-
2 2
measured thermogravimetrically and the oxide scale growth
mechanisms were studied using
* Corresponding author. Tel.: +49 2461 614668; fax: +49 2461
613687.E-mail address: [email protected] (W.J.
Quadakkers).
Corrosion Science 48 (2006) 34283454
www.elsevier.com/locate/corsci0010-938X/$ - see front matter
2006 Elsevier Ltd. All rights reserved.H218O-tracer with subsequent
analyses of oxide scale composition and tracer distribution by
MCs+-SIMS depth proling. The corrosion products were
additionally characterised by light opticalmicroscopy, SEM-EDX and
XRD. It was found that the transition from protective, Cr-rich
oxideformation into non-protective mixed oxide scales is governed
by the ratio H2O
(g)/O2 ratio ratherthan the absolute level of H2O
(g). The results of the tracer studies in combination with the
dataobtained from experiments involving in situ gas changes clearly
illustrated that under the prevailingconditions the penetration of
water vapour molecules triggers the enhanced oxidation and
sustainsthe high growth rates of the poorly protective Fe-rich
oxide scale formed in atmospheres with highH2O
(g)/O2 ratios. The experimental observations can be explained if
one assumes the scale growth tobe governed by a competitive
adsorption of oxygen and water vapour molecules on external
andinternal surfaces of the oxide scales in combination with the
formation of a volatile Fe-hydroxideduring transient oxidation. The
formation of the non-protective Fe-rich oxide scales is
suppressedin atmospheres with low H2O
(g)/O2 -ratios, and the healing of any such scale is promoted.
2006 Elsevier Ltd. All rights reserved.Enhanced oxidation of the
9%Cr steel P91in water vapour containing environments
J. Ehlers a, D.J. Young b, E.J. Smaardijk a, A.K. Tyagi c,H.J.
Penkalla a, L. Singheiser a, W.J. Quadakkers a,*
a Forschungszentrum Julich, Institute for Materials and
Processes in Energy Systems, 52425 Julich, Germanyb University of
New South Wales, Sydney, Australia
c Indira Gandhi Centre for Atomic Research, Kalpakkam, India
Received 11 February 2005; accepted 20 February 2006Available
online 18 April 2006
Abstract
The short term (100 h) oxidation behaviour of the 9%Cr steel P91
was studied at 650 C in N2O H O gas mixtures containing a
relatively low oxygen level of 1%. The oxidation kinetics
weredoi:10.1016/j.corsci.2006.02.002
-
model N2O2H2O gas mixtures at 650 C. In most experiments the
oxygen content was1 vol%, i.e. a value similar to that present in
the combustion gases mentioned above
[2,4,6]. Some new experimental procedures, including the use of
H2
18O tracer, were usedto obtain better insight into the
mechanisms of the enhanced oxidation of chromium steelsKeywords:
Steel; SIMS; TEM; Selective oxidation; Kinetic parameters
1. Introduction
Due to the demand for lower emissions from power generation
systems, a number ofprojects are being carried out world-wide to
improve eciencies of conventional fossilfuel-red power plants
[1,2]. In coal red boilers, eciencies of around 45% can beachieved
if the steam parameters are increased to pressures of 300 bar and
temperaturesof 600650 C [3]. At such high temperatures, the
commonly used low alloy steels andthe higher corrosion resistant
12%Cr steels can no longer be used as construction materialsfor
live steam piping or blading materials in steam turbines, because
of the lack of creepresistance of these materials. Therefore, a
number of modied 9%Cr steels, such as P91,P92 and E911 were
developed to full the new materials requirements in respect to
creepstrength [3]. It has been shown that, in spite of the high
temperatures, these steels alsoshow adequate oxidation resistance
during operation in air [4]. However, it was found thatin simulated
fossil fuel-red power plant combustion gases which contained oxygen
in theorder of 1 vol%, the corrosion rates of these 9%Cr steels can
be several orders of magni-tude higher than in air [46].
Consequently, as thin walled components, they oer nomajor benet
over 12%Cr steels in spite of their signicantly higher creep
strength [6].
The main reason for the high corrosion rates of the 9%Cr steels
in the simulated com-bustion gases was shown [4] to be the presence
of water vapour (typically 715 vol%). Thisdetrimental eect of water
vapour on the oxidation resistance of FeCr alloys has in factbeen
known for many years, [79] and a number of mechanisms have been
proposed byseveral authors to explain the eect. A summary of the
mechanisms proposed in the earlierstudies is given in Ref. [4].
In recent years, the development and construction of power
generation systems withincreased steam parameters has led to a
revival of the research on water vapour eectsin steel oxidation.
The newer studies relate to exposures in mixtures of water vapour
plusoxygen (see e.g. [1017]) as well as to steam-containing
environments to which no oxygenis intentionally added (see e.g.
[1822]). Newer studies on water vapour eects in case ofother types
of metallic materials have been described, e.g. in [23,24].
In spite of these extensive investigations, the numerous
experimental observations haveto date been only partly explained by
the various mechanisms proposed. The diculties innding conclusive
explanations are probably related to the fact that a number of
dierentsteps in the oxidation process may be aected by presence of
water vapour. The rate deter-mining steps in the overall oxidation
process may dier depending on the type of watervapour containing
gas, e.g. depending on the content of free oxygen in the
environment.Thus, the dominant mechanisms in oxygen/water vapour
mixtures may dier from thosein steam or steam/inert gas
mixtures.
In the present paper, the oxidation behaviour of the 9%Cr steel
P91 was studied in
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454 3429due
to the presence of water vapour in frequently encountered service
environments.
-
2. Experimental
The composition of the studied ferritic steel P91 is shown in
Table 1. The two batchcompositions presented were used in various
investigations conducted by the presentauthors into the behaviour
of ferritic 9%Cr steels in water vapour containing gases (seee.g.
[24,6]). Batch A in Table 1 was used in the earlier studies,
whereas, for availabilityreasons, batch B was used in the more
recent studies. The two materials only dier inrespect to minor
element concentrations in the steel. This results in slight
variations in abso-lute growth rate in the various environments,
however, no fundamental dierences inrespect to conditions under
which protective or breakaway type oxidation occurred, werefound.
Most of the studies described in the present paper relate to the
newer batch B. Onlyfew results (as will be indicated in the
respective gures) relate to the earlier batch A.
Rectangular specimens, 20 10 2 mm in size were machined from the
prevailingpieces of the steel P91, and ground to a 1200 grit
surface nish. For a number of shortterm experiments, the specimens
were subsequently polished with 1 lm diamond pasteto suppress the
incubation period frequently encountered under the prevailing
conditions,
3430 J. Ehlers et al. / Corrosion Science 48 (2006) 34283454as
will be further explained in the text and the respective gure
captions. Studies onoxidation kinetics up to exposure times of 100
h were carried out at 650 C inN21 vol%O2x vol%H2O (x = 27) gas
mixtures at a ow rate of 0.15 cm/s, using aSETARAM thermobalance.
To obtain more detailed information on the oxidation kinet-ics,
additional isothermal exposures were performed in an N21 vol%
16O22 vol%H218O
gas mixture at a ow rate of 1 cm/s, for oxidation times ranging
from 1 to 30 h. In thesetests an N21%O2-mixture was bubbled through
a glass container containing the H2
18O atcontrolled temperature. It should be mentioned that
actually the water used was not pureH2
18O but contained a 50%H218O-enrichment. After oxidation, these
specimens were ana-
lysed with respect to composition and oxygen isotope
distribution in the scales by MCs+-SIMS [25] using a CAMECA IMS 4F
secondary ion mass spectrometer (SIMS). The depthproles were
quantied following the procedure described elsewhere [25] and
re-calculatedto results which would have been obtained if pure
H2
18O would have been used. The cor-rosion products on all
specimens were additionally characterised by optical
microscopy,
Table 1Composition of the studied steel batches of P91 in
mass%
Element Batch A Batch B
Fe Base BaseC 0.10 0.10Cr 8.1 8.6Mo 0.92 0.93Mn 0.46 0.41Ni 0.33
0.26V 0.18 0.21Al 0.03 n.d.P 0.02 n.d.Si 0.38 0.36S 0.002 n.d.n.d.:
not determined.
-
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454
3431scanning electron microscopy with energy dispersive X-ray
analysis (SEM-EDX) andX-ray diraction (XRD). For metallographic
cross-section preparation, the oxidized speci-mens were Ni-coated
prior to mounting to protect the oxide scales during
grinding/polish-ing and to reveal a clearer contrast between oxide
scale and mounting material. One of theoxidized specimens was
studied by transmission electron microscopy (TEM). The
cross-section of the oxide scale was made by the common sandwich
preparation method andsubsequent ion milling (PIPS). The analysis
was carried out using a Philips CM200-FTEM.
3. Results
Fig. 1 illustrates, on the basis of earlier [24,6] and more
recent studies [26] typicalexamples of the scale formation on the
steel P91 when exposed in N21 vol%O2 andN21 vol%O2 with water
vapour additions in the range 27 vol%. in the temperature
range600700 C for exposure times of approximately 100 h. After
exposure in the dry N21 vol%O2 mixture the surface scale is
extremely thin. Its morphology and compositionare very similar to
those found after air oxidation [4]. The scales formed in wet gas
consistof four regions (Fig. 1b and c). In the outer part, a thin
Fe2O3 and a thicker Fe3O4 layerhad been formed. The inner scale
consists of an Fe3O4 matrix with (Fe,Cr)3O4 stringers.These
(Fe,Cr)3O4 stringers mainly resulted from oxidation of the
chromium-rich carbideprecipitates in the alloy [21]. Consequently,
they show a morphology and distribution sim-ilar to that of the
alloy carbides. Near the oxide/metal interface an internal
oxidation zoneexists which contained Cr-rich phases such as Cr2O3
or Cr-rich (Fe,Cr)3O4, frequently incombination with FeO. The
latter phase was identied by XRD analysis of sequentiallyground
samples, as will be shown later, and the Cr2O3 by Raman
spectroscopy. The scaleformed in the wet gas shows substantial
porosity and locally, a degree of separationbetween inner and outer
layer is seen to have developed.
Fig. 2 shows the eect of water vapour on the isothermal
oxidation kinetics of P91 inN21 vol%O2 with and without additions
of water vapour at 650 C. In the dry gas, thealloy exhibits
extremely low weight changes. It is evident that addition of water
vapourto the test gas results in an enhanced reaction rate. The
latter is preceded by a short, appar-ent incubation period in which
the weight change rate is relatively small. The switch-overto
enhanced (breakaway) oxidation does, especially in case of ground
specimen sur-faces, not start at the same time over the whole
specimen surface. It was found [2,26] thatnodules of rapidly
growing oxides nucleated and spread over the surface after
extendedexposure times, as has frequently been observed during long
term exposures of similarmaterials in wet gases (see e.g. [19]).
The rates measured in wet gas during these relativelyshort term
exposures after the onset of breakaway (Fig. 2) therefore do not
necessarilyreect the exact, quantitative eect of water vapour on
the oxidation rates of the post-breakaway scales, but rather the
extent of nodule formation. The incubation periodcan be
substantially suppressed by diamond polishing the steel to a
mirror-nish priorto exposure [26] so that the change from
protective to non-protective oxidation in wetgases can be studied
already in experiments with exposure times of only a few hours,
asalready mentioned in Section 2.
The TEM cross-section in Fig. 3 shows that the slowly growing
scale which forms dur-ing exposure in the dry gas, consists of two
layers. EDX analyses in combination with
results from glancing angle X-ray diraction strongly indicate
that at the outer side
-
Fig. 1. Typical examples of metallographic cross-sections of P91
when exposed in the temperature range 600700 C in N21 vol%O2 with
and without water vapour additions of 27 vol%H2O for exposure times
ofapproximately 100 h: (a) dry gas, (b) wet gas and (c) schematic
of typical scale formation in wet gas illustratingnomenclature used
in text.
3432 J. Ehlers et al. / Corrosion Science 48 (2006) 34283454
-
0.0
0.5
1.0
1.5
2.0
2.5
3.0
3.5
0 5 10 15 20 25 30
Wei
ght g
ain
[mg .
cm
-2 ]
650 C;N2 - 1%O2 - H2O
dry gas;ground surface
2% H2O;ground surface
4% H2O;ground surface
4% H2O;polished surface
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454
3433Exposure time [h]Fig. 2. Weight changes during isothermal
oxidation of the 9%Cr steel P91 in N21 vol%O2x vol%H2O mixturesat
650 C.Fe2O3 is present, whereas the inner scale consists of Cr-rich
(Fe,Cr,Mn)3O4 spinel. Smallvoids are seen to have formed at the
oxidemetal interface. These could have resulted fromcondensation of
inwardly diusing cation vacancies at the scale/alloy interface. The
alloy
Fig. 3. TEM cross-section of P91 after 5 h oxidation in N21
vol%O2 at 650 C, (a,b) overview picture plus Cr-distribution, (cf)
larger magnication with corresponding element distributions of Cr,
Mn and Fe, showing two-layered structure and re-oxidation of metal
surface in voids.
-
3434 J. Ehlers et al. / Corrosion Science 48 (2006) 34283454side
of these voids formed new, protective, Cr-rich oxide. The element
mapping in Fig. 3bclearly shows Cr depletion in the sub-scale zone
of the steel beneath the protective scale.
Fig. 4 shows tapered metallographic cross-sections of the
surface scales formed in wetgas (N21%O22% water vapour) after
various oxidation times. In the early stages of oxi-dation (1 h)
the scale mainly consists of a relatively thick, outer Fe2O3 layer
and an innerlayer of Cr2O3-stringers embedded in FeO. The inner and
outer layer are almost com-pletely separated by a gap and, unlike
the scale commonly found after long exposure times[16], hardly any
magnetite is present. The outer hematite exhibits a whisker type
mor-phology and is extremely poorly adherent to the substrate.
Extension of the exposure time(2 h) at rst mainly leads to a
thickening of the inner sub-scale, in which hardly any mag-netite
is found. Relative to the overall scale thickness, the gap seems to
be shifted outward.Upon further exposure, the relative amount of
magnetite strongly increases and the gapbecomes gradually lled with
oxide. After 16 h, the overall scale morphology is very
Fig. 3 (continued)
-
similar to that frequently described for specimens after long
term exposures ([16] andFig. 1b).
Fig. 5 shows weight change data for the P91 steel (batch A)
during isothermal oxidationat 650 C in wet (N21 vol%O24 vol%H2O)
and dry gas (N21 vol%O2), with in situ
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454 3435Fig.
4. Metallographic tapered cross-sections of oxide scales on P91
after oxidation in N21 vol%O22 vol%H2O
at 650 C: (a) 1 h, (b) 2 h, (c)7 h and (d) 16 h. Specimens were
mirror-polished prior to oxidation.
-
3436 J. Ehlers et al. / Corrosion Science 48 (2006)
34283454switching from wet to dry gas and vice versa every 24 h,
i.e. without intermediate cooling.The gravimetric data were
presented earlier in Ref. [11]. In the rst stage (wet gas) fast
Fig. 4 (continued)
80 900
1
2
3
4
5
6
7
8
0 10 20 30 40 50 60 70Time [h]
wet wetdry dry
Mas
s cha
nge
[mg.c
m-2 ]
650C;Isothermal
Fig. 5. Weight gain during isothermal oxidation of P91 (batch A
in Table 1) at 650 C whereby the gas was in situchanged from wet
(N21 vol%O24 vol%H2O) to dry gas (N21 vol%O2) and vice versa every
24 h.
-
oxidation kinetics were observed, similar to those shown in Fig.
1. Switching to dry gasafter 24 h almost immediately decreased the
oxidation rate. Switching back to the wetgas after 48 h again led
to an increase of the oxidation rate after a short incubation
period.Similar observations were made by Narita et al. [16] during
in situ gas changes of an FeAlmodel alloy at 800 C. In Fig. 5, the
oxidation rate at the beginning of the second wet
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454 3437Fig.
6. Metallographic cross-sections of oxide scales on P91 (batch A)
after the various oxidation stages including
in situ gas changes between wet and dry gas indicated in Fig. 5:
(a) 24 h, (b) 48 h (c) 72 h and (d) 96 h.
-
3438 J. Ehlers et al. / Corrosion Science 48 (2006)
34283454oxidation stage was lower than that found at the end of the
rst oxidation stage. Switchingback to dry gas after 72 h resulted
in a very low oxidation rate.
Fig. 6 shows cross-sections of the oxide scales formed after the
four oxidation stages ofFig. 5. After the rst wet gas stage, the
oxide scale consists of three layers plus an inneroxidation zone
and is comparable to that shown in Figs. 1b and 4d. The Fe3O4 in
the outerpart of the scale exhibits substantial porosity. After
subsequent oxidation in dry gas, thisporosity nearly completely
vanished (Fig. 6b) and the Fe2O3 layer increased in thickness.
Fig. 6 (continued)
-
0.00
0.01
0.02
0.03
0.04
0 20 40 60 80 100Time [h]
Mas
s cha
nge
[mg/c
m2]
Mas
s cha
nge
[mg/c
m2]
wetdry
0.00
0.01
0.02
0.04
10 20 30 40 50 60Time [h]
IsothermalChange to Wet Gas
Dry Gas at 650 C
Cooling toRoom Temperature
in Wet Gas
Heating to 650 Cin Wet Gas
0.03
0b
a
Fig. 7. Isothermal oxidation of P91 at 650 C whereby, after 24
h, the gas was switched from dry (N21 vol%O2)to wet gas (N21
vol%O24 vol%H2O): (a) without intermediate temperature change, (b)
cooling to roomtemperature during exposure in wet gas and (c,d) SEM
pictures of oxide morphology after oxidation according toconditions
in (b).
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454 3439
-
3440 J. Ehlers et al. / Corrosion Science 48 (2006) 34283454In
the third stage (wet gas), the Fe2O3 was largely transformed into
an Fe3O4 layer, and alarge amount of porosity again appeared (Fig.
6c). Some remnants of Fe2O3 are apparent,both near the outer Fe2O3
layer and near the inner Fe3O4/(Fe,Cr)3O4 interface. After thelast
stage in dry gas, practically the whole outer layer consisted of
hematite and only smallfragments of Fe3O4 are seen in the outer
scale layer (Fig. 6d).
Fig. 7 shows weight change data measured during an in situ gas
change, after the expo-sure was started in dry gas. The data show
that the protective oxide is up to the maximumexposure time not
destroyed if the dry gas is in situ switched to wet gas. However,
if afterthe change to wet gas, an intermediate cooling to room
temperature is introduced, a tran-sition to rapid oxidation occurs
after an apparent, short incubation period. The externalappearance
of the scale grown in this rapid reaction is shown in Fig. 7c and
d.
To gain more insight into the roles of oxygen and water vapour
species in the scalegrowth processes of P91, a number of
experiments were carried out for dierent oxidationtimes in a N2
16O2H218O gas mixture (1 vol% 16O2 and 2 vol%H2
18O). Isotope distribu-tions in the resulting surface scales
were analysed by MCs+-SIMS (see Section 2 for detailson the test
procedure and quantication method).
After a very short oxidation time of 1 h, the scale seems to
consist of two regions(Fig. 8a). Based on the results in Fig. 4a
and XRD data, the inner Cr-containing part con-sists of a Cr-rich
scale. The outer scale consists of pure Fe-oxide in which no clear
changein oxygen/iron ratio is visible as a function of penetration
depth. This can be explained ifone assumes the iron oxide to nearly
exclusively consist of hematite (compare Fig. 4a). The18O/16O-ratio
diers only slightly as a function of penetration depth, although
the ratio isslightly higher in the outer than in the inner part of
the scale. Similar proles were foundafter 2 and 4 h oxidation.
The SIMS depth proles after 7 h and 30 h oxidation in
N216O2H2
18O are shown inFig. 8b and c. It is obvious that three layers
of dierent compositions exist in the oxidescale. Comparison of the
SIMS depth proles with the metallographic cross-sections(Fig. 4)
and XRD data (Fig. 9) reveal that the outer scale consists of Fe2O3
and Fe3O4,the inner of Fe3O4 + (Fe,Cr)3O4 whereas substantial
amounts of FeO were found in thezone near the scale/alloy
interface. The Cr/Fe-ratio in the inner layer, consisting ofFe3O4 +
(Fe,Cr)3O4, equals approximately 1:4.
In all measured SIMS depth proles, the ratio 18O/16O in the
inner part of the scale issmaller than in the outer scale. After 1
h and 7 h oxidation, the region in which the 18Oconcentration
becomes higher than the 16O concentration is located very near the
interfacebetween outer Fe3O4- and inner Fe3O4 + (Fe,Cr)3O4 layer
(compare Figs. 4 and 8). After30 h oxidation, the cross-over point
occurs approximately in the middle of the outerFe3O4-layer.
Comparison of these SIMS-data with the metallographic
cross-sections inFig. 4 strongly indicate, that the area in which
the concentration of 18O becomes higherthan that of 16O, coincides
with the gap in the scale.
Fig. 10 shows the eect of water vapour and oxygen content on the
oxidation behaviourof the P91 steel at 650 C in N2O2H2O gas
atmosphere. It is seen that in a gas with a lowoxygen content, very
small amounts of water vapour are sucient to initiate the
rapidnon-protective oxidation. The growth rate of the
non-protective oxide scale appears tobe practically independent of
the water vapour content. If the oxygen is increased to20 vol%,
non-protective oxidation does not occur even if the concentration
of watervapour is as high as 10 vol%. Only if the water vapour
content is increased to very high
levels, is the protective oxide destroyed. This result explains
why 9%Cr steels can exhibit
-
020
40
60
80
100
0 0.05 0.1 0.15 0.2Depth [m]
Conc
entra
tion
[at.-%
]Ototal
Fe
Cr
18O
16O
Fe2O3 AlloyCr-rich oxide
.1
0
20
40
60
80
100
0 0.2 0.4 0.6 0.8 1 1.2 1.4Depth [m]
Conc
entra
tion
[at.-%
]
OtotalFe
Cr
18O
16O
Fe2O3 Fe3O4 +(Fe, Cr)3O4
AlloyFe3O4)
0
20
40
60
80
100
0 2 4 6 8 10 12 14 16 18 20Depth [m]
Conc
entra
tion
[at.-%
]
OtotalFe
Cr
16O
18O
Fe2O3 Fe3O4 +(Fe, Cr)3O4
AlloyFe3O4 +)
a
b
c
Fig. 8. MCs+-SIMS depth proles of P91 after oxidation in N21
vol%16O22 vol%H2
18O at 650 C: (a) after 1 h,(b) 7 h and (c) 30 h. Specimens were
mirror-polished prior to oxidation.
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454 3441
-
3442 J. Ehlers et al. / Corrosion Science 48 (2006) 34283454very
low oxidation rates, during exposure in laboratory air up to 10,000
h [4,10]. This veryprotective behaviour occurs in spite of the fact
that during such a long term exposure, the
20 25 30 35 40 45 50 55 60 65 702 theta [degrees]
Inte
nsity
[arbi
trary
units
]
Oxide surface
After 3rd grinding step
After 4th grinding step
Fe2O3 FeO
Fe3O4 Alloy
Fig. 9. XRD patterns of the oxide scale formed on P91 during
oxidation for 30 h in N21 vol%O22 vol%H2O at650 C. The specimen was
mirror-polished prior to oxidation. The rst XRD spectrum was taken
of the as-oxidized specimen (in the gure indicated as oxide
surface). Subsequently the specimen was ground in foursteps before
reaching the metal surface. XRD spectra were take after each
grinding step. Presented are the XRDspectra after grinding steps 3
and 4. The phases which could be detected after steps 1 and 2 were
similar to thoseafter step 3.
0
0.5
1
1.5
2
2.5
3
3.5
4
0 10 20 30 40 50 60H2O concentration [%]
Mas
s cha
nge
[mg/c
m2]
20% O2
2% O21% O2
Fig. 10. Eect of O2 and H2O content on the weight gain after 24
h oxidation of P91 at 650 C in N2O2H2O gasmixtures. Specimens were
mirror-polished prior to oxidation.
-
gram as the diusion path shown. This is seen to be consistent
with local equilibrium, and
Scale development during reaction in wet gas is illustrated by
the series of cross-sectionsin Fig. 4 (N 1 vol%O 2 vol%H O) and the
longer term oxidation product in Fig. 2. The
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454 34432 2
2
scales shown in Fig. 4a and b are seen by reference to Fig. 2 to
correspond to pre-break-away reaction, where the rate is only a
little faster than in dry gas. The outer layers of thesescales are
Fe2O3, just as in the dry gas, but now with an irregular,
whisker-like surface.to imply substantial chromium depletion at the
alloy surface.Selective oxidation of chromium must, of course, lead
to its preferential removal from
the alloy. However, the localisation of this eect to the alloy
sub-surface region is a con-sequence of the relatively low alloy
diusion coecient, DA. This is conrmed by a valuefor an average
diusion coecient DA = 10
13 cm2 s1 which can be extrapolated fromhigh temperature data
for FeCr alloys from Ref. [28]. A similar value can be derived
frommore recent data (see Ref. [29] and compilation in Ref. [12]).
It is clear that any disruptionof the scale would expose an alloy
surface with a low chromium concentration [10]. Re-growth of
continuous, chromium-rich spinel would then be dicult, and
iron-rich oxideformation favoured.water vapour content in the gas
can, for certain time periods, be as high as around 2 vol%,a value
which was shown (see e.g. Figs. 1 and 10) to cause non-protective
oxidation in low-oxygen environments after only very short exposure
times.
4. Discussion
4.1. General remarks
With its chromium level of 9%, the P91 steel has marginal
ability to develop a protec-tive, chromium-rich scale at the
intermediate temperature of 650 C. When such a scaledoes develop,
it has been shown [4] to be capable of providing very long term
protection.If, on the other hand, an iron-rich scale forms, it
grows rapidly, consuming the steel. It isobviously of interest to
identify and understand the processes determining which of
theseoutcomes is arrived at.
4.2. Growth rates and morphologies of protective and
non-protective scales
As seen in Fig. 2, reaction of P91 with dry N21 vol%O2 was
extremely slow. However,addition of even small amounts of water
vapour to the gas led to breakaway reaction,i.e. to a rapid
acceleration in rate after a period of slower reaction. As the
value of p(H2O)was increased, breakaway occurred at shorter times.
Higher p(H2O) values also led tosome acceleration of the initial
pre-breakaway rate.
The detrimental eects of water vapour on FeCr alloy oxidation
have long beenknown [79] but not adequately explained. The
protective scale grown in N21 vol%O2is seen in Fig. 3 to be very
thin, and is similar [4] to those produced by oxidation in air.This
scale consists of an outer layer of Fe2O3 and an inner layer of
mixed chromium spinel,(Fe,Mn,Cr)3O4. If the presence of manganese
is ignored, and if phase equilibrium in theFeCrO system can be
approximated by the high temperature ternary isotherm inFig. 11
[27], then the phase assemblage present in the scale can be mapped
onto the dia-Unlike the continuous chromium spinel formed in dry
gas, the inner layers of these scales
-
O 40
90
80
70
Cr O2 3
50
20
30
O (at %)
-Fe + Cr O 2 3
r O23
+S +
CrO
2
23
+S 1
-
Fe+
FeO +
S
1-x
1
-Fe+
Fe O1-x
Fe O2 3Fe O3 4
S = Fe Cr OS = FeCr O
1 1.5 1.5 4
2 2 4
+S +
CrO
2
23
+S 1
Fe O1-x
Fe O + S1-x 1
3444 J. Ehlers et al. / Corrosion Science 48 (2006)
34283454consist of various phases such as FeO, Cr2O3, (Fe,Cr)2O4
and possibly more stable oxidesof alloying elements such as Si
[18].
The multi-phase mixtures e.g. of FeO + Cr2O3 are, according to
the phase diagramthermodynamically unstable both at high
temperature (Fig. 11), and at 650 C, becausethe reaction
FeO(s) +Cr2O3(s)=FeCr2O4(s) 1is thermodynamically favoured. The
formation of the two-phase reaction product reectsthe low value of
DA relative to the scaling rate, enabling the formation of FeO by
limitingthe chromium availability. Its continued metastability
reects the slow rate of reaction (1).
Scales shown in Fig. 4c and d reect various times of reaction
after breakaway. In allcases, large amounts of Fe3O4 are present,
along with chromium-rich spinel. The magne-tite phase was not
present in the protective scale grown in dry gas, or in the scales
grownin wet gas prior to breakaway. Its presence is characteristic
of breakaway reaction, andreects the failure of the slow diusing
alloy to supply chromium to the scalealloy inter-face. This
interface advances rapidly into the alloy, oxidising the prior
microstructue offerrite plus chromium-rich carbides, and
reproducing it as the Fe3O4 plus chromium-richspinel scale layer
[21]. Thus the interface between Fe3O4 and Fe3O4 + (Fe,Cr)3O4 seen
inFigs. 1 and 4 represents the prior alloy surface location after
oxidation. Additional oxideis formed outside this interface as
result of outward iron diusion through Fe3O4 [21].
Fe 10 20 30 40 50 60 70 80 90 Cr Cr (at %)
10+ Cr O2 3+
C
+
Fig. 11. Phase diagram FeCrO at 1200 C [27]; dotted lines
showing diusion path.
-
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454 3445The
remainder of the post-breakaway scale consists of an outer Fe2O3
layer and aninnermost multi-phase sub-layer. The sequence of oxides
from top to bottom of the scaleis consistent with their increasing
thermodynamic stability and the expected decrease inoxygen activity
from the outside to the interior of the scale. Whilst the
requirements oflocal equilibrium between oxides and oxygen activity
are satised in this sense, true localequilibrium is not achieved.
As already discussed phase mixtures such as e.g. FeO +Cr2O3, or
Fe3O4 + (Fe,Cr)3O4 are metastable; at equilibrium the former would
form aspinel and the latter a single phase.
An important additional feature present in breakaway scales is
the extensive porosityevident in Fig. 4. Whilst ne pores are
present in the inner Fe3O4 + (Fe,Cr)3O4 layer[21], large cavities
exist in the Fe3O4 layer, and these form a more or less
continuousgap at intermediate times. This gap is present also in
pre-breakaway scales grown inwet gas (Fig. 4a and b). The existence
of the gap probably explains the formation ofthe unusually thick
Fe2O3 layers developed in these experiments. Because solid-state
dif-fusion in Fe2O3 is much slower than in the lower iron oxides
[8] it usually develops as onlya very thin layer. However, once the
gap develops and separates the outer scale from theunderlying
material, the supply of iron by outward diusion ceases. Inward
oxygen diu-sion then leads to Fe3O4 oxidation, resulting in Fe2O3
layer thickening. Although the gapblocks solid-state diusion, scale
growth nonetheless continues, both above and below thegap, as seen
in Fig. 4. A similar phenomenon is seen on a much smaller scale in
the pro-tective oxide formed in dry gas (Fig. 3) where voids at the
oxidealloy interface form oxideon the metal surface.
It is clear that water vapour can prevent the formation of
protective chromium-richoxide scale layers on the Fe9Cr steel. The
eect increases in severity with increasedp(H2O), and is associated
with the development of voids and/or a gap within the scale,as well
as the appearance of large amounts of Fe3O4. Continued reaction
despite the pres-ence of this gap must be supported by gaseous mass
transfer, which is now considered.
4.3. Gas phase mass transfer within the scale
In the dry gas reaction, the only relevant vapour species within
the oxide is O2(g). If weassume that pO2 at the interface between
hematite and magnetite is set by the equilibrium
2Fe3O4 + 1/2O2 = 3Fe2O3 2and neglect the presence of chromium,
then it is estimated from thermodynamic data [30]that pO2 equals
approximately 10
13 atm at 650 C. At such a low pressure, the rate ofoxygen
vapourisation through dissociation of Fe2O3 can be estimated from
the HertzLangmuir equation as [8]
ki aipi2pmikT 1=23
where pi is the vapour pressure and mi the mass of the
evaporating molecules; ai is termedthe evaporation coecient. When
pi is expressed in atmospheres and the evaporation rateki in g
cm
2 s1, Eq. (3) takes the form
1=2ki 44 3aipiM=T 3a
-
3446 J. Ehlers et al. / Corrosion Science 48 (2006)
34283454where M is the molecular weight of the evaporating
molecules. If ai is set at unity, then aux of 3 1014 mol cm2 s1 is
calculated. Arrival of this ux at the underlying metalsurface can
support oxide re-growth at a rate of 0.01 nm h1. This dissociation
mecha-nism [7,31,32] is thus seen to provide insucient mass
transfer to account for there-grown oxide, observed in Fig. 3 to be
approximately 5 nm thick in a protective scale.It is therefore
concluded that the oxygen activity must have risen to higher values
asthe detached Fe2O3 approached equilibrium with the ambient
atmosphere.
Such an explanation is not available for mass transport across
the gap formed in break-away scales. As seen in Fig. 4, the gap in
these scales is located within the Fe3O4 layer. Thevalue of pO2
within the gap must therefore be below the equilibrium value for
Eq. (2), andthe oxygen ux due to the O2(g) species is much too low
to account for the rapid growth ofpost-breakaway scale beneath the
gap. Furthermore, it obviously provides no mechanismfor iron
transport to support continued growth of oxide outside the gap. It
is thereforeconcluded that additional transport mechanisms must be
facilitated by the presence ofH2O [8].
If H2O molecules can enter the scale, they can provide a means
of oxygen transport [7,9]through the reaction
H2O=H2 +1/2O2 4as illustrated in Fig. 12. If inward H2O
transport is relatively fast, then the partial pressureof H2O in
the cavity will approach that of the external gas, in the present
case approxi-mately 102 atm. As discussed earlier, local
equilibrium between gas and solid oxideappears to be closely
approached. For local oxygen potentials of 10221013 atm,
corre-sponding to the Fe3O4 existence range at 650 C, it is
calculated from the thermodynamicsof Eq. (4) [4,30] that p(H2)
values lie in the range 5 107 to 102 atm. According to Eq.(3),
these hydrogen pressures could support oxygen transfer rates of 5
107 to102 mol cm2 s1. The breakaway oxidation rate in N2- 1 vol%O22
vol%H2O shownin Fig. 1 corresponds to 1 109 mol cm2 s1 of oxygen
atoms. If approximately halfof this uptake occurs below the gap in
the scale (Figs. 2 and 4), then the available gasphase oxygen
transport rate is more than enough to support it.
An alternative possibility in the presence of water vapour is
the formation of volatile(oxy) hydroxides:
FeOH2O FeOHg2 5Fe3O4 3H2O 3FeOHg2 1=2O2 6Fe2O3 2H2O 2FeOHg2
1=2O2 7Cr2O3 2H2O 3=2O2 2CrO2OHg2 8
The formation of FeOHg2 during pure iron oxidation was proposed
by Surman andCastle [33]. Volatile Fe-species were observed by
Viefhaus [34] during in situ AES studieson steam oxidation of 9Cr
steel and by Jaron et al. [35] during high ow experiments withFe in
steam. The eects of steam on formation of volatile
Cr-oxy-hydroxides Eq. (8) arewell documented [3638] but they seem
not to be directly relevant to the transport of ironacross the
scale gap considered here. Astemann et al. [14,15] clearly showed,
that in high-oxygen/high-water vapour mixtures formation of
volatile Cr-species can trigger break-
away type oxidation. However, comparing the dependence of the
Cr-oxy-hydroxide
-
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454
3447partial pressure on pO2 and pH2O(g) implied by Eq. (8) with the
experimental data in
Fig. 10, leads to the conclusion that chromium volatilisation
was not a critical factor inthe overall oxidation process studied
here. As the process under examination involvedthe growth of iron
oxides, this is not surprising.
Assuming that H2O can enter the scale and that local oxide-gas
equilibrium is achieved,it is seen from Eqs. (6) and (7) that the
lower oxygen potentials in the inner part of thescale will produce
higher FeOHg2 partial pressures than in the outer scale regions.The
resulting gradient in pFeOHg2 leads to outward transport of the
hydroxide. Theprocess is shown schematically in Fig. 13. The FeOHg2
vapour species is unstable athigh-oxygen pressures and deposits as
solid oxide, which is believed to be mainlyFe2O3, but Fe3O4
deposition should also be possible in the outer part of the
magnetitelayers.
Thiele et al. [4] have extrapolated thermodynamic data from much
higher temperatures[30] to calculate FeOHg2 partial pressures at
650 C. They found pFeOHg2 to beapproximately 1011 atm at oxygen
potentials in the Fe3O4 stability range. These valuesare too low to
provide signicant mass transfer, but reliable thermodynamic data
arenot available for this low temperature. Given that the oxide
continues to thicken above
Fig. 12. Schematic illustration showing transport of water
vapour molecules through the scale and oxygentransfer across
in-scale void via H2OH2 bridge (based on Ref. [7]).
-
the gap in the scale (Fig. 4), despite the impossibility of
solid-state diusion, it must beconcluded that vapour phase
transport of iron is occurring. The hydroxide species pro-vides the
vehicle for this transport.
Two mechanisms for gas phase transport within the scale have
been identied. Both
Fig. 13. Schematic illustration showing proposed mechanism for
transport of Fe from inner to outer part of thescale via volatile
specie FeOHg2 .
3448 J. Ehlers et al. / Corrosion Science 48 (2006)
34283454involve H2O(g), and can only operate if that species can
enter the scale. Relevant informa-
tion on the interaction between water vapour and the scale is
provided by the cyclic expo-sure experiments.
4.4. Alternating wet and dry oxidation
As seen in Fig. 5, switching from wet to dry gas after 24 h of
breakaway oxidation led toa rapid decrease in scaling rate.
Comparison of scale cross-sections after these two stages(Fig. 6a
and b) reveals that the dry gas caused an increase in the amount of
Fe2O3 at theexpense of Fe3O4, densication of the oxide and
elimination of the gap. It is clear that dur-ing the second stage
of this experiment, oxygen entered the scale interior where it
con-verted Fe3O4 to Fe2O3. The volume expansion accompanying this
transformation,together with perhaps some additional oxide growth
led to elimination of much of the porespace. The weight of oxygen
uptake measured during the second stage was about0.6 mg cm2,
corresponding to an Fe2O3 thickness of about 4 lm, in reasonable
agreementwith the metallographic evidence of Fig. 6.
For this to occur, the scale originally grown in wet gas (Fig.
6a) must have been perme-able to gas species. It is therefore
concluded that the outer Fe2O3 layer, despite its
compactappearance, allowed inward gas species diusion. Since,
nonetheless, a large gradient inoxygen activity was sustained (as
shown by the distribution of oxide phases), this diusionprocess
must have been much slower than gas phase transport. It is
suggested that molec-ular diusion along internal surfaces provided
the transport mechanism.
-
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454
3449Elimination of the pore spaces indicates that H2O(g) was no
longer present within the
scale. Either H2O(g) diused out into the dry atmosphere, or it
was reduced by reaction
with the scale, and H2 then diused out and/or partly into the
metal as shown to occurduring oxidation of Fe in wet environments
at high temperatures [7].
Subsequent re-introduction of wet gas led to some acceleration
in rate (Fig. 5), butmuch less than was observed in the rst stage.
The acceleration seems to be precededby a short incubation period.
During this third reaction stage, the scale (Fig. 6c) re-devel-oped
porosity and a gap in the interior. These changes were accompanied
by reduction toFe3O4 of part of the thick Fe2O3 layer remaining
from stage 2. These eects can be under-stood if H2O
(g) gained access to the scale interior, causing volatilisation
via reactions (6)and (7), and oxygen transfer via the process shown
in Fig. 12. The observation of irregu-larly distributed Fe2O3
remnants in the outer Fe3O4 layer is consistent with molecular
gastransport within the void space, rather than solid-state oxide
ion diusion. Evidently, theoxidation step in stage 2 did not
completely densify the outer Fe2O3, as some gas accesswas still
possible. The nal oxidation (stage 4) in dry gas led again to
conversion ofFe3O4 in the scale interior, densication and
elimination of the gap, by the same processesas occurred in stage
2.
Commencing the experiment in dry instead of wet gas led to very
dierent results, asshown in Fig. 7a. The protective scale grown in
dry gas was up to the test times used,not aected by subsequent
exposure to H2O
(g), and retained its protective character. If,however, the
scale was cooled and reheated, the coecient of thermal expansion
dierencebetween scale and metal led to scale damage and subsequent
rapid reaction in wet gas(Fig. 7b and c). Clearly the Fe2O3 grown
during isothermal exposure in dry gas is not sub-sequently
permeable to H2O
(g), at least up to the maximum exposure times employed.Unlike
the scale grown in wet gas, the oxide grown in dry gas appears to
be fully denseas long as no scale damage, e.g. by thermal cycling,
is introduced. A similar conclusionwas drawn by Schutze et al. [10]
who found breakaway of the initially formed protectivelayer on P91
to occur during prolonged exposure in air with large amounts of
watervapour. From their acoustic emission analyses, the authors
concluded the occurrence ofscale damage to be related to growth
stresses or thermally induced stresses caused by ther-mal
cycling.
A nal illustration of the gas permeability of breakaway scales
is provided by Fig. 14,taken from Ref. [26]. It is known that long
term oxidation in Ar50 vol%H2O leads to athick, porous scale,
similar to that grown in N21 vol%O24 vol%H2O (Fig. 1) except
that,as long as the scale is suciently dense and no barrier layers
(e.g. of Cr- and/or Si-richoxides) are being formed, an outer Fe2O3
layer is not present [21,26]. A two-stage reactioninvolving
exposure rst to Ar50 vol%H2O and subsequently to air (Fig. 14) led
in the sec-ond stage to oxidation of most of the outer Fe3O4 layer
but not to a change of the innerFe3O4 + (Fe,Mn,Cr)3O4 layer. It is
clear that molecular oxygen penetrated the outer layer,but not the
inner layer. This was presumably a result of their dierent
porosities: large andconnected in the outer layer; small and
isolated in the inner layer [21,22], apart from somelarger voids
near the scale/alloy interface. The observation of isolated Fe3O4
islandsremaining in the oxidized outer layer is also consistent
with molecular transport ratherthan solid-state diusion.
Because no H2O(g) was present in the second stage of this
reaction, no mechanism for
volatilisation was available, and the overall scale growth was
low. Because the rate of
scalealloy interface movement was therefore low, diusion in the
alloy had time in which
-
Fig. 14. Metallographic cross-sections of 9%Cr steel showing
oxide scales after oxidation at 650 C: (a) 1000 hoxidation in Ar50
vol%H2O and (b) 250 h oxidation in Ar50 vol%H2O and subsequent
oxidation in air for750 h.
3450 J. Ehlers et al. / Corrosion Science 48 (2006) 34283454to
deliver chromium to this interface, where a chromium-rich oxide
layer formed(Fig. 14b), i.e. a zone with internal oxidation of
chromia was no longer present.
It is clear from these experiments that the gas permeability of
an iron-rich oxide scaledepends on the gas in which it is grown.
When oxygen is the sole oxidant, the scale is denseand virtually
impermeable, and any prior porous oxide becomes densied. When
oxygen isin the presence of water vapour, this desirable result is
not arrived at. Conversely, watervapour, either alone or in the
presence of oxygen produces a gas-permeable scale. It istherefore
concluded that the volatilisation processes possible in the
presence of H2O
(g)
are responsible for creating the scale defects which permit
molecular gas transport. How-ever, this description does not
explain why the oxygen species present in a mixed gas doesnot
penetrate the scale and lead to oxidation and densication of the
scale interior. Thecompetition between oxygen uptake from O2 and
H2O
(g) is discussed later.
4.5. Distribution of oxygen in scales
It has been deduced that during breakaway oxidation, much of the
reaction is due topenetration of the scale by H2O
(g) which facilitates vapourisation processes. On this basis,it
would be expected that oxygen deriving from H2O
(g) would be distributed dierentlyfrom that coming from O2. Fig.
8 shows the distributions of oxygen within the scale, where18O
derives from H2O
(g) and 16O from molecular oxygen. The phases marked on these
pro-les were identied from the total concentrations of Fe, Cr and
O. It is seen that 16O wasalways more abundant than 18O in the
inner part of the scale. In the outer part of the scale,the two
species were present at approximately equal concentrations in a
pre-breakawayscale (Fig. 8a) but 18O was enriched in this region
after breakaway (Fig. 8b and c). This
-
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454 34514.6.
Eect of oxygen partial pressure on water vapour eects
Fig. 10 shows weight uptake after 24 h oxidation at 650 C in
gases with dierent oxy-gen and water vapour partial pressures. It
is seen that breakaway rates are not very sen-sitive to pH2O
(g), but the value of pH2O(g) at which breakaway initiates
increases with
increasing pO2. To a rst approximation, the data in Fig. 10 can
be summarised by thecondition for breakaway oxidation for the
exposure times used in the present study:
pH2OpO2
P 1 9
Hayashi and Narita [17] also proposed that the change from slow
to rapid oxidation ofFe5%Al alloys at 800 C depended on the
H2O(g)/O2-ratio. Explanations for this ndingare sought on the basis
that the condition for breakaway is that H2O enters the scale
andfacilitates gaseous mass transfer.
Formation of volatile FeOHg2 is dependant on both pH2O(g) and
pO2. In the scaleinterior, where reaction (6) is in eect, local
equilibrium leads to
pFeOH2 K1=36 pH2O pO21=6 10
Alternatively, reaction with Fe2O3 through Eq. (7) leads to
pFeOH2 K1=27 pH2O pO21=4 11nding supports the earlier conclusion
that the H2O(g) species participates in breakaway
oxidation.The oxygen distributions developed in breakaway
scaling indicate that the O2 species
penetrates the outer scale region without being completely
consumed, and reacts preferen-tially with the inner layer. They
also indicate that the H2O
(g) species does react in the outerregion, relatively little of
it reaching the inner scale.
A redistribution process could also occur via FeOHg2
volatilisation in the scale inte-rior and re-deposition at higher
pO2 values toward the scale surface. If the initial transientscale
is assumed, for the sake of argument, to consist of 16O and 18O in
a 1:1 ratio, then thevolatilisation reaction (6) with H2
18O(g) produces FeOHg2 and O2 each with an 16O to18O ratio of
2:5 if mixing is random. Re-deposition through reaction (7) with
16O2 pro-duces Fe2O3 with an
16O to 18O ratio of 3:4. Thus enrichment of 18O in the outer
layeris achieved. This isotopic transfer would not continue
indenitely, because the H2O
(g) pro-duced in reaction (7) is enriched in 16O. To the extent
that H2O
(g) is recycled within thescale, rather than being replaced by
H2
18O(g) from the external gas, the isotopic distribu-tion would
approach a steady state.
The oxygen distribution experiments conrm that when water vapour
is present in suf-cient quantity, oxygen is incorporated into the
scale interior, not merely at the scale sur-face, consistent with
the inward diusion of molecular species. They also conrm thatoxygen
in the presence of water vapour does not react with (and thereby
densify) the outerscale. Instead, reaction with H2O
(g) is favoured in this region. The relative contributions
ofreaction with the two oxidant species is now considered
further.Obviously neither of Eqs. (10) and (11) explains the
condition of Eq. (9).
-
3452 J. Ehlers et al. / Corrosion Science 48 (2006)
34283454Considering now the entry of the molecular species into the
scale, we write the surfaceadsorption equilibria
H2Og S H2O=S 12O2g S O2=S 13
where S represents a vacant surface site, and H2O/S, O2/S
represent adsorbed species. Dis-sociative equilibria are ignored in
light of the fact that isotopic mixing between H2
18O(g)
and 16O2 is extremely slow at these low temperatures. Assuming
that at any instant duringscale development, surface sites are
conserved, then
M S O2=S H2O=S 14Here, M is a constant and square brackets
denote area concentrations. Eliminating [S]between Eqs. (12)(14),
one nds
H2O=S MK12pH2O1K12pH2OK13pO2 15
O2=S MK13pO21K12pH2OK13pO2 16
and it follows immediately that
H2O=SO2=S
K12pH2OK13pO2
17
This competitive adsorption process provides an explanation for
the observation thathigher pO2 values require higher pH2O
(g) values to initiate breakaway oxidation and thecondition of
Eq. (9).
When Eq. (9) is satised, it is likely that K12pH2O > K13pO2,
reecting the preferredadsorption of the polar H2O molecule. Then
Eq. (15) can be approximated as
H2O=S MK12pH2O1K12pH2O 18
If, in addition, K12pH2O > 1, then the surface would saturate
with adsorbed water andthe rate of its inward diusion and
participation in internal mass transfer processes wouldbe
independent of pH2O
(g). This would explain the relative insensitivity of breakaway
ratesto pH2O
(g) (Fig. 10).The competitive adsorption process is also
consistent with the isotope distribution
experiments (Fig. 8), which showed that in the breakaway regime,
oxygen from watervapour was the major species incorporated into the
outer scale and molecular oxygenthe major species taken up by the
inner scale. The preferential adsorption of H2O
(g) inthe outer part of the scale largely excludes the O2
species from the surface and therebyreduces its uptake. Only deep
within the scale, beyond the part at which most of theH2O
(g) has been consumed, is O2 an eective reactant. Finally, the
competitive adsorptionprocess explains the ability of scales formed
in breakaway-inducing atmospheres to resistdensication and retain
their gas-permeability. Adsorbed H2O excludes O2 from the inter-nal
surfaces of the outer scale region, whilst itself reacting only
relatively slowly. Onlywhen H2O
(g) is removed from the gas phase, can O2 gain access to these
surfaces. Finally,
the adsorption model is consistent with the nding that dense,
protective scales grown in
-
H2 couple.
FRG, 1999, ISSN 0944-2952.
J. Ehlers et al. / Corrosion Science 48 (2006) 34283454 3453[3]
P.J. Ennis, in: R. Viswanathan, W.T. Bakker, L.D. Parker (Eds.),
Advances in Material Technology forFossil Power Plants, Institute
of Materials, London, Book no. 0770, 2001, pp. 187196, ISBN
1-86125-145-9.
[4] M. Thiele, H. Teichmann, W. Schwarz, W.J. Quadakkers, VGB
Kraftwerkstechnik 77 (1997) 135;The conclusion that H2O(g) entry
into the scale interior occurs is supported by the
ndings that scales grown in wet gas develop and maintain gas
permeability, that this per-meability is not developed in dry gas,
and that the permeability of a wet gas grown (break-away) scale can
be sealed by subsequent reaction in dry oxygen. The conclusion
isconrmed by the nding that oxidation in a mixture of 16O2 and
H2
18O(g) leads to anon-random distribution of isotopes within the
scale, consistent with gas entry.
A competitive adsorption process in which strongly adsorbed
H2O(g) largely excludes
O2 from internal surfaces is shown to account for the
development in wet gas of gas-per-meable scales and its resistance
to densication by reaction with O2. This process alsoaccounts for
the observations that, in the time-frame examined, a critical
condition forbreakaway is pH2O/pO2 > 1 and that breakaway rates
are relatively insensitive to pH2O
(g).
Acknowledgements
The authors are grateful to their colleagues Mr. Olefs, Mr.
Lersch, Mr. Gutzeit and Mr.Wessel for their assistance in carrying
out the oxidation studies, XRD analyses, opticalmetallography and
SEM/EDX. Prof. B. Gleeson is gratefully acknowledged for his
stim-ulating discussions in the interpretation of the experimental
data. They also acknowledgethe European Commission and Siemens
Power Generation who partially funded thisproject.
References
[1] K.H. Mayer, W. Bendick, R.U. Husemann, T. Kern, R.B.
Scarlin, VGB Power Tech 1 (1998) 22.[2] M. Thiele, W.J. Quadakkers,
F. Schubert, H. Nickel, Report Forschungszentrum Julich, Jul-3712,
Julich,associated with the formation of large amounts of porous
Fe3O4, and the developmentof an essentially continuous gap in the
scale. Under there conditions, rapid scale growthcombined with slow
chromium diusion in the alloy led to loss of local
equilibriumbetween the solid phases of the reacting system.
However, gassolid local equilibriumwas probably still closely
approached.
It is concluded that entry of molecular H2O(g) into the scale
interior was the critical pro-
cess leading to breakaway. The H2O can create porosity by
vapourising the FeOHg2 spe-cies and re-depositing in parts of the
scale where higher pOg2 values exist. The H2O
(g)
species also facilitates oxygen transfer within the scale
through operation of the H2O/dry oxygen are not subsequently
permeated by H2O(g). In the absence of internal surfaces,
adsorption and penetration of molecular H2O(g) is clearly
impossible.
5. Summary and conclusions
The presence of water vapour in oxygen-bearing gas mixtures at
650 C has been shownto provide conditions for breakaway oxidation
of P91 steel. Breakaway was found to beM. Thiele, H. Teichmann, W.
Schwarz, W.J. Quadakkers, VGB Kraftwerkstechnik 2/97 (1997)
129.
-
[5] K. Zabelt, B. Melzer, A. Reuter, in: Conference Korrosion in
Kraftwerken, Wurzburg, FRG, 2930September 1999, 11.
VDI-Jahrestagung Schadensanalyse, VDI Verlag, Dusseldorf, 1999, pp.
99111.
[6] W.J. Quadakkers, M. Thiele, P.J. Ennis, H. Teichmann, W.
Schwarz, in: EUROCORR 97, Trondheim,Norway, 2225 September 1997,
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Enhanced oxidation of the 9%Cr steel P91 in water vapour
containing
environmentsIntroductionExperimentalResultsDiscussionGeneral
remarksGrowth rates and morphologies of protective and
non-protective scalesGas phase mass transfer within the
scaleAlternating wet and dry oxidationDistribution of oxygen in
scalesEffect of oxygen partial pressure on water vapour effects
Summary and conclusionsAcknowledgementsReferences