Oxidation of Stainless Steel in H2O/O2 Environments
– Role of Chromium Evaporation
H. Asteman1, K. Segerdahl1, J.-E. Svensson1, L.-G. Johansson1,
M. Halvarsson2 and J.E. Tang2
1
Department of Environmental Inorganic Chemistry and 2Department of Experimental Physics
The High-Temperature Corrosion Centre
Chalmers University of Technology
S-412 96 Göteborg, Sweden
Keywords: Stainless Steel, Water Vapour, Evaporation, CrO2(OH)2 (g), Breakaway Corrosion
Abstract. We report on the effect of water vapor on the oxidation of ferritic/martensitic and
austenitic chromia-forming steels at 600°C. The samples are investigated by thermogravimetry, GIXRD, SEM/EDX, SAM, GDOES, FIB and TEM. The oxidation Fe/Cr and Fe/Cr/Ni steels in
H2O/O2 environment is strongly affected by the vaporization of CrO2(OH)2. The oxidation behavior
at a certain temperature, gas composition and gas velocity is to a large extent determined by the
ability of the metallic substrate to supply the oxide with chromium to compensate for the losses by
vaporization. The corrosion resistance of chromia-forming steels is enhanced by high Cr
concentration, fast diffusion in the steel bulk (ferrite/martensite rather than austenite), and a high
density of steel grain boundaries (small grain size). The corrosivity of the environment increases
with the concentration of water vapor and oxygen, with the gas velocity and with temperature.
Introduction
Mixtures of water vapour and oxygen are known to be more corrosive towards stainless steel
compared to dry O2 [1]. It is important to understand the reasons for this behaviour since stainless
steels are frequently exposed to H2O/O2 environment at high temperature. One example is the
growing importance of biomass combustion for electricity generation. In this type of environment
the oxygen and water vapour content in the combustion gas is typically 5 and 20-40%, respectively.
The attempts made in the literature to explain the high corrosivity of H2O/O2 mixtures towards
chromia-forming steels may be classified according to the suggested role of water. Some
researchers claim that the transport of iron through the oxide scale is enhanced by the formation of
gaseous iron hydroxide in cracks in the scale [2]. Others propose that oxygen transport through the
scale is enhanced by gaseous water molecules transporting oxygen atoms through cracks and pores
[3]. Still others propose that water vapour changes the solid-state diffusion properties of the oxide.
For example, it was suggested that the diffusion rate of chromium through chromium oxide (or
through a Cr-rich oxide) increases in the presence of water vapour [4]. Recent work by this group
uses a different approach, explaining the corrosivity of H2O/O2 environments in terms of the
volatilisation of chromium from the oxide [5-7]. The poor oxidation behaviour of chromia-forming
alloys in dry oxygen >1000°C is attributed to the loss of chromium from the oxide by CrO3
evaporation [8]. Recently, it was shown that the exposure of FeCrNi alloys in H2O/O2 mixtures
results in chromium evaporation already at 600°C [5,6]. It was concluded that the vaporizing
species is a chromium (VI) oxide hydroxide, probably CrO2(OH)2(g):
1/2Cr2O3(s) + 3/4O2(g) + H2O(g) ! CrO2(OH)2(g)
(1)
The vaporization of chromium was suggested to enhance the oxidation of FeCr alloys because it
depletes the oxide in chromium, resulting in poorer protective properties of the oxide scale. Because
hematite, Fe2O3, and Cr2O3 form a continuous range of solid solutions, a protective Cr-rich oxide is
easily converted into a poorly protective Fe-rich oxide by chromium vaporization. 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 and in the alloy beneath the oxide leads to the loss of its
protective properties. Once the chromium-rich oxide is destroyed, rapid oxidation ensues. The
present paper summarizes our investigations of the effect of H2O/O2 mixtures and chromium
evaporation on the oxidation of three commercially important alloys, the austenitic 304L and 310
and the ferritic/martensitic alloy CrMoV 11 1 (X20) (for compositions, see Table 1).
Experimental
Table 1. Elemental composition of the alloys studied in wt%
X20
304L
310
Cr
11.0
18.5
24.9
Ni
0.65
10.2
19.2
Mn
0.45
1.41
1.55
Si
0.25
0.55
0.45
Mo
0.92
0.49
0.34
Fe
Bal
Bal
Bal
The samples were cut to 15x15x2mm size, finished by polishing with 1!m diamond paste and
carefully cleaned. All exposures were isothermal and were carried out in a horizontal SiO2 glass
reactor. Pure O2 or O2 + 10 or 40% H2O were passed through the reactor with flow velocities from
0.03 to 10 cm/s. Exposure temperature was 600° " 3°C. Three samples were exposed at a time, the
samples being positioned parallel to the direction of flow. Two types of experiments were made,
mass change vs. time and mass change vs. gas velocity for a constant exposure time (168 hours).
Exposed samples were analyzed by X-ray diffraction (XRD) with a grazing incidence beam
attachment and a Göbel mirror. Analytical scanning electron microscopy was carried out using an
Electro-scan 2020 instrument equipped with a Link eXl EDX system. Depth profiling was
performed by a scanning Auger microprobe (PHI660). It was possible to distinguish between the
signals from oxidized and metallic iron and chromium. The concentrations for O, Cr(oxidized) and
Fe(oxidized) were calculated using sensitivity factors determined using standards of pure Cr2O3 and
Fe2O3 and a range of solid solutions ((Cr,Fe)2O3 ) [9]. GDOES-depth profiling was used to provide
information on the chemical composition of the oxide. The GDOES analyses were performed using
a GDS850A instrument with a Grimm-type lamp. The instrument has a 0.75-meter focal length
Paschen-Runge polychromator with 58-channel capability.
Results and discussion
The corrosivity of water vapor/oxygen mixtures is illustrated in Fig. 1, showing mass gain as a
function of exposure time for the ferritic/martensitic CrMoV11 1 (X20) material in O2, O2 +
10%H2O and O2 + 40% H2O at 600°C. In dry oxygen, the sample mass increases slowly during
exposure. In O2 + 10% H2O, a small mass loss is registered (not detectable in the figure). In O2 +
40% H2O, the samples lose mass in the early part of the exposure while there is a huge increase in
mass after longer times corresponding to breakaway corrosion. The mass gain after breakaway is
about 30 times that in the dry O2 runs. The mass loss detected in H2O/O2 environment is explained
by the loss of chromium in the form of CrO2(OH)2 (g) (see reaction (1)).
Mass change (mg/cm2)
3
2.5
O2+40%H2O
O2+10%H2O
Dry O2
2
1.5
1
0.5
0
-0.5
0
100
200
300
400
Time (hours)
Fig. 1. Mass change as a function of exposure time for CrMoV 11 1 oxidized in dry O2, O2 + 10%
H2O and O2 + 40% H2O at 600°C. Gas velocity was 0.5 cms-1.
XRD shows that the alloy forms a corundum-type oxide in dry O2. The GDOES depth profiles in
Fig. 2 show the changes in oxide composition caused by chromium vaporization. In dry O2, the
oxide (<0.1µm) consists of (Fe1-xCrx)2O3 with x = 0.7. The oxide formed in the presence of 10%
H2O is also quite thin. In this case, chromium vaporization depletes the outer part of the oxide in
chromium while the inner part is still rich in Cr. The effect of rapid chromium loss is seen in the O2
+ 40% H2O runs. In this case the thin protective oxide has been replaced by a thick porous scale.
The outer part of the scale consists of hematite (#-Fe2O3) with a few % of chromium. The inner part
consists of Fe-Cr spinel (Fe3-xCrxO4).
100
100
(a)
(b)
80
O
Fe
Cr
Mn
60
40
AC%
AC%
80
20
60
O
Fe
Cr
Mn
40
20
0
0
0.05
0.1
0.15
Depth (µm)
100
0.25
0
0
0.03
0.06
0.09
Depth (µm)
0.12
0.15
(c)
80
AC%
0.2
O
Fe
Cr
Mn
60
40
20
0
0
0.05
0.1
Depth (µm)
0.15
0.2
Fig. 2. GDOES analysis of CrMoV 11 1 steel oxidized in dry (a) O2, (b) O2 + 10% H2O and (c) O2
+ 40% H2O at 600°C for 168 hours.
The observations indicate that the loss of chromium by vaporization in H2O + O2 environment
tends to destroy the protective oxide. If that is the case, the rate of vaporization must be a decisive
factor. One way to test this is to investigate the effect of gas velocity. The rate of evaporation
increases with gas velocity because the thickness of the diffusion layer in the gas close to the
sample surface diminishes with increasing gas velocity. Fig. 3 shows mass change as a function of
gas velocity for X20 steel exposed in O2+40% H2O at 600°C for 168 hours. At low gas velocities,
chromium evaporation (evidenced by a small mass loss) does not trigger the destruction of the
protective oxide. At a certain critical gas velocity, evaporation becomes so rapid that it destroys the
protective chromium-rich corundum-type oxide, replacing it by a thick layered Fe2O3 / Fe3-xCrxO4
scale.
0.05
(a)
2
2.5
Mass change (mg/cm
)
2
Mass change (mg/cm
)
3
O2+40%H2O
2
1.5
1
O2
0.5
0
-0.5
0
2
4
6
8
Gas velocity (cm/s)
10
12
(b)
O2+40%H2O
0.04
0.03
0.02
O2
0.01
0
-0.01
0
2
4
6
8
10
12
-0.02
Gas velocity (cm/s)
Fig. 3. Mass change as a function of flow rate for CrMoV 11 1 oxidized in dry O2, O2 + 10% H2O
and O2 + 40% H2O at 600°C. Exposure time was 168 hours.
Austenitic stainless steels respond to chromium vaporization by local, rather than by a general
failure of the protective oxide. Fig.4 shows alloy 310 after 168 hours at 600°C in dry O2 and O2 +
40% H2O, respectively. After the dry O2 run, the oxide is thin and smooth. Auger depth profiling
shows a 80nm thick protective oxide which is enriched in chromium. In O2 + 40% H2O, the oxide
morphology depends on flow rate. At low flow rates the oxide is thin, smooth and fully protective.
Auger depth profiling reveals the effect of chromium vaporization, showing a strong chromium
gradient through the oxide. At higher flow rates the oxide develops a characteristic “island”
morphology, the oxide islands being 5-10 µm thick and forming on the alloy grain centers. Auger
depth profiling of the thin oxide between the islands shows an even stronger chromium depletion in
this case. The composition and morphology of the oxide islands is essentially the same as in the
case of alloy 304L that is described in detail below.
Fig. 5 shows mass gain as a function of gas velocity for alloy 304L after 168 hours at 600°C in
dry O2 and O2+40%H2O, respectively. SEM images of the oxidized surface are also shown. In
comparison to alloy 310, the main difference in behaviour was that the island-type oxidation is
triggered at lower flow velocities in the case of 304L. The shape of the islands depends on gas
velocity. At 1.25cm/s they are porous while the oxide “islands” formed at 10 cm/s are denser. The
network pattern in the thin oxide between the islands shows the alloy grain boundaries.
0.03
)
2
M ass-change (m g/cm
O2+40%H2O
0.025
0.02
O2
0.015
0.01
0.005
0
-0.005
0
2
4
6
8
10
12
-0.01
Gas-velocity (cm/s)
100
100
Cr
oxide
40
60
40
20
20
0
0
0
50
100
150
200
250
300
Cr
oxide
60
Cr
oxide
40
0
0
50
100
150
200
250
300
0
100
200
300
400
%Fe in substrate
500
600
700
800
Depth (nm)
Depth (nm)
%Cr2O3 in (Cr,Fe)2O3
%Cr2O3 in (Cr,Fe)2O3
%Cr in substrate
%Cr in substrate
!"#$%&%'()*+,#-"#./0%1)2,#3"45"#$6%2//72
)2(#849"#37*)2//72
%Ni in substrate
Cr
alloy
200nm
20
Depth (nm)
%Cr2O3 in (Cr,Fe)2O3
%Fe in substrate
80
AC%
60
100
Cr
alloy
60nm
80
AC%
AC%
Cr
alloy
80nm
80
%Ni in substrate
%Fe in substrate
%Cr in substrate
%Ni in substrate
Fig. 4. Dependence of mass change, oxide morphology and oxide composition on gas velocity for
alloy 310 after 168 hours exposure in dry O2 and in O2 + 40% H2O at 600°C.
2
Mass change (mg/cm )
0.4
0.3
0.2
0.1
0
0
2
4
6
8
10
Flow rate (cm/s)
Fig. 5. Mass gain as a function of gas velocity for alloy 304L exposed 168h at 600°C in dry O2 and
O2 + 40% H2O
Figs. 6 shows FIB images of alloy 304L after 168 hours exposure to O2 + 40% H2O at 600°C at
high gas velocity. Oxide islands of different size have formed on the surface. As in the case of alloy
310, the islands always form on the alloy grain centers. The oxide islands exhibit a layered
structure. Electron diffraction and TEM/EDX reveal that the outer layer consists of hematite
containing small amounts of chromium while the inner oxide layer is made up of iron-chromium
spinel with significant amounts of nickel (Fe3-x-yCrxNiyO4). The thin oxide film that covers the
surface between the oxide islands consists of #-(Fe1-xCrx)2O3. It is similar to that formed in dry O2,
except that its outer parts are chromium depleted.
Cr (atomic% of metal content)
100
t%)
Conc (A
Alloy
Oxide
80
60
40
20
Ni (atomic% of metal content)
0
400
Le
)
nm
100
0
100
De o
th (n
400
500
m)
Fe (atomic% of metal content)
60
40
20
0
400
ng
)
Oxide
80
300
200
)
( nm
100
t%
Conc (A
th
Alloy
100
0
60
Alloy
Oxide
80
Le
(Cr,Fe)3O4
0
200
300
)
(
th
Fe2O3
100
200
t%
Conc (A
ng
300
0
100
200
Dep
300
400
500
m)
th (n
40
20
0
400
Le
ng
300
th
200
)
( nm
100
0
0
100
200
Dep
300
400
500
m)
th (n
Fig. 6. A focussed ion beam image of an oxide island on alloy 304L exposed to O2 + 40% H2O for
168 h at 600°C. Gas velocity 10 cm s-1. The graphs show contour maps for Cr, Fe and Ni (at%
of the metal content) obtained from TEM/EDX linescans.
Fig. 7 shows a TEM image of the alloy-oxide interface in the vicinity of the oxide island shown
in Fig. 6 together with the corresponding elemental distribution from EDX. The elemental maps
show depletion of chromium and enrichment in nickel at the alloy grain boundary and just below
the thin oxide. As in the case of alloy 310, the thin oxide is enriched in chromium. The distribution
of chromium in the steel substrate indicates that the grain boundaries act as rapid diffusion paths for
chromium in the alloy. We propose that the ability of the oxide formed close to an alloy grain
boundary to withstand chromium vaporization without losing its protective properties is explained
by the relatively rapid supply of chromium in this region.
b b
Pt
steel
BFI
Fe
200 nm
Cr
Ni
Fig. 7. Bright field TEM image of the base oxide and a 304L steel region containing a grain
boundary (indicated by arrows) presented with corresponding EDX element maps of Cr, Ni
and Fe.
We propose that the corrosivity of O2/H2O environments towards stainless steel can be
understood in terms of the vaporization of chromium (VI) oxide hydroxide and its effect on the
chromium content of the protective oxide. It is argued that the competition between the loss of
chromium by vaporization and the supply of chromium to the oxide by diffusion in the alloy is one
of the major factors that determine the oxidation behaviour of the alloy.
The stainless steels studied respond to chromium evaporation by three distinct types of oxidation
behavior. In the first case, (type I) the supply of chromium from the substrate to the oxide is rapid
enough to compensate for the loss of chromium by evaporation. As a result of the favorable balance
between chromium loss and supply, the chromium-rich corundum type oxide remains stable and
chromium vaporization does not speed up oxidation. This behavior is exemplified by alloys 310 and
CrMoV11 1 in O2 + 40% H2O at 600°C using low gas velocities. It may be described in terms of
active corrosion. A mass loss is registered that tends to increase with the gas velocity (compare
Figs.1 and 4).
In the second case (type II) chromium evaporation is rapid in comparison to the supply of
chromium from the alloy. This causes the concentration of chromium in the oxide to fall below a
critical limit, resulting in a breakdown of the protective oxide all over the surface. The chromiumrich corundum type oxide is replaced by a layered oxide featuring a Fe3-xCrxO4 spinel oxide covered
by hematite containing only traces of chromium. This behavior is shown by CrMoV11 1 steel at
high gas velocities in O2/H2O environment (Fig. 1).
The third case (type III) may be considered to be intermediate between the previous two. In this
case there is a strong lateral variation over the surface in the supply of chromium to the oxide,
resulting in a local failure of the protective oxide. This results in the characteristic oxide island
morphology exhibited by the austenitic 304L in O2 + 40% H2O at 600°C and by the austenitic alloy
310 in the same environment at high gas velocities (Fig. 4-6). The oxide islands form on top of the
alloy grain centers because the supply of chromium to the oxide is slow. In contrast, the oxide close
to the alloy grain boundaries remains intact because the grain boundaries act as rapid diffusion
paths for chromium in the alloy (see Fig. 7). The oxide islands cover craters in the alloy, the upper
part consisting of iron-rich corundum-type oxide while the “crater” oxide consists of Fe3-xyCrxNiyO4 spinel. It is notable that the layered structure of the oxide islands is very reminiscent of
the thick layered scale formed in type II oxidation.
Conclusion
The vaporization of chromium (VI) oxyhydroxide has important implications for the oxidation of
FeCr alloys in O2/H2O environments. The loss of chromium tends to destroy the protective
chromium rich corundum-type oxide giving rise to rapid oxidation. The mode of failure depends on
alloy type. The ability of a chromia-forming steel to retain its protective properties in O2/H2O
environments is determined by the chromium concentration, by the diffusivity of Cr in the steel
bulk and by the occurrence of rapid diffusion paths for chromium in the alloy.
Acknowledgements
Financial support from the Swedish Research Council Vetenskapsrådet is gratefully acknowledged.
This work was carried out within the Swedish High Temperature Corrosion Center (HTC).
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