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Influence of pre-reduction on microstructure homogeneity and electrical properties of APS Mn1.5Co1.5O4 coatings for SOFC interconnects
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Influence of pre-reduction on microstructure homogeneity and electrical properties of APS Mn1.5Co1.5O4 coatings for SOFC interconnects Ying-Zhen Hu, Shu-Wei Yao, Cheng-Xin Li*, Chang-Jiu Li, Shan-Lin Zhang State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi'an Jiaotong University, Xi'an, Shaanxi, 710049, China
article info
abstract
Article history:
Atmospheric plasma spray is employed to deposit Mn1.5Co1.5O4 (MCO) coatings onto T441
Received 24 May 2017
and porous 430L substrates respectively at room temperature and 400 C. The coatings are
Received in revised form
then subjected to redox treatments to achieve a desirable gas tightness. The gas perme-
29 August 2017
ability, area specific resistance (ASR) and microstructure homogenization of MCO coatings
Accepted 11 September 2017
are evaluated in different stages. The gas permeability of MCO coating deposited at 400 C
Available online 9 October 2017
and exposed to redox treatments can be as low as ~8.2 108 cm4 s1 gf1, which is essential for sufficient protection of metallic substrate against oxidation. The MCO-coated
Keywords:
T441 without pre-reduction presents a nonuniform microstructure characterized by iso-
SOFC
lated precipitates rich in Co and dark matrix rich in Mn. In comparison, the MCO coating
Interconnect
with 5 h pre-reduction presents a uniform and dense spinel structure. The ASR tests in-
Plasma spray
dicates that a homogeneous structure has contributed to an increase in the overall con-
Substrate preheating
ductivity. While bare T441 shows a substantial increase in interfacial ASR during the 200 h
Pre-reduction
test, MCO-coated T441 exhibites a high stability and a much lower ASR value of 13 mU cm2.
Mn1.5Co1.5O4 spinel coating
Besides, the dense MCO coatings effectively inhibit the oxide scale growth and block outward diffusion of Cr during the oxidation tests. Cr-rich oxide scale formed at the MCO/T441 interface is less than 1.2 mm and no further migration of Cr was detected. © 2017 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Solid oxide fuel cells (SOFCs) represent one of most efficient and versatile systems for conversion of chemical fuels to electricity, especially for stationary and distributed power generation [1e6]. As a promising technique, SOFCs have the potential to directly utilize a wide variety of fuels, from hydrogen to natural gas, coal gas, and gasified renewable biofuels [3,4,7e9]. To accumulate the voltage output, multiple
single fuel cells are usually stacked in electrical series via interconnects. Interconnect layer is a key component for construction of SOFC stacks in the form of planar and tubular cell configuration [10e14]; it electrically connects the anode of one cell to the cathode of another while physically separates the fuel from the oxidant gas. In order to perform the intended functions, interconnect should possess the following properties: (a) high electrical conductivity and thermal conductivity, (b) excellent chemical and electrochemical stability in both oxidizing and reducing environments, (c) matched coefficient
* Corresponding author. E-mail address: [email protected] (C.-X. Li). https://doi.org/10.1016/j.ijhydene.2017.09.073 0360-3199/© 2017 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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of thermal expansion (CTE) with other cell components, (d) excellent oxidation and carburization resistances and (e) low production cost [15]. The progresses towards thinner electrolyte membranes [16,17] and the advances in the development of more active and efficient electrode materials and architectures [4,7e9] have successfully reduced the SOFC operating temperature to about 600e800 C, making it possible for some high-temperature oxidation-resistant alloys (e.g., nickelbased and iron-based alloys) to substitute the traditional LaCrO3-based ceramics [18,19]. Ni-Cr-Fe alloy such as Haynes 230 exhibits a relatively high oxidation resistance [20]. However, the high CTE of 15.2 106 K1 [21] results in a low thermal matching with other components of SOFCs (10e13 106 K1) and could lead to the formation of cracks at the interconnect/electrode interface during thermal cycling. In recent years, Cr containing ferritic stainless steel (FSS) such as T441, Crofer22 APU, E-brite and AISI430 has attracted interest of most researchers [22e25] in terms of excellent electrical and thermal conductivity, comparable CTE (~12 106 K1) with other cell components, high ductility and low production cost. However, under the cathode operating environment at relatively high temperature such as 800 C, Cr-rich oxide scales formed at the alloy surface would grow linearly in term of square root of the oxidization time [26], along with the consumption of Mn and Cr in the alloy matrix. The continuous scale growth may lead not only to an increase in the electrical resistance but also to the possibility of coating spallation, which would be further exacerbated by the mismatch in the CTEs of the oxide layer and that of the substrate, since a high compressive stress would develop in the oxide scale during the cooling process due to the CTEs mismatch (9.6 106 K1 for Cr2O3 [27], 7.47 106 K1 for MnCr2O4 [28], 12.3e12.5 106 K1 for T441 [29]). Once the scale reached a critical thickness (eg. 11.2 mm for bare Crofer 22 APU under isothermal cooling [26]) at which interfacial fracture or delamination occurs, the fresh surface would be exposed constantly and furtherly oxidized. Moreover, the transport of highly volatile CrO2(OH)2 originating from the Cr2O3 layer could bring about degradation of the cathode [30,31]. It is necessary to apply a dense and electrically conductive protective coating on the ferritic stainless steel ICs to decrease the oxide scale growth, improving the scale/ICs interfacial bonding and electrical conductivity. Numerous materials have been explored as potential coatings for metallic ICs, such as reactive element oxide (REO) [32], ABO3 perovskite-type oxide [33,34] and AB2O4 spinel-type oxide [35,36]. Through adding small amount of reactive elements (e.g., Y or Ce), REO coatings have shown effectiveness in reducing the oxidation rate and area specific resistance over 500 h oxidation at 750 C [32]. However, due to the rather thin thickness and porous structure, REO coatings are regarded to be insufficient in preventing outward diffusion of Cr to the scale surface [37]. Electronically conductive ABO3 perovskites with appropriate doping, such as La0.8Sr0.2CrO3 [33] and La0.8Sr0.2FeO3 [34], have compatible CTE and sufficient structure stability even in low oxygen partial pressures. Although the perovskite coatings have exhibited certain effect at improving surface stability and electrical performance of the metallic substrates, it is doubtful about the effectiveness in Cr
inhibiting due to their ionically conducting nature. Besides, the difficulty in production of dense perovskite coating also limits its application. In comparison, cubic AB2O4 spinel coatings exhibit better results [38,39] in inhibiting the Cr-rich oxide scale growth. Among variety of binary spinels, Mn-Co system has evoked great interest due to its excellent electrical conductivity, matched CTE, low oxide ion conductivity, high chemical stability and fine sintering property [40]. Several fabrication methods to date have been exploited to deposit MnxCo3xO4 protection coatings, such as screen printing [38], slurry spraying [41], RF-sputtering [42], plasma spraying [43e46], electrodeposition [47e49], and aerosol spraying [29]. As an alternative deposition method, thermal spraying has the advantages of high flexibility, high deposition rate, and easy automation, which makes it particularly suitable for deposition of sophisticated functional layers with thickness ranging from tens to hundreds of micrometers in SOFCs. Currently, some researches have shown the availability of conventional thermal spray techniques such as atmospheric plasma spraying (APS) for the deposition of anode, electrolyte [50] and electrically insulating coatings in SOFC stacks [12]. However, conventional plasma-sprayed coatings typically present a lamellar structure characterized with limited interface bonding and numerous micro-cracks, making it necessary to perform substrate preheating or posttreatment for gas-tight protective coating in SOFC operation. Based on several studies [51e53], the bonding ratio between the splats and subsequent performance were significantly increased by elevating the depositing temperature. On the other hand, the effectiveness of post heat-treatment in densification of plasma-sprayed MnCo2O4 coating was also explored by E. Saoutieff et al. [54] to show that the processed coating exhibited dense structure and stable spinel phase in comparison to the as-sprayed coating. At the same time, it can restrain Cr diffusion to some extent. In addition to dense microstructure for inhibiting oxide scale growth, the other demand for protective coating is good adhesion on metallic substrates, which is particularly important for long-term durability. Although previous works [29,41] have showed the usefulness of short-term reduction treatment in densifying of spinel coatings and in the enhancement of coating/substrate bonding strength, studies concerning the effects of prereduction on microstructure homogenization and electrical property of atmospheric plasma sprayed Mn1.5Co1.5O4 spinel coating is rather limited. Hence, the primary goal of our study was to illuminate the homogenization process of thermalsprayed Mn1.5Co1.5O4 coatings with different post-spray treatments, and in the meantime, to evaluate the effect of coating microstructure on the electrical properties. In the present work, Mn1.5Co1.5O4 (MCO) coatings were prepared by APS process, companied by substrate preheating and reduction treatment aiming at improving the interface bonding, reducing the micro-cracks, and eventually achieving gas-tight and homogeneous spinel coatings. First, the gas permeability of MCO coatings in different preparation stages was measured to intuitively evaluate the density of MCO coatings. Second, area specific resistances (ASR) of MCOcoated T441 prepared under different conditions were tested at 800 C in air to evaluate the influence of microstructure homogenization on electrical property of MCO coatings. The
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interfacial ASRs between La0.8Sr0.2MnO3 (LSM) and T441 with and without MCO coating were also studied to evaluate the stability of MCO/T441 system for IT-SOFC interconnect applications. Third, the microstructure evolution mechanism of MCO coatings during post-treatments was proposed to illuminate the effect of pre-reduction on microstructure homogenization of MCO coatings.
Experimental Materials Two different substrate materials were used in this study: porous 430L was chosen as supporting substrate in order to characterize the gas-permeability of thermally sprayed MCO coatings, and T441 bulk substrate was used for area specific resistance (ASR) test and oxidation resistance evaluation. T441 used herein demonstrates a CTE of 12.3e12.5 106 K1 from 600 to 800 C [29] and has a nominal chemical composition (wt %) of 17.5 Cr, 0.44 Nb, 0.42 Si, 0.24 Mn, 0.19 Ti, 0.011 C, 0.017 P, 0.001 S and balance Fe. The commercially acquired (Mn,Co)3O4 agglomerated powder with a nominal composition of Mn1.5Co1.5O4 was used as starting material. In prior to coating preparation, the powders were sintered in air atmosphere, with the intention of reducing porosity and improving the melting degree. Then the sintered powder underwent a sieving process, and particles with an average size of 30.7 mm measured by laser diffract instrument (Beckman-Coulter, LS 230) were selected for APS process.
Coating preparation Firstly, MCO coatings of ~100 mm were deposited on 430L and T441 substrates with and without preheating. Secondly, redox post-spray treatments consist of reduction and subsequent oxidation were introduced to evaluate the effect of posttreatments on homogenization process of MCO coatings. On the other hand, LSM coating was deposited onto bare T441 and MCO-coated T441, respectively, to simulate the cathode/ interconnect structure in SOFC stack. All the aforementioned coatings were prepared by a commercial plasma spray system (GP-80, 80 kW class, China), the detailed plasma spray parameters and preparation conditions are listed in Tables 1 and 2, respectively.
Gas permeability measurement The gas permeability of MCO coatings deposited on porous 430L with and without substrate preheating was measured. For coatings under different preparation conditions, three coatings in each test group were randomly chosen and tested to improve the measurement accuracy. The effective test area
Table 2 e Sample preparation conditions. Samples
Substrate
Deposition temperature
RT-1 RT-2
430L 430L
RT RT
400e1 400e2
430L 430L
400 C 400 C
T-MCO-1 T-MCO-2 T-MCO-3
T441 T441 T441
RT 400 C 400 C
Treatment conditions as-sprayed with redox post-spray treatments as-sprayed with redox post-spray treatments with pre-reduction as-sprayed with pre-reduction
Note: “with redox post-spray treatments” indicates heattreatments in 14%H2/Ar at 800 C for 5 h and then in air at 800 C for 10 h after coating deposition; “with pre-reduction” indicates heat-treatment in 14%H2/Ar at 800 C for 5 h before ASR test.
for each coating is 1 cm2 and the edges of the specimen was completely sealed with two-component acrylate adhesive to prevent the passage of gas around the edges. The testing apparatus and detailed calculations were reported elsewhere [55]. In essence, it is based on the changes of differential pressure (△P) due to the gas leakage across the coatings. A vacuum pump was employed to reduce the pressure on the coating side and produce a pressure difference across the coating layer. When the maximum vacuum state was achieved, the pump was switched off and the pressure difference decreased due to the leakage of air through the coating. Then the gas permeability can be estimated based on the variation of gas pressure difference versus time. The average value and standard deviation are given at the end. Worth to mention is that the value of △P is affected by two factors: the coating compactness and the leak tightness of the test fixture. While for the latter, it was examined in advance by using a T441 block specimen which exhibited a rather low gas leakage value of 4.96 1010 cm4 s1 gf1, indicating that the possible gas permeability of the test fixture is lower than 1010 order. Thus, the measuring error introduced by the external factor can be neglected in present measurement.
Area specific resistance test T441 specimens with MCO coatings on both sides (T-MCO-1, T-MCO-2 and T-MCO-3) were oxidized in air for 200 h at 800 C, the ASRs were measured in-situ using a standard 4-probe technique. Besides, the interfacial ASRs between LSM and T441 with and without MCO coating were also studied to evaluate the effectiveness of spinel coating. The MCO coating herein mentioned was deposited at 400 C and exposed to redox post-spray treatments (as noted in Table 2) before interfacial ASR test. A Solartron 1260/1287 electrochemical system was employed to characterize the electrical
Table 1 e Spray processing parameters. Coatings MCO LSM
Arc current (A)
Arc power (kW)
Ar (slpm)
H2 (slpm)
N2 (slpm)
Spray distance (mm)
Traverse speed (mm s1)
600 600
36 30
50 50
2.5 2.5
2.5 8.5
80 80
600 1200
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Fig. 1 e Apparatus for ASR measurement.
performance of the abovementioned MCO/T441/MCO and LSM/T441/LSM test units. The apparatus for ASR measurement is schematically shown in Fig. 1. Platinum layer in diameter of F8 mm was firstly sputtered onto the coatings at pressure of 7 Pa and current of 20 mA for 30 min, in order to improve the contact between the samples and the current collector. Pt paste used as the current collector was then applied onto both sides of the coated specimen. Silver wires were attached to Pt layer by spot welding.
Fig. 2 e Polished cross-sectional and surface views of (a) original Mn1.5Co1.5O4 powders, and after sintering in air at (b) 1000 C, (c)-(d) 1100 C for 2 h, (e)-(f) 1100 C for 10 h, (g)-(h) 1200 C for 10 h respectively.
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Fig. 3 e XRD patterns for Mn1.5Co1.5O4 powders: (a) original, and after sintering in air at (b) 1000 C, (c) 1100 C, (d) 1200 C for 10 h, respectively.
Fig. 4 e Cross-sectional views of (a) porous 430L substrate and (b) as-sprayed MCO coating deposited at 400 C: (c) surface layer, (d) with good splat bonding.
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based on EDS spot scan, the relative atom percent of Mn and Co is close to 50/50 for the original powder, and powders sintered at 1000 C, 1100 C, 1200 C. Accordingly, the XRD analysis shown in Fig. 3 indicates dual phases of cubic MnCo2O4 (PDF#23-1237) and tetragonal CoMn2O4 (PDF#18-0408) spinels for the original and sintered powders. It is worth mentioning that the intensities of reflections of the tetragonal phase are relatively weak for 1200 C sintered powder as shown in Fig. 3(d), which indicates relatively lower phase content. After comprehensive consideration, the densified Mn1.5Co1.5O4 spinel powder sintered at 1100 C for 10 h is chosen as the optimum feedstock to achieve a fully-molten state during the subsequent spraying process.
Gas permeability of MCO coatings Fig. 5 e The gas leakage rate of MCO coatings under different preparation conditions.
Fig. 6 e Area specific resistance (ASR) of MCO coatings prepared under different conditions.
Results and discussion Powders Fig. 2(a) indicates that the original agglomerated powder presents a porous structure with nanometer or sub-micrometer scale primary particles, which is liable to cause gas entrapment and leads to the formation of pores inside the sprayed coating. After sintering at 1000 C for 10 h, the initial fine particles have grown up to micrometer particles characterized with grown sintering necks and through-holes as shown in Fig. 2(b). When the sintering temperature raised to 1100 C, the seal and spheroidization of pores have be en clearly observed from the polished section presented in Fig. 2(c). The surface view of spherical particle in Fig. 2(d) demonstrates closely arranged fine spinel crystals. By prolonging the duration of sintering to 10 h, the porosity was further decreased, accompanied by the merge and growth of grains shown in Fig. 2(f). At 1200 C, the powder presents dense structure and coarse grains (see Fig. 2(g) and (h)). According to statistical analysis
To effectively inhibit the O-diffusion and improve the oxidationresistance of the steel interconnect during IT-SOFC operation process, the MCO protective coating must have low porosity and good splat bonding. But above all, penetrating cracks as the primary path for gas transport must be avoided. The thickness of coatings prepared in this work was set as 100 mm. According to the cross-sectional view shown in Fig. 4(b) and (d), no penetrating crack is detected, instead, only few randomly located micro-cracks is observed. During plasma spraying, fast cooling of the splats from a high temperature near the melting point to a relatively low temperature inevitably results in a quenching stress. Once the cumulative stress exceeds the tensile stress of the deposited material, it can only be released in the form of cracking. In general, the thermally sprayed Mn1.5Co1.5O4 coating deposited at 400 C shows a high density with a relatively uniform thickness and good splat bonding which is attributed to elevated deposition temperature [53]. To evaluate the influence of substrate preheating and redox post-spray treatments on the density of MCO coatings, a comparison testing of the gas permeability which can directly reflect the porosity of through pores was carried out. Fig. 4(a) exhibits the cross-sectional microstructure of 430L supporting substrate with an open porosity of ~17.7% which was measured by image analysis software Image J. According to the gas permeability test, a rather high gas permeability value of 1.07 102 cm4 s1 gf1 was obtained, manifesting that the porous substrate used herein could provide enough communicating pores for gas transport. The gas permeability values for MCO coatings under different preparation conditions are then calculated and the standard deviation is also presented. As indicated in Fig. 5, the as-sprayed MCO coating without substrate preheating presents a high gas permeability of 1.21 106 cm4 s1 gf1, which is slightly decreased after subsequent redox treatment. In comparison, for coating deposited at 400 C, the gas permeability is only 1.23 107 cm4 s1 gf1, and a lower value of 8.2 108 cm4 s1 gf1 is achieved after 5 h reduction and 10 h oxidation, implying that a sufficiently dense coating was produced.
Influence of microstructure homogeneity on electrical performance of MCO coatings Fig. 6 shows the ASR of T441 coated with MCO coatings prepared under different treatment conditions. The ASRs were
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Fig. 7 e Sectional morphology of (a)-(b) T-MCO-1, (c)-(d) T-MCO-2, (e)-(f) T-MCO-3 samples after ASR test in air at 800 C for 200 h.
measured at 800 C in air as a function of time, under a current density of 200 mA cm2 flowing through the specimen. The ASR values exhibit a significant decline in the initial stage. This might be related to the increasing contact area and densifying of spinel coatings, since the ASR testing process was actually equivalent to a separate post-spray treatment under the oxidizing atmosphere. Besides, T-MCO-1 and TMCO-3 with pre-reduction show higher initial resistances, which is linked to the rapid generation of Cr-rich scale in the initial oxidizing process. Nevertheless, of the three coatings, T-MCO-3 exhibits the optimum performance and shows a final ASR value of 12 mU cm2. Once the ASR tests were completed, the samples were epoxymounted and polished for microstructure analysis. As can be seen, the MCO samples shown in Fig. 7 present typical sintered structures. For ASR tested T-MCO-1 (see Fig. 7(a) and (b)), a few interior cracks and numerous closed-pores are clearly observed, despite good interface bonding after redox treatments.
MCO coating without pre-reduction For T-MCO-2 sample without reduction pre-treatment, the tested coating presents distinct two-layer structure with a dense top layer and a nonuniform inner layer which is characterized by light-coloured precipitates and dark matrix, see Fig. 7(c) and (d). In addition, it can be seen that numerous pores are concentrated in the inner layer, especially near the coating/substrate interface. In order to reveal the phase structure and elementary composition of sample T-MCO-2, the top layer with a thickness of 50 mm was carefully removed, then energy spectrum and x-ray diffraction analysis were respectively performed on the polished nonuniform section. As indicated by XRD results shown in Fig. 8(a), MCO coating in the as-deposited state is comprised of tetragonal and cubic spinel phases, in addition to metal-oxide mixed phase MO (where M indicates Mn or Co) which was caused by partially spinel decomposition during thermal spraying [56,57]. According to the phase analysis
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Fig. 8 e Crystallographic phase analysis on MCO coatings (a) as-deposited, and after 200 h ASR test: (b) nonuniform region and (c) surface region of T-MCO-2 coating; (d) T-MCO-3 coating.
shown in Fig. 8(b), the inner region of oxidized (ASR tested) TMCO-2 is composed of tetragonal CoMn2O4, cubic MnCo2O4 and metal-oxide phases which were also detected in the asdeposited coating. The increase in the intensity of spinel phases is inferred to be connected with the transformation of metastable MO into spinel during oxidation treatment between 400 C to 800 C [58]. For the top layer, most of the diffraction peaks is identified to be corresponding to MnCo2O4 and CoMn2O4, see Fig. 8(c). As a semiquantitative analytical method, the measurement error for EDS point scan is generally less than 1% when heavy element such as Mn, Co is tested. In our present investigation, the work distance is set as 15 mm and the EDS analyse indicates that the average ratio of Mn/Co is 50/50, 51/49 for the dense top layer and the whole inner layer, respectively. However, for the inner layer, map scanning results shown in Fig. 9 indicate that the light isolated precipitates marked as spot 1 is in rich of Co (corresponding to the cubic CoxOy or MnCo2O4 phase), and the dark matrix marked as spot 2 is in rich of Mn (corresponding to the tetragonal MnxOy or CoMn2O4 phase). Similar microstructure composed of isolated precipitates was observed in the 1150 C sintered Mn1.75Co1.25O4 bulk sample, where the isolated precipitates were identified as the cubic MnCo2O4 phase and the matrix as the tetragonal phase [59].
MCO coating with pre-reduction Unlike the two-layer structure, the oxidized T-MCO-3 shown in Fig. 7(e) and (f) presents a uniform and dense structure with fewer pores, cracks and good interface bonding which would contribute to an increased electrical conductivity [51]. The XRD result in Fig. 8(d) shows a composition corresponding to MnCo2O4 and CoMn2O4. According to EDS data, the Mn/Co ratio is 50/50, 51/49 and 51/49 respectively for the top, middle and inner layers, manifesting that selective evaporation of elements did not occur and a homogeneous composition was obtained for the whole coating. Besides, it is worthwhile mentioning that a relatively denser spinel layer is observed in surface area of MCO coating, which is attributed to the high sinterability of MCO after redox treatment at a temperature no higher than 800 C [29]. In this work, the sintering of metallic cobalt under the reducing environment was essentially necessary and beneficial for the subsequent homogenization and densifying process during the oxidizing treatment. According to Brylewski [59], the electrical conductivity of Mn3xCoxO4 spinels was significantly influenced by phase structures and the highest conductivity was obtained in range of 1.5 < x < 2. The relatively higher resistance of T-MCO-2 coating could be related to the metastable phase MO detected in the nonuniform layer. In contrast, T-MCO-3 comprised of
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Fig. 9 e SEM views of EDS map scan based on nonuniform region of ASR tested T-MCO-2 sample. tetragonal CoMn2O4 and remaining cubic MnCo2O4 presented a lower resistance. To simulate the cathode/interconnect structure in SOFC stacks, porous LSM cathode was deposited onto bare T441 and MCO-coated T441, respectively. Interfacial ASRs between LSM and T441 with and without MCO protection layer were then studied to evaluate the effectiveness of MCO coatings. Prior to the resistance measurements, MCO coating deposited at 400 C was processed with redox post-spray treatments as noted in Table 2. According to Fig. 10, LSM/T441 exhibits a substantial increase in interfacial ASR during the test (to a final ASR value of 42 mU cm2). The poor oxidation resistance performance and Crrich spinel scale with low electric conductivity in the oxidizing atmosphere have negative effects on the power output when bare T441 is used as interconnects. It was reported that the bare 441 showed a high ASR of 75 mU cm2 after long-term oxidation in air at 800 C [38]. In comparison, LSM/MCO-T441 shows a significant decline in the initial stage and exhibits a high stability during the continuous testing process. A much lower value of 13 mU cm2 was obtained at the end of ASR test. Fig. 11 presents the cross-sectional microstructure of LSM/ MCO-T441 after ASR test. As can be seen, LSM/MCO exhibits a good interface bonding which would contribute to a lower interface resistance. Moreover, the MCO protection coating also shows a good adhesion to T441 substrate, despite a few isolated micro-voids are visible at the interface region. The EDS analysis (see Fig. 12) indicates that the Cr-rich oxide scale formed at the MCO/T441 interface is less than 1.2 mm and no further migration of Cr is detected, demonstrating the
effectiveness of MCO protection coating in inhibiting subscale growth. Dense spinel microstructure achieved by APS and redox treatments makes up the most important factor to provide adequate protection against Cr diffussion and minimize the electrical resistance of T441 interconnects.
Mechanism of microstructure homogenization during posttreatments The microstructure homogenization process of MCO coating during post-treatments is schematically presented in Fig. 13.
Fig. 10 e Interfacial ASRs between a LSM cathode and T441 with and without the spinel protection layer.
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Fig. 11 e Sectional morphology of LSM/MCO-T441 after ASR test at 800 C in air for 200 h. The as-sprayed coating characterized with a few closed pores and randomly located micro-cracks is indicated as Fig. 13(a). As previously mentioned, partially spinel decomposition was detected during the thermal spraying. Therefore, mixed phases of spinel and metal-oxides with partially amorphous characters were formed. At the initial oxidation stage, the molecular oxygen rapidly diffused into the coating surface and brought to the re-formation of spinel structure, accompanied by the crack self-healing. Unlike the free surface, the formation rate of spinels in the inner coating was significantly influenced by the diffusion rate of molecular oxygen. With the re-formation of spinel structure in the top layer, the inner diffusion of O/O2 was retarded to some extent, leading to an oxygen deficiency and a nonuniform region within the coating. Thus, a distinct two-layer structure was presented, as indicated in Fig. 13(b). Nevertheless, the interface between the top spinel layer and the inner nonuniform layer was still
assumed to gradually evolve towards the substrate side in prolonged oxidation process, accompanied by the formation of Cr-rich scale. It is plausible to speculate that the homogenization process would proceed along with the re-formation of Mn-Co spinel. This assumption was also confirmed by the existence of metastable phase at the inner nonuniform region of T-MCO-2 after 200 h oxidation test. For MCO coating with pre-reduction, different microstructure evolution mechanism was presented. The introduction of reduction pre-treatment could enhance the interfacial adhesion on micro level and promote the closure of micro-cracks originally found in the as-sprayed coating. When the porous MnO-Co coating was subsequently exposed to oxidizing atmosphere, O2 rapidly diffused into the whole coating, accompanied by in-situ oxidizing reaction. Assuming that the molecular oxygen distributed evenly, the formation rate of primary spinel should be similar throughout the whole
Fig. 12 e EDS line scan based on line AB shown in Fig. 11(b).
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Fig. 13 e Schematic demonstrating the microstructure evolution mechanism of MCO/T441 under the following states: (a) assprayed, (b) at the initial oxidation, (c) in subsequent oxidation, respectively; and (d) MCO with pre-reduction, and then (e) at the initial oxidation, (f) in subsequent oxidation, respectively.
coating, which would promote the homogenization process of Mn-Co spinel. In the meantime, it was speculated that the inner diffusion of O/O2 advanced more rapidly than the in-situ oxidizing reaction, and the molecular oxygen reached the coating/steel interface within a relatively short time [38,60]. Thus, the oxygen activity at the coating/steel interface could be well above the critical value of Cr2O3(1028 atm), leading to the formation of a thin and compact Cr2O3 scale inevitably, as shown in Fig. 13(e). Nevertheless, the diffussion rate of O/O2 gradually decreased with the growth of spinel grains and the reducing of grain boundaries. A homogeneous and stable spinel was produced as the oxidation sintering process went on. Thermodynamically, the basic driving force for shrinkage of ceramic coating in sintering process is the decrease of surface energy. The surface energy minimization leads to the elimination or coalescence of small isolated pores and a sufficiently dense spinel coating which is impermeable to molecular oxygen can be eventually achieved, see Fig. 13(f).
Conclusions In this study, Mn1.5Co1.5O4 coatings with a relatively uniform thickness were deposited at room temperature and 400 C via atmospheric plasma spraying, then followed by redox postspray treatments. The gas tightness of the as-sprayed MCO coating deposited at 400 C was 1.23 107 cm4 s1 gf1, and that of the as-treated coatings can be as low as ~8.2 108 cm4 s1 gf1, which is essential for sufficient protection of metallic substrate against oxidation. Metal-oxide mixed phase MO (where M indicates Mn or Co) from partially spinel decomposition during thermal spraying was detected in the as-sprayed MCO coatings. For ASR tested coating without pre-reduction, a distinct two-layer structure composed of dense top layer and nonuniform inner layer was observed. The inner layer was characterized by isolated precipitates rich in Co and dark matrix rich in Mn. While for
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coating with pre-reduction, the Mn/Co ratio was closed to 50/ 50, manifesting that selective evaporation of elements did not occur and a uniform composition was obtained. ASR tests also confirmed that a homogeneous structure had contributed to an increase in the overall conductivity. Of the three coatings, T-MCO-3 with pre-reduction exhibited the optimum performance and a final ASR of 12 mU cm2 was obtained at the end of 200 h duration. Besides, LSM/MCO-T441 system exhibited a high stability during the 200 h test, and a final interfacial ASR of 13 mU cm2 was obtained. SEM/EDS analysis indicated that the Cr-rich oxide scale formed at the MCO/T441 interface was less than 1.2 mm and no further migration of Cr was detected, demonstrating the effectiveness of MCO protection coating in inhibiting the subscale growth and Cr diffussion. In comparision, the bare T441 exhibited a substantial increase in interfacial ASR during the test (to a final ASR value of 42 mU cm2). Once properly treated, APS-prepared Mn1.5Co1.5O4 coatings have desirable gas-tightness and sufficient high electrical conductivity as protective coatings for IT-SOFC interconnects.
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Available online at www.sciencedirect.com
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Influence of pre-reduction on microstructure homogeneity and electrical properties of APS Mn1.5Co1.5O4 coatings for SOFC interconnects Ying-Zhen Hu, Shu-Wei Yao, Cheng-Xin Li*, Chang-Jiu Li, Shan-Lin Zhang State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi'an Jiaotong University, Xi'an, Shaanxi, 710049, China
article info
abstract
Article history:
Atmospheric plasma spray is employed to deposit Mn1.5Co1.5O4 (MCO) coatings onto T441
Received 24 May 2017
and porous 430L substrates respectively at room temperature and 400 C. The coatings are
Received in revised form
then subjected to redox treatments to achieve a desirable gas tightness. The gas perme-
29 August 2017
ability, area specific resistance (ASR) and microstructure homogenization of MCO coatings
Accepted 11 September 2017
are evaluated in different stages. The gas permeability of MCO coating deposited at 400 C
Available online 9 October 2017
and exposed to redox treatments can be as low as ~8.2 108 cm4 s1 gf1, which is essential for sufficient protection of metallic substrate against oxidation. The MCO-coated
Keywords:
T441 without pre-reduction presents a nonuniform microstructure characterized by iso-
SOFC
lated precipitates rich in Co and dark matrix rich in Mn. In comparison, the MCO coating
Interconnect
with 5 h pre-reduction presents a uniform and dense spinel structure. The ASR tests in-
Plasma spray
dicates that a homogeneous structure has contributed to an increase in the overall con-
Substrate preheating
ductivity. While bare T441 shows a substantial increase in interfacial ASR during the 200 h
Pre-reduction
test, MCO-coated T441 exhibites a high stability and a much lower ASR value of 13 mU cm2.
Mn1.5Co1.5O4 spinel coating
Besides, the dense MCO coatings effectively inhibit the oxide scale growth and block outward diffusion of Cr during the oxidation tests. Cr-rich oxide scale formed at the MCO/T441 interface is less than 1.2 mm and no further migration of Cr was detected. © 2017 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Solid oxide fuel cells (SOFCs) represent one of most efficient and versatile systems for conversion of chemical fuels to electricity, especially for stationary and distributed power generation [1e6]. As a promising technique, SOFCs have the potential to directly utilize a wide variety of fuels, from hydrogen to natural gas, coal gas, and gasified renewable biofuels [3,4,7e9]. To accumulate the voltage output, multiple
single fuel cells are usually stacked in electrical series via interconnects. Interconnect layer is a key component for construction of SOFC stacks in the form of planar and tubular cell configuration [10e14]; it electrically connects the anode of one cell to the cathode of another while physically separates the fuel from the oxidant gas. In order to perform the intended functions, interconnect should possess the following properties: (a) high electrical conductivity and thermal conductivity, (b) excellent chemical and electrochemical stability in both oxidizing and reducing environments, (c) matched coefficient
* Corresponding author. E-mail address: [email protected] (C.-X. Li). https://doi.org/10.1016/j.ijhydene.2017.09.073 0360-3199/© 2017 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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of thermal expansion (CTE) with other cell components, (d) excellent oxidation and carburization resistances and (e) low production cost [15]. The progresses towards thinner electrolyte membranes [16,17] and the advances in the development of more active and efficient electrode materials and architectures [4,7e9] have successfully reduced the SOFC operating temperature to about 600e800 C, making it possible for some high-temperature oxidation-resistant alloys (e.g., nickelbased and iron-based alloys) to substitute the traditional LaCrO3-based ceramics [18,19]. Ni-Cr-Fe alloy such as Haynes 230 exhibits a relatively high oxidation resistance [20]. However, the high CTE of 15.2 106 K1 [21] results in a low thermal matching with other components of SOFCs (10e13 106 K1) and could lead to the formation of cracks at the interconnect/electrode interface during thermal cycling. In recent years, Cr containing ferritic stainless steel (FSS) such as T441, Crofer22 APU, E-brite and AISI430 has attracted interest of most researchers [22e25] in terms of excellent electrical and thermal conductivity, comparable CTE (~12 106 K1) with other cell components, high ductility and low production cost. However, under the cathode operating environment at relatively high temperature such as 800 C, Cr-rich oxide scales formed at the alloy surface would grow linearly in term of square root of the oxidization time [26], along with the consumption of Mn and Cr in the alloy matrix. The continuous scale growth may lead not only to an increase in the electrical resistance but also to the possibility of coating spallation, which would be further exacerbated by the mismatch in the CTEs of the oxide layer and that of the substrate, since a high compressive stress would develop in the oxide scale during the cooling process due to the CTEs mismatch (9.6 106 K1 for Cr2O3 [27], 7.47 106 K1 for MnCr2O4 [28], 12.3e12.5 106 K1 for T441 [29]). Once the scale reached a critical thickness (eg. 11.2 mm for bare Crofer 22 APU under isothermal cooling [26]) at which interfacial fracture or delamination occurs, the fresh surface would be exposed constantly and furtherly oxidized. Moreover, the transport of highly volatile CrO2(OH)2 originating from the Cr2O3 layer could bring about degradation of the cathode [30,31]. It is necessary to apply a dense and electrically conductive protective coating on the ferritic stainless steel ICs to decrease the oxide scale growth, improving the scale/ICs interfacial bonding and electrical conductivity. Numerous materials have been explored as potential coatings for metallic ICs, such as reactive element oxide (REO) [32], ABO3 perovskite-type oxide [33,34] and AB2O4 spinel-type oxide [35,36]. Through adding small amount of reactive elements (e.g., Y or Ce), REO coatings have shown effectiveness in reducing the oxidation rate and area specific resistance over 500 h oxidation at 750 C [32]. However, due to the rather thin thickness and porous structure, REO coatings are regarded to be insufficient in preventing outward diffusion of Cr to the scale surface [37]. Electronically conductive ABO3 perovskites with appropriate doping, such as La0.8Sr0.2CrO3 [33] and La0.8Sr0.2FeO3 [34], have compatible CTE and sufficient structure stability even in low oxygen partial pressures. Although the perovskite coatings have exhibited certain effect at improving surface stability and electrical performance of the metallic substrates, it is doubtful about the effectiveness in Cr
inhibiting due to their ionically conducting nature. Besides, the difficulty in production of dense perovskite coating also limits its application. In comparison, cubic AB2O4 spinel coatings exhibit better results [38,39] in inhibiting the Cr-rich oxide scale growth. Among variety of binary spinels, Mn-Co system has evoked great interest due to its excellent electrical conductivity, matched CTE, low oxide ion conductivity, high chemical stability and fine sintering property [40]. Several fabrication methods to date have been exploited to deposit MnxCo3xO4 protection coatings, such as screen printing [38], slurry spraying [41], RF-sputtering [42], plasma spraying [43e46], electrodeposition [47e49], and aerosol spraying [29]. As an alternative deposition method, thermal spraying has the advantages of high flexibility, high deposition rate, and easy automation, which makes it particularly suitable for deposition of sophisticated functional layers with thickness ranging from tens to hundreds of micrometers in SOFCs. Currently, some researches have shown the availability of conventional thermal spray techniques such as atmospheric plasma spraying (APS) for the deposition of anode, electrolyte [50] and electrically insulating coatings in SOFC stacks [12]. However, conventional plasma-sprayed coatings typically present a lamellar structure characterized with limited interface bonding and numerous micro-cracks, making it necessary to perform substrate preheating or posttreatment for gas-tight protective coating in SOFC operation. Based on several studies [51e53], the bonding ratio between the splats and subsequent performance were significantly increased by elevating the depositing temperature. On the other hand, the effectiveness of post heat-treatment in densification of plasma-sprayed MnCo2O4 coating was also explored by E. Saoutieff et al. [54] to show that the processed coating exhibited dense structure and stable spinel phase in comparison to the as-sprayed coating. At the same time, it can restrain Cr diffusion to some extent. In addition to dense microstructure for inhibiting oxide scale growth, the other demand for protective coating is good adhesion on metallic substrates, which is particularly important for long-term durability. Although previous works [29,41] have showed the usefulness of short-term reduction treatment in densifying of spinel coatings and in the enhancement of coating/substrate bonding strength, studies concerning the effects of prereduction on microstructure homogenization and electrical property of atmospheric plasma sprayed Mn1.5Co1.5O4 spinel coating is rather limited. Hence, the primary goal of our study was to illuminate the homogenization process of thermalsprayed Mn1.5Co1.5O4 coatings with different post-spray treatments, and in the meantime, to evaluate the effect of coating microstructure on the electrical properties. In the present work, Mn1.5Co1.5O4 (MCO) coatings were prepared by APS process, companied by substrate preheating and reduction treatment aiming at improving the interface bonding, reducing the micro-cracks, and eventually achieving gas-tight and homogeneous spinel coatings. First, the gas permeability of MCO coatings in different preparation stages was measured to intuitively evaluate the density of MCO coatings. Second, area specific resistances (ASR) of MCOcoated T441 prepared under different conditions were tested at 800 C in air to evaluate the influence of microstructure homogenization on electrical property of MCO coatings. The
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interfacial ASRs between La0.8Sr0.2MnO3 (LSM) and T441 with and without MCO coating were also studied to evaluate the stability of MCO/T441 system for IT-SOFC interconnect applications. Third, the microstructure evolution mechanism of MCO coatings during post-treatments was proposed to illuminate the effect of pre-reduction on microstructure homogenization of MCO coatings.
Experimental Materials Two different substrate materials were used in this study: porous 430L was chosen as supporting substrate in order to characterize the gas-permeability of thermally sprayed MCO coatings, and T441 bulk substrate was used for area specific resistance (ASR) test and oxidation resistance evaluation. T441 used herein demonstrates a CTE of 12.3e12.5 106 K1 from 600 to 800 C [29] and has a nominal chemical composition (wt %) of 17.5 Cr, 0.44 Nb, 0.42 Si, 0.24 Mn, 0.19 Ti, 0.011 C, 0.017 P, 0.001 S and balance Fe. The commercially acquired (Mn,Co)3O4 agglomerated powder with a nominal composition of Mn1.5Co1.5O4 was used as starting material. In prior to coating preparation, the powders were sintered in air atmosphere, with the intention of reducing porosity and improving the melting degree. Then the sintered powder underwent a sieving process, and particles with an average size of 30.7 mm measured by laser diffract instrument (Beckman-Coulter, LS 230) were selected for APS process.
Coating preparation Firstly, MCO coatings of ~100 mm were deposited on 430L and T441 substrates with and without preheating. Secondly, redox post-spray treatments consist of reduction and subsequent oxidation were introduced to evaluate the effect of posttreatments on homogenization process of MCO coatings. On the other hand, LSM coating was deposited onto bare T441 and MCO-coated T441, respectively, to simulate the cathode/ interconnect structure in SOFC stack. All the aforementioned coatings were prepared by a commercial plasma spray system (GP-80, 80 kW class, China), the detailed plasma spray parameters and preparation conditions are listed in Tables 1 and 2, respectively.
Gas permeability measurement The gas permeability of MCO coatings deposited on porous 430L with and without substrate preheating was measured. For coatings under different preparation conditions, three coatings in each test group were randomly chosen and tested to improve the measurement accuracy. The effective test area
Table 2 e Sample preparation conditions. Samples
Substrate
Deposition temperature
RT-1 RT-2
430L 430L
RT RT
400e1 400e2
430L 430L
400 C 400 C
T-MCO-1 T-MCO-2 T-MCO-3
T441 T441 T441
RT 400 C 400 C
Treatment conditions as-sprayed with redox post-spray treatments as-sprayed with redox post-spray treatments with pre-reduction as-sprayed with pre-reduction
Note: “with redox post-spray treatments” indicates heattreatments in 14%H2/Ar at 800 C for 5 h and then in air at 800 C for 10 h after coating deposition; “with pre-reduction” indicates heat-treatment in 14%H2/Ar at 800 C for 5 h before ASR test.
for each coating is 1 cm2 and the edges of the specimen was completely sealed with two-component acrylate adhesive to prevent the passage of gas around the edges. The testing apparatus and detailed calculations were reported elsewhere [55]. In essence, it is based on the changes of differential pressure (△P) due to the gas leakage across the coatings. A vacuum pump was employed to reduce the pressure on the coating side and produce a pressure difference across the coating layer. When the maximum vacuum state was achieved, the pump was switched off and the pressure difference decreased due to the leakage of air through the coating. Then the gas permeability can be estimated based on the variation of gas pressure difference versus time. The average value and standard deviation are given at the end. Worth to mention is that the value of △P is affected by two factors: the coating compactness and the leak tightness of the test fixture. While for the latter, it was examined in advance by using a T441 block specimen which exhibited a rather low gas leakage value of 4.96 1010 cm4 s1 gf1, indicating that the possible gas permeability of the test fixture is lower than 1010 order. Thus, the measuring error introduced by the external factor can be neglected in present measurement.
Area specific resistance test T441 specimens with MCO coatings on both sides (T-MCO-1, T-MCO-2 and T-MCO-3) were oxidized in air for 200 h at 800 C, the ASRs were measured in-situ using a standard 4-probe technique. Besides, the interfacial ASRs between LSM and T441 with and without MCO coating were also studied to evaluate the effectiveness of spinel coating. The MCO coating herein mentioned was deposited at 400 C and exposed to redox post-spray treatments (as noted in Table 2) before interfacial ASR test. A Solartron 1260/1287 electrochemical system was employed to characterize the electrical
Table 1 e Spray processing parameters. Coatings MCO LSM
Arc current (A)
Arc power (kW)
Ar (slpm)
H2 (slpm)
N2 (slpm)
Spray distance (mm)
Traverse speed (mm s1)
600 600
36 30
50 50
2.5 2.5
2.5 8.5
80 80
600 1200
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Fig. 1 e Apparatus for ASR measurement.
performance of the abovementioned MCO/T441/MCO and LSM/T441/LSM test units. The apparatus for ASR measurement is schematically shown in Fig. 1. Platinum layer in diameter of F8 mm was firstly sputtered onto the coatings at pressure of 7 Pa and current of 20 mA for 30 min, in order to improve the contact between the samples and the current collector. Pt paste used as the current collector was then applied onto both sides of the coated specimen. Silver wires were attached to Pt layer by spot welding.
Fig. 2 e Polished cross-sectional and surface views of (a) original Mn1.5Co1.5O4 powders, and after sintering in air at (b) 1000 C, (c)-(d) 1100 C for 2 h, (e)-(f) 1100 C for 10 h, (g)-(h) 1200 C for 10 h respectively.
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Fig. 3 e XRD patterns for Mn1.5Co1.5O4 powders: (a) original, and after sintering in air at (b) 1000 C, (c) 1100 C, (d) 1200 C for 10 h, respectively.
Fig. 4 e Cross-sectional views of (a) porous 430L substrate and (b) as-sprayed MCO coating deposited at 400 C: (c) surface layer, (d) with good splat bonding.
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based on EDS spot scan, the relative atom percent of Mn and Co is close to 50/50 for the original powder, and powders sintered at 1000 C, 1100 C, 1200 C. Accordingly, the XRD analysis shown in Fig. 3 indicates dual phases of cubic MnCo2O4 (PDF#23-1237) and tetragonal CoMn2O4 (PDF#18-0408) spinels for the original and sintered powders. It is worth mentioning that the intensities of reflections of the tetragonal phase are relatively weak for 1200 C sintered powder as shown in Fig. 3(d), which indicates relatively lower phase content. After comprehensive consideration, the densified Mn1.5Co1.5O4 spinel powder sintered at 1100 C for 10 h is chosen as the optimum feedstock to achieve a fully-molten state during the subsequent spraying process.
Gas permeability of MCO coatings Fig. 5 e The gas leakage rate of MCO coatings under different preparation conditions.
Fig. 6 e Area specific resistance (ASR) of MCO coatings prepared under different conditions.
Results and discussion Powders Fig. 2(a) indicates that the original agglomerated powder presents a porous structure with nanometer or sub-micrometer scale primary particles, which is liable to cause gas entrapment and leads to the formation of pores inside the sprayed coating. After sintering at 1000 C for 10 h, the initial fine particles have grown up to micrometer particles characterized with grown sintering necks and through-holes as shown in Fig. 2(b). When the sintering temperature raised to 1100 C, the seal and spheroidization of pores have be en clearly observed from the polished section presented in Fig. 2(c). The surface view of spherical particle in Fig. 2(d) demonstrates closely arranged fine spinel crystals. By prolonging the duration of sintering to 10 h, the porosity was further decreased, accompanied by the merge and growth of grains shown in Fig. 2(f). At 1200 C, the powder presents dense structure and coarse grains (see Fig. 2(g) and (h)). According to statistical analysis
To effectively inhibit the O-diffusion and improve the oxidationresistance of the steel interconnect during IT-SOFC operation process, the MCO protective coating must have low porosity and good splat bonding. But above all, penetrating cracks as the primary path for gas transport must be avoided. The thickness of coatings prepared in this work was set as 100 mm. According to the cross-sectional view shown in Fig. 4(b) and (d), no penetrating crack is detected, instead, only few randomly located micro-cracks is observed. During plasma spraying, fast cooling of the splats from a high temperature near the melting point to a relatively low temperature inevitably results in a quenching stress. Once the cumulative stress exceeds the tensile stress of the deposited material, it can only be released in the form of cracking. In general, the thermally sprayed Mn1.5Co1.5O4 coating deposited at 400 C shows a high density with a relatively uniform thickness and good splat bonding which is attributed to elevated deposition temperature [53]. To evaluate the influence of substrate preheating and redox post-spray treatments on the density of MCO coatings, a comparison testing of the gas permeability which can directly reflect the porosity of through pores was carried out. Fig. 4(a) exhibits the cross-sectional microstructure of 430L supporting substrate with an open porosity of ~17.7% which was measured by image analysis software Image J. According to the gas permeability test, a rather high gas permeability value of 1.07 102 cm4 s1 gf1 was obtained, manifesting that the porous substrate used herein could provide enough communicating pores for gas transport. The gas permeability values for MCO coatings under different preparation conditions are then calculated and the standard deviation is also presented. As indicated in Fig. 5, the as-sprayed MCO coating without substrate preheating presents a high gas permeability of 1.21 106 cm4 s1 gf1, which is slightly decreased after subsequent redox treatment. In comparison, for coating deposited at 400 C, the gas permeability is only 1.23 107 cm4 s1 gf1, and a lower value of 8.2 108 cm4 s1 gf1 is achieved after 5 h reduction and 10 h oxidation, implying that a sufficiently dense coating was produced.
Influence of microstructure homogeneity on electrical performance of MCO coatings Fig. 6 shows the ASR of T441 coated with MCO coatings prepared under different treatment conditions. The ASRs were
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Fig. 7 e Sectional morphology of (a)-(b) T-MCO-1, (c)-(d) T-MCO-2, (e)-(f) T-MCO-3 samples after ASR test in air at 800 C for 200 h.
measured at 800 C in air as a function of time, under a current density of 200 mA cm2 flowing through the specimen. The ASR values exhibit a significant decline in the initial stage. This might be related to the increasing contact area and densifying of spinel coatings, since the ASR testing process was actually equivalent to a separate post-spray treatment under the oxidizing atmosphere. Besides, T-MCO-1 and TMCO-3 with pre-reduction show higher initial resistances, which is linked to the rapid generation of Cr-rich scale in the initial oxidizing process. Nevertheless, of the three coatings, T-MCO-3 exhibits the optimum performance and shows a final ASR value of 12 mU cm2. Once the ASR tests were completed, the samples were epoxymounted and polished for microstructure analysis. As can be seen, the MCO samples shown in Fig. 7 present typical sintered structures. For ASR tested T-MCO-1 (see Fig. 7(a) and (b)), a few interior cracks and numerous closed-pores are clearly observed, despite good interface bonding after redox treatments.
MCO coating without pre-reduction For T-MCO-2 sample without reduction pre-treatment, the tested coating presents distinct two-layer structure with a dense top layer and a nonuniform inner layer which is characterized by light-coloured precipitates and dark matrix, see Fig. 7(c) and (d). In addition, it can be seen that numerous pores are concentrated in the inner layer, especially near the coating/substrate interface. In order to reveal the phase structure and elementary composition of sample T-MCO-2, the top layer with a thickness of 50 mm was carefully removed, then energy spectrum and x-ray diffraction analysis were respectively performed on the polished nonuniform section. As indicated by XRD results shown in Fig. 8(a), MCO coating in the as-deposited state is comprised of tetragonal and cubic spinel phases, in addition to metal-oxide mixed phase MO (where M indicates Mn or Co) which was caused by partially spinel decomposition during thermal spraying [56,57]. According to the phase analysis
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Fig. 8 e Crystallographic phase analysis on MCO coatings (a) as-deposited, and after 200 h ASR test: (b) nonuniform region and (c) surface region of T-MCO-2 coating; (d) T-MCO-3 coating.
shown in Fig. 8(b), the inner region of oxidized (ASR tested) TMCO-2 is composed of tetragonal CoMn2O4, cubic MnCo2O4 and metal-oxide phases which were also detected in the asdeposited coating. The increase in the intensity of spinel phases is inferred to be connected with the transformation of metastable MO into spinel during oxidation treatment between 400 C to 800 C [58]. For the top layer, most of the diffraction peaks is identified to be corresponding to MnCo2O4 and CoMn2O4, see Fig. 8(c). As a semiquantitative analytical method, the measurement error for EDS point scan is generally less than 1% when heavy element such as Mn, Co is tested. In our present investigation, the work distance is set as 15 mm and the EDS analyse indicates that the average ratio of Mn/Co is 50/50, 51/49 for the dense top layer and the whole inner layer, respectively. However, for the inner layer, map scanning results shown in Fig. 9 indicate that the light isolated precipitates marked as spot 1 is in rich of Co (corresponding to the cubic CoxOy or MnCo2O4 phase), and the dark matrix marked as spot 2 is in rich of Mn (corresponding to the tetragonal MnxOy or CoMn2O4 phase). Similar microstructure composed of isolated precipitates was observed in the 1150 C sintered Mn1.75Co1.25O4 bulk sample, where the isolated precipitates were identified as the cubic MnCo2O4 phase and the matrix as the tetragonal phase [59].
MCO coating with pre-reduction Unlike the two-layer structure, the oxidized T-MCO-3 shown in Fig. 7(e) and (f) presents a uniform and dense structure with fewer pores, cracks and good interface bonding which would contribute to an increased electrical conductivity [51]. The XRD result in Fig. 8(d) shows a composition corresponding to MnCo2O4 and CoMn2O4. According to EDS data, the Mn/Co ratio is 50/50, 51/49 and 51/49 respectively for the top, middle and inner layers, manifesting that selective evaporation of elements did not occur and a homogeneous composition was obtained for the whole coating. Besides, it is worthwhile mentioning that a relatively denser spinel layer is observed in surface area of MCO coating, which is attributed to the high sinterability of MCO after redox treatment at a temperature no higher than 800 C [29]. In this work, the sintering of metallic cobalt under the reducing environment was essentially necessary and beneficial for the subsequent homogenization and densifying process during the oxidizing treatment. According to Brylewski [59], the electrical conductivity of Mn3xCoxO4 spinels was significantly influenced by phase structures and the highest conductivity was obtained in range of 1.5 < x < 2. The relatively higher resistance of T-MCO-2 coating could be related to the metastable phase MO detected in the nonuniform layer. In contrast, T-MCO-3 comprised of
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Fig. 9 e SEM views of EDS map scan based on nonuniform region of ASR tested T-MCO-2 sample. tetragonal CoMn2O4 and remaining cubic MnCo2O4 presented a lower resistance. To simulate the cathode/interconnect structure in SOFC stacks, porous LSM cathode was deposited onto bare T441 and MCO-coated T441, respectively. Interfacial ASRs between LSM and T441 with and without MCO protection layer were then studied to evaluate the effectiveness of MCO coatings. Prior to the resistance measurements, MCO coating deposited at 400 C was processed with redox post-spray treatments as noted in Table 2. According to Fig. 10, LSM/T441 exhibits a substantial increase in interfacial ASR during the test (to a final ASR value of 42 mU cm2). The poor oxidation resistance performance and Crrich spinel scale with low electric conductivity in the oxidizing atmosphere have negative effects on the power output when bare T441 is used as interconnects. It was reported that the bare 441 showed a high ASR of 75 mU cm2 after long-term oxidation in air at 800 C [38]. In comparison, LSM/MCO-T441 shows a significant decline in the initial stage and exhibits a high stability during the continuous testing process. A much lower value of 13 mU cm2 was obtained at the end of ASR test. Fig. 11 presents the cross-sectional microstructure of LSM/ MCO-T441 after ASR test. As can be seen, LSM/MCO exhibits a good interface bonding which would contribute to a lower interface resistance. Moreover, the MCO protection coating also shows a good adhesion to T441 substrate, despite a few isolated micro-voids are visible at the interface region. The EDS analysis (see Fig. 12) indicates that the Cr-rich oxide scale formed at the MCO/T441 interface is less than 1.2 mm and no further migration of Cr is detected, demonstrating the
effectiveness of MCO protection coating in inhibiting subscale growth. Dense spinel microstructure achieved by APS and redox treatments makes up the most important factor to provide adequate protection against Cr diffussion and minimize the electrical resistance of T441 interconnects.
Mechanism of microstructure homogenization during posttreatments The microstructure homogenization process of MCO coating during post-treatments is schematically presented in Fig. 13.
Fig. 10 e Interfacial ASRs between a LSM cathode and T441 with and without the spinel protection layer.
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Fig. 11 e Sectional morphology of LSM/MCO-T441 after ASR test at 800 C in air for 200 h. The as-sprayed coating characterized with a few closed pores and randomly located micro-cracks is indicated as Fig. 13(a). As previously mentioned, partially spinel decomposition was detected during the thermal spraying. Therefore, mixed phases of spinel and metal-oxides with partially amorphous characters were formed. At the initial oxidation stage, the molecular oxygen rapidly diffused into the coating surface and brought to the re-formation of spinel structure, accompanied by the crack self-healing. Unlike the free surface, the formation rate of spinels in the inner coating was significantly influenced by the diffusion rate of molecular oxygen. With the re-formation of spinel structure in the top layer, the inner diffusion of O/O2 was retarded to some extent, leading to an oxygen deficiency and a nonuniform region within the coating. Thus, a distinct two-layer structure was presented, as indicated in Fig. 13(b). Nevertheless, the interface between the top spinel layer and the inner nonuniform layer was still
assumed to gradually evolve towards the substrate side in prolonged oxidation process, accompanied by the formation of Cr-rich scale. It is plausible to speculate that the homogenization process would proceed along with the re-formation of Mn-Co spinel. This assumption was also confirmed by the existence of metastable phase at the inner nonuniform region of T-MCO-2 after 200 h oxidation test. For MCO coating with pre-reduction, different microstructure evolution mechanism was presented. The introduction of reduction pre-treatment could enhance the interfacial adhesion on micro level and promote the closure of micro-cracks originally found in the as-sprayed coating. When the porous MnO-Co coating was subsequently exposed to oxidizing atmosphere, O2 rapidly diffused into the whole coating, accompanied by in-situ oxidizing reaction. Assuming that the molecular oxygen distributed evenly, the formation rate of primary spinel should be similar throughout the whole
Fig. 12 e EDS line scan based on line AB shown in Fig. 11(b).
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Fig. 13 e Schematic demonstrating the microstructure evolution mechanism of MCO/T441 under the following states: (a) assprayed, (b) at the initial oxidation, (c) in subsequent oxidation, respectively; and (d) MCO with pre-reduction, and then (e) at the initial oxidation, (f) in subsequent oxidation, respectively.
coating, which would promote the homogenization process of Mn-Co spinel. In the meantime, it was speculated that the inner diffusion of O/O2 advanced more rapidly than the in-situ oxidizing reaction, and the molecular oxygen reached the coating/steel interface within a relatively short time [38,60]. Thus, the oxygen activity at the coating/steel interface could be well above the critical value of Cr2O3(1028 atm), leading to the formation of a thin and compact Cr2O3 scale inevitably, as shown in Fig. 13(e). Nevertheless, the diffussion rate of O/O2 gradually decreased with the growth of spinel grains and the reducing of grain boundaries. A homogeneous and stable spinel was produced as the oxidation sintering process went on. Thermodynamically, the basic driving force for shrinkage of ceramic coating in sintering process is the decrease of surface energy. The surface energy minimization leads to the elimination or coalescence of small isolated pores and a sufficiently dense spinel coating which is impermeable to molecular oxygen can be eventually achieved, see Fig. 13(f).
Conclusions In this study, Mn1.5Co1.5O4 coatings with a relatively uniform thickness were deposited at room temperature and 400 C via atmospheric plasma spraying, then followed by redox postspray treatments. The gas tightness of the as-sprayed MCO coating deposited at 400 C was 1.23 107 cm4 s1 gf1, and that of the as-treated coatings can be as low as ~8.2 108 cm4 s1 gf1, which is essential for sufficient protection of metallic substrate against oxidation. Metal-oxide mixed phase MO (where M indicates Mn or Co) from partially spinel decomposition during thermal spraying was detected in the as-sprayed MCO coatings. For ASR tested coating without pre-reduction, a distinct two-layer structure composed of dense top layer and nonuniform inner layer was observed. The inner layer was characterized by isolated precipitates rich in Co and dark matrix rich in Mn. While for
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coating with pre-reduction, the Mn/Co ratio was closed to 50/ 50, manifesting that selective evaporation of elements did not occur and a uniform composition was obtained. ASR tests also confirmed that a homogeneous structure had contributed to an increase in the overall conductivity. Of the three coatings, T-MCO-3 with pre-reduction exhibited the optimum performance and a final ASR of 12 mU cm2 was obtained at the end of 200 h duration. Besides, LSM/MCO-T441 system exhibited a high stability during the 200 h test, and a final interfacial ASR of 13 mU cm2 was obtained. SEM/EDS analysis indicated that the Cr-rich oxide scale formed at the MCO/T441 interface was less than 1.2 mm and no further migration of Cr was detected, demonstrating the effectiveness of MCO protection coating in inhibiting the subscale growth and Cr diffussion. In comparision, the bare T441 exhibited a substantial increase in interfacial ASR during the test (to a final ASR value of 42 mU cm2). Once properly treated, APS-prepared Mn1.5Co1.5O4 coatings have desirable gas-tightness and sufficient high electrical conductivity as protective coatings for IT-SOFC interconnects.
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