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Oxidation behavior of ferritic stainless steels in simulated automotive exhaust gas containing 5 vol.% water vapor
Accepted Manuscript Oxidation Behavior of Ferritic Stainless Steels in Simulated Automotive Exhaust Gas Containing 5 vol.% Water Vapor
L.L. Wei, L.Q. Chen, M.Y. Ma, H.L. Liu, R.D.K. Misra PII:
S0254-0584(17)30931-8
DOI:
10.1016/j.matchemphys.2017.11.051
Reference:
MAC 20171
To appear in:
Materials Chemistry and Physics
Received Date:
27 June 2017
Revised Date:
25 October 2017
Accepted Date:
25 November 2017
Please cite this article as: L.L. Wei, L.Q. Chen, M.Y. Ma, H.L. Liu, R.D.K. Misra, Oxidation Behavior of Ferritic Stainless Steels in Simulated Automotive Exhaust Gas Containing 5 vol.% Water Vapor, Materials Chemistry and Physics (2017), doi: 10.1016/j.matchemphys.2017.11.051
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ACCEPTED MANUSCRIPT
Highlights
Ce and W were added in Fe-Cr alloy to improve its high temperature performance.
Oxidation kinetics displayed significantly different with adding Ce and W at 950–1100 ºC.
The property of oxide film was obvious improved with adding ~0.05 wt% Ce in steel.
Laves phase and porous SiO2 caused the oxide film spalling when adding ~1.0 wt% W.
ACCEPTED MANUSCRIPT
Oxidation Behavior of Ferritic Stainless Steels in Simulated Automotive Exhaust Gas Containing 5 vol.% Water Vapor L.L. Wei1), L.Q. Chen1, *), M.Y. Ma1), H.L. Liu1) and R.D.K. Misra2) 1) State Key Laboratory of Rolling and Automation, Northeastern University, 3-11 Wenhua Road, Shenyang 110819, China 2) Department of Metallurgical, Materials and Biomedical Engineering, University of Texas at El Paso, El Paso, Texas 79968, USA. * Corresponding author. Tel.: +86-24-83681819, Fax: +86-24-23906472, E-mail address: [email protected]
Abstract In this study, the effect of alloying elements, cerium (Ce) and tungsten (W), on the oxidation behavior of medium chromium ferritic stainless steel in simulated automotive exhaust gases containing 5 vol.% H2O was studied in the temperature range of 950–1100 °C for 5 h. The oxidation kinetics and oxide film characteristics were analyzed by means of thermogravimetric analysis (TGA), X-ray diffraction (XRD), scanning electron microscope (SEM) and elemental probe micro-analyzer (EPMA). The growth rate of oxidation and oxidation mass gain were significantly decreased on the addition of Ce. A dense, uniform, and thin oxide scale formed on Ce-containing steels. The steel containing Ce and ~0.5 wt% W displayed similar oxidation behavior compared to the steel containing only Ce. A large number of cracks and pores existed in the oxide film and oxide/metal interface leading to the spallation of oxide film, when the addition of W reached ~1.0 wt%. Key words: Ferritic stainless steel; Oxidation; Rare earth element; Tungsten; Water vapor
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ACCEPTED MANUSCRIPT 1. Introduction Ferritic stainless steels with single phase ferritic structure, and generally free of nickel, have
excellent corrosion resistance and low cost making them a promising material to
substitute austenitic stainless steel, for use in storage vessels, petroleum refining equipment and furnace parts, etc. In addition, the characteristic of low thermal expansion coefficient, good thermal conductivity and oxidation resistance render ferritic stainless steels as viable materials for manufacturing automotive exhaust pipes. Automobile exhaust manifold is the hot-end of exhaust system connected directly with the engine. The operating environment is generally ~900 °C. However, in order to satisfy the increasing stringent auto emissions standards, the temperature of automobile exhaust gas may reach 950–1100 °C [1, 2]. Thus, materials used for manifold are required to have excellent high temperature oxidation resistance and thermal fatigue resistance. According to the previous study, adding rare earth element is an effective way to improve the high temperature oxidation resistance of ferritic stainless steel because of its “reactive element effect” [3, 4]. The creep rupture strength of the steel can be enhanced by adding tungsten (W) through precipitation and/or solid-solution strengthening effect [5–7]. Kim et al. [8] reported that the addition of W can also retard nucleation and growth of σ phase during aging, thereby delaying degradation of corrosion resistance and mechanical properties of 25% chromium duplex stainless steel. Based on this idea, the chemical composition of ferritic stainless steel was redesigned by adding alloying elements, cerium (Ce) and tungsten (W). Ferritic stainless steels have excellent high temperature oxidation resistance, because preferential oxidation of chromium leads to a thin layer of Cr-rich oxide that inhibits oxidation reaction. The formation of oxide scale is influenced by temperature, alloy 2
ACCEPTED MANUSCRIPT composition, exposure time and especially the atmosphere. The atmosphere in the internal surface of manifold is automotive exhaust gas. Gasoline blend is a mixture of paraffin and aromatic hydrocarbons undergoes combustion in air by the following reaction [9]: gasoline + O2 (in air) → CO2 + H2O + heat
(1)
In addition to water vapor and CO2, automotive exhaust emission also contains a small amount of O2, sulfur dioxide (SO2), nitrogen oxides (NOx), carbon monoxide (CO), and unburned hydrocarbon (HC) [10]. But among of them, water vapor and O2 play the most important role in the high temperature oxidation of ferritic stainless steel. Previous studies have demonstrated that the presence of water vapor in the exposure environment can have a dramatic effect on the oxidation behavior of chromia-forming alloys as compared to the dry condition, involving rapid and catastrophic oxidation. Various mechanisms have been proposed to account for the degradation behavior of Fe-Cr alloys oxidized in water vapor containing environment. At high partial pressure of oxygen and water vapor, the formation of vaporized oxyhydroxides CrO2(OH)2 leads to less protection of oxide film [11, 12]. Rahmel and Tobolski [13] proposed that H2-H2O mixture is formed in voids at the scale/metal interface and/or within the scale, which dissociates leading to inward transport oxygen. Essuman et al. [14] suggested that when the water vapor is the source of oxygen, the internal oxidation is accelerated, and water vapor affects the solubility and/or diffusivity of oxygen in the alloy. Yuan et al. [15] proposed that the stress induced by accumulation of molecular hydrogen in the pores introduces cracking of chromia scale in pure steam. Water vapor can also enhance scale growth stress, the scale experiences frequent cracking and healing processes [16, 17]. In order to improve the high temperature properties of the manifold used ferritic stainless 3
ACCEPTED MANUSCRIPT steel, alloying elements Ce and W are added. But the role of alloying elements on high temperature oxidation resistance of the new developed ferritic stainless steel in water vapor containing automotive exhaust gas environment continue to be unclear. Thus, the aim of this study is to explore the influence of different content of alloying elements, Ce and W, on the high temperature corrosion behavior of ferritic stainless steel at 950–1100 ºC in simulated automotive exhaust gas environment. 2. Experimental method 2.1. Materials Four ferritic stainless steels were used in this study, and are named as F1, F2, F3 and F4. The F1 steel is an existing ferritic stainless steel (grade B444M2). In F2, F3 and F4 steels, different content of alloying elements, Ce and W, were added to F1 steel. The nominal chemical composition of these steels is presented in Table. 1. The difference among them is the content of W and Ce. Specimens of dimensions of 8 mm×12 mm×1 mm were fabricated from a cold-rolled plate with a small hole of 2 mm in diameter near one edge. F1 steel was annealed at 1020 °C for 2 min and F2, F3, and F4 steels were annealed at 1050 °C for 2 min, followed by air cooling. Prior to oxidation experiments, all the specimens were metallographically polished with SiC paper up to 1200 grit, ultrasonically cleaned, and stored in a desiccator. 2.2. Oxidation experiments The oxidation kinetics of ferritic stainless steels was investigated using thermo-gravimetric analysis (SETSYS Evolution 1750, SETARAM) with a sensibility of 5×10–5 g. Before heating the furnace to the test temperatures, the chamber was filled with pure nitrogen. When the actual temperature reached the set temperature, the synthetic exhaust gas of N2-10CO2-5H2O4
ACCEPTED MANUSCRIPT 0.5O2 (in vol.%) was introduced into the furnace at a total flow rate of 50 mlmin1 and a total pressure of 1 atm. The composition of synthetic exhaust gas is shown in Table 2, which is nearly the typical gasoline engine exhaust gas [9, 10, 18]. Samples were placed in the vertical tube furnace and isothermally oxidized for 5 h. The mass gain of the sample was continuously measured by a microbalance and the data was transmitted to the computer in real time. Finally, pure nitrogen was introduced in the furnace at a flow rate of 200 mlmin1, and the specimens cooled to room temperature. The water vapor was generated by water bath heated to 35 ºC, and the relative humidity of the mixed gas was 95%. All these parameters could be set and real-time monitored. The gas inlet line was heated to prevent condensation of water. A schematic of the experimental system is shown in Fig. 1. 2.3. Surface analysis The surface morphology of oxide scale was observed by scanning electron microscopy (SEM, FEI Quanta 600) equipped with an energy dispersive spectrometer (EDS). Surface 3D profiles of the oxide scales were examined by confocal laser scanning microscopy (VL2000DX SVF17SP/15FTC). The oxidation products were identified by using X-ray diffraction (XRD, MPDDY2094) with Cu Kα radiation source. The cross-section morphology and the elemental distribution in the oxide layer was studied by electron microprobe analysis (EPMA, JEOL JXA8530F). 3. Results and discussion 3.1 Oxidation kinetics Thermo-gravimetric kinetic curves of the experimental steels oxidized in simulated automotive exhaust gases are shown in Fig. 2. Oxidation temperature had a significant impact on the oxidation behavior of steels and the weight gain of experimental steels was increased 5
ACCEPTED MANUSCRIPT with test temperature. When the oxidation temperature was 950 ºC, the curve of F1 steel obeyed parabolic law during the first 125 min, then followed by a linear law. F2 steel also followed the parabolic law during the first 75 min and then obeyed linear law. F3 steel followed linear law during the period of 150–300 min. When the temperature reached 1000– 1050 ºC, F1, F2 and F3 steel followed parabolic law during the entire oxidation process. F4 steel obeyed parabolic law during the entire oxidation process in the temperature range of 950 ºC to 1050 ºC (Fig. 2a-c). However, the oxidation weight gain curves became complex (Fig. 2d) when the temperature approached 1100 ºC. The entire oxidation process can be divided into three different stages, including initial oxidation, rapid oxidation, and slow oxidation. But in the initial oxidation stage, the mass gain curves followed parabolic law, implying that the growth of oxide film was diffusion controlled [19]. The Ce-containing F2 steel displayed better oxidation resistance than F1 steel, especially when the temperature was above 1050 °C, such that the mass gain of F2 steel was nearly half of F1 steel. F3 steel had lower oxidation rate than F2 steel when the oxidation was below 1000 °C, and the mass gain was only 0.086 mg/cm2 at 950 °C. Nevertheless, F2 and F3 steels showed similar mass gain curves when tested at 1000–1050 °C. When the addition of W was 1.0 wt%, F4 steel did not exhibit superior oxidation resistance than F1 steel at 950–1050 °C. The high temperature oxidation kinetics of metals or alloys is generally controlled by the diffusion of cationic or anionic species through the oxide scale, and the oxide film growth process generally obeys parabolic law, which is defined by [20]: (Δm/s)2= C +kpt
(2)
where Δm is the weight gain in mg, s is surface area of specimen in cm2, kp is parabolic rate constant, expressed in mg2·cm4·s1，C is a constant, and t is the oxidation time with unit s. 6
ACCEPTED MANUSCRIPT However, the relationship of specific weight gain vs. time followed linear law when the oxide scale formed on the metal surface is not protective. This is may be due to oxide cracking or spalling, volatile or molten oxidation products formed during the oxide process. The linear law is given by the following relationship [20]: Δm/s= klt
(3)
where kl is linear rate constant with unit mg·cm2·s1. Fig. 3 shows detailed oxidation kinetics of coupons oxidized at 950 °C. Fig. 3a, c, d shows the parabolic plot of the square of the mass gain as a function of time. The parabolic rate constant kp was not always the same during the entire oxidation process. This is due to the change in the oxidation mechanism during the oxidation process. In statistics, R2 is the coefficient of determination, which represents how well statistical model fits a set of observations. Thus, the value of R2 was used to evaluate the fitness of the model to the data. When the value of R2 is close to 1, it suggests that the model fits the observed values very well. The oxidation kinetics of F1 and F2 steels followed a mixed parabolic-linear law. In the initial oxidation stage, the kp value of F1 and F2 steels was 1.40×106 mg2·cm4·s1 and 8.05×107 mg2·cm4·s1, respectively. After the initial period, the oxidation kinetics of F1, F2 steels obeyed linear law (Fig. 3b), but the corresponding kl value of F2 steel was nearly half of F1 steel. For F3 steel, the oxidation kinetics followed logarithmic law when the oxidation time was less than 5500 s, and subsequently followed linear law. We can conclude that F3 steel had best oxidation resistance capability at 950 °C. The parabolic or linear rate constant of experimental steels oxidized at 1000 °C and 1050 °C was also calculated. However, the oxide process was complex, when the temperature approached 1100 °C, and the corresponding parameters could not be derived. Table 3 shows 7
ACCEPTED MANUSCRIPT parabolic or linear oxidation rate constant for the tested steels oxidized at 9501050 °C in simulated exhaust gas. The results show that the oxidation rate constant of Ce-bearing F2 steel was about one order of magnitude lower than F1 steel, when the oxidation temperature was above 1000 °C. The reason is the “reactive element effect” on the addition of rare earth element. Many theories have been proposed to explain the various aspects of the reactive element effect. It is reported that rare earth element in the alloy can enhance the selective oxidation of chromium, reduce the growth rate of oxide film, change the transport mechanism and increase the resistance of scale to spallation [21]. Hou and Stringer [22] proposed that the segregation of RE along the oxide grain boundaries greatly retards outward transport cations through the scale. Seo et al. [23] put forward that the major role of REs is to make the oxidation process of chromia-rich scale the dominant inward diffusion of anions. In conclusion, the addition of rare earth element can significantly reduce the oxidation rate, such that the addition of Cerium in F2 steel led to its lower oxidation rate compared to Ce-free F1 steel. 3.2. Oxide scale surface morphology The macro-surface morphology of the specimens was observed to evaluate the degree of oxidation. After 5 h oxidation in simulated exhaust gas, all the specimen surfaces turned black. The oxide film of F1 and F4 steels showed spalling phenomenon. Fig. 4 shows micro-surface morphology of the oxide scale formed on the test steels oxidized at 950 °C for 5 h. The oxide film indicated spalling on F1 and F4 steel (Fig. 4a, d). The residual oxide film had a number of cracks and almost peeled off from the substrate. F2 and F3 steels did not experience spalling (Fig. 4b, c), and the oxide film displayed better adhesion. Moreover, the oxide particles were fine, dense and uniform, which could effectively 8
ACCEPTED MANUSCRIPT block the inward/outward diffusion of elements. As shown in the magnified images in Fig. 4a, the oxide on F1 steel (~8.9μm) was greater than F2, F3 and F4 steels (~3.6 μm, 4.1 μm and 4.9 μm, respectively), and in the spalled area, there existed many particles. Kuhn [24] studied microstructural features, mechanical and steam oxidation properties of ferritic stainless steel containing 18–23% chromium and addition of niobium and tungsten. They determined the chemical composition of the precipitates, which were rich in niobium and tungsten, and suggested that the phase is probably a (Fe, Cr, Si)2(Nb, W) Laves phase. Silva et al. [25] studied the in heat affected zone of AISI 444 ferritic stainless steel. They also proposed that the phase present in the HAZ was Laves phase. Thus, from the EDS analysis of these particles, we deduced that they were mainly Laves phase (Fe2(Nb, Mo, W)). The precipitates formed in F4 steel were larger and more in number than in F1 steel. Laves phase is intermetallic phase of the type AB2, and in some instances other alloying elements such as Cr, Si or Mo may also be present in the phase [26, 27]. Laves phase is very hard and brittle at ambient temperature, and the difference in the thermal expansion coefficient between Laves phase and matrix is large [28-30]. Thus, the precipitation of a large number of Laves phase at the oxide scale/metal interface is not beneficial for the adhesion of oxide film. Fig. 5 shows the micro-surface morphology of the oxide film formed at 1100 °C for 5 h. The size of oxide was increased compared to oxidation at 950 °C. Oxide formed on F2 and F3 steels was spherical in shape and uniformly distributed. However, oxide formed on F1 and F4 steel had strip-like morphology. The three dimensional profile of oxide scale surface after 5 h oxidation at 1100 °C is shown in Fig. 6. The surface fluctuations on F1 and F4 steels was large, especially F4 steel. The change in the thickness of oxide film is not beneficial to release growth stresses and thermal stresses. The accumulation of strain energy at the oxide/metal 9
ACCEPTED MANUSCRIPT interface will lead to the spallation of oxide film [31]. Morphology at high magnification is presented in the inset of Fig. 5. According to the XRD data in Fig. 7, the surface oxide mainly consisted of Cr-Mn spinel oxide. The morphology of MnCr2O4 spinel was octahedral in shape, which is shown as red circle in Fig. 5d. But in F1 steel tetragonal prism-shaped crystallites (marked by red circle in Fig. 5a) was present. This phenomenon was also observed in other studies [32, 33], and is due to MnxCr3-xO4 spinel crystal system that changes from cubic to tetragonal when x˃1.17. Thus, the MnCr2O4 spinel oxide in F1 steel may have partially decomposed. 3.3 Identification of oxide scale phase XRD analysis was conducted to determine phases present in the surface oxide film. Fig. 7 shows the XRD pattern of oxidized layer after the ferritic stainless steels was subjected to isothermal oxidation for 5h at 1100 ºC in simulated exhaust gas. The results revealed that the surface oxide film mainly composed of corundum M2O3 and spinel M3O4 type oxides, where M may be Fe, Cr, Mn or combination of them. Because the ionic radii of Fe, Cr and Mn are very similar, it is therefore difficult to determine the exact composition of the oxide film by the position of the peaks in the XRD patterns [34, 35]. But from the EDS analysis and EPMA elemental mapping of the cross-section of oxide film, we concluded that the oxide products mainly comprised of Cr2O3, (Mn, Cr)3O4 and small amount of SiO2. Fe-Cr substrate was also detected and is due to the oxide film being thin and the X-rays penetrated the film. 3.4 Cross-section analysis 3.4.1 Cross-sectional morphology Fig. 8 shows the cross-sectional morphology of the experimental steels after 5 h of oxidation in simulated exhaust gas obtained by back-scattered electron (BSE). The oxide film formed 10
ACCEPTED MANUSCRIPT on all the specimens can be divided into two parts: an outer layer of (Mn, Cr, Fe)3O4 spinel oxides and inner layer of Cr-rich (mainly Cr2O3) oxides. The oxide film formed on F1 steel is shown in Fig. 8a. Transverse cracks and small longitudinal cracks were present in the oxide film, which can lead to spallation of oxide film. In addition, many small pores existed in the oxide film (black circle in Fig. 8a). The cluster of pores coalesce and facilitate initiation of cracks [36]. In addition, at the oxide/metal interface gaps were present, which may reduce the adhesion between oxide scale and the matrix. The oxide film formed on F2 and F3 steels was dense, uniform and straight, and there was near absence of defects in the oxide film. Spherical oxide particle could be clearly seen on the oxide film surface of F2 steel, which was mainly composed of Cr-Mn spinel oxides. Fig. 8d shows the oxide film formed on F4 steel, which displayed notable difference compared to other steels. The oxide film was porous and undulated, a number of large cavities (marked by black arrow in Fig. 8d) and cracks were present in the oxide scale. It is worth noting that a thin layer of oxide near the matrix of F4 steel was compact, but the outer layer was loose. The pores and micro-cracks act as diffusivity paths allowing H2O molecules or O2 to easily penetrate in the oxide film, and H2O can provide a means of oxygen transport by dissociation through the following reaction [37, 38] H2O=H2+1/2O2
(4)
The partial pressure of H2O in the cavity approaches the external gas when the inward H2O transport is relatively fast, and the recycling of H2/H2O couple provides continuous transport of O2 [39]. In the case of F4 steel (Fig. 8d), the oxide film had many cracks and cavities which may lead to continuous oxidation and formation of a less protective oxide film. A large number of studies indicated that the presence of water vapor in the exposure environment can significantly influence the oxidation behavior of the chromia-forming 11
ACCEPTED MANUSCRIPT materials as compared to in dry atmosphere, e.g. oxygen or air [14]. In the case of Fe-Cr alloys with intermediate Cr content (10–20 wt%), wet gases can lead to degradation behavior of the oxide film by growing Fe-rich oxide scales [40]. However, there is still not a complete understanding of the mechanism behind this phenomenon. Majority of the studies attribute the breakaway oxidation of Fe-Cr alloys in water vapor containing environment to the formation of volatile chromium oxohydroxide CrO2(OH)2 [11,12]. Ehlers et al. [40] studied the oxidation behavior of 9% Cr steels in N2-O2-H2O gas mixture. They proposed that the penetration of water vapor molecules triggers enhanced oxidation and the tendency of the formation of non-protective oxide scale increased with increasing H2O(g)/O2 ratio. In our study, the partial pressure of O2 and H2O was 0.005 and 0.05 atm, respectively. All the experimental steels did not experience breakaway oxidation behavior in the temperature range of 950–1100ºC. This may due to the formation of (Mn, Cr, Fe)3O4 spinel oxide layer on the Cr2O3 scale. Fig. 9 shows the thickness and the mass gain per unit volume of the oxide film after 5 h oxidation at 1100 ºC. The oxide film thickness (Fig. 9a) can be obtained from the crosssectional morphology, and the thickness of oxide scale formed on F1, F2, F3 and F4 steels was ~8.1, 3.5, 4.7 and 10.1 μm, respectively. Combining with the mass gain data (Fig. 2d), the average mass gain per unit volume of the oxide film was calculated (Fig. 9b), and the results are 3.2, 4.3, 4.5 and 2.0 g·cm-3, respectively. If the formation of volatile chromium oxide and other situations are ignored, the weight gain is mainly from oxygen. Based on the results, we can say that the addition of small amount of Ce can significantly reduce the thickness and improve the density of oxide film. This phenomenon has been previously observed [3, 41]. 3.4.2 Elemental distribution in the oxide film 12
ACCEPTED MANUSCRIPT Elemental distribution involving the oxide layer and the substrate obtained by EPMA elemental mapping is shown in Fig. 10. The variation in elemental concentration is represented by change of color. Manganese was distributed in the outermost surface of the oxide film, and existed mainly in the form of spinel oxide (Mn, Cr)3O4. A compact and continuous layer of (Mn, Cr)3O4 spinel oxide doped with a small amount of iron formed in the top layer of Cr2O3, which effectively prevents the volatilization of chromium. Previous studies also showed that the outermost Cr-Mn spinel oxide can strongly reduce CrO2(OH)2 volatilization [42, 43]. Sachitanand et al. and Stanislowski et al. [44, 45] have shown that steel containing 0.30.5 wt% Mn develops an adherent (Mn, Cr)3O4 top layer above the Cr2O3 layer at 800 ºC and 850 ºC, and that the rate of Cr vaporization for these alloys was 2-3 times lower than alloys that form pure Cr2O3 scale or a non-continuous (Mn, Cr)3O4 top layer. The Mn content of the experimental steels in our study was greater than 0.3 wt. %, and results showed that a compact (Mn, Cr)3O4 layer formed on all the steels. Silicon was distributed at the oxide/metal interface and can be seen in Fig. 10. The presence of Si in stainless steel generally existed in the form of a silica interface between the external chromia layer and the base metal and acts as a barrier for cation transport. However, the existence of silicon layer in stainless steel enhances scale spallation during cyclic tests [46]. As shown in Fig. 10, silicon layer in F1 and F4 steels was more obvious than in F2 and F3 steel, and the discontinuous silicon layer was porous, which may lead to spallation of oxide film on F1 and F4 steels. The oxygen distributed in the oxide film had a concentration gradient, especially when the oxide film was thick. Near the oxide/metal interface, the oxygen concentration was lower. This implied that the oxidation process was diffusion-controlled. In the oxidation process, the 13
ACCEPTED MANUSCRIPT O–2 diffused through the oxide film from the gas/oxide interface to the oxide/metal interface and the metal cations diffused toward the opposite direction. The growth rate of oxide film was affected by concentration gradient and diffusion rate of the reactive elements. When the oxide scale becomes thicker, the diffusion velocity as well as the growth rate of the oxide film will decline. Chromium was distributed between the surface and the substrate. Cr depletion zone was not observed near the oxide/metal interface because of large residual Cr in the alloy. The rare-earth element in the oxide scale could not be detected because of very small addition of Ce (0.05 wt%) in the steels. The mechanism of Ce on the significant improvement of the high temperature oxidation resistance of Fe-Cr alloys is still unclear. But from the results of the experiments, it can be concluded that the addition of Ce can reduce the oxidation rate of the alloys at higher temperatures (above 1000 ºC) and improve the density of the oxide film. The oxide scale displayed better adhesion and protective nature in wet gas. The content of W in the alloys also had a dramatic influence on the high temperature oxidation behavior of the alloys. The addition of Ce together with ~0.5 wt% W can improve the oxidation resistance of the alloy compared to only addition Ce at temperatures below 1000 ºC. However, the oxide film experienced substantial spalling phenomenon during cooling to room temperature, when the addition was ~1.0 wt% W. The high content of W promoted the coarsing Laves phase precipitate at the oxide/scale interface, which can significantly reduce the adhesion of oxide scale. In addition, the higher content of W also led to the formation of porous oxide film on the alloy. All of these factors together resulted in the poor oxidation resistance of the oxide scale. 4. Conclusions The isothermal oxidation behavior of ferritic stainless steel was studied at 9501100 ºC in 14
ACCEPTED MANUSCRIPT simulated exhaust gas environment. Based on the results, the following are the conclusions: (1) The oxidation growth rate of B444M2 steel was significantly reduced on the addition of Ce or combination of Ce and 0.5 wt% W. They followed a mixed parabolic-linear law at 950 ºC and parabolic law at 1000 or 1050 ºC. (2) Multi-layer oxide film formed on all the steels. The surface layer was mainly (Mn, Cr)3O4 spinel oxide, which can effectively reduce Cr volatilization, and the layer beneath was Cr2O3 oxide. (3) The oxide formed on the steels was fine, dense and uniform with the addition of Ce, and the oxide film was thin, smooth and displayed better adhesion. The number of defects in the oxide scale was decreased. (4) Many cracks and pores were present in the oxide film of F1 and F4 steels. The silicon layer at the oxide/metal interface was porous, and there were many Laves phases precipitated at the oxide/metal interface at 950 ºC. All these factors may lead to severe spallation of the oxide film on F1 and F4 steels. The addition of 1.0 wt% W was not beneficial to the oxidation resistance.
Acknowledgements This work was financially supported by the National Natural Science Foundation of China (Grant No. U1660205). R.D.K. Misra gratefully acknowledges continued collaboration with Northeastern University as an Honorary Professor in providing guidance to the students.
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ACCEPTED MANUSCRIPT [26] T. Juuti, L. Rovatti, A. Mäkelä, L.P. Karjalainen, D. Porter, Influence of long heat treatments on the Laves phase nucleation in a type 444 ferritic stainless steel, J. Alloys Compd. 616 (2014) 250-256. [27] N. Takata, H. Ghassemi-Armaki, M. Takeyama, S. Kumar, Nanoindentation study on solid solution softening of Fe-rich Fe2Nb Laves phase by Ni in Fe-Nb-Ni ternary alloys, Intermetallics 70 (2016) 7-16. [28] Y. Kimura, D.E. Luzzi, D.P. Pope, Deformation twinning in a HfV2+Nb-based Laves phase alloy, Mater. Sci. Eng. A 329 (2002) 241-248. [29] B. Mayer, H. Anton, E. Bott, M. Methfessel, J. Sticht, J. Harris, P.C. Schmidt, Ab-initio calculation of the elastic constants and thermal expansion coefficients of Laves phases, Intermetallics 11 (2003) 23-32. [30] A.F. Padiliha, I.F. Machado, R.L. Plaut. Microstructures and mechanical properties of Fe-15% Cr-15% Ni austenitic stainless steels containing different levels of niobium additions submitted to various processing stages, J. Mater. Process. Tech. 170 (2005) 89-96. [31] A. Hayashi, N. Hiraide, Y. Inoue, Spallation Behavior of Oxide Scale on Stainless Steels, Oxid. Met. 85 (2016) 87-101. [32] H. Li, W.X. Chen, Stability of MnCr2O4 spinel and Cr2O3 in high temperature carbonaceous environments with varied oxygen partial pressures, Corros. Sci. 52 (2010) 2481-2488. [33] A. Naumidis, H.A. Schulze, W. Jungen, P. Lersch, Phase studies in the chromiummanganese-titanium oxide system at different oxygen partial pressures, J. Eur. Ceram. Soc. 7 (1991) 55-63. [34] Z.G. Yang, G.G. Xia, P. Singh, J.W. Stevenson, Effects of water vapor on oxidation behavior of ferritic stainless steels under solid oxide fuel cell interconnect exposure conditions, Solid State Ionics 176 (2005) 1495-1503. [35] X.W. Cheng, Z.Y. Jiang, D.B. Wei, J.W. Zhao, B.J. Monaghan, R.J. Longbottom, L.Z. Jiang, Charastics of oxide scale formed on ferritic stainless steels in simulated reheating atmosphere, Surf. Coat. Technol. 258 (2014) 257-267. [36] J.S. Zheng, X.M. Hou, X.B. Wang, Y. Meng, X. Zheng, L. Zheng, Isothermal oxidation mechanism of Nb-Ti-V-Al-Zr alloy at 700-1200 °C: diffusion and interface reaction, Corros. Sci. 96 (2015) 186-195. [37] P.L. Surman, J.E. Castle, Gas phase transport in the oxidation of Fe and steel, Corros. 18
ACCEPTED MANUSCRIPT Sci. 9 (1969) 771-777. [38] C.T. Fuji, R.A. Meussner, Oxide structure produced on iron chromium alloys by a dissociative mechanism, J. Electrochem. Soc. 110 (1963) 1195-1204. [39] N.K. Othman , N. Othman , J. Zhang , D.J. Young, Effects of water vapour on isothermal oxidation of chromia-forming alloys in Ar/O2 and Ar/H2 atmospheres, Corros. Sci. 51 (2009) 3039-3049. [40] J. Ehlers, D.J. Young, E.J. Smaardijk, A.K. Tyagi, H.J. Penkalla, L. Singheisera, W.J. Quadakkers, Enhanced oxidation of the 9%Cr steel P91 in water vapour containing environments, Corros. Sci. 48 (2006) 3428-3454. [41] J. Piekoszewski, B. Sartowska, M. Barlak, P. Konarski, L. Dąbrowski, L. Dąbrowski, W. Starosta, L. Walis, Z. Werner, C. Pochrybniak, K. Bochenska, P. Stoch, W. Szymczyk, Improvement of high temperature oxidation resistance of AISI 316L stainless steel by incorporation of Ce-La elements using intense pulsed plasma beams, Surf. Coat. Technol. 206 (2011) 854-858. [42] X.W. Cheng, Z.Y. Jiang, B.J. Monaghana, D.B. Wei, R.J. Longbottoma, J.W. Zhao, J.P. Peng, M. Luo, L. Ma, S.Z. Luo, L.Z. Jiang, Breakaway oxidation behaviour of ferritic stainless steels at 1150 ºC in humid air, Corros. Sci. 108 (2016) 11-12. [43] G.R. Holcomb, D.E. Alman, The effect of manganese additions on the reactive evaporation of chromium in Ni-Cr alloys, Scripta Mater. 54 (2006) 1821-1825. [44] R. Sachitanand, M. Sattari, J.E. Svensson, J. Froitzheim, Evaluation of the oxidation and Cr evaporation properties of selected FeCr alloys used as SOFC interconnects, Int. J. Hydrogen Energy 38 (2013) 15328-15334. [45] M. Stanislowski, E. Wessel, K. Hilpert, T. Markus, L. Singheiser, Chromium vaporization from high-temperature alloys, J.
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ACCEPTED MANUSCRIPT Table captions Table 1.
Chemical composition of ferritic stainless steels (in wt%).
Table 2.
The composition of the synthetic exhaust gas.
Table 3. gases.
Values of kp obtained at 950 – 1050 °C for test steels oxidized in simulated exhaust
20
ACCEPTED MANUSCRIPT Figure captions Fig. 1.
Schematic of the experimental system for oxidation studies.
Fig. 2.
Mass change versus time curves for the test steels oxidized for 300 min at different
temperatures in simulated synthetic exhaust gases. (a) 950 °C, (b) 1000 °C, (c) 1050 °C, and (d) 1100 °C. Fig. 3.
TGA curves of the test coupons oxidized at 950 °C for 18000 s in simulated
synthetic exhaust gas. (a) parabolic plot (Δm/s)2–t of F1 steel, (b) Δm/s–t plots of the F2 and F3 steels, (c) parabolic plot (Δm/s)2–t of F2 steel, and (d) parabolic plot (Δm/s)2–t of F4 steel. Fig. 4.
Surface morphology of the specimens after 5 h oxidation at 950 °C in synthetic
exhaust gases. (a) F1, (b) F2, (c) F3, and (d) F4. Fig. 5.
Surface morphology of the specimens after 5 h oxidation at 1100 °C in synthetic
exhaust gas. (a) F1, (b) F2, (c) F3, and (d) F4. Fig. 6.
Three dimensional profiles of the oxide scale surfaces after 5 h oxidation at 1100 °C
in synthetic exhaust gas. (a) F1, (b) F2, (c) F3, and (d) F4. Fig. 7.
XRD patterns of oxidation layer after the ferritic stainless steels isothermal oxidation
for 5 h at 1100 ºC in simulated exhaust gas. Fig. 8.
BSE morphology of the cross-section of the test steels oxidized at 1100 ºC for 5 h.
(a) F1, (b) F2, (c) F3, and (d) F4. Fig. 9.
The thickness and the mass gain per unit volume of the oxide films after 5 h
oxidation at 1100 ºC in simulated exhaust gas. Fig. 10.
BSE micrographs and EPMA element maps of oxide scales of the experimental
steels oxidized at 1100 ºC for 5 h in simulated exhaust gas.
21
ACCEPTED MANUSCRIPT Table 1. Chemical composition of ferritic stainless steels (in wt%). Steel
C
Si
Mn
P
S
Cr
Nb
Ti
N
Mo
Ce
W
Fe
F1
0.006
0.54
0.33
0.006
0.006
19.7
0.436
0.174
0.072
2.09
Bal.
F2
0.009
0.52
0.32
0.008
0.008
19.5
0.450
0.155
0.072
1.97
0.048
Bal.
F3
0.007
0.50
0.32
0.010
0.009
19.1
0.441
0.137
0.071
1.95
0.056
0.58
Bal.
F4
0.010
0.53
0.35
0.008
0.010
19.9
0.443
0.113
0.071
2.03
0.053
1.12
Bal.
22
ACCEPTED MANUSCRIPT Table 2. The composition of the synthetic exhaust gas. Gas
O2
H 2O
CO2
N2
Vol.%
0.5
5
10
Bal.
23
ACCEPTED MANUSCRIPT Table 3.
Values of kp obtained at 950–1050 °C for the experimental steels oxidized in
simulated exhaust gases. Steel grades
Temperature (°C)
Kinetics of oxidation
Rate constant values
R2
Parabolic (kp) (0–7500s)
1.40×10–12 g2·cm4·s1
0.98
Linear (kl) (7500–18000s)
1.57×10–8 g·cm2·s1
0.99
1000
Parabolic (kp)
3.21×10–12 g2·cm4·s1
0.99
1050
Parabolic (kp)
1.97×10–8 g2·cm4·s1
0.99
Parabolic (kp) (0–4500s)
8.05×10–13 g2·cm4·s1
0.99
Linear (kl) (4500–18000s)
7.68×10–9 g·cm2·s1
0.99
1000
Parabolic (kp)
1.29×10–11 g2·cm4·s1
0.99
1050
Parabolic (kp)
5.30×10–9 g2·cm4·s1
0.99
950
Linear (kl) (9000–18000s)
5.28×10–9 g·cm2·s1
0.99
1000
Parabolic (kp)
1.74×10–11 g2·cm4·s1
0.99
1050
Parabolic (kp)
6.70×10–9 g2·cm4·s1
0.99
Parabolic (kp)
1.50×10–12 g2·cm4·s1
0.99
Parabolic (kp)
4.23×10–12 g2·cm4·s1
0.99
1000
Parabolic (kp)
3.04×10–11 g2·cm4·s1
0.99
1050
Parabolic (kp)
7.50×10–9 g2·cm4·s1
0.99
950 F1
950 F2
F3
950 F4
Note: The time intervals 0–7500s and 0–4500s represent that during this period of time the curves obey parabolic law. The time intervals 7500–18000s, 4500–18000s and 9000–18000s represent that the curves follow linear law during these period.
24
ACCEPTED MANUSCRIPT
Fig. 1.
Schematic of the experimental system for oxidation studies.
25
ACCEPTED MANUSCRIPT (a)
(b)
(d)
(c)
Fig. 2.
Mass change versus time curves for the test steels oxidized for 300 min at different
temperatures in simulated synthetic exhaust gases. (a) 950 °C, (b) 1000 °C, (c) 1050 °C, and (d) 1100 °C.
26
ACCEPTED MANUSCRIPT
Fig. 3.
(a)
(b)
(c)
(d)
TGA curves of the test coupons oxidized at 950 °C for 18000 s in simulated
synthetic exhaust gas. (a) parabolic plot (Δm/s)2–t of F1 steel, (b) Δm/s–t plots of the F2 and F3 steels, (c) parabolic plot (Δm/s)2–t of F2 steel, and (d) parabolic plot (Δm/s)2–t of F4 steel.
27
ACCEPTED MANUSCRIPT (a)
(b)
20 μm
4 μm
Spalled area
40 μm
400 μm
(c)
(d)
4 μm
20 μm
Spalled area 400 μm
40 μm
Fig. 4.
Surface morphology of the specimens after 5 h oxidation at 950 °C in synthetic
exhaust gases. (a) F1, (b) F2, (c) F3, and (d) F4.
28
ACCEPTED MANUSCRIPT (a)
(b)
2 μm
2 μm
40 μm
40 μm
(c)
(d)
2 μm
2 μm
. . 2 2 7 7 0 40 μm 40 μm 0 . . Fig. 5. Surface morphology2 of the specimens after 5 h oxidation2 at 1100 °C in synthetic 7 exhaust gas. (a) F1, (b) F2, (c)7 F3, and (d) F4.
29
ACCEPTED MANUSCRIPT (a)
(b)
(c)
(d)
Fig. 6.
Three dimensional profiles of the oxide scale surfaces for the experimental steels
after 5 h oxidation at 1100 °C in synthetic exhaust gas. (a) F1, (b) F2, (c) F3, and (d) F4.
30
ACCEPTED MANUSCRIPT
Fig. 7.
XRD patterns of oxidation layer after the ferritic stainless steels isothermal oxidation for 5 h at 1100 ºC in simulated exhaust gas.
31
ACCEPTED MANUSCRIPT (b)
(a) Resin
Cr-Mn spinel Layer Transversal Cracks
Cr2O3 Layer
Longitudinal Cracks
Resin
Cr-Mn spinel Layer
Cr2O3 Layer
gap
Substrate
Substrate 10 μm
10 μm
(d)
(c) Resin
Resin
Cr-Mn spinel Layer
Cr2O3 Layer
Cr2O3 Layer Substrate
Fig. 8.
Cr-Mn spinel Layer
Substrate
10 μm
20 μm
BSE morphology of the cross-section of the experimental steels oxidized at 1100 ºC
for 5 h. (a) F1, (b) F2, (c) F3, and (d) F4.
32
ACCEPTED MANUSCRIPT (b)
(a)
Fig. 9.
The thickness (a) and the mass gain per unit volume (b) of the oxide film on the
surface of the steels after 5 h oxidation at 1100 ºC in simulated exhaust gas.
33
ACCEPTED MANUSCRIPT
(a)
O
Cr
Mn
Fe
Si
O
Cr
Mn
Fe
Si
O
Cr
Mn
Fe
Si
O
Cr
Mn
Fe
Si
10 μm
(b)
10 μm
(c)
10 μm
(d)
10 μm
Low
High Element Concentration
Fig. 10.
BSE micrographs and EPMA element maps of oxide scales of the experimental steels oxidized at 1100 ºC for 5 h in simulated exhaust gas.
34
L.L. Wei, L.Q. Chen, M.Y. Ma, H.L. Liu, R.D.K. Misra PII:
S0254-0584(17)30931-8
DOI:
10.1016/j.matchemphys.2017.11.051
Reference:
MAC 20171
To appear in:
Materials Chemistry and Physics
Received Date:
27 June 2017
Revised Date:
25 October 2017
Accepted Date:
25 November 2017
Please cite this article as: L.L. Wei, L.Q. Chen, M.Y. Ma, H.L. Liu, R.D.K. Misra, Oxidation Behavior of Ferritic Stainless Steels in Simulated Automotive Exhaust Gas Containing 5 vol.% Water Vapor, Materials Chemistry and Physics (2017), doi: 10.1016/j.matchemphys.2017.11.051
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ACCEPTED MANUSCRIPT
Highlights
Ce and W were added in Fe-Cr alloy to improve its high temperature performance.
Oxidation kinetics displayed significantly different with adding Ce and W at 950–1100 ºC.
The property of oxide film was obvious improved with adding ~0.05 wt% Ce in steel.
Laves phase and porous SiO2 caused the oxide film spalling when adding ~1.0 wt% W.
ACCEPTED MANUSCRIPT
Oxidation Behavior of Ferritic Stainless Steels in Simulated Automotive Exhaust Gas Containing 5 vol.% Water Vapor L.L. Wei1), L.Q. Chen1, *), M.Y. Ma1), H.L. Liu1) and R.D.K. Misra2) 1) State Key Laboratory of Rolling and Automation, Northeastern University, 3-11 Wenhua Road, Shenyang 110819, China 2) Department of Metallurgical, Materials and Biomedical Engineering, University of Texas at El Paso, El Paso, Texas 79968, USA. * Corresponding author. Tel.: +86-24-83681819, Fax: +86-24-23906472, E-mail address: [email protected]
Abstract In this study, the effect of alloying elements, cerium (Ce) and tungsten (W), on the oxidation behavior of medium chromium ferritic stainless steel in simulated automotive exhaust gases containing 5 vol.% H2O was studied in the temperature range of 950–1100 °C for 5 h. The oxidation kinetics and oxide film characteristics were analyzed by means of thermogravimetric analysis (TGA), X-ray diffraction (XRD), scanning electron microscope (SEM) and elemental probe micro-analyzer (EPMA). The growth rate of oxidation and oxidation mass gain were significantly decreased on the addition of Ce. A dense, uniform, and thin oxide scale formed on Ce-containing steels. The steel containing Ce and ~0.5 wt% W displayed similar oxidation behavior compared to the steel containing only Ce. A large number of cracks and pores existed in the oxide film and oxide/metal interface leading to the spallation of oxide film, when the addition of W reached ~1.0 wt%. Key words: Ferritic stainless steel; Oxidation; Rare earth element; Tungsten; Water vapor
1
ACCEPTED MANUSCRIPT 1. Introduction Ferritic stainless steels with single phase ferritic structure, and generally free of nickel, have
excellent corrosion resistance and low cost making them a promising material to
substitute austenitic stainless steel, for use in storage vessels, petroleum refining equipment and furnace parts, etc. In addition, the characteristic of low thermal expansion coefficient, good thermal conductivity and oxidation resistance render ferritic stainless steels as viable materials for manufacturing automotive exhaust pipes. Automobile exhaust manifold is the hot-end of exhaust system connected directly with the engine. The operating environment is generally ~900 °C. However, in order to satisfy the increasing stringent auto emissions standards, the temperature of automobile exhaust gas may reach 950–1100 °C [1, 2]. Thus, materials used for manifold are required to have excellent high temperature oxidation resistance and thermal fatigue resistance. According to the previous study, adding rare earth element is an effective way to improve the high temperature oxidation resistance of ferritic stainless steel because of its “reactive element effect” [3, 4]. The creep rupture strength of the steel can be enhanced by adding tungsten (W) through precipitation and/or solid-solution strengthening effect [5–7]. Kim et al. [8] reported that the addition of W can also retard nucleation and growth of σ phase during aging, thereby delaying degradation of corrosion resistance and mechanical properties of 25% chromium duplex stainless steel. Based on this idea, the chemical composition of ferritic stainless steel was redesigned by adding alloying elements, cerium (Ce) and tungsten (W). Ferritic stainless steels have excellent high temperature oxidation resistance, because preferential oxidation of chromium leads to a thin layer of Cr-rich oxide that inhibits oxidation reaction. The formation of oxide scale is influenced by temperature, alloy 2
ACCEPTED MANUSCRIPT composition, exposure time and especially the atmosphere. The atmosphere in the internal surface of manifold is automotive exhaust gas. Gasoline blend is a mixture of paraffin and aromatic hydrocarbons undergoes combustion in air by the following reaction [9]: gasoline + O2 (in air) → CO2 + H2O + heat
(1)
In addition to water vapor and CO2, automotive exhaust emission also contains a small amount of O2, sulfur dioxide (SO2), nitrogen oxides (NOx), carbon monoxide (CO), and unburned hydrocarbon (HC) [10]. But among of them, water vapor and O2 play the most important role in the high temperature oxidation of ferritic stainless steel. Previous studies have demonstrated that the presence of water vapor in the exposure environment can have a dramatic effect on the oxidation behavior of chromia-forming alloys as compared to the dry condition, involving rapid and catastrophic oxidation. Various mechanisms have been proposed to account for the degradation behavior of Fe-Cr alloys oxidized in water vapor containing environment. At high partial pressure of oxygen and water vapor, the formation of vaporized oxyhydroxides CrO2(OH)2 leads to less protection of oxide film [11, 12]. Rahmel and Tobolski [13] proposed that H2-H2O mixture is formed in voids at the scale/metal interface and/or within the scale, which dissociates leading to inward transport oxygen. Essuman et al. [14] suggested that when the water vapor is the source of oxygen, the internal oxidation is accelerated, and water vapor affects the solubility and/or diffusivity of oxygen in the alloy. Yuan et al. [15] proposed that the stress induced by accumulation of molecular hydrogen in the pores introduces cracking of chromia scale in pure steam. Water vapor can also enhance scale growth stress, the scale experiences frequent cracking and healing processes [16, 17]. In order to improve the high temperature properties of the manifold used ferritic stainless 3
ACCEPTED MANUSCRIPT steel, alloying elements Ce and W are added. But the role of alloying elements on high temperature oxidation resistance of the new developed ferritic stainless steel in water vapor containing automotive exhaust gas environment continue to be unclear. Thus, the aim of this study is to explore the influence of different content of alloying elements, Ce and W, on the high temperature corrosion behavior of ferritic stainless steel at 950–1100 ºC in simulated automotive exhaust gas environment. 2. Experimental method 2.1. Materials Four ferritic stainless steels were used in this study, and are named as F1, F2, F3 and F4. The F1 steel is an existing ferritic stainless steel (grade B444M2). In F2, F3 and F4 steels, different content of alloying elements, Ce and W, were added to F1 steel. The nominal chemical composition of these steels is presented in Table. 1. The difference among them is the content of W and Ce. Specimens of dimensions of 8 mm×12 mm×1 mm were fabricated from a cold-rolled plate with a small hole of 2 mm in diameter near one edge. F1 steel was annealed at 1020 °C for 2 min and F2, F3, and F4 steels were annealed at 1050 °C for 2 min, followed by air cooling. Prior to oxidation experiments, all the specimens were metallographically polished with SiC paper up to 1200 grit, ultrasonically cleaned, and stored in a desiccator. 2.2. Oxidation experiments The oxidation kinetics of ferritic stainless steels was investigated using thermo-gravimetric analysis (SETSYS Evolution 1750, SETARAM) with a sensibility of 5×10–5 g. Before heating the furnace to the test temperatures, the chamber was filled with pure nitrogen. When the actual temperature reached the set temperature, the synthetic exhaust gas of N2-10CO2-5H2O4
ACCEPTED MANUSCRIPT 0.5O2 (in vol.%) was introduced into the furnace at a total flow rate of 50 mlmin1 and a total pressure of 1 atm. The composition of synthetic exhaust gas is shown in Table 2, which is nearly the typical gasoline engine exhaust gas [9, 10, 18]. Samples were placed in the vertical tube furnace and isothermally oxidized for 5 h. The mass gain of the sample was continuously measured by a microbalance and the data was transmitted to the computer in real time. Finally, pure nitrogen was introduced in the furnace at a flow rate of 200 mlmin1, and the specimens cooled to room temperature. The water vapor was generated by water bath heated to 35 ºC, and the relative humidity of the mixed gas was 95%. All these parameters could be set and real-time monitored. The gas inlet line was heated to prevent condensation of water. A schematic of the experimental system is shown in Fig. 1. 2.3. Surface analysis The surface morphology of oxide scale was observed by scanning electron microscopy (SEM, FEI Quanta 600) equipped with an energy dispersive spectrometer (EDS). Surface 3D profiles of the oxide scales were examined by confocal laser scanning microscopy (VL2000DX SVF17SP/15FTC). The oxidation products were identified by using X-ray diffraction (XRD, MPDDY2094) with Cu Kα radiation source. The cross-section morphology and the elemental distribution in the oxide layer was studied by electron microprobe analysis (EPMA, JEOL JXA8530F). 3. Results and discussion 3.1 Oxidation kinetics Thermo-gravimetric kinetic curves of the experimental steels oxidized in simulated automotive exhaust gases are shown in Fig. 2. Oxidation temperature had a significant impact on the oxidation behavior of steels and the weight gain of experimental steels was increased 5
ACCEPTED MANUSCRIPT with test temperature. When the oxidation temperature was 950 ºC, the curve of F1 steel obeyed parabolic law during the first 125 min, then followed by a linear law. F2 steel also followed the parabolic law during the first 75 min and then obeyed linear law. F3 steel followed linear law during the period of 150–300 min. When the temperature reached 1000– 1050 ºC, F1, F2 and F3 steel followed parabolic law during the entire oxidation process. F4 steel obeyed parabolic law during the entire oxidation process in the temperature range of 950 ºC to 1050 ºC (Fig. 2a-c). However, the oxidation weight gain curves became complex (Fig. 2d) when the temperature approached 1100 ºC. The entire oxidation process can be divided into three different stages, including initial oxidation, rapid oxidation, and slow oxidation. But in the initial oxidation stage, the mass gain curves followed parabolic law, implying that the growth of oxide film was diffusion controlled [19]. The Ce-containing F2 steel displayed better oxidation resistance than F1 steel, especially when the temperature was above 1050 °C, such that the mass gain of F2 steel was nearly half of F1 steel. F3 steel had lower oxidation rate than F2 steel when the oxidation was below 1000 °C, and the mass gain was only 0.086 mg/cm2 at 950 °C. Nevertheless, F2 and F3 steels showed similar mass gain curves when tested at 1000–1050 °C. When the addition of W was 1.0 wt%, F4 steel did not exhibit superior oxidation resistance than F1 steel at 950–1050 °C. The high temperature oxidation kinetics of metals or alloys is generally controlled by the diffusion of cationic or anionic species through the oxide scale, and the oxide film growth process generally obeys parabolic law, which is defined by [20]: (Δm/s)2= C +kpt
(2)
where Δm is the weight gain in mg, s is surface area of specimen in cm2, kp is parabolic rate constant, expressed in mg2·cm4·s1，C is a constant, and t is the oxidation time with unit s. 6
ACCEPTED MANUSCRIPT However, the relationship of specific weight gain vs. time followed linear law when the oxide scale formed on the metal surface is not protective. This is may be due to oxide cracking or spalling, volatile or molten oxidation products formed during the oxide process. The linear law is given by the following relationship [20]: Δm/s= klt
(3)
where kl is linear rate constant with unit mg·cm2·s1. Fig. 3 shows detailed oxidation kinetics of coupons oxidized at 950 °C. Fig. 3a, c, d shows the parabolic plot of the square of the mass gain as a function of time. The parabolic rate constant kp was not always the same during the entire oxidation process. This is due to the change in the oxidation mechanism during the oxidation process. In statistics, R2 is the coefficient of determination, which represents how well statistical model fits a set of observations. Thus, the value of R2 was used to evaluate the fitness of the model to the data. When the value of R2 is close to 1, it suggests that the model fits the observed values very well. The oxidation kinetics of F1 and F2 steels followed a mixed parabolic-linear law. In the initial oxidation stage, the kp value of F1 and F2 steels was 1.40×106 mg2·cm4·s1 and 8.05×107 mg2·cm4·s1, respectively. After the initial period, the oxidation kinetics of F1, F2 steels obeyed linear law (Fig. 3b), but the corresponding kl value of F2 steel was nearly half of F1 steel. For F3 steel, the oxidation kinetics followed logarithmic law when the oxidation time was less than 5500 s, and subsequently followed linear law. We can conclude that F3 steel had best oxidation resistance capability at 950 °C. The parabolic or linear rate constant of experimental steels oxidized at 1000 °C and 1050 °C was also calculated. However, the oxide process was complex, when the temperature approached 1100 °C, and the corresponding parameters could not be derived. Table 3 shows 7
ACCEPTED MANUSCRIPT parabolic or linear oxidation rate constant for the tested steels oxidized at 9501050 °C in simulated exhaust gas. The results show that the oxidation rate constant of Ce-bearing F2 steel was about one order of magnitude lower than F1 steel, when the oxidation temperature was above 1000 °C. The reason is the “reactive element effect” on the addition of rare earth element. Many theories have been proposed to explain the various aspects of the reactive element effect. It is reported that rare earth element in the alloy can enhance the selective oxidation of chromium, reduce the growth rate of oxide film, change the transport mechanism and increase the resistance of scale to spallation [21]. Hou and Stringer [22] proposed that the segregation of RE along the oxide grain boundaries greatly retards outward transport cations through the scale. Seo et al. [23] put forward that the major role of REs is to make the oxidation process of chromia-rich scale the dominant inward diffusion of anions. In conclusion, the addition of rare earth element can significantly reduce the oxidation rate, such that the addition of Cerium in F2 steel led to its lower oxidation rate compared to Ce-free F1 steel. 3.2. Oxide scale surface morphology The macro-surface morphology of the specimens was observed to evaluate the degree of oxidation. After 5 h oxidation in simulated exhaust gas, all the specimen surfaces turned black. The oxide film of F1 and F4 steels showed spalling phenomenon. Fig. 4 shows micro-surface morphology of the oxide scale formed on the test steels oxidized at 950 °C for 5 h. The oxide film indicated spalling on F1 and F4 steel (Fig. 4a, d). The residual oxide film had a number of cracks and almost peeled off from the substrate. F2 and F3 steels did not experience spalling (Fig. 4b, c), and the oxide film displayed better adhesion. Moreover, the oxide particles were fine, dense and uniform, which could effectively 8
ACCEPTED MANUSCRIPT block the inward/outward diffusion of elements. As shown in the magnified images in Fig. 4a, the oxide on F1 steel (~8.9μm) was greater than F2, F3 and F4 steels (~3.6 μm, 4.1 μm and 4.9 μm, respectively), and in the spalled area, there existed many particles. Kuhn [24] studied microstructural features, mechanical and steam oxidation properties of ferritic stainless steel containing 18–23% chromium and addition of niobium and tungsten. They determined the chemical composition of the precipitates, which were rich in niobium and tungsten, and suggested that the phase is probably a (Fe, Cr, Si)2(Nb, W) Laves phase. Silva et al. [25] studied the in heat affected zone of AISI 444 ferritic stainless steel. They also proposed that the phase present in the HAZ was Laves phase. Thus, from the EDS analysis of these particles, we deduced that they were mainly Laves phase (Fe2(Nb, Mo, W)). The precipitates formed in F4 steel were larger and more in number than in F1 steel. Laves phase is intermetallic phase of the type AB2, and in some instances other alloying elements such as Cr, Si or Mo may also be present in the phase [26, 27]. Laves phase is very hard and brittle at ambient temperature, and the difference in the thermal expansion coefficient between Laves phase and matrix is large [28-30]. Thus, the precipitation of a large number of Laves phase at the oxide scale/metal interface is not beneficial for the adhesion of oxide film. Fig. 5 shows the micro-surface morphology of the oxide film formed at 1100 °C for 5 h. The size of oxide was increased compared to oxidation at 950 °C. Oxide formed on F2 and F3 steels was spherical in shape and uniformly distributed. However, oxide formed on F1 and F4 steel had strip-like morphology. The three dimensional profile of oxide scale surface after 5 h oxidation at 1100 °C is shown in Fig. 6. The surface fluctuations on F1 and F4 steels was large, especially F4 steel. The change in the thickness of oxide film is not beneficial to release growth stresses and thermal stresses. The accumulation of strain energy at the oxide/metal 9
ACCEPTED MANUSCRIPT interface will lead to the spallation of oxide film [31]. Morphology at high magnification is presented in the inset of Fig. 5. According to the XRD data in Fig. 7, the surface oxide mainly consisted of Cr-Mn spinel oxide. The morphology of MnCr2O4 spinel was octahedral in shape, which is shown as red circle in Fig. 5d. But in F1 steel tetragonal prism-shaped crystallites (marked by red circle in Fig. 5a) was present. This phenomenon was also observed in other studies [32, 33], and is due to MnxCr3-xO4 spinel crystal system that changes from cubic to tetragonal when x˃1.17. Thus, the MnCr2O4 spinel oxide in F1 steel may have partially decomposed. 3.3 Identification of oxide scale phase XRD analysis was conducted to determine phases present in the surface oxide film. Fig. 7 shows the XRD pattern of oxidized layer after the ferritic stainless steels was subjected to isothermal oxidation for 5h at 1100 ºC in simulated exhaust gas. The results revealed that the surface oxide film mainly composed of corundum M2O3 and spinel M3O4 type oxides, where M may be Fe, Cr, Mn or combination of them. Because the ionic radii of Fe, Cr and Mn are very similar, it is therefore difficult to determine the exact composition of the oxide film by the position of the peaks in the XRD patterns [34, 35]. But from the EDS analysis and EPMA elemental mapping of the cross-section of oxide film, we concluded that the oxide products mainly comprised of Cr2O3, (Mn, Cr)3O4 and small amount of SiO2. Fe-Cr substrate was also detected and is due to the oxide film being thin and the X-rays penetrated the film. 3.4 Cross-section analysis 3.4.1 Cross-sectional morphology Fig. 8 shows the cross-sectional morphology of the experimental steels after 5 h of oxidation in simulated exhaust gas obtained by back-scattered electron (BSE). The oxide film formed 10
ACCEPTED MANUSCRIPT on all the specimens can be divided into two parts: an outer layer of (Mn, Cr, Fe)3O4 spinel oxides and inner layer of Cr-rich (mainly Cr2O3) oxides. The oxide film formed on F1 steel is shown in Fig. 8a. Transverse cracks and small longitudinal cracks were present in the oxide film, which can lead to spallation of oxide film. In addition, many small pores existed in the oxide film (black circle in Fig. 8a). The cluster of pores coalesce and facilitate initiation of cracks [36]. In addition, at the oxide/metal interface gaps were present, which may reduce the adhesion between oxide scale and the matrix. The oxide film formed on F2 and F3 steels was dense, uniform and straight, and there was near absence of defects in the oxide film. Spherical oxide particle could be clearly seen on the oxide film surface of F2 steel, which was mainly composed of Cr-Mn spinel oxides. Fig. 8d shows the oxide film formed on F4 steel, which displayed notable difference compared to other steels. The oxide film was porous and undulated, a number of large cavities (marked by black arrow in Fig. 8d) and cracks were present in the oxide scale. It is worth noting that a thin layer of oxide near the matrix of F4 steel was compact, but the outer layer was loose. The pores and micro-cracks act as diffusivity paths allowing H2O molecules or O2 to easily penetrate in the oxide film, and H2O can provide a means of oxygen transport by dissociation through the following reaction [37, 38] H2O=H2+1/2O2
(4)
The partial pressure of H2O in the cavity approaches the external gas when the inward H2O transport is relatively fast, and the recycling of H2/H2O couple provides continuous transport of O2 [39]. In the case of F4 steel (Fig. 8d), the oxide film had many cracks and cavities which may lead to continuous oxidation and formation of a less protective oxide film. A large number of studies indicated that the presence of water vapor in the exposure environment can significantly influence the oxidation behavior of the chromia-forming 11
ACCEPTED MANUSCRIPT materials as compared to in dry atmosphere, e.g. oxygen or air [14]. In the case of Fe-Cr alloys with intermediate Cr content (10–20 wt%), wet gases can lead to degradation behavior of the oxide film by growing Fe-rich oxide scales [40]. However, there is still not a complete understanding of the mechanism behind this phenomenon. Majority of the studies attribute the breakaway oxidation of Fe-Cr alloys in water vapor containing environment to the formation of volatile chromium oxohydroxide CrO2(OH)2 [11,12]. Ehlers et al. [40] studied the oxidation behavior of 9% Cr steels in N2-O2-H2O gas mixture. They proposed that the penetration of water vapor molecules triggers enhanced oxidation and the tendency of the formation of non-protective oxide scale increased with increasing H2O(g)/O2 ratio. In our study, the partial pressure of O2 and H2O was 0.005 and 0.05 atm, respectively. All the experimental steels did not experience breakaway oxidation behavior in the temperature range of 950–1100ºC. This may due to the formation of (Mn, Cr, Fe)3O4 spinel oxide layer on the Cr2O3 scale. Fig. 9 shows the thickness and the mass gain per unit volume of the oxide film after 5 h oxidation at 1100 ºC. The oxide film thickness (Fig. 9a) can be obtained from the crosssectional morphology, and the thickness of oxide scale formed on F1, F2, F3 and F4 steels was ~8.1, 3.5, 4.7 and 10.1 μm, respectively. Combining with the mass gain data (Fig. 2d), the average mass gain per unit volume of the oxide film was calculated (Fig. 9b), and the results are 3.2, 4.3, 4.5 and 2.0 g·cm-3, respectively. If the formation of volatile chromium oxide and other situations are ignored, the weight gain is mainly from oxygen. Based on the results, we can say that the addition of small amount of Ce can significantly reduce the thickness and improve the density of oxide film. This phenomenon has been previously observed [3, 41]. 3.4.2 Elemental distribution in the oxide film 12
ACCEPTED MANUSCRIPT Elemental distribution involving the oxide layer and the substrate obtained by EPMA elemental mapping is shown in Fig. 10. The variation in elemental concentration is represented by change of color. Manganese was distributed in the outermost surface of the oxide film, and existed mainly in the form of spinel oxide (Mn, Cr)3O4. A compact and continuous layer of (Mn, Cr)3O4 spinel oxide doped with a small amount of iron formed in the top layer of Cr2O3, which effectively prevents the volatilization of chromium. Previous studies also showed that the outermost Cr-Mn spinel oxide can strongly reduce CrO2(OH)2 volatilization [42, 43]. Sachitanand et al. and Stanislowski et al. [44, 45] have shown that steel containing 0.30.5 wt% Mn develops an adherent (Mn, Cr)3O4 top layer above the Cr2O3 layer at 800 ºC and 850 ºC, and that the rate of Cr vaporization for these alloys was 2-3 times lower than alloys that form pure Cr2O3 scale or a non-continuous (Mn, Cr)3O4 top layer. The Mn content of the experimental steels in our study was greater than 0.3 wt. %, and results showed that a compact (Mn, Cr)3O4 layer formed on all the steels. Silicon was distributed at the oxide/metal interface and can be seen in Fig. 10. The presence of Si in stainless steel generally existed in the form of a silica interface between the external chromia layer and the base metal and acts as a barrier for cation transport. However, the existence of silicon layer in stainless steel enhances scale spallation during cyclic tests [46]. As shown in Fig. 10, silicon layer in F1 and F4 steels was more obvious than in F2 and F3 steel, and the discontinuous silicon layer was porous, which may lead to spallation of oxide film on F1 and F4 steels. The oxygen distributed in the oxide film had a concentration gradient, especially when the oxide film was thick. Near the oxide/metal interface, the oxygen concentration was lower. This implied that the oxidation process was diffusion-controlled. In the oxidation process, the 13
ACCEPTED MANUSCRIPT O–2 diffused through the oxide film from the gas/oxide interface to the oxide/metal interface and the metal cations diffused toward the opposite direction. The growth rate of oxide film was affected by concentration gradient and diffusion rate of the reactive elements. When the oxide scale becomes thicker, the diffusion velocity as well as the growth rate of the oxide film will decline. Chromium was distributed between the surface and the substrate. Cr depletion zone was not observed near the oxide/metal interface because of large residual Cr in the alloy. The rare-earth element in the oxide scale could not be detected because of very small addition of Ce (0.05 wt%) in the steels. The mechanism of Ce on the significant improvement of the high temperature oxidation resistance of Fe-Cr alloys is still unclear. But from the results of the experiments, it can be concluded that the addition of Ce can reduce the oxidation rate of the alloys at higher temperatures (above 1000 ºC) and improve the density of the oxide film. The oxide scale displayed better adhesion and protective nature in wet gas. The content of W in the alloys also had a dramatic influence on the high temperature oxidation behavior of the alloys. The addition of Ce together with ~0.5 wt% W can improve the oxidation resistance of the alloy compared to only addition Ce at temperatures below 1000 ºC. However, the oxide film experienced substantial spalling phenomenon during cooling to room temperature, when the addition was ~1.0 wt% W. The high content of W promoted the coarsing Laves phase precipitate at the oxide/scale interface, which can significantly reduce the adhesion of oxide scale. In addition, the higher content of W also led to the formation of porous oxide film on the alloy. All of these factors together resulted in the poor oxidation resistance of the oxide scale. 4. Conclusions The isothermal oxidation behavior of ferritic stainless steel was studied at 9501100 ºC in 14
ACCEPTED MANUSCRIPT simulated exhaust gas environment. Based on the results, the following are the conclusions: (1) The oxidation growth rate of B444M2 steel was significantly reduced on the addition of Ce or combination of Ce and 0.5 wt% W. They followed a mixed parabolic-linear law at 950 ºC and parabolic law at 1000 or 1050 ºC. (2) Multi-layer oxide film formed on all the steels. The surface layer was mainly (Mn, Cr)3O4 spinel oxide, which can effectively reduce Cr volatilization, and the layer beneath was Cr2O3 oxide. (3) The oxide formed on the steels was fine, dense and uniform with the addition of Ce, and the oxide film was thin, smooth and displayed better adhesion. The number of defects in the oxide scale was decreased. (4) Many cracks and pores were present in the oxide film of F1 and F4 steels. The silicon layer at the oxide/metal interface was porous, and there were many Laves phases precipitated at the oxide/metal interface at 950 ºC. All these factors may lead to severe spallation of the oxide film on F1 and F4 steels. The addition of 1.0 wt% W was not beneficial to the oxidation resistance.
Acknowledgements This work was financially supported by the National Natural Science Foundation of China (Grant No. U1660205). R.D.K. Misra gratefully acknowledges continued collaboration with Northeastern University as an Honorary Professor in providing guidance to the students.
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ACCEPTED MANUSCRIPT Table captions Table 1.
Chemical composition of ferritic stainless steels (in wt%).
Table 2.
The composition of the synthetic exhaust gas.
Table 3. gases.
Values of kp obtained at 950 – 1050 °C for test steels oxidized in simulated exhaust
20
ACCEPTED MANUSCRIPT Figure captions Fig. 1.
Schematic of the experimental system for oxidation studies.
Fig. 2.
Mass change versus time curves for the test steels oxidized for 300 min at different
temperatures in simulated synthetic exhaust gases. (a) 950 °C, (b) 1000 °C, (c) 1050 °C, and (d) 1100 °C. Fig. 3.
TGA curves of the test coupons oxidized at 950 °C for 18000 s in simulated
synthetic exhaust gas. (a) parabolic plot (Δm/s)2–t of F1 steel, (b) Δm/s–t plots of the F2 and F3 steels, (c) parabolic plot (Δm/s)2–t of F2 steel, and (d) parabolic plot (Δm/s)2–t of F4 steel. Fig. 4.
Surface morphology of the specimens after 5 h oxidation at 950 °C in synthetic
exhaust gases. (a) F1, (b) F2, (c) F3, and (d) F4. Fig. 5.
Surface morphology of the specimens after 5 h oxidation at 1100 °C in synthetic
exhaust gas. (a) F1, (b) F2, (c) F3, and (d) F4. Fig. 6.
Three dimensional profiles of the oxide scale surfaces after 5 h oxidation at 1100 °C
in synthetic exhaust gas. (a) F1, (b) F2, (c) F3, and (d) F4. Fig. 7.
XRD patterns of oxidation layer after the ferritic stainless steels isothermal oxidation
for 5 h at 1100 ºC in simulated exhaust gas. Fig. 8.
BSE morphology of the cross-section of the test steels oxidized at 1100 ºC for 5 h.
(a) F1, (b) F2, (c) F3, and (d) F4. Fig. 9.
The thickness and the mass gain per unit volume of the oxide films after 5 h
oxidation at 1100 ºC in simulated exhaust gas. Fig. 10.
BSE micrographs and EPMA element maps of oxide scales of the experimental
steels oxidized at 1100 ºC for 5 h in simulated exhaust gas.
21
ACCEPTED MANUSCRIPT Table 1. Chemical composition of ferritic stainless steels (in wt%). Steel
C
Si
Mn
P
S
Cr
Nb
Ti
N
Mo
Ce
W
Fe
F1
0.006
0.54
0.33
0.006
0.006
19.7
0.436
0.174
0.072
2.09
Bal.
F2
0.009
0.52
0.32
0.008
0.008
19.5
0.450
0.155
0.072
1.97
0.048
Bal.
F3
0.007
0.50
0.32
0.010
0.009
19.1
0.441
0.137
0.071
1.95
0.056
0.58
Bal.
F4
0.010
0.53
0.35
0.008
0.010
19.9
0.443
0.113
0.071
2.03
0.053
1.12
Bal.
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ACCEPTED MANUSCRIPT Table 2. The composition of the synthetic exhaust gas. Gas
O2
H 2O
CO2
N2
Vol.%
0.5
5
10
Bal.
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ACCEPTED MANUSCRIPT Table 3.
Values of kp obtained at 950–1050 °C for the experimental steels oxidized in
simulated exhaust gases. Steel grades
Temperature (°C)
Kinetics of oxidation
Rate constant values
R2
Parabolic (kp) (0–7500s)
1.40×10–12 g2·cm4·s1
0.98
Linear (kl) (7500–18000s)
1.57×10–8 g·cm2·s1
0.99
1000
Parabolic (kp)
3.21×10–12 g2·cm4·s1
0.99
1050
Parabolic (kp)
1.97×10–8 g2·cm4·s1
0.99
Parabolic (kp) (0–4500s)
8.05×10–13 g2·cm4·s1
0.99
Linear (kl) (4500–18000s)
7.68×10–9 g·cm2·s1
0.99
1000
Parabolic (kp)
1.29×10–11 g2·cm4·s1
0.99
1050
Parabolic (kp)
5.30×10–9 g2·cm4·s1
0.99
950
Linear (kl) (9000–18000s)
5.28×10–9 g·cm2·s1
0.99
1000
Parabolic (kp)
1.74×10–11 g2·cm4·s1
0.99
1050
Parabolic (kp)
6.70×10–9 g2·cm4·s1
0.99
Parabolic (kp)
1.50×10–12 g2·cm4·s1
0.99
Parabolic (kp)
4.23×10–12 g2·cm4·s1
0.99
1000
Parabolic (kp)
3.04×10–11 g2·cm4·s1
0.99
1050
Parabolic (kp)
7.50×10–9 g2·cm4·s1
0.99
950 F1
950 F2
F3
950 F4
Note: The time intervals 0–7500s and 0–4500s represent that during this period of time the curves obey parabolic law. The time intervals 7500–18000s, 4500–18000s and 9000–18000s represent that the curves follow linear law during these period.
24
ACCEPTED MANUSCRIPT
Fig. 1.
Schematic of the experimental system for oxidation studies.
25
ACCEPTED MANUSCRIPT (a)
(b)
(d)
(c)
Fig. 2.
Mass change versus time curves for the test steels oxidized for 300 min at different
temperatures in simulated synthetic exhaust gases. (a) 950 °C, (b) 1000 °C, (c) 1050 °C, and (d) 1100 °C.
26
ACCEPTED MANUSCRIPT
Fig. 3.
(a)
(b)
(c)
(d)
TGA curves of the test coupons oxidized at 950 °C for 18000 s in simulated
synthetic exhaust gas. (a) parabolic plot (Δm/s)2–t of F1 steel, (b) Δm/s–t plots of the F2 and F3 steels, (c) parabolic plot (Δm/s)2–t of F2 steel, and (d) parabolic plot (Δm/s)2–t of F4 steel.
27
ACCEPTED MANUSCRIPT (a)
(b)
20 μm
4 μm
Spalled area
40 μm
400 μm
(c)
(d)
4 μm
20 μm
Spalled area 400 μm
40 μm
Fig. 4.
Surface morphology of the specimens after 5 h oxidation at 950 °C in synthetic
exhaust gases. (a) F1, (b) F2, (c) F3, and (d) F4.
28
ACCEPTED MANUSCRIPT (a)
(b)
2 μm
2 μm
40 μm
40 μm
(c)
(d)
2 μm
2 μm
. . 2 2 7 7 0 40 μm 40 μm 0 . . Fig. 5. Surface morphology2 of the specimens after 5 h oxidation2 at 1100 °C in synthetic 7 exhaust gas. (a) F1, (b) F2, (c)7 F3, and (d) F4.
29
ACCEPTED MANUSCRIPT (a)
(b)
(c)
(d)
Fig. 6.
Three dimensional profiles of the oxide scale surfaces for the experimental steels
after 5 h oxidation at 1100 °C in synthetic exhaust gas. (a) F1, (b) F2, (c) F3, and (d) F4.
30
ACCEPTED MANUSCRIPT
Fig. 7.
XRD patterns of oxidation layer after the ferritic stainless steels isothermal oxidation for 5 h at 1100 ºC in simulated exhaust gas.
31
ACCEPTED MANUSCRIPT (b)
(a) Resin
Cr-Mn spinel Layer Transversal Cracks
Cr2O3 Layer
Longitudinal Cracks
Resin
Cr-Mn spinel Layer
Cr2O3 Layer
gap
Substrate
Substrate 10 μm
10 μm
(d)
(c) Resin
Resin
Cr-Mn spinel Layer
Cr2O3 Layer
Cr2O3 Layer Substrate
Fig. 8.
Cr-Mn spinel Layer
Substrate
10 μm
20 μm
BSE morphology of the cross-section of the experimental steels oxidized at 1100 ºC
for 5 h. (a) F1, (b) F2, (c) F3, and (d) F4.
32
ACCEPTED MANUSCRIPT (b)
(a)
Fig. 9.
The thickness (a) and the mass gain per unit volume (b) of the oxide film on the
surface of the steels after 5 h oxidation at 1100 ºC in simulated exhaust gas.
33
ACCEPTED MANUSCRIPT
(a)
O
Cr
Mn
Fe
Si
O
Cr
Mn
Fe
Si
O
Cr
Mn
Fe
Si
O
Cr
Mn
Fe
Si
10 μm
(b)
10 μm
(c)
10 μm
(d)
10 μm
Low
High Element Concentration
Fig. 10.
BSE micrographs and EPMA element maps of oxide scales of the experimental steels oxidized at 1100 ºC for 5 h in simulated exhaust gas.
34