Effect of W and Ti addition on oxidation behavior and area-specific resistance of Fe–22Cr–0.5Mn ferritic stainless steel for SOFCs interconnect

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Effect of W and Ti addition on oxidation behavior and area-specific resistance of Fee22Cre0.5Mn ferritic stainless steel for SOFCs interconnect Azim Safikhani*, Mohammad Aminfard 1 Department of Materials Sciences and Engineering, Islamic Azad University (Saveh Branch), Saveh, Iran

article info

abstract

Article history:

Ferritic stainless steel (FSS) is an appropriate material to be used in solid oxide fuel cells

Received 11 October 2013

(SOFCs) interconnect. Chromium in FSS evaporates at high temperature and it is deposited

Received in revised form

on the cathode surface of SOFC and makes the cell poisonous. So the present study, the

20 November 2013

effect of W and Ti addition on oxidation behavior, electrical property and Cr evaporation

Accepted 25 November 2013

resistance were discussed in terms of microstructure property of the oxide scale by using

Available online 27 December 2013

SEM analysis. W and Ti addition reduce the oxidation rate, Cr evaporation rate and areaspecific resistance of the FSS. Presence of W in this steel formed chi (c) phase in scale/

Keywords:

alloy interface and prevented the diffusion of cation to the oxide scale. It is also a barrier for

Solid oxide fuel cell

the influence of the oxygen anion to the inward of the FSS. The presence of high amount of

Interconnect

W and low amount of Ti were effective in improving oxidation and electrical conductivity.

Oxidation

However, to enhance the Cr evaporation resistance, high Ti content is required.

Chi phase

Copyright ª 2013, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights

Tungsten

reserved.

Titanium

1.

Introduction

By rising energy consumption, diminishing quantities of fossil fuels, and the more and more severe environment pollution, the demand for more efficient and environment-friendly devices is growing rapidly. The problem which has been presented to the scientific community is to develop technology to allow the power generation industries use a variety of clean fuels with higher conversion efficiencies and low emissions than is currently possible [1]. Considerable effort has been and continues to be exerted in the research of a fuel cell as an

alternative to fossil fuel in hopes of providing more environmentally friendly and clean energy sources [2]. Fuel cells are electrochemical devices that convert chemical energy in fuels into electrical energy directly, promising power generation with high efficiency and low environmental impact [3]. Significant progress in the research and development of solid oxide electrochemical devices has been made for several years to convert chemical reaction energy directly into electrical energy without combustion process [4]. Among various types of fuel cells, solid oxide fuel cells (SOFCs) have extra merits including acceptable energy output and fuel flexibility [5]. A SOFC is composed of two porous electrodes, an anode and a

* Corresponding author. Home address: Block 2, Apartment of Housing Foundation, SqKhorramshahr, Bojnord City 9417764958, North Khorasan Province, Iran. Tel.: þ98 584 2411476, þ98 935 7990637 (mobile). E-mail addresses: [email protected] (A. Safikhani), [email protected] (M. Aminfard). 1 Home address: Iran, Alborz Province, Karaj City, Keshavarz Street, Resalat Lane, No. 338, Unit 6. Tel.: þ98 263 6103815, þ98 919 5515728 (mobile). 0360-3199/$ e see front matter Copyright ª 2013, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2013.11.100

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cathode, separated by a gas imperviousness ion-conducting electrolyte. The interconnect provides electrical connection between anode of one individual cell to the cathode of the neighboring one. It also acts as a physical barrier to avoid any contact between the reducing and the oxidizing atmospheres [4]. It is well-known that if fuel cells want to reach the practical power levels, individual cells must be repeated to form a fuel cell stack. Interconnects are the components that separate individual cells and provide the means to complete the electrical circuit. Thus interconnects maintain the uniform fuel and air flow. They also play a critical role in efficiency and power density of fuel cell stacks [6]. Each interconnect contains an electrolyte, an anode and a cathode. Since the SOFC system operates at high temperature (650e800  C) and oxygen is supplied for chemical reaction with fuel (mainly hydrogen or hydrocarbon compounds) in order to generate electricity, oxidation of interconnect is anticipated. The requirements for interconnect materials are stringent, and a satisfactory solution has not been found so far [3]. Interconnects must be suitable electrical conductors and exhibit good oxidation resistance in the SOFC operating environment. Ceramic interconnects such as La1x (Sr, Ca)x CrO3 have been used in a SOFC system operating above 800  C; but metallic interconnects such as stainless steel have been included in interconnect research due to the decrease in the operating temperature to an intermediate range (650e800  C) [7]. Metallic interconnects have some advantages: low material cost, suitable mechanical properties, high thermal conductivity, high electrical conductivity, and easy manufacturing process in a large area, etc. [8]. Among the investigated alloys as interconnect material, ferritic Stainless Steels (FSSs) have attracted more attention due to their superior properties such as good TEC matches in comparison to other cell components, low cost, chemical stability, corrosion resistance, thermal properties, mechanical strength, competitive price and high oxidation resistance which is caused by chromia scale formed on the FSS at temperatures around 650  Ce800  C in both oxidizing and reducing atmospheres [9,10]. In light of these requirements, chromia forming ferritic stainless steels such as Allegheny Ludlum (AL) 441-HP or Crofer 22 APU (ThyssenKrupp VDM) and ZMG232 (Hitachi Metals) [11] and also Ni-base alloys such as Haynes 230 or Haynes 242 [31] are the most promising materials of interconnects in solid oxide fuel cells. A significant drawback to chromia-forming interconnects is chromium degradation or chromium poisoning. Depending on the atmosphere at the cathode, chromium rich alloys form chromium three oxides CrO3, or chromium hydroxides CrO2(OH)2. Upon combining with oxygen ions in the cathode active area, the chromium compounds reduce back to chromia scale, Cr2O3. Thus the chromia scale formation results in decreasing cathode active area. In addition, the electrically insulating nature of chromia scales increases the contact area specific resistance (ASR) between interconnect and cathode. Consequently, the performance and efficiency are significantly degraded [6,12,13]. In order to solve this problem, protective coatings can be useful; such as perovskite coating materials: (LaSr) CoO3, (LaSr)CrO3, (LaSr)MnO3, have been studied extensively and in general because of diffusion of chromium through the

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coatings and potential formation of thick interfacial reaction layers they are not very effective. Several conductive spinel coatings, such as MnCo2O4, Mn1.5Co1.5O4, (Cu,Mn)3O4, etc. were found to be more promising to block Cr migration [14]. Also NiFe2O4 spinel coating, which is nonvolatile, conductive and better chemical compatibility to the Fe-based and Nibased metallic interconnects, can be employed not only to reduce the Cr2O3 evaporation effectively but also enhance the hot corrosion resistance of the Cr2O3-forming interconnects significantly, i.e., limiting the growth of the Cr-based oxide scale, by either smartly controlling the Ni content in its precursor material or improving the coating density and thickness [32]. The first important parameter in the appropriate design for manufacturing SOFC interconnect is to choose a suitable alloy. A suitable alloy must be inexpensive, and have TEC matches with that of other cell components. By choosing a suitable alloy and a good coating, desired properties can be achieved. The Ni-base alloys Haynes 230 and Haynes 242 exhibited better oxidation resistance than the Fe-base alloys Ebrite and Crofer 22 APU. The mass gain of Haynes 242 was the lowest, while that of Crofer 22 APU was the highest. The scale ASR increased faster by increasing scale thickness for the Fe-base alloys in comparison to the Ni-base alloys which is also consistent with their oxidation resistance [31]. Although the results of the Ni-base alloys are better than Fe-base alloys. But the Ni-base alloys are more expensive. A designer should consider the all aspects. So one of the best options for SOFC metallic interconnect is the Fe-based alloys such as ferritic stainless steel. The reactionary elements, such as Y, La, Ce [7] and Zr, in the alloy are also known for promoting oxidation resistance, scale adherence and conductivity [15]. Recently, Nb is commonly added to FSS to improve the high-temperature strength, creep resistance, and prevent sensitization [9e11]. But if the high amount of Nb (about 0.16 wt%) is added to the FSS, the Nb will precipitate as Fe2Nb in the grain boundaries [35]. Adding a low amount of laves phase formation element (Nb, Mo, W, or Ti) to traditional FeeCr alloys (such as SUS430) is effective to control the elemental diffusivity at the alloy grain boundaries by forming a laves type phase in air [16]. Reported that Nb, W and Mo are commonly added to FSS to improve the mechanical property of the high temperature strength, oxidation resistance, creep resistance and reduction of ASR [17,18]. But these elements have disadvantages as production of brittle intermetallic secondary phases. These brittle intermetallic phases are generally known to degrade impact toughness and corrosion resistance of the alloys [25]. In Haynes 242, Mo forms intermetallic phase with Cr and Ni to promote the formation of NiO at the surface, which can reduce Cr evaporation rate [16]. However, it is suggested that the service temperature of Haynes 242 be lower than 700  C for the potential harm of volatile MoO3 [1]. It is reported that if more than 4 wt% Mo is added to the FSS, the formation of Ferich spinel will promote. Furthermore, Mo addition of more than 4 wt% can be detrimental to the long-term stability of protective oxide layers [17]. It is also reported that presence 3.8wt% of the Mo in 27.4Cre3.8Moe2.1Ni ferritic stainless steel formed c-phase precipitates. Therefore, it decreases the ductility and increases the hardness in the ferritic stainless

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Table 1 e Chemical composition of specimens (wt.%). Base F.Ti F.W F.Ti1W1 F.Ti1W2 F.Ti2W1 F.Ti2W2

Fe

Cr

Mn

C

Si

Al

Ti

W

Bal. Bal. Bal. Bal. Bal. Bal. Bal.

21.8 21.9 20.9 20.6 19.8 20.7 19.6

0.49 0.49 0.46 0.47 0.44 0.46 0.43

0.20 0.19 0.19 0.20 0.18 0.19 0.19

0.01 0.03 0.01 0.02 0.01 0.01 0.02

0.002 0.002 0.004 0.003 0.002 0.002 0.004

0.001 0.26 e 0.25 0.23 0.52 0.48

0.002 e 1.98 2.01 3.98 2.00 3.95

steel [34]. If the amount of tungsten is too high, it can be more easily accumulated at the oxide/metal interface and can form voids. Void formation can reduce the oxidation rate for a short time by disconnecting diffusion path from metal to oxide [24]. It is reported that Ti addition to interconnect material was effective to form conductive oxide scale [19,20]. Addition of Ti was also effective to reduce Cr evaporation of ferritic stainless steel [21,22]. But if the amount of Ti is too high, oxidation rate and ASR will increase [21]. In this study, the effect of W and Ti additions are investigated to the Fee22Cre0.5Mn FSS separately and together, on the oxidation behavior, electrical property in the oxide scale, and Cr-evaporation. The results clearly show that the simultaneously addition of W and Ti to ferritic stainless steel causes the using of FSS in metallic interconnect to be possible.

2.

Experimental

2.1.

Materials and sample preparation

Various amounts of Tungsten (W) and titanium (Ti) were added to the base steel. Each specimen was prepared by vacuum induction melting (VIM) and metals with purity higher than 99.9% were used. VIM technique is an effective method to provide controlled alloy composition and homogeneity. To prevent the oxidation of the elements, the chamber was evacuated under 105 torr, and then filled with pure argon gas. Melted specimens were homogenized by heat treating at 1200( C) for 24 (h) and then water quenched. The chemical composition of the specimens is listed in Table 1. The chemical composition was confirmed by chemical analysis equipment.

2.2.

X2 ¼ kp t þ C

(1)

X2 ¼ ðDm=AÞ2

(2)

Dm ¼ m2  m1

(3)

Where in equation (1) X ¼ weight gain, kp ¼ parabolic rate constant, t ¼ oxidation time, C ¼ integration constant. Where in Equation (2) Dm ¼ weight change, A ¼ sample surface area. Where in Equation (3) m2 ¼ weight of sample after each 100 h oxidation and m1 ¼ Initial weight of sample before oxidation process.

2.3.

Area specific resistance (ASR) measurement

To measure the electrical resistance of the oxide scale, samples were prepared with dimension of 15 mm diameter and 2 mm thickness with the circular shape area. The surface was polished up to P2000 grade by SiC paper. A popular method for determining the resistivity of materials is a test as the fourpoint probes. In this test, firstly the samples must be oxidized at 800 ( C) for 100 (h). Then for ASR measurement, the four-point probes method with DC current must be used. Platinum paste was painted on both sides of the sample and then platinum meshes were attached to both sides (Fig. 1). The

Oxidation experiments

For the oxidation test, 7 samples were provided. The samples were cut into 20 mm diameter and 5 mm thickness with the circular shape area. The surface was polished up to P2000 grade with SiC paper. After rinsing with ethanol and acetone, the samples were oxidized at 800 ( C) for 500 (h) in a box furnace with ambient air. The specimens were cooled down to room temperature to measure the weight change in every 100 h. To avoid the spallation of thermal shock, the temperature increasing rate and decreasing rate were controlled to be fewer than 5 ( C/min) to minimize thermal stresses on the oxide scale of samples during the operation. The parabolic rate constant of each sample for quantitative analysis on the oxidation rate was evaluated by Equations (1)e(3).

Fig. 1 e A schematic of sample design prepared for ASR measurement.

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Fig. 2 e Experimental set-up for measuring the electrical contact resistance of oxide scales [23].

specimen and the platinum sheets were placed in a ceramic holder (Fig. 2). The current was applied to use a Jandel current source, at a fixed constant current density 0.33 A cm2. The voltage drop across the specimen was also measured by the Keithely Multimeter. The data were used to calculate the resistance according to Ohm’s law by Equation (4) and the ASR by Equation (5). R ¼ V=I

(4)

ASR ¼ R$A

(5)

In above equations, R is the resistance in Ohms (U), V is the voltage (Volt), I is the current (Amp) and A is the sample surface area.

2.4.

Cr evaporation rate measurement

To investigate the Cr-evaporation rate, circular shape samples with 15 mm diameter and 2 mm thickness were prepared. Samples were polished up to P2000 grade with SiC paper. In this test, firstly the samples were oxidized at 800  C for 100 h, and then they were placed at the flowing air through the quartz tube at 800  C for 24 h. The relative humidity (RH) in the test zone was controlled to supply the air saturated with water vapor at room temperature which is equivalent to 92% RH (with 1% error rate), and the gas flow rate was controlled at 300 ml.min1. The quartz tube was rinsed by a solution of 2 mol HCl (with 35% concentration per mole) {((35 g HCl/100 g solution)  100 ¼ 35%)}þ1 mol HNO3 (with 60% concentration

Fig. 3 e A schematic of test apparatus for evaporation process [24].

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Fig. 4 e Weight gain of the FSS for various samples after oxidation at 800  C in ambient air as a function of oxidation time.

per mole){((60 g HNO3/100 g solution)  100 ¼ 60%)}. The amount of Cr vaporizable samples was measured by inductively coupled plasma (ICP) analysis. Schematic of the Cr evaporation process is shown in Fig. 3.

2.5.

Analysis method for oxide scale of samples

The oxide phases were formed on the alloy surface and secondary phases were identified by energy-dispersive X-ray spectroscopy. A cross-section of the scale was examined by scanning electron microscopy (SEM), and distribution of maximum and minimum alloying elements in the oxide scale was analyzed by glow discharge optical emission spectrometry (GDS).

3.

Results and discussion

3.1.

Oxidation test results

Fig. 4 shows the weight gain as a function of oxidation time for each specimen that was listed in Table 1. The test was conducted at 800 C in ambient air for 500 h. For quantitative analysis of the oxidation kinetics, the parabolic rate constant of each sample was evaluated and listed in Table 2. Referring to Fig. 4, it is clearly observed that the sample with code F.Ti1W2 has the best oxidation resistance. According to Table 1, this sample contains 3.98 wt% of W and 0.23 wt % of Ti. On the other hand, this sample has a minimum amount of parabolic rate constant (Table 2). The highest weight change was obtained in the base sample. W and Ti alloying elements were caused reduction of the oxidation rate in the steel. The weight gain in the sample with code F.W is

higher than F.Ti. However, the weight gain in the sample with code F.Ti2W1 is higher than both of them. Therefore, the oxidation rate can be increased by adding the amount of Titanium and decreasing the amount of Tungsten. But in sample with code F.Ti1W2, the oxidation rate decreased by adding the amount of W and decreasing the amount of Ti. It is reported that 1wt% Ti addition increases the oxidation rate of ferritic stainless steel, but the addition of 0.05 wt% Ti suppressed it. The presence of the small amount of Ti in Cr2O3 matrix retarded the oxidation rate by changing the oxidation kinetics. However, the high amount of Ti in Cr2O3 matrix generated an excess ionic defect and reduced the oxidation resistance [21].Ti can also induce the creation of additional cationic vacancies in the Cr2O3 matrix and subsequently promote fast diffusion of ion through Cr2O3 layer [19,30]. Consequently, the effect of Ti on the oxidation rate is changed depending on the amount of Ti addition. A large amount of Ti addition in the FSS i) accelerated the growth of Cr2O3 layer, ii) suppressed the growth of inner MneCr spinel and iii) accelerated the oxidation of outermost MneCr spinel as reported in Ref. [21]. The growth of Cr2O3 layer was promoted by changing the defect structure of Ti addition rather than volume expansion of Ti oxidation. The amount of Ti in the Cr2O3 matrix was too low to increase the volume of oxide. However, diffusion property of Cr2O3 can be changed by a small amount of Ti addition. For example, Atkinson et al. reported that 2 mol% of TiO2 in the Cr2O3 generated VCr and increased the electrical conductivity of both electronic and ionic ones [20]. Although a large amount of Ti addition increased both Cr and Mn diffusivity in the oxide scale, a small amount of Ti suppressed diffusion of these cations. On the other hand, according to Table 1, an increasing amount of tungsten from (2 wt%) to (4 wt%) in ferritic stainless steel improved the oxidation rate. It is reported that (4e6 wt%) of W reduce the oxidation rate in ferritic stainless steel by chi (c) phase formation [24]. Though c phase is stable in the thermodynamic point of view, the atomic size of the tungsten is so large that its diffusion will be slow. This means that another mechanism other than diffusion of W may affect the formation of c phase at the interface. In according to Fig. 4, increasing the amount of W and decreasing the amount of Ti in Fee22Cre0.5Mn reduced the oxidation rate.

3.2.

GDS and SEM analysis results

Fig. 5 shows GDS depth profiles for samples of (base, F.Ti, F.W and F.Ti1W2) after 100 h oxidation and Fig. 6 shows SEM images of them. The GDS depth profile results are consistent with the SEM results. The oxide scale of each specimen mainly consisted of CreMn spinel at the outermost layer and Cr2O3 in the inner layer, which matches with oxide structure. It seems that this inner spinel layer be the result of the reaction between Cr2O3 and MnO, which can be formed underneath the

Table 2 e The oxidation rate constant of the specimens tested at 800  C in ambient air. Sample Kp  10

13

2

(g mm

4

1

h )

Base

F.Ti

F.W

F.Ti1W1

F.Ti1W2

F.Ti2W1

F.Ti2W2

10.8

3.10

4.50

1.72

1.10

6.8

1.34

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Fig. 5 e GDS depth profile of distribution elements near interface between scale/alloy after 100 h oxidation at 800  C in air: (a) Base, (b) F.Ti, (c) F.W and (d) F.Ti1W2.

Cr2O3 layer due to lower equilibrium oxygen partial pressure of MnO (1.6  1025 Pa for MnO, 1.4  1023 Pa for Cr2O3) [24]. The metal/oxide interface is divided into three regions: “A” is the outermost region of the oxide surface (spinel region), “B” is the main body of the oxide, and “C” is the alloy substrate region. The outermost layer of the oxide scale of each sample in spinel region is enriched by Cr and Mn. And Cr2O3 is formed in the inner layer of the region B. The Mn content decreases rapidly and Cr becomes dominant in the outer part of region B. So, according to GDS results, the peak of Mn is formed below

the Cr2O3 substrate. Below the spinel, Mn content decreases rapidly and Cr becomes dominant in the outer part of region B. This is an indication that Cr2O3 formed mainly in this region. In the lower side of region B, the Cr content decreases as it moves from the oxide to the alloy. This means that the peak of Cr is obtained in the middle of region B. In accordance with Fig. 5(b) for sample F.Ti, Ti is the outermost oxide scale surface and the scale/alloy interface were segregated. On the scale surface shown in Fig. 5(b), there are Ti-oxides in some parts of MneCr spinel and this is also observed in the SEM results in

Fig. 6 e Cross-section SEM image of the specimens after 100 h oxidation at 800  C in air: (a) Base, (b) F.Ti, (c) F.W, (d) F.Ti1.W2.

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Fig. 7 e Ellingham diagram for standard reaction of the various elements.

Fig. 6(b). In the Cr2O3 layer (region B), Ti distribution is localized near the scale/alloy interface dominantly. However, according to Figs. 5(b) and 6(b), Ti distribution in Cr2O3 decreases significantly when Ti content is near 0.26 wt%. The presence of Ti addition in F.Ti suppressed the growth of the inner MneCr spinel compared to the base sample. And by reduction of oxidation rate, the oxide scale thickness is reduced when compared to steel base. But the gap value in F.Ti sample at the scale/alloy interface is lower than the base sample as shown in Fig. 6. The presence of gap in the scale/alloy interface of the base sample is resulting from the difference in thermal expansion coefficient of the substrate metal and chromium oxide. According to the Ellingham diagram in Fig. 7, Ti is an oxygen active element. As shown in Fig. 4, in the F.Ti sample, the effect of Tiaddition is similar to the effect of reactive element addition which reduces the oxidation rate of stainless steel and the growth rate of Cr2O3. The reactive element reduces the oxidation rate of the steel through the reduction of ion mobility in the scale by retarding ion diffusion through the interface poisoning and grain boundary segregation [7,33]. It has been reported that the reactive element should be finely and uniformly distributed and the proper amount of the reactive element be typically between 0.1 and 0.2 at% to maximize its effect. However, when the amount of reactive element is larger than the appropriate amount, there are some disadvantages such as deterioration in the scale adhesion [7,21].Ti is an effective element to increase the electric conductivity and oxidation resistance of the oxide scale, and the effect of Ti is related to distribution of Ti oxide in the scale [11]. In F.W

sample as shown in Fig. 5(c), the effect of W-addition is intensity reduction of Fe content in the spinel layer and the oxide scale. In F.W specimen, the thickness of the CreMn spinel layer and the oxide scale were thinner than the Base and F.Ti specimens. The peak of manganese was observed in the spinel layer and then declined sharply in the oxide scale. According to Figs. 5(c) and 6(c), the peak of W was observed in the scale/alloy interface. As shown in Fig. 6 (c) and (d), W rich second phase (cphase) was beneath the oxide layer in F.W and F.Ti1W2 specimens. This can promote the formation of W rich second phase at the scale/alloy interface as well as the metal grain boundary. These results are similar to the results of reference [24,25] for 4 and 8 wt% of W respectively. As Dae Won Yun et al [24]. have reported that the c phase particles were also formed inside the grains. This seems to be related with the large grain size, which is about a few millimeters. Although the grain boundary is a favorable site for both thermodynamically and kinetically, a large number of particles is formed inside the grains since the grain boundary area is so small. The increase in W amount from 2 wt% (in F.W sample) to 4 wt% (in F.Ti1W2 sample), increased c phase particles in the scale/alloy interface and adhesion of oxide scale to the substrate. According to Figs. 5(d) and 6(d), the thickness of spinel layer and the oxide scale were decreased compared to the other samples and it shows the improvement of the oxidation resistance. These results are similar to the results of Fig. 4.The reduced Cr concentration in the oxide scale at the F.W and F.Ti1W2 specimens can be explained by this second phase effect. This reduced Cr diffusion was responsible for the lower oxidation rate of F.W and F.Ti1W2 specimens. The amount of W added to the ferritic stainless steel is an important because a large number of voids can be observed in the high amount of W. This high amount is higher than 4wt% of W. According to ref [24], the highest amount of the voids was observed in specimen with 6wt% of W. This void formation is mainly occurred by the accumulation of vacancies. Cation vacancies diffuse toward the oxide/metal interface, and annihilated by vacancy sinks such as dislocations and grain boundaries at the oxide/metal interface. However, void formation will decrease the oxidation resistance in a long term exposure by dissociative transport across the voids and development of micro-channels. Moreover, a large number of void formations will eventually lead to spallation of the oxide layer, which is detrimental to the oxidation resistance of interconnects and electric contact with cells. The peak of W observed at the scale/alloy interface. Dark particles in the metal substrate were identified as TiO2 by energydispersive X-ray spectroscopy. These oxide inclusions were formed during the melting process. The higher amount of TiO2 observed under the W-rich region. The presence of c-phase particles in the interface prevented from the cation outward diffusion and the anion inward diffusion in the oxide scale.

Table 3 e Contents of principal metallic elements in cphase as determined by energy-dispersive X-ray spectroscopy after oxidation at 800  C for 100 h in Fee22Cre0.5Mn ferritic stainless steel (wt.%). W 21.8 9.6

Cr

Fe

Elements

17.9 18.4

Bal. Bal.

c-phase (F.Ti1W2) c-phase (F.W)

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Fig. 8 e ASR of each specimen oxide scale measured by 4point probe method with direct current after 100 h preoxidation at 800  C.

Escriba and et al. [26] reported the c-phase precipitation was found to decorate ferrite/ferrite interface in duplex stainless steel after 1 h at 750  C and after 2 h at 700  C during homogenization. Chi (c) secondary phase was formed during the oxidation test at 800  C after 100 h Table 3 shows the chemical composition of Chi (c) phase for F.Ti1W2 and F.W samples that observed by energy-dispersive X-ray spectroscopy.

3.3.

ASR results

Fig. 8 shows the area specific resistance changes for each specimen that was measured during the oxidation from 600 to 800  C after 100 h pre-oxidation at 800  C in ambient air. The ASR at 800  C is listed in Table 4.The oxidation of metallic interconnects in the SOFC operating environment is inevitable, and the thermally grown oxides with semi-conductor behavior will significantly deteriorate the cell performance after long term oxidation. A low and stable electrical resistance is required for metallic alloys as interconnects in SOFCs. The ASR of commercially available steel interconnects increases as the oxide scale grows [27,28]. There are various factors that can contribute to this increased resistance. The main factor is attributed to the growth of two spinel layer and chromia oxide scales. There are also other contributions to increased ASR that are related to the metal/oxide scale interface. Interfacial imperfections, including voids and cavities, as well as impurity segregation, reduce scale-to-metal adhesion and actual surface area of intimate contact between the metal and oxide scale. Consequently, this increases the interconnect ASR [23].

Zhenguo Yang et al. [29] reported that the oxidation behavior and ASR in ferritic stainless steel for SOFC interconnect applications with the minor alloying addition (Nb and Ti) lead to laves phase formation both inside grains and along grain boundaries during exposure in the intermediate SOFC operating temperature range. So it can improve the oxidation behavior and the electrical resistance. The effect of dissolved Ti ions is constant at all oxygen activities and this in turn, causes changes in the oxygen pressure dependence of the various defects at reduced oxygen activities. The Ti-doped chromia is an electronic conductor; this model predicts that the oxide is an n-conductor at reduced oxygen activities and a p-conductor at high oxygen activities [30]. Alan Atkinson et al. [20] reported that doping of chromia by Ti causes compensation Cr vacancies and electrical conductivity results has shown that high Ti doping concentration in this steel increases the oxygen activity. It is reported that addition of 1 wt% Ti increases ASR value significantly due to thick oxide scale and thus by long electric conduction path. But in case 0.05e0.07 wt% Ti decreases ASR value significantly due to thin oxide scale [21]. According to Table 4 and Fig. 8 the lowest amount of ASR is for the F.Ti1W2 sample that is 0.85  102 (U cm2). By comparing Figs. 4 and 8 for the Base, F.Ti, F.W and F.Ti1W2 samples, it is apparent that the electrical conductivity decreases and the ASR value increases with increasing the oxide scale thickness. So, the oxide scale on metal surface reduces the electrical conductivity. The presence of W was formed cphase in scale/alloy interface and caused adhesion increasing of the oxide to the substrate. Maximum concentration of W (3.98 wt%) and minimum concentration of Ti (0.23 wt%) is the best amount in ferritic stainless steel in this study. The high amount of W and the low amount of Ti were reduced the penetration of oxygen into the oxide scale by suppression of the oxygen activation and also reduced the cation mobility and prevented the moving of cation to the outside of the oxide scale. Thus, both the oxide scale thickness and ASR were decreased. Increasing the Ti amount in this steel, while Ti oxide is dissolved in the Cr-oxide, creates additional conductible species such as electron, hole and vacancy in oxide scale and these barriers causes the higher conductivity in the oxide scale. There were gaps in the scale/alloy interface of the base sample and there were also the porous in oxide scale. So, the ASR increased in the base sample compared to F.Ti, F.W and F.Ti1W2 samples.

3.4.

Cr evaporation rate

Fig. 9 shows Cr evaporation rate of each specimen that were measured in humid air at 800  C after 100 h pre-oxidation at 800  C in ambient air. The lowest Cr evaporation rate observed in the F.Ti2W2 sample and the highest amount of it observed in the base sample. However, F.Ti sample has also lower Cr

Table 4 e ASR of the oxide scale for each specimen measured after oxidation at 800  C in ambient air for 100 h preoxidation. Sample ASR  10

2

2

(U cm )

Base

F.Ti

F.W

F.Ti1W1

F.Ti1W2

F.Ti2W1

F.Ti2W2

2.69

2.46

1.32

1.91

0.85

3.55

1.04

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regeneration of the passive film even if it fails locally and the addition of Ti is also effective in the steel. Ti is a reactive element. One beneficial effect of dispersed reactive element oxides is their improvement in scale adherence, so these oxide additions are particularly effective in improving resistance to thermal cycling [38]. This is a good property of W and Ti as alloying elements to FSS since the electrolyte environment of the proton exchange membrane fuel cell (PEMFC) that is strongly acidic. With improved high-temperature strength by W addition in the FSS [39] and with improved resistance to thermal cycling by Ti addition the FSS and also with beneficial results obtained from the present work such as improve the oxidation rate, conductivity and Cr evaporation resistance in the ferritic stainless steel, the alloy that was designed, can be implemented in planar SOFC stack systems. Fig. 9 e Cr evaporation rate of ferritic stainless steel at 800  C in humid air after 100 h pre-oxidation.

5. evaporation rate than Base sample though it has thinner outer CreMn spinel layer than Base sample. It suggests that the reduction of Cr evaporation by Ti addition be not just due to physical property of the outer MneCr spinel layer. At first, Ti oxide formed on the oxide scale surface acts as a barrier to Cr evaporation reaction. In case of F.Ti2W2, this assumption is reasonable because it has high Ti content near oxide scale surface and it forms TiO2 mixed with the outermost MneCr spinel. However, in case of F.Ti1W2, Ti concentration in the oxide scale is much lower than Cr and Mn concentration and only a few Ti oxides are observed on the surface as presented in Fig. 6. Therefore, Cr evaporation reduction by Ti addition to these steels is not by forming physical barrier of Ti oxide but by changing chemical property of the oxide scale. So, the presence of high amount of Ti near the oxide scale surface acts as a barrier to Cr evaporation reaction. The presence of high amount of W by c-phase formation in the scale/alloy interface acts as a barrier for moving the Cr atoms and reduces the Cr evaporation rate. Formation of c phase during oxidation seems to affect Cr evaporation rate by altering the diffusion process. To improve the Cr evaporation rate, Ti concentration is important and the role of Ti is more important than W. According to the results of the oxidation and ASR, the F.Ti1W2 sample is better than the other samples. However, in the test of Cr evaporation rate, the F.Ti2W2 sample is the best result compared with the other samples.

4.

Use cases of this alloy in future research

At the end, it can be stated that the new alloy design concept for FSS can be possible by adding W and Ti, which can increase both the corrosion resistance and electrical conductivity [36]. According to the Pourbaix (or pH-potential) diagram at ref [37], W may form a stable WO3 passive layer even in a strong acidic solution. In fact, the addition of W to various grades of stainless steel has proved an increase in the corrosion resistance in strong acidic solutions but not in alkali solution. The addition of W to stainless steel has improved the repassivation kinetics during scratch testing so that it can assure

Conclusions

The effect of Ti and W addition to Fee22Cre0.5Mn ferritic stainless steel were investigated on the oxidation behavior, electrical properties and Cr evaporation rate and the results were discussed in terms of distribution of alloying elements in the oxide scale by GDS analysis and SEM. The tests were conducted at operating temperature of SOFCs (800  C). The results showed that the effect of Ti and W addition to the ferritic stainless steel can be useful to provide the operation of SOFC interconnects at high temperatures (800  C) by decreasing of oxidation rate and increasing the electrical conductivity. The results are summarized as follows: 1) The oxidation rate and the ASR value increased with the increase in oxide scale thickness. 2) In the base sample, some gaps were created in the scale/ alloy interface. The presence of Ti and W alloying elements, separately and together, in the ferritic stainless steel reduced the gaps and the porosity values. 3) The presence of 0.26 wt% Ti (in F.Ti sample) and 1.98 wt% W (in F.W sample) reduced both the oxidation rate and the thickness of oxide scale and also improved the ASR compared with the base sample. So the high amount of W and low amount of Ti can improve these properties. 4) The increase in the amount of W from 2 wt% to 4wt% in the F.Ti1W2 sample obtained a higher amount of c-phase in scale/alloy interface. 5) The presence of the c-phase particles in the scale/alloy interface prevented the diffusion of cation to the oxide scale. It was also a barrier for the influence of the oxygen anion to move from oxide scale to the FSS and improved the internal and the external oxidation and reduced the ASR. 6) The increase in the weight percentage of Ti in this steel when the Ti oxides dissolved in the Cr-oxide, created additional conductible species such as electron, hole and vacancy in oxide scale and these barriers caused the higher conductivity in the oxide scale. 7) W and Ti alloying elements had a considerable effect on the Cr evaporation rate in the F.Ti2W2 sample. The lowest Cr evaporation rate was obtained by the high amounts of W and Ti about 4 wt% and 0.5 wt% respectively.

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 2 2 8 6 e2 2 9 6

8) Finally, the F.Ti1W2 sample with 3.98 wt% of W and 0.23 wt % of Ti is the best sample in comparison with other samples because the highest oxidation resistance and the lowest electrical resistance were obtained in this sample.

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