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: j.quadakkers@fz-juelich.de (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, Proceedings, European Federation of Corrosion, vol. II, pp. 3540.

[7] A. Rahmel, J. Tobolski, Corrosion Science 5 (1965) 333.[8] P. Kofstad, High Temperature Corrosion, Elsevier Applied Science, London, 1988.[9] C.T. Fujii, R.A. Meussner, Journal of the Electrochemical Society 110 (12) (1963) 11951204.

3454 J. Ehlers et al. / Corrosion Science 48 (2006) 34283454[10] M. Schutze, D. Renusch, M. Schorr, Corrosion Engineering, Science and Technology 39 (2004) 157166.[11] C. Williams, M. Thiele, W.J. Quadakkers, in: EUROCORR 96, 2426 September 1996, Nice, F.,

Proceedings, vol. III, pp. 10/110/4.[12] R. Peraldi, B.A. Pint, Oxidation of Metals 61 (56) (2004) 463483.[13] K. Nakagawa, Y. Matsunaga, T. Yanagisawa, Materials at High Temperature 18 (1) (2001) 5156.[14] H. Asteman, J.-E. Svensson, L.-G. Johansson, M. Norell, Oxidation of Metals 52 (12) (1999) 95111.[15] H. Asteman, K. Segerdahl, J.-E. Svensson, L.-G. Johansson, Material Science Forum 369372 (2001) 277

286.[16] M. Fukumoto, S. Maeda, S. Hayashi, T. Narita, Oxidation of Metals 55 (56) (2001) 401421.[17] S. Hayashi, T. Narita, Oxidation of Metals 56 (34) (2001) 251270.[18] M. Ueda, Y. Oyama, K. Kawamura, T. Maruyama, Materials at High Temperature 22 (12) (2005) 7985.[19] I.G. Wright, B.A. Pint, in: Proceedings, NACE CORROSION/2002, Denver, CO, USA, April 811, 2002;

paper no. 02377.[20] A. Aguero, R. Muelas, Materials Science Forum 461464 (2004) 957964.[21] R.J. Ehlers, P.J. Ennis, L. Singheiser, W.J. Quadakkers, T. Link, in: M. Schutze, W.J. Quadakkers, J.

Nicholls (Eds.), European Federation of Corrosion Monograph, no. 34, The Institute of Materials, London,2001, pp. 178193, ISSN 1354-5116.

[22] J. Zurek, E. Wessel, L. Niewolak, F. Schmitz, T.K. Kern, L. Singheiser, W.J. Quadakkers, Corrosion Science46 (2004) 23012317.

[23] M. Hansel, W.J. Quadakkers, D.J. Young, Oxidation of Metals 59 (34) (2003) 285301.[24] A. Galerie, Y. Wouters, M. Caillet, Materials Science Forum 369372 (2001) 231238.[25] J.-P. Pfeifer, H. Holzbrecher, W.J. Quadakkers, W. Speier, Fresenius Journal for Analytical Chemistry 346

(1993) 186191.[26] J. Ehlers, W.J. Quadakkers, Report Forschungszentrum Julich, Julich, FRG, Jul-3883, June 2001, ISSN

0944-2952.[27] N. Birks, G.H. Meier, Introduction to High Temperature Oxidation of Metals, Edward Arnold, London,

UK, 1983.[28] W. Bowen, G.M. Leak, Metallurgical Transactions 1 (1970) 2767.[29] C. Piehl, Zs. Tokei, H. Grabke, Materials at High Temperature 17 (2) (2000) 243246.[30] Program Factsage, Data bank SGTE, GTT Technologies, Herzogenrath, FRG.[31] A. Rahmel, W. Schwenk, Korrosion und Korrosionsschutz von Stahlen, Verlag Chemie, Weinheim, New

York, 1977.[32] S. Mrowec, in: Proc. 3rd JIM Int. Sympos. High Temperature Mater. Chem. 99-38, 2000, p. 115.[33] P.L. Surman, J.E. Castle, Corrosion Science 9 (1969) 77.[34] H. Viefhaus, Max Planck Institute, Dusseldorf, FRG, 2001, unpublished results.[35] A. Jaron, Z. Zurek, A. Stawiarski, J. Zurek, M. Homa, Ochrona przed Korozja 11 s/A/, 2003, pp. 97101.[36] D.L. Douglass, P. Kofstad, A. Rahmel, G.C. Wood, Oxidation of Metals 45 (1996) 529.[37] K. Hilpert, D. Das, M. Miller, D.H. Peck, R. Wei, Journal of Electrochemical Society 143 (1996) 3642.[38] B.B. Ebbinghaus, Combustion and Flame 93 (1993) 119.

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