted for 100 h at 800 � C to thermally grow the layer of oxide on the pellet’s surface. After the pre-oxidation, platinum (Pt) wire was connected to the ASR samples for current collection, as shown in Fig. 1 (a). Pt paste was used to minimize the interfacial resistance between the current collectors and the samples [53]. The ASR samples were placed in a specially designed jig, which provided a compressive pressure of 50 kPa on the sample by the load (Fig. 1(b–f)) [54]. The prepared alloys were tested to determine their long-term ASR behavior by exposing the alloys SUS430, SUS430þ3 wt% LaCrO3, SUS430 þ3 wt Table 1 Chemical composition of the ferritic stainless steel alloys (SUS430 and Crofer 22APU). Steel alloy SUS430 Crofer 22APU a
a
Element Composition (wt. %) Fe Cr S P Balance 18.08 0.028 0.035 Balance 20–24 0.02 0.05
C 0.079 0.03
Si 1.02 0.5
Mn 0.26 0.3–0.8
ThyssenKrupp Crofer® 22 APU Datasheet No. 4046. 3
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3. Results and discussion
43 μm for SUS430, SUS430 þ 3 wt % LaCrO3, SUS430 þ 3 wt % CeO2, and SUS430þ 3 wt % Y2O3, respectively.
3.1. Microstructure analysis of the sintered nano-oxide dispersed steel alloys
3.2. Oxidation behaviors of the nano-oxide dispersed steel alloys during long-term ASR tests
Fig. 2 shows SEM images of the nano-oxide added SUS430 powder during the processing. The supplied SUS430 powder have a particle size between 20 and 60 μm, as shown in Fig. 2(a) and 2(b) shows the morphology of the particles after 12 h ball milling in the high energy ball mill, yielded a clear decrease in the particle size. Particle size analysis of the powder after high energy ball milling process is shown in Fig. S2. Fig. 2(c and d) present the surface morphology of the developed ODS alloy with a SUS430-3 wt % CeO2 composition and it can be noticed that the added nano-oxides are agglomerated at the grain boundaries with an average particle size of 1–5 μm. The quantitative results of the EDS are given in Fig. 2(e). To find the normalized mass percentage of the ele ments present in the alloy, point EDS was used and presence of CeO2 at the grain boundaries is confirmed via EDS analysis. The morphology of the sintered alloy shown in Fig. 2(c) indicates that the specimens were fully sintered and furthermore no pores were observed; the relative density of all the developed alloys was found to be more than 95%. For the sintered alloys, the average grain sizes were ~160, 30, 34, and
The morphology of the oxide scale grown on the substrate during 1000 h isothermal oxidation at 800 � C is studied by SEM and EDS ana lyses. Fig. 3 provides a comparison of the surfaces of the oxide scales for SUS430, and SUS430þnano-oxide alloys after 1000 h. It can be noted that the morphologies of SUS430þ3 wt% CeO2 and SUS430þ3 wt% LaCrO3 resemble that of Crofer 22 APU reported by Qu et al. [30]. However, the particle shapes of the oxide scales with CeO2 and LaCrO3 nano oxides vary from tetrahedral to diamond-shapes and had crystal shape scales with distinctive facets. The oxide scales formed on the surface of SUS430 and SUS430þ3 wt% Y2O3 appear to have round-faced particles. In the literatures, it is reported that prism and diamond type particle shapes belong to (Mn, Cr)3O4 spinel and that Cr2O3 oxide scale forms corundum shaped structures [35,55]. Although it is evident from the SEM micrographs given in Fig. 3 that the respective additions of nano-CeO2 and -LaCrO3 have a clear influ ence on the morphology of the oxide scale, we performed a detailed EDS
Fig. 2. (a) As-received SUS430 alloy powder from Metal Player®, Korea (b) after mixing of nano-oxides and SUS430 powder by using high energy ball-milling for 12 h, (c, d) Backscattered SEM analysis of sintered pellet of the SUS430þ3 wt% CeO2 alloy showing presence of the agglomerated nanoparticles of CeO2 at the grain boundaries (e) EDS analysis to confirm the presence of agglomerated nano-particles at the grain boundaries of SUS430þ3 wt % CeO2 alloy. 4
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Fig. 3. Surface morphology of the oxide scale after 1000 h isothermal oxidation of SUS430, and SUS430 þ 3 wt% nano-oxides: (a) SUS430, (b) SUS430þ3 wt% LaCrO3, (c) SUS430þ3 wt%CeO2, (d) SUS430þ3 wt%Y2O3. The arrows with numbers show the points of quantitative EDS analyses.
analysis of the different particles present on the oxide scale (marked in Fig. 3) to quantify the nature of these particles. The results are presented in Table 2. From the EDS results, we can deduce that the compositions of the two types of particles are almost identical in the samples with nanoCeO2 and -LaCrO3, although they appear to have different sizes and shapes in the SEM micrograph. From Fig. 3 (a), we can see that, at point 4, the particle appears to be overgrown and large, while the area near point 5 has many small particles. From the EDS results at points 3 and 5, both types of particles composed mainly of CrO2 and Mn appear to be absent. However, at point 4, Mn is present, but in a very small amount (~0.66%). This means that the surface of the oxide scale of SUS430 after 1000 h isothermal oxidation is predominantly CrO2, along with small traces of Mn-Cr. That is, the oxidation process of SUS430 still might be controlled by the outward migration of Cr ions from the Fr-Cr matrix to the CrO2-Air interface. The large-sized grains of corundum shape may provide evidence of continued growth and Cr2þ diffusion dominated oxidation kinetics in SUS430 [35]. In the SUS430þ3 wt% CeO2 oxide scale, three points were selected for analysis (referred to as points 8, 9, and 10 in Fig. 3 (c)). Point 8 marked particle is overgrown relative to others and exhibits a distinct morphology of Mn-Cr spinel oxides from its composition analysis.
Although the particle 9 and 10 shows Cr-rich phase, most of the observed particles in the scale have a Mn-Cr spinel morphology. In the case of nano-LaCrO3 addition, Mn in the oxide scale was present in a lesser amount in comparison with the SUS430þ3 wt% CeO2 sample. In contrast, it can be seen from the EDS data that with the addition of 3 wt % Y2O3, the morphology and the composition of the oxide scale were not altered to a great extent. The composition of the oxide scale is similar to that of SUS430 without nano-oxide additives. From the EDS data and the morphology of the oxide scale after 1000 h, it is considered that the addition of nano-CeO2 and nano-LaCrO3 might have affected the mechanism of oxidation and the formation of the oxide scale. The higher conductive spinel on the top surface of the oxide scale helps to reduce inward migration of oxide ion and outward diffusion of Cr element, thus modifying the oxidation rates and ASR characteristics. The oxide scale structure, thickness, and adherence to the substrate were also studied using SEM and EDS analyses. The conductivity of the oxide scale is controlled by the scale thickness and, with an increase in the oxidation time, the thickness of the chromia scale increases [4]. Fig. 4(a–h) shows SEM/EDS images of the cross-section of the oxide scale after the 1000 h ASR test. The oxide scale thickness was deter mined by measuring the thickness at 4–5 different locations, after which the average value was calculated. The average oxide scale thickness values for SUS430, SUS430 þ 3 wt% LaCrO2, SUS430 þ3 wt% CeO2, and SUS430 þ3 wt% Y2O3 were 1.823, 0.592, 0.386, and 0.939 μm, respectively. Another aspect of the oxide scales that can be observed from the SEM analysis of the cross-section is that there are no areas of delamination or spallation except SUS 430. As shown in Fig. 4 (a), the oxide scale grown over the SUS430 substrate was not perfectly attached to the substrate. Some gaps can be observed in the SEM image. This may indicate that scale adherence to the substrate might be poor and this might contribute to loss of conductivity of the sample. However, in the SUS430 þ 3 wt% CeO2, a uniform and dense oxide scale is attached tightly with the substrate, as shown in Fig. 4 (e). The oxide scale thickness in the nano-CeO2 added alloy was very small (~0.386 μm) during the 1000 h oxidation test. Similarly, for the case of SUS430 þ 3 wt% LaCrO3, the oxide scale was also similar to that of the ceria added alloy, with a slightly greater thickness (~0.592 μm). For the case of the SUS430þ3 wt% Y2O3 sample, the oxide scale thickness was greater
Table 2 EDS point analysis of the surface of SUS430 and nano-oxide dispersed alloys after 1000 h oxidation in air at 800 � C to determine the composition of the various microstructures of the oxide scale. Sample
Point
SUS 430
3 4 5 6 7 8 9 10 11 12
SUS430 þ 3 wt% LaCrO3 SUS430 þ 3 wt% CeO2 SUS430 þ 3 wt% Y2O3
Elemental Composition (Atomic %) C
O
Cr
Mn
Fe
11.02 6.32 5.28 3.89 4.78 10.34 11.13 12.85 4.79 6.37
66.28 65.55 63.48 64.10 64.75 60.45 60.8 61.65 61.90 49.55
22.5 27.05 30.63 29.04 22.58 16.27 18.12 22.18 32.86 43.04
– 0.66 – 1.63 7.39 12.41 9.62 2.87 – –
0.23 0.43 0.61 1.34 0.50 0.53 0.44 0.46 0.45 1.03
5
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Fig. 4. SEM images of the cross-section of the oxide scale after 1000 h ASR test and corresponding concentration profiles of different elements determined by EDS line scan across the oxide scale, (a, b) SUS430 (c, d) SUS430 þ 3 wt% LaCrO3 (e, f) SUS430 þ3 wt% CeO2 (g, h) SUS430 þ3 wt% Y2O3.
(~0.939 μm) and the morphology of the oxide scale was comparable to that of SUS430 [56]. In Fig. 4 (b, d, f, and h), the EDS line scan results are provided to highlight the composition of the oxide scale along the thickness. From the EDS line scans, it can be easily seen that SUS430þ3 wt% CeO2, and SUS430þ3 wt% LaCrO3 show higher concentrations of Mn toward the outer region of the oxide scale, while Mn was present in the middle part of the oxide layer in the other samples. As discussed above, the presence of a high concentration of Mn in the outer section of the oxide layer is significant in determining the thickness of the oxide layer, as well as its conductivity. The Mn-Cr spinel enhances the conductivity of the oxide scale and might reduce the ASR values during long-term oxidation testing. Fig. S5 (a-d) shows SEM scale thickness and morphology of the oxide scale after 1000 h long-term tests as a function of CeO2 amount. It is noted that the oxide scale thickness is decreased with an increasing amount of nano-CeO2. With addition of 0.5 wt% CeO2, the average thickness of the oxide scale after 1000 h was 0.919 μm which reduced to 0.386 μm for SUS430þ3 wt% CeO2. The XRD analyses of the surface of the oxide scale (Fig. 5) after 100, 500, and 1000 h revealed the different nature of the phases present
during the oxidation test. Fig. 5 (a) shows the XRD analysis results of the surface scale of SUS430 after 100, 500, and 1000 h continuous isothermal oxidation at 800 � C in air. The XRD peaks are matched with JCPDS cards for Fe, Fe2O3, Cr2Fe2O3, Cr2O3, and Mn1.5Cr1.5O4. It can be observed that, after 100 h, small intensity peaks started appearing; these mostly correspond to Cr2O3 (JCPS # 38–1479). There are no matching peaks with Mn-Cr phases for the 100 h sample. However, the 500 h test sample for SUS430 shows one or two small peaks matching Mn1.5Cr1.5O4 (JCPDS# 33–0892) and, after 1000 h oxidation, the Mn-Cr peaks dis appeared and only Cr2O3 peaks were present. This trend has two possible explanations. First, the overall oxidation of SUS430 is domi nated by the formation of Cr2O3 scale and only a smaller amount of MnCr spinel is present but undetectable in the samples at 500 and 1000 h. The other possible reason is the preferential oxidation of Cr and Mn in the steel. Hence, during the first 100 h, the Cr present near the surface of the substrate was oxidized and formed a chromia layer. After the for mation of the Cr2O3 layer during the initial oxidation, the depletion of Cr ions near the surface might have caused the diffusion of Mn ions, which have faster diffusion rate than that of Cr ions. However, in the late stages of oxidation, after 1000 h, when most of the Mn near the substrate 6
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Fig. 5. XRD analysis of surface scale of SUS430 alloys after 100, 500, and 1000 h continuous isothermal oxidation at 800 � C in air. (a) SUS430 (b) SUS430þ3 wt% LaCrO3 (c) SUS430þ3 wt% CeO2 and (d) SUS430þ3 wt% Y2O3.
surface had been consumed, Cr diffusion might have started again and formed a layer of Cr2O3 on the top. Fig. 5 (b – d) show XRD analysis results for the surface of the oxide scale developed on the SUS430þnano-oxide alloy during oxidation at 800 � C. During the initial 100 h of the oxidation, the XRD phases are similar in all three samples, which indicate that the initial oxidations of the alloys are almost identical and the addition of the nano-oxide does not change this initial oxidation significantly. However, when the XRD peaks at 500 h and 1000 h are compared for all three developed alloys, it should be noted that the Mn-Cr phase is present in the SUS430 þ CeO2 and SUS430 þ LaCrO3 samples, but not in the SUS430 þ Y2O3 samples. The evolution of the oxide scale with the oxidation time shows that the presence of nano-oxides within the matrix of ferritic stainless steel might have altered the growth kinetics of the oxide scale, resulting in different morphologies and phases at different times. It is clear that the samples with stable long-term ASR characteristics have Mn-Cr spinel layers on the top and that the spinel structures are present together with chromia throughout the scale. It has been reported in the literature that Mn-Cr spinel has very high conductivity, resulting in a smaller ASR increase during a 1000 h test [57]. However, for the case of SUS430þ3 wt% Y2O3, only Cr2O3 was present.
oxidation test. However, the addition of 3 wt% Y2O3 did not result in significant improvement in the conductivity or the ASR properties of the interconnect alloys. Furthermore, it was observed, as shown in Fig. 6(a), that the SUS430þ3 wt% Y2O3 sample showed an large increase in ASR during initial 300 h and, after that, the rate of increase in the ASR is slow as compared with the SUS430 sample. The oxide scale of the Y con taining Steel is known to consist of Cr2O3, Y2O3, YCrO3, and the for mation of conductive YCrO3 is a very slow process [62]. Thus the initially large ASR increase of the SUS430þ3 wt% Y2O3 sample is considered to be due to presence of low conductive Y2O3in the oxide scale, and then its subsequent oxidation forms the stable and relatively thin oxide scale probably with small amount of YCrO3. At the end of the 1000 h test, the ASR of the SUS430þ3 wt% Y2O3 was 57 mΩ cm2. The increase in the ASR values of the SUS430þ3 wt% Y2O3 during 1000 h is very high in comparison with other nano-oxide added alloys. Long-term ASR graphs for the developed alloys indicate that the addition of nano CeO2 and LaCrO3 oxides to SUS430 significantly improves the electrical properties of these materials for use as a SOFC interconnect. Fig. 6(b) shows the effect of temperature on the ASR of the oxide scale measured at 650, 700, 750, and 800 � C after 1000 h continuous oxidation. The ASR values of the oxide scale formed on the SUS430 during continuous oxidation decrease with increasing temperature, which is typical behavior of the Cr2O3 based oxides. As the conductivity in the metallic interconnect is dominated by the oxide layer, the higher ASR value for SUS430 and the highest slope of the ASR vs. 1000/T plot indicate that the oxide layer might have been overgrown in the SUS430, causing a reduction in the conductivity. The minimum dependence of the ASR on the temperature was observed in the SUS430þ3 wt% CeO2 samples. This means that the dominant oxidation mechanism is changed due to formation of the stable thin oxide scale including Mn-Cr spinel layers. The effect of the nano-CeO2 amount on the long-term ASR properties of the SUS430 alloys was also studied. Samples were prepared with SUS430 þ xwt.% CeO2 (x ¼ 0.5, 1, 3, and 5) and subjected to a long-term
3.3. ASR behavior of the nano-oxide dispersed SUS430 alloys Fig. 6(a) shows the change in the ASR values of various samples during the 1000 h continuous test. The SUS430 sample shows a continuous increase in the ASR value. The measured ASR values depend on the oxide scale because the electrical conductivity of the metal sub strate is very high in several order. The initial ASR value at 800 � C for SUS430 was 21 mΩ cm2. However, after 1000 h test at 800 � C, the ASR value reached 89 mΩ cm2. Similarly, when we compared the behavior of the nano-oxide dispersed alloy during the long-term ASR test, it was found that the addition of 3 wt% nano-CeO2 and -LaCrO3 resulted in substantially improved ASR characteristics during the long-term 7
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Fig. 6. (a) Long-term ASR values of the SUS430 samples with 3 wt% nano CeO2, 3 wt% nano LaCrO3, and 3 wt% nano Y2O3, measured at 800 � C in air. (b) Effect of temperature on the ASR of the samples after 1000 h test. The SUS 430 is used as a reference sample. (c) Long-term ASR values of the SUS430 with 0.5, 1, 3, and 5 wt % nano-CeO2, (d) ASR behavior of the nano CeO2 dispersed alloys during 10 thermal cycles between 400 and 800 � C in air, (e) Change in cathodic overpotentials of the LSCF/GDC/Pt half-cells with different interconnects, measured at a constant current density of 50 mA cm 2.
ASR test under the same conditions as described in section 2.3. Fig. Fig. 6 (c) shows the trends of ASR of the specimen measured during the longterm test at 800 � C. On increasing the amount of nano-CeO2, the ASR property of the SUS430 improved significantly. A lower ASR value in the 5 wt% CeO2 addition is observed due to the formation of a thinner oxide layer. The CeO2 addition of smaller amounts such as 0.5 and 1 wt % also decreased sufficiently the ASR during the initial hours of operation. However, at longer times, the ASR was increased rapidly. This might be attributed to the small quantity of nano CeO2 particles, which failed to mitigate the transport of cations for longer time periods. Longer time tests in the low content of CeO2, therefore, produced the formation of more Cr2O3 and Fe3O4, resulting in a fast increase in the ASR values. Significant amounts, such as 3 wt % and 5 wt % of CeO2, appear to have promisingly stopped the cation transport and thus their ASR values maintained to be quite low for long time test. Fig. 6(d) shows that the nano-CeO2 dispersed SUS430 alloys show stable behavior during the
thermal cycles. The samples were exposed to 10 consecutive thermal cycles between 400 and 800 � C; the ASR values were measured at 800 � C after each thermal cycle. The ASR values were almost stable for all the nano-CeO2 added alloys, without any apparent delamination of the oxide layer. Fig. 6(e) shows the change in cathodic overpotential of the LSCF electrode measured in the presence of the developed interconnect alloys. For comparison, we used SUS430 and Crofer 22 APU interconnects as references. All the half cells showed an increase in the cathodic over potential with time, which is due to LSCF degradation in Cr environ ment. The half cells with SUS430þ3 wt.CeO2 and Crofer 22 APU showed a relatively low increase in the cathodic overpotential with time. It is well known that Cr species in gaseous phase are released from chromia scales on the interconnect surface and then are electrochemically deposited at the LSCF electrode. The Cr oxide deposits not only reduce the active sites for oxygen reduction, but also block the migration of 8
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oxygen ions from the electrode to the electrolyte, which results in the performance degradation [39]. As compared to SUS430, the cell with SUS430 þ 3 wt% CeO2 interconnect showed lower change in cathodic overpotential. This indicates that the stable oxide scale formation in the nano CeO2-dispersed SUS430 retards a vaporization of gaseous Cr spe cies from the interconnect, and thus amount of the Cr oxide deposit in the LSCF cathode decreases. Further detailed investigations are required to make it clear.
segregation model might be applicable. The addition of dispersed nano-CeO2 has demonstrated superior properties regarding the oxida tion resistance of newly developed alloys; this might be due to the better oxygen affinity of CeO2, which can influence the cation and anion diffusion across the oxide scale [60]. Moreover, CeO2 has a different oxidation state from that of Cr2O3, and it improves intermetallic contact and remains at the grain boundary. This means that even during long-term oxidation, CeO2 nanoparticles might be present at the grain boundary, thereby hindering the diffusion of Cr ions outward. Form Fig. 2(c and d), we can see that the CeO2 nanoparticles are present at the grain boundaries in the form of 1–2 micron-sized agglomerates. The presence of the agglomerated CeO2 particles at the grain boundary of the alloy might have improved the oxidation properties of the CeO2 added SUS430 alloy due to retarding the Cr ion diffusion outward. The detailed investigation of the oxide scale formation indicates that, due to the addition of the nano-oxides in the alloy, the oxidation kinetics might have been altered, as depicted by schematics in Fig. 7. Upon startup of oxidation, the favorable oxidation of chromium near the free surface of the alloy can be enhanced by rare earth oxide ceria, and an integrated and densely packed Cr2O3 oxide layer can then be obtained. Ce shows good affinity towards oxygen and the cerium oxides accelerate the formation of protective oxidation layers by providing the nucleation sites that rapidly form Cr2O3 oxides layers [44]. Fig. S5 shows a sig nificant decrease in oxide layer thickness with an increment of Ce con tent. At the higher content of CeO2, the more incorporation of CeO2 into the oxide scale is clear to produce more thin stable oxide scale, and the more Ce at grain boundaries of the oxide scale prevents the growth of oxide scale due to retarded cation diffusion during long-term oxidation. Grain boundaries also affect the high-temperature oxidation resistance of metals. Upon decreasing the grain size, grain boundary area increases, which makes the diffusion of Cr better, leading to a decrease of the critical concentration of chromium to form a Cr2O3 layer [14]. Actually, more CeO2 addition produced finer grain size in the nano oxide-dispersed SUS 430. Although there is grain size effect in the oxidation kinetics, in the present work the effect of CeO2 content was greater. Nevertheless, it is still needed to find out the precise manner of interaction of these materials with cations and anions across the oxide layer. We are studying this further and hoping to report following this work.
3.4. Oxidation mechanism of the nano-oxide dispersed SUS 430s Hou and Stinger [40] reported that, for Cr2O3-forming alloys, the presence of a reactive element [RE] or its oxide can reduce the growth rate of the chromia scale and improve the adhesion of the scale to the substrate. The grain boundary segregation model has proven to be one of the most credible theories describing the proposed mechanisms for RE effect [43,44]. The growth of chromia scale with outward diffusion of Cr cations is carried out below 1000 � C temperature [58]. From the experimentally monitored rates of oxidation, chromium cation diffusion in Cr2O3 bulk form is found to be quite slow [59]. Therefore, it is considered that Cr cation diffusion during oxidation takes place through ‘short circuits’, i.e. at the points of grain boundaries and dislocations [60]. In case of the SUS 430 with the RE, the grain boundary segregation model explains that the reactive elements create hindrance at the grain boundaries and thus retard cations to diffuse through the grain bound aries. Yurek et al. explained that there is significant effect on oxidation of Ni–Cr alloy at 1000 � C when Y2O3 is added. They searched for the oxide scale, which mainly consists of NiO layer and chromia layer. Ev idence for the presence of Y was found at the grain boundaries of the NiO and chromia layer upon performing the STEM and EDS [61]. Czerwinski and Szpunar [59] reported that when Ni and Cr based superalloys are coated by the CeO2 prepared by sol-gel method, changes in oxidation are observed. The oxide-gas interface region was found to be rich in Ce for the oxide layer of Ni formed at 700 � C. They also found the presence of Ce at the NiO grain boundaries by EDS. In the present study, it has been shown that upon adding the nano particles of RE oxides, oxide scale thickness get reduced and adhesion of the oxide scale becomes better. Whittle and Stringer [43] argued that the higher number of RE oxide particles present at the alloy surface assist the initial oxidation of Cr at the start of oxidation and a homogeneous oxide layer is formed. In the later stages of oxidation, the grain boundary
Fig. 7. Schematics of the mechanism of formation of oxide scale on nano-oxide dispersed SUS430 alloy. 9
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4. Conclusions [8]
The addition of nano-oxides of reactive elements in SUS430 has significant effects on the oxidation and electrical characteristics of ferritic stainless steel during long-term test for SOFC interconnect application. We conducted a detailed investigation of the oxidation and electrical properties of different SUS430 based alloys modified with various nano-oxides. The sintered samples shows a uniform distribution of the nano-oxide throughout the pellet. XRD data indicate that the added oxide particles are quite stable and do not react with the base steel. Oxidation properties of the newly prepared alloys were deter mined by long-term ASR test at 800 � C for 1000 h. It was observed that higher amount of nano CeO2 greatly improves the performance of SUS430 by lowering the ASR value to as low as 20 mΩ cm2, which is comparable with the commercially used Crofer 22 APU. In the ASR test, the maximum increases in the ASR due to oxide scale growth were just 15, 29, and 56 mΩcm2 for CeO2, LaCrO3, and Y2O3 additions, respec tively. However, the ASR of SUS430 without any nano-oxide reached 100 mΩcm2. It was also found that for the case of the nano-CeO2 dispersed alloy, the top layer of the oxide scale contains more Mn-Cr rich spinel compared to the other developed alloys. The addition of nanooxides was found to improve the oxide scale adherence with the sub strate, forming a thin dense oxide layer. Therefore, the lower increase in the ASR for the dispersed nano-CeO2 alloys was attributed to formation of conductive Mn-Cr spinel layer, Ce incorporation into the oxide scale and retarding the cation diffusion at the grain boundary, which produces the lower thickness (~0.4 μm) and the improved adhesion of oxide scale to the substrate. Finally, it is concluded that the nano-oxide dispersed alloys can be successfully employed as SOFC interconnects because of their excellent chemical, electrical, and oxidation properties.
[9] [10] [11] [12]
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[14] [15]
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Acknowledgments
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This work was supported by the Technology Innovation Program, (No. 10080607) funded by the Ministry of Trade, Industry and Energy, Korea, and also supported by a National Research Foundation (NRF) grant from the Ministry of Science and ICT, Korea under Technology Development Program to Solve Climate Change (No. NRF2017M1A2A2044926).
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Appendix A. Supplementary data [24]
Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227109.
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Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
Highly durable nano-oxide dispersed ferritic stainless steel interconnects for intermediate temperature solid oxide fuel cells Muhammad Taqi Mehran a, b, Tae-Hun Kim a, c, Muhammad Zubair Khan a, Seung-Bok Lee a, Tak-Hyoung Lim a, Rak-Hyun Song a, c, * a b c
Fuel Cell Research Center, Korea Institute of Energy Research (KIER), 152 Gajeong-ro, Yuseong-gu, Daejeon, 34129, South Korea School of Chemical and Materials Engineering (SCME), National University of Sciences and Technology (NUST), Main Campus, H-12, Islamabad, 44000, Pakistan Korea University of Science and Technology, 217 Gajeong-ro, Yuseong-gu, Daejeon, 34113, South Korea
H I G H L I G H T S
G R A P H I C A L A B S T R A C T
� Nano-CeO2 dispersed SUS430 shows better oxidation properties for long-term testing. � Nano oxide dispersion in steel improves oxide-scale adhesion and thickness. � 3 wt% CeO2 dispersed steel alloy results 15 mΩcm2 ASR increase in 1000 h at 800 � C. � Presence of dispersed oxides at grain boundary influences the oxidation kinetics.
A R T I C L E I N F O
A B S T R A C T
Keywords: Metallic interconnects Reactive element Solid oxide fuel cell Area-specific resistance
Herein, ferritic stainless steel alloys are developed with the dispersion of nano-oxides of reactive elements to improve their oxidation characteristics for solid oxide fuel cell (SOFC) interconnects. Nano-oxides of yttrium, cerium, and lanthanum chromite are homogeneously dispersed in a ferritic stainless steel base powder and specimen pellets are prepared by pressing the powder and sintering it in hydrogen at 1400 � C. The nano-oxide dispersed alloys are tested for oxidation characteristics and electrical properties at 800 � C for 1000 h in air. The results show that the ferritic stainless steel with 3 wt% nano-ceria exhibits significantly improved areaspecific resistance (ASR) characteristics. The increase in the ASR for this alloy is found to be only 15 mΩ cm2/kh and is ascribed to the thin dense oxide layer (~0.4 μm), the presence of spinel (MnCrO4) at the outer side of the oxide layer, and the improved oxide layer adherence to the metal substrate. The addition of reactive element oxides altered the oxidation kinetics of the Fe-Cr alloys, thus making nano-oxide dispersed ferritic stainless steel highly durable and most suitable for SOFC interconnects. The oxidation kinetics of the nano-oxide dispersed steel was also discussed as related to cation diffusion.
* Corresponding author. Fuel Cell Research Center, Korea Institute of Energy Research (KIER), 152 Gajeong-ro, Yuseong-gu, Daejeon, 34129, South Korea. E-mail address: [email protected] (R.-H. Song). https://doi.org/10.1016/j.jpowsour.2019.227109 Received 11 July 2019; Received in revised form 26 August 2019; Accepted 3 September 2019 Available online 5 September 2019 0378-7753/© 2019 Elsevier B.V. All rights reserved.
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~10 mΩcm2. To control the scale composition and conductivity, various commercial alloys named as Crofer 22 APU, ZMG232, etc. have been structured by modifying the Fe-Cr alloy [24,34]. During recent years, the development of the alloy with the oxide dispersion has made considerable progress [36,37]. The better adher ence of oxide layer to base metal and deceleration of cations to the interface were caused by these oxide particles as a result of improved oxidation resistance of the base alloys [11,38]. When RE’s are added, in one of the following two forms of metallic or oxide dispersion, proved to be quite beneficial in enhancing the scale adhesion to the substrate [39–42]. Y, La, Ce, and Zr are mostly used as REs. Quadakkers et al. and Whittle et al. mentioned that, besides developing the adhesion proper ties, the REs also play a vital role in giving a considerable decrement in the scale growth of oxides and promoting the selective oxidization of Cr in low-Cr alloys [6,43]. The scale growth can be suppressed by the introduction of REs which reduces outward Cr transport. The oxidation resistance of steel alloys along with the better mechanical properties can be achieved by adding RE-oxides in dispersed phase [44,45]. Linderoth et al. reported that the oxide dispersed chromium alloys is having a very lower rate of oxidation than that of chromium alloys with availability of no reactive element oxides [46]. Quadakker et al. presented his work on improving the oxidation resistance of alloy by adding rare-earth ele ments/reactive elements (REs) that form RE oxides at the boundaries of grain [47]. The coatings of reactive element exceptionally improved oxidation properties of Iron–Chromium steels for long run applications of SOFC [15–17]. Fontana et al. [17] investigated the effect of coatings (thickness: 102– 2 � 102nm) of La2O3, Nd2O3, and Y2O3 on the proper ties of oxidation and conductivity for three different ferritic stainless steels, Crofer 22 APU, AL 453, and Haynes 230. The coated samples showed lower growth rates for oxide and improved scale adhesion to the substrate. Furthermore, the best result for conductivity was showed by La2O3 coated Crofer 22 APU. Fontana et al. in another study [48] investigated long-term exposure for more than two years at 800 � C in air of La2O3 and Y2O3 coated Crofer 22 APU. They reported enhanced oxidation resistance, slower growth rates of oxides, and reduced elec trical resistance within the coated samples compared to the uncoated material [49]. Commercial SUS430 has proved to be a cost-effective and cheap interconnect material in comparison to Crofer 22 APU, and when nanooxides of REs are alloyed in the SUS430 base material, gave the results for possible improvement in oxidation protection, electronic conduc tion, and chromium retention for SOFC systems [50,51]. In our previous study [52], we reported that the addition of nano-oxides significantly improves the properties of SUS430 as an interconnect. However, the oxide layer development and the influence of the addition of nano-oxides on the oxidation properties during the long-term test should be further investigated. The main focus of this work was to configure different ferritic steel-based nano-oxide dispersed alloys (nano-CeO2, -Y2O3, and -LaCrO3 dispersed SUS430 steel) by using an improved powder mixing method. The dispersion of nano-oxides of the reactive elements influenced the oxide scale thickness, adhesion, and chemistry, and we have therefore conducted a detailed study of the oxide scale and the electrical properties of RE-oxide dispersed SUS430 alloys during a 1000 h ASR test at 800 � C in air. The thermally grown oxide scale was accordingly characterized based on the composition, microstructure, and electrical properties. The evaluation of grown oxide was done to determine the suitability of the oxide dispersed alloys as interconnect materials for IT-SOFCs.
1. Introduction The solid oxide fuel cell (SOFC) stack is formed as a lamellar struc ture of planar cells in order to reduce the internal resistance and to in crease the effective electrode area per unit volume. To separate two adjacent cells in a stack, a separator or interconnector is employed [1,2]. To apply such a device as a SOFC interconnect, the material of the separator should have a thermal expansion coefficient (TEC) closer to that of the other cell components, as well as high conductivity. Because the SOFC’s operating temperature is high, the interconnect material must possess good corrosion resistance in fuel and air electrode envi ronments [3–5]. Typical interconnects are formed of electrically conductive ceramic materials such as LaCrO3. However, these ceramic materials are relatively expensive and difficult to fabricate. Also, due to recent improvements that have reduced the operating temperature of the SOFCs below 800 � C, many research groups are extensively search ing for metallic interconnects materials. Metallic materials exhibit higher thermal and electrical conductivities as compared to ceramic interconnects and are less costly and can be fabricated easily [1,6–8]. Ferritic stainless steel (Fe-Cr) based alloys are considered to be a suitable metallic interconnect material where preferential oxidation of Chromium to Chromium oxide would configure a continuous protective oxide layer. However, during long-term exposure to the high tempera ture oxidizing and reducing atmosphere in the SOFC, the Cr-containing steel experiences continuous growth of an oxide layer on the steel sub strate [9]. This oxide layer has lesser conductivity and a different TEC than those of the base material, thus causing a substantial increase in the area-specific resistance (ASR) and spallation of the oxide layer. The ASR values of the currently developed ferritic stainless steel alloys are pro jected to be 150 mΩ-cm2 after 40,000 h operation at 850 � C; the oxide thickness after long-term exposure is expected to be more than 30 μm, which can cause delamination and cracking during the thermal cycling operation of the SOFC [10,11]. In addition to the oxide scale growth, Cr poisoning is also a significant issue related to the use of Cr2O3-forming alloys. Cr-containing gaseous species are formed from the oxide scales, which results in poisoning of the electrodes and performance degrada tion of the SOFCs [8,12–14]. In order to improve the properties of Cr2O3-forming alloys such that they will be suitable for application in high-temperature SOFC service conditions, surface modification and coatings are applied. Reactive element (RE) [15–17], spinel [18,19], and perovskite [20–23] oxide coatings significantly mitigate the Cr-poisoning and the conductivity of the oxide scale is also reported to be improved. In particular, Cr-free spinel coatings such as (Mn, Co)3O4 [24,25], (Cu,Mn)3O4 [19], Cu-Fe spinel [26], and NiCo2O3 [27] have shown promising results for SOFC interconnects. Similarly, many reports suggest that the application of a protective coating of reactive element oxides (such as La [28], Ce [29], and/or Y [30]) on the metallic interconnects enhances the electrical properties [31]. However, while applying surface coatings can effec tively inhibit Cr poisoning, practical improvement of the electrical conductivity and oxidation resistance of the metallic interconnect cannot be achieved for long-term operation. Moreover, the coating process increases the manufacturing cost of the metallic interconnects. In order to meet the requirements of an appropriate TEC and suitable oxide scale properties for the metallic interconnects, modifi cation of Fe-Cr alloys is considered to be a viable approach [9,24,32,33]. Horita et al. [34] added Nb (0.35 wt%) and Mo (0.5 wt%) to Fe-20Cr alloy and investigated the oxidation properties in an H2/H2O atmo sphere at 800 � C for more than 1000 h. They reported that, in addition to frequently observed Cr2O3 and Cr-Mn spinel in the scale and at the grain boundary, an Nb lava phase is also formed. During a long-term test, the grain boundary lava-phase was found which resulted in an increment in the oxidation resistance and a decrease in the ASR. Similarly, Hua et al. [35] developed a promising ferritic steel alloy Fe-17Cr-1Mn-0.5Ti with the addition of La, Y, and Zr; After running 1000 h isothermal oxidation test at 750 � C in air, this alloy showed an ASR value as low as
2. Experimental 2.1. Preparation of powders and pellets Commercial nano-CeO2 and -Y2O3 powders (Sigma-Aldrich, USA, particle size >25 nm) were used for blending with the steel powder for commercial use such as SUS430 (Metal Player, Korea) to develop the 2
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steel with dispersed oxide; nano-LaCrO3 (Kanto Chemical Co., Japan) had an average particle size of < 50 nm. Crofer 22APU Fe-Cr alloy (ThyssenKrupp VDM, Germany) was used as a reference material. The chemical compositions of the starting alloys (SUS430 and Crofer 22 APU) are presented in Table 1. The composition of SUS430 was deter mined by using an Inductively Coupled Plasma Mass Spectrometer (ICPMS, model: NexION 300D, PerkinElmer). SUS430, SUS430þ3 wt% CeO2, SUS430þ3 wt% Y2O3, and SUS430þ3 wt% LaCrO3 were homogenized by mixing in a high energy planetary ball mill for several hours (12–48 h) using polyvinyl butyral (PVB) which binds a mixture of isopropyl alcohol and toluene in a stainless steel jar using stainless steel balls. The high-energy ball mill ensured homogenous mixing of the SUS430 and nano-oxide powders. The powder mixtures were then compacted at 600 MPa and 200 MPa via uniaxial and cold isostatic pressing, respectively, to yield specimen pellets of 5 mm � 8 mm � 25 mm dimensions. The pellets were sintered in hydrogen for 10 h at 1400 � C to obtain dense pellet samples. The details of sample processing and sintering are given in Fig. S1 in the supplementary material. 2.2. Material characterization The sintered pellets were first ground using 1000 grit SiC paper and dried at 100 � C for 2 h before density measurement by the Archimedes method. For each measurement, at least three pellets were used and the average relative density was determined. The sintered pellets were characterized by SEM and XRD analyses. The surface and cross-section of the selected specimens were analyzed by a scanning electron micro scope (SEM, Hitachi S4800) equipped with an energy dispersive spec trometer (EDS), which was used to assess the thickness, structural morphology, and element distribution of the oxide scales. For observa tion of the grain structure, the polished specimens were etched with aqua regia þ glycerol for 30 s before the SEM analysis. The crystalline phases of the oxidation products on the surfaces of these specimens were identified by X-Ray diffraction (XRD). Additionally, cross-sections of the specimens were also examined using SEM/EDS. The thermal expansion coefficient (TEC) of the alloys was evaluated from room temperature to 800 � C at a rate of 5 � C per minute. Air flowed into the chamber at 100 cc.min 1. To measure the TEC, the device was first calibrated with the standard Al2O3 specimen and the calibration was used to determine the thermal expansion behavior of the SUS430 alloys up to 800 � C.
Fig. 1. (a) Detailed schematics showing the positions of Pt wires and mesh attached to the pellet for ASR measurement test; (b–d) Jig and application of load on the sample using steel slabs to ensure appropriate contact. (e,f) Actual image of final ASR test samples.
% CeO2, and SUS430þ3 wt% Y2O3 to the SOFC cathode environment (800 � C, air) during 1000 h, and the ASR was measured every 20 h by the four-probe method [53]. To measure the ASR of the samples, 50–300 mA current was applied by a current source (Toyotech, DP 30-03 TP) at an interval of 50 mA and the respective values of voltage were measured using an Agilent Mul timeter (model: 34450A). The ASR was calculated for each sample using Equation (1) for each current applied and the average ASR value was determined. 2.4. Cr poisoning experiment
2.3. Long-term ASR testing set-up
To study relative Cr poisoning of a cathode in the developed alloy based interconnects, a half-cell was fabricated as reported elsewhere [19]. Gd2O3-doped CeO2 (GDC) pellet (thickness: 1 mm) was used as the electrolyte and a slurry of La0.6Sr0.4Co0.2Fe0.8O3 (LSCF) cathode mate rial was coated on one side of the sintered GDC coin cell using a screen-printing and sintered in air for 3 h at 1150 � C. The area of the coated LSCF working electrode was 1.0 cm2. The reference and counter electrodes were prepared by applying a conductive Pt paste on the other side and at the edge of GDC coin cell, respectively, and fired in air at for 2 h 1000 � C. Pt meshes and wires were used as current collectors and voltage probes. Three metallic interconnect specimen SUS430, SUS430þ3 wt%CeO2 and Crofer 22 APU were prepared with area of 1.5 cm2, thus entirely covering the cathode active area. The alloy specimens were placed onto the Pt mesh current collector for the working electrode. Humidified air (3% H2O) was supplied to the cell at 800 � C. Half-cell voltage was measured at a constant current density of 50 mA cm 2 and AC-impedance spectra were periodically measured at an OCV. Electrochemical experiments were carried out using Biologic Potentiostat (SP-240).
To prepare the samples for the long-term ASR test, the surfaces of the specimens were polished mechanically with SiC paper up to 2000 grade. The samples were then cleaned ultrasonically in acetone for 30 min. Preoxidation was conducted for 100 h at 800 � C to thermally grow the layer of oxide on the pellet’s surface. After the pre-oxidation, platinum (Pt) wire was connected to the ASR samples for current collection, as shown in Fig. 1 (a). Pt paste was used to minimize the interfacial resistance between the current collectors and the samples [53]. The ASR samples were placed in a specially designed jig, which provided a compressive pressure of 50 kPa on the sample by the load (Fig. 1(b–f)) [54]. The prepared alloys were tested to determine their long-term ASR behavior by exposing the alloys SUS430, SUS430þ3 wt% LaCrO3, SUS430 þ3 wt Table 1 Chemical composition of the ferritic stainless steel alloys (SUS430 and Crofer 22APU). Steel alloy SUS430 Crofer 22APU a
a
Element Composition (wt. %) Fe Cr S P Balance 18.08 0.028 0.035 Balance 20–24 0.02 0.05
C 0.079 0.03
Si 1.02 0.5
Mn 0.26 0.3–0.8
ThyssenKrupp Crofer® 22 APU Datasheet No. 4046. 3
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3. Results and discussion
43 μm for SUS430, SUS430 þ 3 wt % LaCrO3, SUS430 þ 3 wt % CeO2, and SUS430þ 3 wt % Y2O3, respectively.
3.1. Microstructure analysis of the sintered nano-oxide dispersed steel alloys
3.2. Oxidation behaviors of the nano-oxide dispersed steel alloys during long-term ASR tests
Fig. 2 shows SEM images of the nano-oxide added SUS430 powder during the processing. The supplied SUS430 powder have a particle size between 20 and 60 μm, as shown in Fig. 2(a) and 2(b) shows the morphology of the particles after 12 h ball milling in the high energy ball mill, yielded a clear decrease in the particle size. Particle size analysis of the powder after high energy ball milling process is shown in Fig. S2. Fig. 2(c and d) present the surface morphology of the developed ODS alloy with a SUS430-3 wt % CeO2 composition and it can be noticed that the added nano-oxides are agglomerated at the grain boundaries with an average particle size of 1–5 μm. The quantitative results of the EDS are given in Fig. 2(e). To find the normalized mass percentage of the ele ments present in the alloy, point EDS was used and presence of CeO2 at the grain boundaries is confirmed via EDS analysis. The morphology of the sintered alloy shown in Fig. 2(c) indicates that the specimens were fully sintered and furthermore no pores were observed; the relative density of all the developed alloys was found to be more than 95%. For the sintered alloys, the average grain sizes were ~160, 30, 34, and
The morphology of the oxide scale grown on the substrate during 1000 h isothermal oxidation at 800 � C is studied by SEM and EDS ana lyses. Fig. 3 provides a comparison of the surfaces of the oxide scales for SUS430, and SUS430þnano-oxide alloys after 1000 h. It can be noted that the morphologies of SUS430þ3 wt% CeO2 and SUS430þ3 wt% LaCrO3 resemble that of Crofer 22 APU reported by Qu et al. [30]. However, the particle shapes of the oxide scales with CeO2 and LaCrO3 nano oxides vary from tetrahedral to diamond-shapes and had crystal shape scales with distinctive facets. The oxide scales formed on the surface of SUS430 and SUS430þ3 wt% Y2O3 appear to have round-faced particles. In the literatures, it is reported that prism and diamond type particle shapes belong to (Mn, Cr)3O4 spinel and that Cr2O3 oxide scale forms corundum shaped structures [35,55]. Although it is evident from the SEM micrographs given in Fig. 3 that the respective additions of nano-CeO2 and -LaCrO3 have a clear influ ence on the morphology of the oxide scale, we performed a detailed EDS
Fig. 2. (a) As-received SUS430 alloy powder from Metal Player®, Korea (b) after mixing of nano-oxides and SUS430 powder by using high energy ball-milling for 12 h, (c, d) Backscattered SEM analysis of sintered pellet of the SUS430þ3 wt% CeO2 alloy showing presence of the agglomerated nanoparticles of CeO2 at the grain boundaries (e) EDS analysis to confirm the presence of agglomerated nano-particles at the grain boundaries of SUS430þ3 wt % CeO2 alloy. 4
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Fig. 3. Surface morphology of the oxide scale after 1000 h isothermal oxidation of SUS430, and SUS430 þ 3 wt% nano-oxides: (a) SUS430, (b) SUS430þ3 wt% LaCrO3, (c) SUS430þ3 wt%CeO2, (d) SUS430þ3 wt%Y2O3. The arrows with numbers show the points of quantitative EDS analyses.
analysis of the different particles present on the oxide scale (marked in Fig. 3) to quantify the nature of these particles. The results are presented in Table 2. From the EDS results, we can deduce that the compositions of the two types of particles are almost identical in the samples with nanoCeO2 and -LaCrO3, although they appear to have different sizes and shapes in the SEM micrograph. From Fig. 3 (a), we can see that, at point 4, the particle appears to be overgrown and large, while the area near point 5 has many small particles. From the EDS results at points 3 and 5, both types of particles composed mainly of CrO2 and Mn appear to be absent. However, at point 4, Mn is present, but in a very small amount (~0.66%). This means that the surface of the oxide scale of SUS430 after 1000 h isothermal oxidation is predominantly CrO2, along with small traces of Mn-Cr. That is, the oxidation process of SUS430 still might be controlled by the outward migration of Cr ions from the Fr-Cr matrix to the CrO2-Air interface. The large-sized grains of corundum shape may provide evidence of continued growth and Cr2þ diffusion dominated oxidation kinetics in SUS430 [35]. In the SUS430þ3 wt% CeO2 oxide scale, three points were selected for analysis (referred to as points 8, 9, and 10 in Fig. 3 (c)). Point 8 marked particle is overgrown relative to others and exhibits a distinct morphology of Mn-Cr spinel oxides from its composition analysis.
Although the particle 9 and 10 shows Cr-rich phase, most of the observed particles in the scale have a Mn-Cr spinel morphology. In the case of nano-LaCrO3 addition, Mn in the oxide scale was present in a lesser amount in comparison with the SUS430þ3 wt% CeO2 sample. In contrast, it can be seen from the EDS data that with the addition of 3 wt % Y2O3, the morphology and the composition of the oxide scale were not altered to a great extent. The composition of the oxide scale is similar to that of SUS430 without nano-oxide additives. From the EDS data and the morphology of the oxide scale after 1000 h, it is considered that the addition of nano-CeO2 and nano-LaCrO3 might have affected the mechanism of oxidation and the formation of the oxide scale. The higher conductive spinel on the top surface of the oxide scale helps to reduce inward migration of oxide ion and outward diffusion of Cr element, thus modifying the oxidation rates and ASR characteristics. The oxide scale structure, thickness, and adherence to the substrate were also studied using SEM and EDS analyses. The conductivity of the oxide scale is controlled by the scale thickness and, with an increase in the oxidation time, the thickness of the chromia scale increases [4]. Fig. 4(a–h) shows SEM/EDS images of the cross-section of the oxide scale after the 1000 h ASR test. The oxide scale thickness was deter mined by measuring the thickness at 4–5 different locations, after which the average value was calculated. The average oxide scale thickness values for SUS430, SUS430 þ 3 wt% LaCrO2, SUS430 þ3 wt% CeO2, and SUS430 þ3 wt% Y2O3 were 1.823, 0.592, 0.386, and 0.939 μm, respectively. Another aspect of the oxide scales that can be observed from the SEM analysis of the cross-section is that there are no areas of delamination or spallation except SUS 430. As shown in Fig. 4 (a), the oxide scale grown over the SUS430 substrate was not perfectly attached to the substrate. Some gaps can be observed in the SEM image. This may indicate that scale adherence to the substrate might be poor and this might contribute to loss of conductivity of the sample. However, in the SUS430 þ 3 wt% CeO2, a uniform and dense oxide scale is attached tightly with the substrate, as shown in Fig. 4 (e). The oxide scale thickness in the nano-CeO2 added alloy was very small (~0.386 μm) during the 1000 h oxidation test. Similarly, for the case of SUS430 þ 3 wt% LaCrO3, the oxide scale was also similar to that of the ceria added alloy, with a slightly greater thickness (~0.592 μm). For the case of the SUS430þ3 wt% Y2O3 sample, the oxide scale thickness was greater
Table 2 EDS point analysis of the surface of SUS430 and nano-oxide dispersed alloys after 1000 h oxidation in air at 800 � C to determine the composition of the various microstructures of the oxide scale. Sample
Point
SUS 430
3 4 5 6 7 8 9 10 11 12
SUS430 þ 3 wt% LaCrO3 SUS430 þ 3 wt% CeO2 SUS430 þ 3 wt% Y2O3
Elemental Composition (Atomic %) C
O
Cr
Mn
Fe
11.02 6.32 5.28 3.89 4.78 10.34 11.13 12.85 4.79 6.37
66.28 65.55 63.48 64.10 64.75 60.45 60.8 61.65 61.90 49.55
22.5 27.05 30.63 29.04 22.58 16.27 18.12 22.18 32.86 43.04
– 0.66 – 1.63 7.39 12.41 9.62 2.87 – –
0.23 0.43 0.61 1.34 0.50 0.53 0.44 0.46 0.45 1.03
5
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Fig. 4. SEM images of the cross-section of the oxide scale after 1000 h ASR test and corresponding concentration profiles of different elements determined by EDS line scan across the oxide scale, (a, b) SUS430 (c, d) SUS430 þ 3 wt% LaCrO3 (e, f) SUS430 þ3 wt% CeO2 (g, h) SUS430 þ3 wt% Y2O3.
(~0.939 μm) and the morphology of the oxide scale was comparable to that of SUS430 [56]. In Fig. 4 (b, d, f, and h), the EDS line scan results are provided to highlight the composition of the oxide scale along the thickness. From the EDS line scans, it can be easily seen that SUS430þ3 wt% CeO2, and SUS430þ3 wt% LaCrO3 show higher concentrations of Mn toward the outer region of the oxide scale, while Mn was present in the middle part of the oxide layer in the other samples. As discussed above, the presence of a high concentration of Mn in the outer section of the oxide layer is significant in determining the thickness of the oxide layer, as well as its conductivity. The Mn-Cr spinel enhances the conductivity of the oxide scale and might reduce the ASR values during long-term oxidation testing. Fig. S5 (a-d) shows SEM scale thickness and morphology of the oxide scale after 1000 h long-term tests as a function of CeO2 amount. It is noted that the oxide scale thickness is decreased with an increasing amount of nano-CeO2. With addition of 0.5 wt% CeO2, the average thickness of the oxide scale after 1000 h was 0.919 μm which reduced to 0.386 μm for SUS430þ3 wt% CeO2. The XRD analyses of the surface of the oxide scale (Fig. 5) after 100, 500, and 1000 h revealed the different nature of the phases present
during the oxidation test. Fig. 5 (a) shows the XRD analysis results of the surface scale of SUS430 after 100, 500, and 1000 h continuous isothermal oxidation at 800 � C in air. The XRD peaks are matched with JCPDS cards for Fe, Fe2O3, Cr2Fe2O3, Cr2O3, and Mn1.5Cr1.5O4. It can be observed that, after 100 h, small intensity peaks started appearing; these mostly correspond to Cr2O3 (JCPS # 38–1479). There are no matching peaks with Mn-Cr phases for the 100 h sample. However, the 500 h test sample for SUS430 shows one or two small peaks matching Mn1.5Cr1.5O4 (JCPDS# 33–0892) and, after 1000 h oxidation, the Mn-Cr peaks dis appeared and only Cr2O3 peaks were present. This trend has two possible explanations. First, the overall oxidation of SUS430 is domi nated by the formation of Cr2O3 scale and only a smaller amount of MnCr spinel is present but undetectable in the samples at 500 and 1000 h. The other possible reason is the preferential oxidation of Cr and Mn in the steel. Hence, during the first 100 h, the Cr present near the surface of the substrate was oxidized and formed a chromia layer. After the for mation of the Cr2O3 layer during the initial oxidation, the depletion of Cr ions near the surface might have caused the diffusion of Mn ions, which have faster diffusion rate than that of Cr ions. However, in the late stages of oxidation, after 1000 h, when most of the Mn near the substrate 6
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Fig. 5. XRD analysis of surface scale of SUS430 alloys after 100, 500, and 1000 h continuous isothermal oxidation at 800 � C in air. (a) SUS430 (b) SUS430þ3 wt% LaCrO3 (c) SUS430þ3 wt% CeO2 and (d) SUS430þ3 wt% Y2O3.
surface had been consumed, Cr diffusion might have started again and formed a layer of Cr2O3 on the top. Fig. 5 (b – d) show XRD analysis results for the surface of the oxide scale developed on the SUS430þnano-oxide alloy during oxidation at 800 � C. During the initial 100 h of the oxidation, the XRD phases are similar in all three samples, which indicate that the initial oxidations of the alloys are almost identical and the addition of the nano-oxide does not change this initial oxidation significantly. However, when the XRD peaks at 500 h and 1000 h are compared for all three developed alloys, it should be noted that the Mn-Cr phase is present in the SUS430 þ CeO2 and SUS430 þ LaCrO3 samples, but not in the SUS430 þ Y2O3 samples. The evolution of the oxide scale with the oxidation time shows that the presence of nano-oxides within the matrix of ferritic stainless steel might have altered the growth kinetics of the oxide scale, resulting in different morphologies and phases at different times. It is clear that the samples with stable long-term ASR characteristics have Mn-Cr spinel layers on the top and that the spinel structures are present together with chromia throughout the scale. It has been reported in the literature that Mn-Cr spinel has very high conductivity, resulting in a smaller ASR increase during a 1000 h test [57]. However, for the case of SUS430þ3 wt% Y2O3, only Cr2O3 was present.
oxidation test. However, the addition of 3 wt% Y2O3 did not result in significant improvement in the conductivity or the ASR properties of the interconnect alloys. Furthermore, it was observed, as shown in Fig. 6(a), that the SUS430þ3 wt% Y2O3 sample showed an large increase in ASR during initial 300 h and, after that, the rate of increase in the ASR is slow as compared with the SUS430 sample. The oxide scale of the Y con taining Steel is known to consist of Cr2O3, Y2O3, YCrO3, and the for mation of conductive YCrO3 is a very slow process [62]. Thus the initially large ASR increase of the SUS430þ3 wt% Y2O3 sample is considered to be due to presence of low conductive Y2O3in the oxide scale, and then its subsequent oxidation forms the stable and relatively thin oxide scale probably with small amount of YCrO3. At the end of the 1000 h test, the ASR of the SUS430þ3 wt% Y2O3 was 57 mΩ cm2. The increase in the ASR values of the SUS430þ3 wt% Y2O3 during 1000 h is very high in comparison with other nano-oxide added alloys. Long-term ASR graphs for the developed alloys indicate that the addition of nano CeO2 and LaCrO3 oxides to SUS430 significantly improves the electrical properties of these materials for use as a SOFC interconnect. Fig. 6(b) shows the effect of temperature on the ASR of the oxide scale measured at 650, 700, 750, and 800 � C after 1000 h continuous oxidation. The ASR values of the oxide scale formed on the SUS430 during continuous oxidation decrease with increasing temperature, which is typical behavior of the Cr2O3 based oxides. As the conductivity in the metallic interconnect is dominated by the oxide layer, the higher ASR value for SUS430 and the highest slope of the ASR vs. 1000/T plot indicate that the oxide layer might have been overgrown in the SUS430, causing a reduction in the conductivity. The minimum dependence of the ASR on the temperature was observed in the SUS430þ3 wt% CeO2 samples. This means that the dominant oxidation mechanism is changed due to formation of the stable thin oxide scale including Mn-Cr spinel layers. The effect of the nano-CeO2 amount on the long-term ASR properties of the SUS430 alloys was also studied. Samples were prepared with SUS430 þ xwt.% CeO2 (x ¼ 0.5, 1, 3, and 5) and subjected to a long-term
3.3. ASR behavior of the nano-oxide dispersed SUS430 alloys Fig. 6(a) shows the change in the ASR values of various samples during the 1000 h continuous test. The SUS430 sample shows a continuous increase in the ASR value. The measured ASR values depend on the oxide scale because the electrical conductivity of the metal sub strate is very high in several order. The initial ASR value at 800 � C for SUS430 was 21 mΩ cm2. However, after 1000 h test at 800 � C, the ASR value reached 89 mΩ cm2. Similarly, when we compared the behavior of the nano-oxide dispersed alloy during the long-term ASR test, it was found that the addition of 3 wt% nano-CeO2 and -LaCrO3 resulted in substantially improved ASR characteristics during the long-term 7
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Fig. 6. (a) Long-term ASR values of the SUS430 samples with 3 wt% nano CeO2, 3 wt% nano LaCrO3, and 3 wt% nano Y2O3, measured at 800 � C in air. (b) Effect of temperature on the ASR of the samples after 1000 h test. The SUS 430 is used as a reference sample. (c) Long-term ASR values of the SUS430 with 0.5, 1, 3, and 5 wt % nano-CeO2, (d) ASR behavior of the nano CeO2 dispersed alloys during 10 thermal cycles between 400 and 800 � C in air, (e) Change in cathodic overpotentials of the LSCF/GDC/Pt half-cells with different interconnects, measured at a constant current density of 50 mA cm 2.
ASR test under the same conditions as described in section 2.3. Fig. Fig. 6 (c) shows the trends of ASR of the specimen measured during the longterm test at 800 � C. On increasing the amount of nano-CeO2, the ASR property of the SUS430 improved significantly. A lower ASR value in the 5 wt% CeO2 addition is observed due to the formation of a thinner oxide layer. The CeO2 addition of smaller amounts such as 0.5 and 1 wt % also decreased sufficiently the ASR during the initial hours of operation. However, at longer times, the ASR was increased rapidly. This might be attributed to the small quantity of nano CeO2 particles, which failed to mitigate the transport of cations for longer time periods. Longer time tests in the low content of CeO2, therefore, produced the formation of more Cr2O3 and Fe3O4, resulting in a fast increase in the ASR values. Significant amounts, such as 3 wt % and 5 wt % of CeO2, appear to have promisingly stopped the cation transport and thus their ASR values maintained to be quite low for long time test. Fig. 6(d) shows that the nano-CeO2 dispersed SUS430 alloys show stable behavior during the
thermal cycles. The samples were exposed to 10 consecutive thermal cycles between 400 and 800 � C; the ASR values were measured at 800 � C after each thermal cycle. The ASR values were almost stable for all the nano-CeO2 added alloys, without any apparent delamination of the oxide layer. Fig. 6(e) shows the change in cathodic overpotential of the LSCF electrode measured in the presence of the developed interconnect alloys. For comparison, we used SUS430 and Crofer 22 APU interconnects as references. All the half cells showed an increase in the cathodic over potential with time, which is due to LSCF degradation in Cr environ ment. The half cells with SUS430þ3 wt.CeO2 and Crofer 22 APU showed a relatively low increase in the cathodic overpotential with time. It is well known that Cr species in gaseous phase are released from chromia scales on the interconnect surface and then are electrochemically deposited at the LSCF electrode. The Cr oxide deposits not only reduce the active sites for oxygen reduction, but also block the migration of 8
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oxygen ions from the electrode to the electrolyte, which results in the performance degradation [39]. As compared to SUS430, the cell with SUS430 þ 3 wt% CeO2 interconnect showed lower change in cathodic overpotential. This indicates that the stable oxide scale formation in the nano CeO2-dispersed SUS430 retards a vaporization of gaseous Cr spe cies from the interconnect, and thus amount of the Cr oxide deposit in the LSCF cathode decreases. Further detailed investigations are required to make it clear.
segregation model might be applicable. The addition of dispersed nano-CeO2 has demonstrated superior properties regarding the oxida tion resistance of newly developed alloys; this might be due to the better oxygen affinity of CeO2, which can influence the cation and anion diffusion across the oxide scale [60]. Moreover, CeO2 has a different oxidation state from that of Cr2O3, and it improves intermetallic contact and remains at the grain boundary. This means that even during long-term oxidation, CeO2 nanoparticles might be present at the grain boundary, thereby hindering the diffusion of Cr ions outward. Form Fig. 2(c and d), we can see that the CeO2 nanoparticles are present at the grain boundaries in the form of 1–2 micron-sized agglomerates. The presence of the agglomerated CeO2 particles at the grain boundary of the alloy might have improved the oxidation properties of the CeO2 added SUS430 alloy due to retarding the Cr ion diffusion outward. The detailed investigation of the oxide scale formation indicates that, due to the addition of the nano-oxides in the alloy, the oxidation kinetics might have been altered, as depicted by schematics in Fig. 7. Upon startup of oxidation, the favorable oxidation of chromium near the free surface of the alloy can be enhanced by rare earth oxide ceria, and an integrated and densely packed Cr2O3 oxide layer can then be obtained. Ce shows good affinity towards oxygen and the cerium oxides accelerate the formation of protective oxidation layers by providing the nucleation sites that rapidly form Cr2O3 oxides layers [44]. Fig. S5 shows a sig nificant decrease in oxide layer thickness with an increment of Ce con tent. At the higher content of CeO2, the more incorporation of CeO2 into the oxide scale is clear to produce more thin stable oxide scale, and the more Ce at grain boundaries of the oxide scale prevents the growth of oxide scale due to retarded cation diffusion during long-term oxidation. Grain boundaries also affect the high-temperature oxidation resistance of metals. Upon decreasing the grain size, grain boundary area increases, which makes the diffusion of Cr better, leading to a decrease of the critical concentration of chromium to form a Cr2O3 layer [14]. Actually, more CeO2 addition produced finer grain size in the nano oxide-dispersed SUS 430. Although there is grain size effect in the oxidation kinetics, in the present work the effect of CeO2 content was greater. Nevertheless, it is still needed to find out the precise manner of interaction of these materials with cations and anions across the oxide layer. We are studying this further and hoping to report following this work.
3.4. Oxidation mechanism of the nano-oxide dispersed SUS 430s Hou and Stinger [40] reported that, for Cr2O3-forming alloys, the presence of a reactive element [RE] or its oxide can reduce the growth rate of the chromia scale and improve the adhesion of the scale to the substrate. The grain boundary segregation model has proven to be one of the most credible theories describing the proposed mechanisms for RE effect [43,44]. The growth of chromia scale with outward diffusion of Cr cations is carried out below 1000 � C temperature [58]. From the experimentally monitored rates of oxidation, chromium cation diffusion in Cr2O3 bulk form is found to be quite slow [59]. Therefore, it is considered that Cr cation diffusion during oxidation takes place through ‘short circuits’, i.e. at the points of grain boundaries and dislocations [60]. In case of the SUS 430 with the RE, the grain boundary segregation model explains that the reactive elements create hindrance at the grain boundaries and thus retard cations to diffuse through the grain bound aries. Yurek et al. explained that there is significant effect on oxidation of Ni–Cr alloy at 1000 � C when Y2O3 is added. They searched for the oxide scale, which mainly consists of NiO layer and chromia layer. Ev idence for the presence of Y was found at the grain boundaries of the NiO and chromia layer upon performing the STEM and EDS [61]. Czerwinski and Szpunar [59] reported that when Ni and Cr based superalloys are coated by the CeO2 prepared by sol-gel method, changes in oxidation are observed. The oxide-gas interface region was found to be rich in Ce for the oxide layer of Ni formed at 700 � C. They also found the presence of Ce at the NiO grain boundaries by EDS. In the present study, it has been shown that upon adding the nano particles of RE oxides, oxide scale thickness get reduced and adhesion of the oxide scale becomes better. Whittle and Stringer [43] argued that the higher number of RE oxide particles present at the alloy surface assist the initial oxidation of Cr at the start of oxidation and a homogeneous oxide layer is formed. In the later stages of oxidation, the grain boundary
Fig. 7. Schematics of the mechanism of formation of oxide scale on nano-oxide dispersed SUS430 alloy. 9
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4. Conclusions [8]
The addition of nano-oxides of reactive elements in SUS430 has significant effects on the oxidation and electrical characteristics of ferritic stainless steel during long-term test for SOFC interconnect application. We conducted a detailed investigation of the oxidation and electrical properties of different SUS430 based alloys modified with various nano-oxides. The sintered samples shows a uniform distribution of the nano-oxide throughout the pellet. XRD data indicate that the added oxide particles are quite stable and do not react with the base steel. Oxidation properties of the newly prepared alloys were deter mined by long-term ASR test at 800 � C for 1000 h. It was observed that higher amount of nano CeO2 greatly improves the performance of SUS430 by lowering the ASR value to as low as 20 mΩ cm2, which is comparable with the commercially used Crofer 22 APU. In the ASR test, the maximum increases in the ASR due to oxide scale growth were just 15, 29, and 56 mΩcm2 for CeO2, LaCrO3, and Y2O3 additions, respec tively. However, the ASR of SUS430 without any nano-oxide reached 100 mΩcm2. It was also found that for the case of the nano-CeO2 dispersed alloy, the top layer of the oxide scale contains more Mn-Cr rich spinel compared to the other developed alloys. The addition of nanooxides was found to improve the oxide scale adherence with the sub strate, forming a thin dense oxide layer. Therefore, the lower increase in the ASR for the dispersed nano-CeO2 alloys was attributed to formation of conductive Mn-Cr spinel layer, Ce incorporation into the oxide scale and retarding the cation diffusion at the grain boundary, which produces the lower thickness (~0.4 μm) and the improved adhesion of oxide scale to the substrate. Finally, it is concluded that the nano-oxide dispersed alloys can be successfully employed as SOFC interconnects because of their excellent chemical, electrical, and oxidation properties.
[9] [10] [11] [12]
[13]
[14] [15]
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[17] [18]
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Acknowledgments
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This work was supported by the Technology Innovation Program, (No. 10080607) funded by the Ministry of Trade, Industry and Energy, Korea, and also supported by a National Research Foundation (NRF) grant from the Ministry of Science and ICT, Korea under Technology Development Program to Solve Climate Change (No. NRF2017M1A2A2044926).
[22]
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Appendix A. Supplementary data [24]
Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227109.
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