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O2 mixtures at 600 °C: the effect of water-vapour-enhanced chromium evaporation
Corrosion Science 44 (2002) 2635–2649 www.elsevier.com/locate/corsci
Oxidation of 310 steel in H2O/O2 mixtures at 600 °C: the effect of water-vapour-enhanced chromium evaporation H. Asteman
a,*
, J.-E. Svensson b, L.-G. Johansson
a
a
b
Department of Chemistry, G€oteborg University, S-412 96 G€oteborg, Sweden Department of Environmental Inorganic Chemistry, Chalmers University of Technology, S-412 96, Sweden Received 4 February 2002; accepted 15 February 2002
Abstract The oxidation of type 310 stainless steel was investigated at 600 °C in the presence of O2 and O2 þ 10% and 40% H2 O. The effect of gas velocity was studied. The oxidized samples were investigated by grazing angle X-ray diffraction, SEM/EDX and SAM. The addition of H2 O to O2 resulted in a change of oxidation behaviour. A strong dependence on flow rate was observed in O2 /H2 O mixtures. At low flow rates a thin (30–50 nm) protective a-(Cr,Fe)2 O3 formed, the outer part being depleted in chromium. When the flow rate was increased beyond a critical value the protective oxide failed. Under these conditions P 5 lm thick a-Fe2 O3 / (Cr,Fe)3 O4 , oxide islands formed on the part of the surface corresponding to the centre of the alloy grains. The effect of water vapour is attributed to the water-vapour-assisted evaporation of chromium from the oxide, in the form of a chromium oxide hydroxide, probably CrO2 (OH)2 . The oxidation behaviour is rationalized using a qualitative mechanism proposed previously and parallels that of the 304L alloy. Ó 2002 Published by Elsevier Science Ltd. Keywords: Stainless steel; AES; SEM; High temperature corrosion; Selective oxidation
1. Introduction The influence of water vapour on the high temperature oxidation of Fe, Cr, and Fe–Cr and Fe–Cr–Ni alloys, has received considerable attention [1–12,16]. It has
*
Corresponding author. E-mail addresses: [email protected] (H. Asteman), [email protected] (J.-E. Svensson), [email protected] (L.-G. Johansson). 0010-938X/02/$ - see front matter Ó 2002 Published by Elsevier Science Ltd. PII: S 0 0 1 0 - 9 3 8 X ( 0 2 ) 0 0 0 5 6 - 2
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been observed that H2 O/O2 mixtures are far more corrosive towards this group of alloys compared to dry O2 . Several attempts have been made to explain the apparent corrosivity of water vapour. The explanations put forward can be classified according to the role proposed for water: (1) increasing the mass transport of oxygen through iron oxide [2], (2) forming gaseous iron hydroxide leading to increased mass transport of iron through the scale [4], (3) increasing the diffusion rate of chromium through chromium oxide (or through a Cr-rich oxide) [12] and (4) causing the vapourization of chromium oxide hydroxide [6,7,9–11]. Explanations belonging to groups 1 and 2 may explain why iron and other iron oxide-formers oxidize faster in the presence of water vapour. However, they do not explain why water vapour causes the protective properties of chromia-forming Fe– Cr and Fe–Cr–Ni alloys to deteriorate. Similarly, explanations in group 3, involving changes in the electronic and ionic conductivity of chromium oxide, may explain why pure chromium or Cr–Ni alloys oxidize faster in the presence of water vapour. However, they do not explain why water vapour tends to cause Fe–Cr and Fe–Cr–Ni alloys to form iron-rich oxides rather than chromium-rich oxides. In fact, a more rapid diffusion of Cr in the oxide is expected to promote the formation of a protective Cr-rich scale. It is well-known that Cr-rich oxide scales lose chromium by CrO3 evaporation in oxygen-containing environments >1000 °C resulting in ‘‘active corrosion’’ [13,14]. As a result, alloys that form protective chromia layers in oxygen at lower temperatures tend to fail doing this >1000 °C. Recently, it was shown that the exposure of Fe–Cr–Ni alloys in H2 O/O2 mixtures results in chromium evaporation already at 600 °C [6]. It was concluded that the vapourizing species is a chromium (VI) oxide hydroxide, probably CrO2 (OH)2 (g): 1=2Cr2 O3 ðsÞ þ 3=4O2 ðgÞ þ H2 O ! CrO2 ðOHÞ2 ðgÞ
ð1Þ
It is argued that water-vapour-enhanced chromium evaporation is an attractive explanation of the corrosivity of water vapour in oxygen towards Fe–Cr alloys <1000 °C. Recently, we showed that the oxidation of 304L steel (Cr18Ni10) in H2 O/ O2 mixtures is flow dependent, high flow velocities destabilizing the protective Crrich oxide [7,9]. The dependence of oxidation on flow velocity shows that oxidation depends on transport processes in the gas phase. Gas velocity is expected to influence evaporation because the volatile species formed at the surface have to diffuse through a stagnant layer of gas separating the surface from the bulk gas in order to be carried away by the flowing gas. The thickness of the diffusion layer diminishes with increasing flow rate, resulting in an increased evaporation rate. As there is no conceivable other way by which gas velocity can influence oxide properties, the fact that oxidation depends on gas velocity proves that vapourization of chromium significantly influences the corrosion process. In our previous studies it was proposed that the vapourization of chromium speeds up the oxidation of Fe–Cr alloys in H2 O/O2 environments because it tends to deplete the oxide in chromium, resulting in poorer protective properties of the oxide scale. The argument is based on the fact that Fe2 O3 and Cr2 O3 form a continuous
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range of solid solutions. This means that a protective Cr-rich oxide is easily converted into poorly protective Fe-rich oxide by chromium vapourization. At low evaporation rates, the metal is able to sustain a high enough flux of chromium to the oxide in order for the oxide to retain its protective properties. When the evaporation rate exceeds a critical value, excessive chromium depletion of the oxide leads to the loss of its protective properties. Once the chromium-rich oxide is destroyed, rapid oxidation ensues. A qualitative mechanism was presented, explaining the differences in oxidation behaviour of Fe–Cr and Fe–Cr–Ni alloys in H2 O/O2 mixtures in terms of the supply of chromium from the substrate to the oxide. It was proposed that the resistance towards high temperature corrosion in this type of environment is enhanced by high Cr concentration, fast diffusion of Cr in the steel bulk (ferrite rather than austenite) and by a high density of grain boundaries in the alloy (small grain size). The previous investigations also showed that the corrosivity of the environment increases as the gas composition approaches the stoichiometry of reaction (1), and that watervapour-enhanced oxidation is promoted by high gas velocity and high temperature [7,9]. The present paper extends our investigation of the effect of H2 O/O2 mixtures on oxidation to the commercially important 310 alloy (for composition, see Table 1). This alloy has essentially the same microstructure, e.g., grain size, as 304L but a considerably higher chromium content (25% as opposed to 18%). We focus especially on whether the differences between this material and the previously studied 304L can be rationalized on the basis of the mechanism proposed.
2. Experimental In all experiments, 310-coupons with the size 15 15 2 mm were exposed (for chemical composition, see Table 1). The samples were ground on SiC-paper to 1000 mesh, polished with 1 lm diamond paste and then cleaned in acetone and ethanol using ultrasonic agitation.
2.1. Experimental set-up for the isothermal exposures The exposures were carried out in a horizontal furnace fitted with a 48 mm diameter SiO2 -glass tube. All exposures were isothermal, the temperature was kept at 600 3°. The samples were mounted three at a time using a sample holder consisting of a dense alumina-plate with the dimensions 40 50 5 mm, featuring three 1 mm deep and 2.1 mm broad slits. The distance between the slits was 12 mm. The samples were positioned parallel to the direction of flow. This set-up enables the samples to be exposed to the reaction gas with minimum contact with the sample holder and with optimal flow pattern and minimizes the disturbance caused by the presence of multiple samples in the reactor.
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Table 1 Chemical composition of 310 steel (wt%) 310
Cr (%)
Ni (%)
Mn (%)
Si (%)
Mo (%)
Fe (%)
24.9
19.2
1.55
0.45
0.34
Balance
2.2. Experimental conditions/flow rate dependence Exposures were made in a furnace system fitted with a humidifier producing a reaction gas consisting of O2 þ 10% H2 O or O2 þ 40% H2 O. After each exposure the samples were allowed to cool in dry air. The mass changes were recorded using a fivedecimal balance. Two types of experiments were made, mass change vs. time and mass change vs. flow rate for a constant exposure time. In the time-resolved experiments two different flow rates were used, 200 ml/min (NTP) and 2000 ml/min (NTP), corresponding to 0.50 and 5.0 cm/s, respectively. For the investigation of the effect of flow rate, a standard exposure time of 168 h was selected. Flow rates were 10, 200, 500, 1000, 2000 and 4000 ml/min (NTP). These flow rates correspond to an average net gas velocity inside the furnace of 0.03, 0.50, 1.25, 2.5, 5.0 and 10.0 cm/s, respectively. 2.3. Corrosion product characterization 2.3.1. X-ray diffraction (GI-XRD) Crystalline corrosion products were analyzed by X-ray diffraction (XRD) using a Siemens D5000 powder diffractometer equipped with grazing incidence beam attachment and a G€ obel mirror. CuKa radiation was used and the angle of incidence was 0.50°. The detector measures between 20° < 2h < 80°. 2.4. Analytical environmental scanning electron microscopy (ESEM/EDX) Analytical scanning electron microscopy was carried out using an Electro-scan 2020 equipped with a Link eXl EDX system. The microscope was operated at 20 kV for secondary electron imaging and at 10–20 kV for EDX analyses. 2.5. Scanning Auger electron microscopy (SAM) The Auger analyses were performed using a scanning Auger microprobe (PHI660). The electron beam voltage was 10 kV and the beam current was about 160 nA. The depth profiles were obtained using a differentially pumped ion gun (Arþ ) with acceleration voltages between 1.5 and 4 kV. The etch rates were calibrated on . The collected rawflat samples of Ta2 O5 with well-known oxide thickness of 1000 A data was refined into the form presented in this paper, by using MultiPakÓ v.6.0 software. The software makes it possible to evaluate peak shapes of differentiated Auger-peaks. This feature can be used in order to distinguish between oxidized and metallic iron and chromium signal, as these elements exhibit significant chemical
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shifts in oxidized form. The concentrations for O, Cr(oxidized) and Fe(oxidized) were calculated using sensitivity factors determined using standards of pure Cr2 O3 and Fe2 O3 and a range of solid solutions ((Cr,Fe)2 O3 ) [15]. 3. Results 3.1. Gravimetry Fig. 1 shows the marked flow rate dependence of the oxidation of 310 steel after 168 h in O2 , O2 þ 10% and 40% H2 O at 600 °C. In O2 þ 40% H2 O at 10 ml/min (0.03 cm/s), a small mass gain is registered. With increasing flow rate, mass gain decreases and at 500 (1.25 cm/s) and 1000 ml/min (2.5 cm/s) a small but significant mass loss is registered. Increasing the flow rate 2000 (5.0 cm/s) and 4000 ml/min (10.0 cm/s) resulted in a mass gain. The transition from mass loss to mass gain is abrupt and at high flow rates the mass gain is greater than in dry O2 . The dependence of mass change on flow rate in O2 þ 10% H2 O environment resembles that in O2 þ 40% H2 O. For example, the transition from mass loss to mass gain appears at the same flow rate. In comparison, the dry O2 exposures resulted in a mass gain that did not depend on flow rate. The results show that, in O2 þ H2 O environment, material is being lost from the samples by vapourization. Moreover it appears that, depending on flow rate, vapourization results in two distinct types of behaviour. At low flow velocities vapourization mainly results in a mass loss and does not seem to increase the rate of oxidation. In contrast, at high flow rates, mass increases indicating accelerated corrosion. In order to investigate the mass gain in the two regimes mentioned, a series of time-resolved exposures were carried out. The time dependence of mass change in O2 þ 40% H2 O was investigated at low flow (200 ml/min, 0.50 cm/s) and at high flow (2000 ml/min, 5.0 cm/s) (see Fig. 2). In the set of exposures at low flow, a mass gain is
Fig. 1. Mass-change vs. flow rate; 310 oxidized for 168 h at 600 °C in dry O2 , O2 þ 10% H2 O and O2 þ 40% H2 O atmosphere.
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Fig. 2. Mass-change vs. exposure time; 310 oxidized for up to 336 h at 600 °C in O2 þ 40% H2 O atmosphere with low gas-flow velocity 0.50 cm/s and high gas-flow velocity 5.0 cm/s.
registered that increases with time up to 72 h. After longer exposure times mass gain starts to decrease and at 336 h the samples exhibit a mass loss. In the experiments involving higher flow rates mass gains are registered at all exposure times. Mass gain increases with exposure time, levelling out after 168 h. 3.2. Phase analysis (GI-XRD) The phase analysis by GI-XRD indicated only one crystalline oxide, i.e., corundum-type M2 O3 corresponding to a-Fe2 O3 , Cr2 O3 or the solid solution (Cr,Fe)2 O3 . 3.3. Morphology (ESEM) In Fig. 3a a typical ESEM image of a sample oxidized at 600 °C in dry O2 for 168 h is presented. A thin and smooth oxide film covers the surface. The network pattern that can be seen is a projection of the grain boundaries in the alloy and can be seen clearly after sputtering when performing AES-depth profiling. Fig. 3b shows a sample exposed for 168 h in O2 þ 40% H2 O at low flow rate (200 ml/min corresponding to 0.50 cm/s). The morphology is very similar to the dry O2 exposure. It may be noted that the interference colours developed by the samples in Fig. 3a and b were different. The samples oxidized in dry O2 were deep blue while the ones oxidized in O2 þ 40% H2 O ranged from bluish grey to pale lilac. This indicates that the oxide films differ in thickness or in chemical composition. The morphology changes dramatically when samples are exposed to higher flow rates in O2 þ 40% H2 O environment (2000 ml/min, 5.0 cm/s) (see Fig. 3c). On top of the thin oxide, ‘‘islands’’ occur, consisting of thick oxide. It may be noted that the oxide islands never cross the network pattern mentioned above. This means that the oxide islands only form on the part of the sample surface that corresponds to the
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Fig. 3. ESEM pictures of 310 oxidized for 168 h at 600 °C: (a) O2 , (b) O2 þ 40% H2 O flow: 200 ml/min (0.50 cm/s), (c) O2 þ 40% H2 O flow: 2000 ml/min (5.0 cm/s) and (d) same as (c) but at higher magnification.
interior of the alloy grains. Oxide islands are not formed on every steel grain. However, a closer look at some of the steel grains lacking oxide islands, shows that a thicker oxide has started to form (see Fig. 3d). When samples are exposed at even higher flow rate in this environment (>2000 ml/min, 10.0 cm/s), the sample morphology remains qualitatively the same while the number density of oxide islands increases. At flow rates up to 1000 ml/min (2.5 cm/s), the surface morphologies developed in O2 þ 10% H2 O were essentially the same as described above for O2 þ 40% H2 O (see Fig. 4a). At higher flow rates, a transition from protective behaviour to a local breakdown of protective properties occurs in O2 þ 10% H2 O. However, in this case the oxide islands only occur sporadically (see Fig. 4b). 3.4. Chemical composition (AES) Figs. 5–8 show AES-depth profiles of samples exposed for 168 h. The ESEM images of the samples in Figs. 5–7 are not included as the surfaces were smooth and featureless. Figs. 5a, 6a and 7a show the chemical composition of the oxide film on samples oxidized in dry O2 , and in O2 þ 40% H2 O at 10 and 1000 ml/min (0.03 and
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Fig. 4. ESEM pictures of 310 oxidized for 168 h at 600 °C: (a) O2 þ 40% H2 O flow: 1000 ml/min (2.5 cm/s) and (b) O2 þ 10% H2 O flow: 2000 ml/min (5.0 cm/s), respectively.
2.5 cm/s) flow rate, respectively. Figs. 5b, 6b and 7b show the Cr/cation ratio in the oxide and the composition of the alloy. The chromium concentration in a(Cr,Fe)2 O3 , is calculated from CðCr2 O3 Þ ¼ ðCCrðoxÞ =ðCCrðoxÞ þ CFeðoxÞ Þ 100%Þ, where CMðoxÞ is the concentration of the Fe(ox) and Cr(ox) signal separated from the Fe(metal) and Cr(metal) by MultipakTM v.6.0 software as described in Section 2. The Fe(metal) and Cr(metal) signal, in turn, are combined with the nickel signal (Ni is assumed to be only in non-oxidized form) using the formula CMðalloyÞ ¼ ðCMðmeÞ = ðCCrðmeÞ þ CFeðmeÞ þ CNi Þ 100%Þ (M ¼ Cr, Fe and Ni). This makes it possible to plot the alloy composition (CCrðalloyÞ , CFeðalloyÞ and CNiðalloyÞ ) as a function of the distance from the oxide–gas interface. Fig. 5a and b show that the 80 nm oxide film formed in dry O2 consists of a a(Cr,Fe)2 O3 solid solution with about 70% Cr and that the steel substrate beneath the oxide is only marginally depleted in chromium. The oxides formed in O2 þ 40% H2 O exposures have a different chemical composition (see Figs. 6 and 7). Even at the lowest flow rate (10 ml/min, 0.03 cm/s), the outer part of the oxide is depleted in chromium (see Fig. 6a and b). The Cr=Cr þ Fe ratio in the oxide is seen to be about 30% at the oxide–gas interface while the inner parts of the oxide remains Cr-rich. In addition, there is a marked depletion of chromium in the alloy while iron and nickel
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Fig. 5. AES profile of the oxide formed on 310 oxidized for 168 h at 600 °C in dry O2 : (a) the pure depthprofile and (b) the Cr concentration in (Cr,Fe)2 O3 , and the metallic Cr, Fe and Ni as a function of depth.
are enriched. When the flow rate is increased to 1000 ml/min (2.5 cm/s) (see Fig. 7), the chromium depletion in the oxide and in the substrate becomes more marked. As a consequence, the Ni concentration in the substrate immediately below the oxide reaches about 30 at.%, (compare the bulk value in the steel of 25%). Film thickness is also affected by water vapour. In dry O2 film thickness is about 80 nm. In comparison, the samples exposed in O2 þ 40% H2 O at a flow rate of 10 ml/min, produced a 60 nm oxide while the oxide formed in the same environment at a flow rate of 1000 ml/min was only 40 nm. This indicates that the presence of water vapour tends to etch the oxide and produce a thinner oxide rather than accelerating the growth of a(Cr,Fe)2 O3 . In O2 þ 40% H2 O and very high flow rates (P 2000 ml/min in the case of O2 þ 40% H2 O) the morphology, thickness and chemical composition of the oxidized surface changes dramatically. In Fig. 8 an ESEM image of a sample that was
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Fig. 6. AES profile of the oxide formed on 310 oxidized for 168 h at 600 °C in O2 þ 40% H2 O flow: 10 ml/ min (0.03 cm/s): (a) the pure depth-profile and (b) the Cr concentration in (Cr,Fe)2 O3 and the metallic Cr, Fe and Ni as a function of depth.
oxidized at 600 °C for 168 h in O2 þ 40% H2 O at a flow rate of 2000 ml/min (5.0 cm/ s) is shown. In addition, AES profiles are shown, representing different characteristic features on the surface. It can clearly be seen that the outer part of the oxide between the oxide islands is almost completely depleted in chromium. In addition, the inner part of the oxide is also considerably depleted in chromium compared to the previously described oxides formed in O2 þ H2 O environments at low flow rate. The oxide between the oxide islands has a thickness of about 250 nm, which is about three times the value in dry O2 and six times thicker compared to the oxide formed in O2 þ 40% H2 O at 1000 ml/min. The thickness of the oxide islands is about 5 lm. The islands exhibit a layered structure, the outer part being made up of iron-rich oxide while the inner part contains considerable amounts of chromium. The combined evidence from XRD and Auger spectroscopy shows that the outer part of the islands is iron-rich (Fe,Cr)2 O3 . Even though XRD does not indicate that other oxides are
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Fig. 7. AES profile of the oxide formed on 310 oxidized for 168 h at 600 °C in O2 þ 40% H2 O flow: 1000 ml/min (2.5 cm/s): (a) the pure depth-profile and (b) the Cr concentration in (Cr,Fe)2 O3 and the metallic Cr, Fe and Ni as a function of depth.
present, it is suggested that the inner, chromium-rich part of the oxide consists of iron-chromium spinel (Fe,Cr)3 O4 .
4. Discussion Before turning to the present results concerning 310 steel, it is useful to recapitulate the recently reported oxidation behaviour of the 304L alloy in O2 /H2 O environment and the mechanism proposed to rationalize the observations made [6,7,9]. Similarly to the 310 alloy, the 304L material forms a thin (50–100 nm) protective oxide consisting of chromium-rich a-(Cr,Fe)2 O3 when exposed to dry O2 at 600 °C for 1–4 weeks. In O2 /H2 O environment at the same temperature, the oxidation behaviour depended on flow rate. At low flow rates, the oxide showed signs of chromium depletion while the protective properties remained intact. At higher flow rates a gradual and localized breakdown of the protective properties was observed. The flow rate dependence was explained in terms of the water-vapour-enhanced chromium vapourization mentioned in the introduction. The evaporation of chromium
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oxy-hydroxides from the surface depletes the protective oxide in chromium. As already noted, the evaporation rate is expected to depend on the equilibrium partial pressure of the volatile species and on gas velocity. At low flow rates, corresponding to low evaporation rates, the diffusion of chromium from the steel substrate is sufficiently rapid so that the chromium content in the oxide is not seriously lowered. As the evaporation rate increases, the diffusion of chromium from the alloy eventually becomes insufficient and the chromium concentration in the oxide drops below a critical value. When this happens, the oxide loses part of its protective properties and starts to behave similar to a-Fe2 O3 , resulting in a dramatic increase in oxidation rate. It was observed that the breakdown of protective properties of 304L always occurred on the parts of the surface corresponding to the middle of the alloy grains, resulting in the formation of hematite islands of several lm thickness. This was attributed to a poor supply of chromium from the alloy. Because the diffusion rate is high in the grain boundaries compared to the interior of a crystal, the oxide formed on the parts of the surface corresponding to the interior of the alloy grains tends to be more susceptible to breakdown compared to the oxide in the vicinity of a grain boundary. Similarly, the tenacity of the thin protective chromium-rich a-(Cr,Fe)2 O3 , layer separating the oxide islands was attributed to the greater supply of chromium close to the alloy grain boundaries. At still higher flow rates the 304L material also suffered a partial degradation of the protective oxide between the oxide islands, causing this part of the oxide to grow to about 700 nm thickness. Based on these observations it was suggested that the ability of Fe–Cr alloys to withstand oxidation in H2 O/O2 environment is connected to the capacity of the alloy to compensate for the losses of chromium from the oxide by vapourization. This implies that oxidation resistance is related to chromium content and to the diffusivity of chromium in the alloy. The 310 and 304L steels are quite similar, both alloys being austenitic and having almost the same microstructure. Therefore we expect a relatively low bulk diffusivity of chromium as well as the same density of rapid diffusion paths (primarily grain boundaries) in both materials. This implies that we expect qualitatively the same oxidation behaviour for the two alloys in H2 O/O2 environment. However, as the 310 alloy contains more chromium (25% as compared to 18% for 304L) it is predicted to be somewhat less susceptible to oxidation. This means that while we expect a similar localized breakdown of the protective oxide on the 310 alloy due to water-vapourenhanced chromium evaporation, a higher flow rate is expected to be necessary in order to trigger this behaviour. The present study clearly shows the effect of water-vapour-enhanced chromium evaporation (see Figs. 1, 5 and 6). At low flow rate, the samples exposed in H2 O/O2 environment develop an oxide, the outer part of which is markedly depleted in chromium compared to the dry O2 exposure. The outer part of the oxide consists of almost pure a-(Fe2 O3 ), sometimes with a sporadic presence of Mn. Still, the thin and smooth oxide lacking in morphological features reveals that, in spite of the partial chromium depletion, the oxide formed on 310 in H2 O/O2 environment at low flow rate is completely protective, at least up to 168 h. The oxide formed on the 310 alloy in H2 O/O2 environment at low flow rates has the same chemical composition and
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thickness as the thin and protective oxide that formed close to the alloy grain boundaries on 304L [7]. As the flow rate increases it causes chromium vapourization from the 310 alloy to accelerate (see Fig. 1). This is also evident from the Auger depth profiling (see Figs. 6 and 7) As predicted, the flow rate necessary for triggering the deterioration of protective properties of the oxide is greater than in the case of the 304 alloy [7]. In the case of the 310 alloy the critical flow rate that causes the oxide to loose its protective properties in O2 þ 40% H2 O at 600 °C is between 1000 and 2000 ml/min (see Figs. 1 and 8). For alloy 304L in the same environment, the transition from protective to partly non-protective behaviour occurred close to 200 ml/min [7]. The transition from protective to non-protective behaviour is more sudden for the 310 alloy than for the 304L material. In the former case local breakaway corrosion, resulting in the formation of oxide islands, and the partial loss of the protective properties of the thin oxide between islands occurs at the same flow rate. The oxide islands that form have a thickness in the micrometer range (304L, 5–20 lm and 310, 5 lm). The islands exhibit a layered structure with an outer iron-rich corundum-type (Cr,Fe)2 O3 and an inner, more Cr-rich layer possibly consisting of (Fe,Cr)3 O4 . The (Cr,Fe)2 O3 , oxide separating the oxide islands on the 310 alloy exposed to H2 O/O2 environment at high flow rates is about 250 nm on the 310 alloy (compared to about 50 nm at low flow rates). The partial loss of the protective properties of this part of the oxide is
Fig. 8. AES profiles of the oxide formed on 310 oxidized for 168 h at 600 °C in O2 þ 40% H2 O flow: 2000 ml/min (5.0 cm/s).
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also attributed to the chromium depletion caused by water-vapour-enhanced chromium vapourization (see Fig. 8). However, due to the greater supply of chromium to this part of the oxide by grain boundary diffusion in the alloy, chromium loss only results in a partial loss of protective properties. The oxide morphology developed on the 310 alloy in O2 þ 10% H2 O is essentially the same as in O2 þ 40% H2 O. The main difference is that in O2 þ 10% H2 O environment, a higher flow rate is needed to cause the protective properties of the oxide to breakdown resulting in the formation of oxide islands. To summarize, the oxidation behaviour of the 310 material is very similar to that of the previously studied 304L alloy. However, the higher chromium content makes the 310 alloy less sensitive towards the destabilization of the protective properties of the oxide by water-vapour-enhanced chromium vapourization. When chromium vapourization is sufficiently rapid (high water vapour content and high flow rate) the 310 alloy suffers a partial breakdown of the protective oxide. It is suggested that when chromium vapourization becomes sufficiently rapid, (high water vapour content and flow rates in the m/s range) both materials will suffer a more or less complete breakdown of the protective oxide at 600 °C. Such conditions may appear in some industrial applications, e.g., in biomass combustion.
5. Conclusions The addition of H2 O to O2 results in a change of oxidation behaviour of the 310 alloy. A strong dependence on flow rate was observed in O2 /H2 O mixtures. At low flow rates a thin (30–50 nm) protective a-(Cr,Fe)2 O3 , forms, the outer part being depleted in chromium. When the flow rate is increased beyond a critical value the protective oxide fails. Under these conditions P 5 lm thick a-Fe2 O3 /(Cr,Fe)3 O4 , oxide islands forms on the part of the surface corresponding to the centre of the alloy grains. The a-(Cr,Fe)2 O3 oxide close to the alloy grain boundaries remained partially protective. The effect of water vapour is attributed to the water-vapour-assisted evaporation of chromium from the oxide, in the form of a chromium oxide hydroxide, probably CrO2 (OH)2 . The oxidation behaviour is rationalized using a qualitative mechanism proposed previously and parallels that of the 304L alloy. The mechanism implies that the ability of Fe–Cr alloys to withstand oxidation in H2 O/O2 environment is connected to the capacity of the alloy to compensate for the losses of chromium from the oxide by vapourization. The slightly better resistance of the 310 alloy as compared to 304L is attributed to the higher chromium content of 310. The present results are suggested to be relevant to environments with high concentrations of H2 O and O2 , e.g., in biomass combustion.
Acknowledgements This work was carried out within the Swedish High Temperature Corrosion Center (HTC).
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Oxidation of 310 steel in H2O/O2 mixtures at 600 °C: the effect of water-vapour-enhanced chromium evaporation H. Asteman
a,*
, J.-E. Svensson b, L.-G. Johansson
a
a
b
Department of Chemistry, G€oteborg University, S-412 96 G€oteborg, Sweden Department of Environmental Inorganic Chemistry, Chalmers University of Technology, S-412 96, Sweden Received 4 February 2002; accepted 15 February 2002
Abstract The oxidation of type 310 stainless steel was investigated at 600 °C in the presence of O2 and O2 þ 10% and 40% H2 O. The effect of gas velocity was studied. The oxidized samples were investigated by grazing angle X-ray diffraction, SEM/EDX and SAM. The addition of H2 O to O2 resulted in a change of oxidation behaviour. A strong dependence on flow rate was observed in O2 /H2 O mixtures. At low flow rates a thin (30–50 nm) protective a-(Cr,Fe)2 O3 formed, the outer part being depleted in chromium. When the flow rate was increased beyond a critical value the protective oxide failed. Under these conditions P 5 lm thick a-Fe2 O3 / (Cr,Fe)3 O4 , oxide islands formed on the part of the surface corresponding to the centre of the alloy grains. The effect of water vapour is attributed to the water-vapour-assisted evaporation of chromium from the oxide, in the form of a chromium oxide hydroxide, probably CrO2 (OH)2 . The oxidation behaviour is rationalized using a qualitative mechanism proposed previously and parallels that of the 304L alloy. Ó 2002 Published by Elsevier Science Ltd. Keywords: Stainless steel; AES; SEM; High temperature corrosion; Selective oxidation
1. Introduction The influence of water vapour on the high temperature oxidation of Fe, Cr, and Fe–Cr and Fe–Cr–Ni alloys, has received considerable attention [1–12,16]. It has
*
Corresponding author. E-mail addresses: [email protected] (H. Asteman), [email protected] (J.-E. Svensson), [email protected] (L.-G. Johansson). 0010-938X/02/$ - see front matter Ó 2002 Published by Elsevier Science Ltd. PII: S 0 0 1 0 - 9 3 8 X ( 0 2 ) 0 0 0 5 6 - 2
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been observed that H2 O/O2 mixtures are far more corrosive towards this group of alloys compared to dry O2 . Several attempts have been made to explain the apparent corrosivity of water vapour. The explanations put forward can be classified according to the role proposed for water: (1) increasing the mass transport of oxygen through iron oxide [2], (2) forming gaseous iron hydroxide leading to increased mass transport of iron through the scale [4], (3) increasing the diffusion rate of chromium through chromium oxide (or through a Cr-rich oxide) [12] and (4) causing the vapourization of chromium oxide hydroxide [6,7,9–11]. Explanations belonging to groups 1 and 2 may explain why iron and other iron oxide-formers oxidize faster in the presence of water vapour. However, they do not explain why water vapour causes the protective properties of chromia-forming Fe– Cr and Fe–Cr–Ni alloys to deteriorate. Similarly, explanations in group 3, involving changes in the electronic and ionic conductivity of chromium oxide, may explain why pure chromium or Cr–Ni alloys oxidize faster in the presence of water vapour. However, they do not explain why water vapour tends to cause Fe–Cr and Fe–Cr–Ni alloys to form iron-rich oxides rather than chromium-rich oxides. In fact, a more rapid diffusion of Cr in the oxide is expected to promote the formation of a protective Cr-rich scale. It is well-known that Cr-rich oxide scales lose chromium by CrO3 evaporation in oxygen-containing environments >1000 °C resulting in ‘‘active corrosion’’ [13,14]. As a result, alloys that form protective chromia layers in oxygen at lower temperatures tend to fail doing this >1000 °C. Recently, it was shown that the exposure of Fe–Cr–Ni alloys in H2 O/O2 mixtures results in chromium evaporation already at 600 °C [6]. It was concluded that the vapourizing species is a chromium (VI) oxide hydroxide, probably CrO2 (OH)2 (g): 1=2Cr2 O3 ðsÞ þ 3=4O2 ðgÞ þ H2 O ! CrO2 ðOHÞ2 ðgÞ
ð1Þ
It is argued that water-vapour-enhanced chromium evaporation is an attractive explanation of the corrosivity of water vapour in oxygen towards Fe–Cr alloys <1000 °C. Recently, we showed that the oxidation of 304L steel (Cr18Ni10) in H2 O/ O2 mixtures is flow dependent, high flow velocities destabilizing the protective Crrich oxide [7,9]. The dependence of oxidation on flow velocity shows that oxidation depends on transport processes in the gas phase. Gas velocity is expected to influence evaporation because the volatile species formed at the surface have to diffuse through a stagnant layer of gas separating the surface from the bulk gas in order to be carried away by the flowing gas. The thickness of the diffusion layer diminishes with increasing flow rate, resulting in an increased evaporation rate. As there is no conceivable other way by which gas velocity can influence oxide properties, the fact that oxidation depends on gas velocity proves that vapourization of chromium significantly influences the corrosion process. In our previous studies it was proposed that the vapourization of chromium speeds up the oxidation of Fe–Cr alloys in H2 O/O2 environments because it tends to deplete the oxide in chromium, resulting in poorer protective properties of the oxide scale. The argument is based on the fact that Fe2 O3 and Cr2 O3 form a continuous
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range of solid solutions. This means that a protective Cr-rich oxide is easily converted into poorly protective Fe-rich oxide by chromium vapourization. At low evaporation rates, the metal is able to sustain a high enough flux of chromium to the oxide in order for the oxide to retain its protective properties. When the evaporation rate exceeds a critical value, excessive chromium depletion of the oxide leads to the loss of its protective properties. Once the chromium-rich oxide is destroyed, rapid oxidation ensues. A qualitative mechanism was presented, explaining the differences in oxidation behaviour of Fe–Cr and Fe–Cr–Ni alloys in H2 O/O2 mixtures in terms of the supply of chromium from the substrate to the oxide. It was proposed that the resistance towards high temperature corrosion in this type of environment is enhanced by high Cr concentration, fast diffusion of Cr in the steel bulk (ferrite rather than austenite) and by a high density of grain boundaries in the alloy (small grain size). The previous investigations also showed that the corrosivity of the environment increases as the gas composition approaches the stoichiometry of reaction (1), and that watervapour-enhanced oxidation is promoted by high gas velocity and high temperature [7,9]. The present paper extends our investigation of the effect of H2 O/O2 mixtures on oxidation to the commercially important 310 alloy (for composition, see Table 1). This alloy has essentially the same microstructure, e.g., grain size, as 304L but a considerably higher chromium content (25% as opposed to 18%). We focus especially on whether the differences between this material and the previously studied 304L can be rationalized on the basis of the mechanism proposed.
2. Experimental In all experiments, 310-coupons with the size 15 15 2 mm were exposed (for chemical composition, see Table 1). The samples were ground on SiC-paper to 1000 mesh, polished with 1 lm diamond paste and then cleaned in acetone and ethanol using ultrasonic agitation.
2.1. Experimental set-up for the isothermal exposures The exposures were carried out in a horizontal furnace fitted with a 48 mm diameter SiO2 -glass tube. All exposures were isothermal, the temperature was kept at 600 3°. The samples were mounted three at a time using a sample holder consisting of a dense alumina-plate with the dimensions 40 50 5 mm, featuring three 1 mm deep and 2.1 mm broad slits. The distance between the slits was 12 mm. The samples were positioned parallel to the direction of flow. This set-up enables the samples to be exposed to the reaction gas with minimum contact with the sample holder and with optimal flow pattern and minimizes the disturbance caused by the presence of multiple samples in the reactor.
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Table 1 Chemical composition of 310 steel (wt%) 310
Cr (%)
Ni (%)
Mn (%)
Si (%)
Mo (%)
Fe (%)
24.9
19.2
1.55
0.45
0.34
Balance
2.2. Experimental conditions/flow rate dependence Exposures were made in a furnace system fitted with a humidifier producing a reaction gas consisting of O2 þ 10% H2 O or O2 þ 40% H2 O. After each exposure the samples were allowed to cool in dry air. The mass changes were recorded using a fivedecimal balance. Two types of experiments were made, mass change vs. time and mass change vs. flow rate for a constant exposure time. In the time-resolved experiments two different flow rates were used, 200 ml/min (NTP) and 2000 ml/min (NTP), corresponding to 0.50 and 5.0 cm/s, respectively. For the investigation of the effect of flow rate, a standard exposure time of 168 h was selected. Flow rates were 10, 200, 500, 1000, 2000 and 4000 ml/min (NTP). These flow rates correspond to an average net gas velocity inside the furnace of 0.03, 0.50, 1.25, 2.5, 5.0 and 10.0 cm/s, respectively. 2.3. Corrosion product characterization 2.3.1. X-ray diffraction (GI-XRD) Crystalline corrosion products were analyzed by X-ray diffraction (XRD) using a Siemens D5000 powder diffractometer equipped with grazing incidence beam attachment and a G€ obel mirror. CuKa radiation was used and the angle of incidence was 0.50°. The detector measures between 20° < 2h < 80°. 2.4. Analytical environmental scanning electron microscopy (ESEM/EDX) Analytical scanning electron microscopy was carried out using an Electro-scan 2020 equipped with a Link eXl EDX system. The microscope was operated at 20 kV for secondary electron imaging and at 10–20 kV for EDX analyses. 2.5. Scanning Auger electron microscopy (SAM) The Auger analyses were performed using a scanning Auger microprobe (PHI660). The electron beam voltage was 10 kV and the beam current was about 160 nA. The depth profiles were obtained using a differentially pumped ion gun (Arþ ) with acceleration voltages between 1.5 and 4 kV. The etch rates were calibrated on . The collected rawflat samples of Ta2 O5 with well-known oxide thickness of 1000 A data was refined into the form presented in this paper, by using MultiPakÓ v.6.0 software. The software makes it possible to evaluate peak shapes of differentiated Auger-peaks. This feature can be used in order to distinguish between oxidized and metallic iron and chromium signal, as these elements exhibit significant chemical
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shifts in oxidized form. The concentrations for O, Cr(oxidized) and Fe(oxidized) were calculated using sensitivity factors determined using standards of pure Cr2 O3 and Fe2 O3 and a range of solid solutions ((Cr,Fe)2 O3 ) [15]. 3. Results 3.1. Gravimetry Fig. 1 shows the marked flow rate dependence of the oxidation of 310 steel after 168 h in O2 , O2 þ 10% and 40% H2 O at 600 °C. In O2 þ 40% H2 O at 10 ml/min (0.03 cm/s), a small mass gain is registered. With increasing flow rate, mass gain decreases and at 500 (1.25 cm/s) and 1000 ml/min (2.5 cm/s) a small but significant mass loss is registered. Increasing the flow rate 2000 (5.0 cm/s) and 4000 ml/min (10.0 cm/s) resulted in a mass gain. The transition from mass loss to mass gain is abrupt and at high flow rates the mass gain is greater than in dry O2 . The dependence of mass change on flow rate in O2 þ 10% H2 O environment resembles that in O2 þ 40% H2 O. For example, the transition from mass loss to mass gain appears at the same flow rate. In comparison, the dry O2 exposures resulted in a mass gain that did not depend on flow rate. The results show that, in O2 þ H2 O environment, material is being lost from the samples by vapourization. Moreover it appears that, depending on flow rate, vapourization results in two distinct types of behaviour. At low flow velocities vapourization mainly results in a mass loss and does not seem to increase the rate of oxidation. In contrast, at high flow rates, mass increases indicating accelerated corrosion. In order to investigate the mass gain in the two regimes mentioned, a series of time-resolved exposures were carried out. The time dependence of mass change in O2 þ 40% H2 O was investigated at low flow (200 ml/min, 0.50 cm/s) and at high flow (2000 ml/min, 5.0 cm/s) (see Fig. 2). In the set of exposures at low flow, a mass gain is
Fig. 1. Mass-change vs. flow rate; 310 oxidized for 168 h at 600 °C in dry O2 , O2 þ 10% H2 O and O2 þ 40% H2 O atmosphere.
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Fig. 2. Mass-change vs. exposure time; 310 oxidized for up to 336 h at 600 °C in O2 þ 40% H2 O atmosphere with low gas-flow velocity 0.50 cm/s and high gas-flow velocity 5.0 cm/s.
registered that increases with time up to 72 h. After longer exposure times mass gain starts to decrease and at 336 h the samples exhibit a mass loss. In the experiments involving higher flow rates mass gains are registered at all exposure times. Mass gain increases with exposure time, levelling out after 168 h. 3.2. Phase analysis (GI-XRD) The phase analysis by GI-XRD indicated only one crystalline oxide, i.e., corundum-type M2 O3 corresponding to a-Fe2 O3 , Cr2 O3 or the solid solution (Cr,Fe)2 O3 . 3.3. Morphology (ESEM) In Fig. 3a a typical ESEM image of a sample oxidized at 600 °C in dry O2 for 168 h is presented. A thin and smooth oxide film covers the surface. The network pattern that can be seen is a projection of the grain boundaries in the alloy and can be seen clearly after sputtering when performing AES-depth profiling. Fig. 3b shows a sample exposed for 168 h in O2 þ 40% H2 O at low flow rate (200 ml/min corresponding to 0.50 cm/s). The morphology is very similar to the dry O2 exposure. It may be noted that the interference colours developed by the samples in Fig. 3a and b were different. The samples oxidized in dry O2 were deep blue while the ones oxidized in O2 þ 40% H2 O ranged from bluish grey to pale lilac. This indicates that the oxide films differ in thickness or in chemical composition. The morphology changes dramatically when samples are exposed to higher flow rates in O2 þ 40% H2 O environment (2000 ml/min, 5.0 cm/s) (see Fig. 3c). On top of the thin oxide, ‘‘islands’’ occur, consisting of thick oxide. It may be noted that the oxide islands never cross the network pattern mentioned above. This means that the oxide islands only form on the part of the sample surface that corresponds to the
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Fig. 3. ESEM pictures of 310 oxidized for 168 h at 600 °C: (a) O2 , (b) O2 þ 40% H2 O flow: 200 ml/min (0.50 cm/s), (c) O2 þ 40% H2 O flow: 2000 ml/min (5.0 cm/s) and (d) same as (c) but at higher magnification.
interior of the alloy grains. Oxide islands are not formed on every steel grain. However, a closer look at some of the steel grains lacking oxide islands, shows that a thicker oxide has started to form (see Fig. 3d). When samples are exposed at even higher flow rate in this environment (>2000 ml/min, 10.0 cm/s), the sample morphology remains qualitatively the same while the number density of oxide islands increases. At flow rates up to 1000 ml/min (2.5 cm/s), the surface morphologies developed in O2 þ 10% H2 O were essentially the same as described above for O2 þ 40% H2 O (see Fig. 4a). At higher flow rates, a transition from protective behaviour to a local breakdown of protective properties occurs in O2 þ 10% H2 O. However, in this case the oxide islands only occur sporadically (see Fig. 4b). 3.4. Chemical composition (AES) Figs. 5–8 show AES-depth profiles of samples exposed for 168 h. The ESEM images of the samples in Figs. 5–7 are not included as the surfaces were smooth and featureless. Figs. 5a, 6a and 7a show the chemical composition of the oxide film on samples oxidized in dry O2 , and in O2 þ 40% H2 O at 10 and 1000 ml/min (0.03 and
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Fig. 4. ESEM pictures of 310 oxidized for 168 h at 600 °C: (a) O2 þ 40% H2 O flow: 1000 ml/min (2.5 cm/s) and (b) O2 þ 10% H2 O flow: 2000 ml/min (5.0 cm/s), respectively.
2.5 cm/s) flow rate, respectively. Figs. 5b, 6b and 7b show the Cr/cation ratio in the oxide and the composition of the alloy. The chromium concentration in a(Cr,Fe)2 O3 , is calculated from CðCr2 O3 Þ ¼ ðCCrðoxÞ =ðCCrðoxÞ þ CFeðoxÞ Þ 100%Þ, where CMðoxÞ is the concentration of the Fe(ox) and Cr(ox) signal separated from the Fe(metal) and Cr(metal) by MultipakTM v.6.0 software as described in Section 2. The Fe(metal) and Cr(metal) signal, in turn, are combined with the nickel signal (Ni is assumed to be only in non-oxidized form) using the formula CMðalloyÞ ¼ ðCMðmeÞ = ðCCrðmeÞ þ CFeðmeÞ þ CNi Þ 100%Þ (M ¼ Cr, Fe and Ni). This makes it possible to plot the alloy composition (CCrðalloyÞ , CFeðalloyÞ and CNiðalloyÞ ) as a function of the distance from the oxide–gas interface. Fig. 5a and b show that the 80 nm oxide film formed in dry O2 consists of a a(Cr,Fe)2 O3 solid solution with about 70% Cr and that the steel substrate beneath the oxide is only marginally depleted in chromium. The oxides formed in O2 þ 40% H2 O exposures have a different chemical composition (see Figs. 6 and 7). Even at the lowest flow rate (10 ml/min, 0.03 cm/s), the outer part of the oxide is depleted in chromium (see Fig. 6a and b). The Cr=Cr þ Fe ratio in the oxide is seen to be about 30% at the oxide–gas interface while the inner parts of the oxide remains Cr-rich. In addition, there is a marked depletion of chromium in the alloy while iron and nickel
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Fig. 5. AES profile of the oxide formed on 310 oxidized for 168 h at 600 °C in dry O2 : (a) the pure depthprofile and (b) the Cr concentration in (Cr,Fe)2 O3 , and the metallic Cr, Fe and Ni as a function of depth.
are enriched. When the flow rate is increased to 1000 ml/min (2.5 cm/s) (see Fig. 7), the chromium depletion in the oxide and in the substrate becomes more marked. As a consequence, the Ni concentration in the substrate immediately below the oxide reaches about 30 at.%, (compare the bulk value in the steel of 25%). Film thickness is also affected by water vapour. In dry O2 film thickness is about 80 nm. In comparison, the samples exposed in O2 þ 40% H2 O at a flow rate of 10 ml/min, produced a 60 nm oxide while the oxide formed in the same environment at a flow rate of 1000 ml/min was only 40 nm. This indicates that the presence of water vapour tends to etch the oxide and produce a thinner oxide rather than accelerating the growth of a(Cr,Fe)2 O3 . In O2 þ 40% H2 O and very high flow rates (P 2000 ml/min in the case of O2 þ 40% H2 O) the morphology, thickness and chemical composition of the oxidized surface changes dramatically. In Fig. 8 an ESEM image of a sample that was
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Fig. 6. AES profile of the oxide formed on 310 oxidized for 168 h at 600 °C in O2 þ 40% H2 O flow: 10 ml/ min (0.03 cm/s): (a) the pure depth-profile and (b) the Cr concentration in (Cr,Fe)2 O3 and the metallic Cr, Fe and Ni as a function of depth.
oxidized at 600 °C for 168 h in O2 þ 40% H2 O at a flow rate of 2000 ml/min (5.0 cm/ s) is shown. In addition, AES profiles are shown, representing different characteristic features on the surface. It can clearly be seen that the outer part of the oxide between the oxide islands is almost completely depleted in chromium. In addition, the inner part of the oxide is also considerably depleted in chromium compared to the previously described oxides formed in O2 þ H2 O environments at low flow rate. The oxide between the oxide islands has a thickness of about 250 nm, which is about three times the value in dry O2 and six times thicker compared to the oxide formed in O2 þ 40% H2 O at 1000 ml/min. The thickness of the oxide islands is about 5 lm. The islands exhibit a layered structure, the outer part being made up of iron-rich oxide while the inner part contains considerable amounts of chromium. The combined evidence from XRD and Auger spectroscopy shows that the outer part of the islands is iron-rich (Fe,Cr)2 O3 . Even though XRD does not indicate that other oxides are
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Fig. 7. AES profile of the oxide formed on 310 oxidized for 168 h at 600 °C in O2 þ 40% H2 O flow: 1000 ml/min (2.5 cm/s): (a) the pure depth-profile and (b) the Cr concentration in (Cr,Fe)2 O3 and the metallic Cr, Fe and Ni as a function of depth.
present, it is suggested that the inner, chromium-rich part of the oxide consists of iron-chromium spinel (Fe,Cr)3 O4 .
4. Discussion Before turning to the present results concerning 310 steel, it is useful to recapitulate the recently reported oxidation behaviour of the 304L alloy in O2 /H2 O environment and the mechanism proposed to rationalize the observations made [6,7,9]. Similarly to the 310 alloy, the 304L material forms a thin (50–100 nm) protective oxide consisting of chromium-rich a-(Cr,Fe)2 O3 when exposed to dry O2 at 600 °C for 1–4 weeks. In O2 /H2 O environment at the same temperature, the oxidation behaviour depended on flow rate. At low flow rates, the oxide showed signs of chromium depletion while the protective properties remained intact. At higher flow rates a gradual and localized breakdown of the protective properties was observed. The flow rate dependence was explained in terms of the water-vapour-enhanced chromium vapourization mentioned in the introduction. The evaporation of chromium
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oxy-hydroxides from the surface depletes the protective oxide in chromium. As already noted, the evaporation rate is expected to depend on the equilibrium partial pressure of the volatile species and on gas velocity. At low flow rates, corresponding to low evaporation rates, the diffusion of chromium from the steel substrate is sufficiently rapid so that the chromium content in the oxide is not seriously lowered. As the evaporation rate increases, the diffusion of chromium from the alloy eventually becomes insufficient and the chromium concentration in the oxide drops below a critical value. When this happens, the oxide loses part of its protective properties and starts to behave similar to a-Fe2 O3 , resulting in a dramatic increase in oxidation rate. It was observed that the breakdown of protective properties of 304L always occurred on the parts of the surface corresponding to the middle of the alloy grains, resulting in the formation of hematite islands of several lm thickness. This was attributed to a poor supply of chromium from the alloy. Because the diffusion rate is high in the grain boundaries compared to the interior of a crystal, the oxide formed on the parts of the surface corresponding to the interior of the alloy grains tends to be more susceptible to breakdown compared to the oxide in the vicinity of a grain boundary. Similarly, the tenacity of the thin protective chromium-rich a-(Cr,Fe)2 O3 , layer separating the oxide islands was attributed to the greater supply of chromium close to the alloy grain boundaries. At still higher flow rates the 304L material also suffered a partial degradation of the protective oxide between the oxide islands, causing this part of the oxide to grow to about 700 nm thickness. Based on these observations it was suggested that the ability of Fe–Cr alloys to withstand oxidation in H2 O/O2 environment is connected to the capacity of the alloy to compensate for the losses of chromium from the oxide by vapourization. This implies that oxidation resistance is related to chromium content and to the diffusivity of chromium in the alloy. The 310 and 304L steels are quite similar, both alloys being austenitic and having almost the same microstructure. Therefore we expect a relatively low bulk diffusivity of chromium as well as the same density of rapid diffusion paths (primarily grain boundaries) in both materials. This implies that we expect qualitatively the same oxidation behaviour for the two alloys in H2 O/O2 environment. However, as the 310 alloy contains more chromium (25% as compared to 18% for 304L) it is predicted to be somewhat less susceptible to oxidation. This means that while we expect a similar localized breakdown of the protective oxide on the 310 alloy due to water-vapourenhanced chromium evaporation, a higher flow rate is expected to be necessary in order to trigger this behaviour. The present study clearly shows the effect of water-vapour-enhanced chromium evaporation (see Figs. 1, 5 and 6). At low flow rate, the samples exposed in H2 O/O2 environment develop an oxide, the outer part of which is markedly depleted in chromium compared to the dry O2 exposure. The outer part of the oxide consists of almost pure a-(Fe2 O3 ), sometimes with a sporadic presence of Mn. Still, the thin and smooth oxide lacking in morphological features reveals that, in spite of the partial chromium depletion, the oxide formed on 310 in H2 O/O2 environment at low flow rate is completely protective, at least up to 168 h. The oxide formed on the 310 alloy in H2 O/O2 environment at low flow rates has the same chemical composition and
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thickness as the thin and protective oxide that formed close to the alloy grain boundaries on 304L [7]. As the flow rate increases it causes chromium vapourization from the 310 alloy to accelerate (see Fig. 1). This is also evident from the Auger depth profiling (see Figs. 6 and 7) As predicted, the flow rate necessary for triggering the deterioration of protective properties of the oxide is greater than in the case of the 304 alloy [7]. In the case of the 310 alloy the critical flow rate that causes the oxide to loose its protective properties in O2 þ 40% H2 O at 600 °C is between 1000 and 2000 ml/min (see Figs. 1 and 8). For alloy 304L in the same environment, the transition from protective to partly non-protective behaviour occurred close to 200 ml/min [7]. The transition from protective to non-protective behaviour is more sudden for the 310 alloy than for the 304L material. In the former case local breakaway corrosion, resulting in the formation of oxide islands, and the partial loss of the protective properties of the thin oxide between islands occurs at the same flow rate. The oxide islands that form have a thickness in the micrometer range (304L, 5–20 lm and 310, 5 lm). The islands exhibit a layered structure with an outer iron-rich corundum-type (Cr,Fe)2 O3 and an inner, more Cr-rich layer possibly consisting of (Fe,Cr)3 O4 . The (Cr,Fe)2 O3 , oxide separating the oxide islands on the 310 alloy exposed to H2 O/O2 environment at high flow rates is about 250 nm on the 310 alloy (compared to about 50 nm at low flow rates). The partial loss of the protective properties of this part of the oxide is
Fig. 8. AES profiles of the oxide formed on 310 oxidized for 168 h at 600 °C in O2 þ 40% H2 O flow: 2000 ml/min (5.0 cm/s).
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also attributed to the chromium depletion caused by water-vapour-enhanced chromium vapourization (see Fig. 8). However, due to the greater supply of chromium to this part of the oxide by grain boundary diffusion in the alloy, chromium loss only results in a partial loss of protective properties. The oxide morphology developed on the 310 alloy in O2 þ 10% H2 O is essentially the same as in O2 þ 40% H2 O. The main difference is that in O2 þ 10% H2 O environment, a higher flow rate is needed to cause the protective properties of the oxide to breakdown resulting in the formation of oxide islands. To summarize, the oxidation behaviour of the 310 material is very similar to that of the previously studied 304L alloy. However, the higher chromium content makes the 310 alloy less sensitive towards the destabilization of the protective properties of the oxide by water-vapour-enhanced chromium vapourization. When chromium vapourization is sufficiently rapid (high water vapour content and high flow rate) the 310 alloy suffers a partial breakdown of the protective oxide. It is suggested that when chromium vapourization becomes sufficiently rapid, (high water vapour content and flow rates in the m/s range) both materials will suffer a more or less complete breakdown of the protective oxide at 600 °C. Such conditions may appear in some industrial applications, e.g., in biomass combustion.
5. Conclusions The addition of H2 O to O2 results in a change of oxidation behaviour of the 310 alloy. A strong dependence on flow rate was observed in O2 /H2 O mixtures. At low flow rates a thin (30–50 nm) protective a-(Cr,Fe)2 O3 , forms, the outer part being depleted in chromium. When the flow rate is increased beyond a critical value the protective oxide fails. Under these conditions P 5 lm thick a-Fe2 O3 /(Cr,Fe)3 O4 , oxide islands forms on the part of the surface corresponding to the centre of the alloy grains. The a-(Cr,Fe)2 O3 oxide close to the alloy grain boundaries remained partially protective. The effect of water vapour is attributed to the water-vapour-assisted evaporation of chromium from the oxide, in the form of a chromium oxide hydroxide, probably CrO2 (OH)2 . The oxidation behaviour is rationalized using a qualitative mechanism proposed previously and parallels that of the 304L alloy. The mechanism implies that the ability of Fe–Cr alloys to withstand oxidation in H2 O/O2 environment is connected to the capacity of the alloy to compensate for the losses of chromium from the oxide by vapourization. The slightly better resistance of the 310 alloy as compared to 304L is attributed to the higher chromium content of 310. The present results are suggested to be relevant to environments with high concentrations of H2 O and O2 , e.g., in biomass combustion.
Acknowledgements This work was carried out within the Swedish High Temperature Corrosion Center (HTC).
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