acom
2 - 2006
A corrosion management and applications engineering magazine from Outokumpu Stainless
Oxidation of S35315 in
water vapor containing
atmospheres under
cyclic and isothermal
conditions
page 2
Passive Films on
Stainless Steel – Recent
Nano-Range Research
page 15
Dear Reader
The main article of this Acom covers a subject that could be very dear
to a Finn or a Swede since hot and humid conditions are by natural
reasons synonymous with a sauna to these people. However, the paper
is on more hostile conditions than any sauna lover could stand, steam
at up to 1000°C.
The article describes how heat resistant stainless steels tolerate
different moisture contents, which is depending on the type of fuel.
It is easy to realize that moisture contents differ between wet
household waste, coal, oil and gas. If the moisture content is too
high volatile chromium species evaporate from the steel surface and
increase the mass loss and risk of failure.
In the second paper you can learn more about the by far most
important part of a stainless steel, the passive film.
Enjoy the reading!
Yours sincerely
Jan Olsson
Technical editor of Acom
www.outokumpu.com
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2
Oxidation of S35315 in
water vapor containing
atmospheres under cyclic
and isothermal conditions
Dr. Pascale Vangeli
Outokumpu Stainless AB
Avesta Research Center, High Temperature Steels
SE-77422 Avesta, Sweden
Abstract
The effects of air in combination with up to 40% of water and temperatures up to 1000°C
on S35315 are reported. The exposure time varied from 20 h to 1000 h in a tube furnace,
and up to 200 h in a thermobalance, under cyclic and isothermal conditions. The
results show that the oxidation behavior of S35315 in water vapor is characterized by a
protective chromium rich oxide scale and no breakaway oxidation was observed for this
steel. The mechanism of oxidation in an atmosphere containing water vapor is discussed,
linked to the ability for chromium to diffuse to the surface and some assumptions are
proposed regarding the effects of different alloying elements such as Ce, Si, Ni and Cr.
Introduction
Keywords: austenitic stainless steels,
high temperature corrosion, evaporation,
breakaway oxidation, oxidation, water vapor,
isothermal oxidation, cyclic oxidation.
1
153 MATM is a trademark of Outokumpu
2
253 MA® and 353 MA® are registered
trademarks of Outokumpu
Outokumpu 153 MATM1 (S30415), 253 MA®2 (S30815) and 353 MA®2 (S35315) is a
family of heat resistant austenitic stainless steels with increased contents of silicon and
nitrogen, and micro alloyed (MA) with rare earth metals (REM, reactive elements, RE).
S35315 has a significantly higher nickel content than the other steels. The alloying
concept has resulted in characteristic family features:
• High mechanical strength at elevated temperatures, i.e. creep strength.
• Excellent isothermal oxidation resistance and, above all, excellent cyclic oxidation
resistance and oxidation resistance under erosive/abrasive conditions.
Furthermore, MA steels are optimized to complement each other:
• S30415 for medium-high to high temperatures and moderately aggressive, mainly
oxidizing atmospheres. This grade is extra resistant against embrittlement after service
at medium-high temperatures.
• S30815 for high to very high temperatures and/or rather aggressive, mainly oxidizing
atmospheres. This grade is the “work-horse” of the family.
• S35315 for the highest temperatures and/or toughest conditions, which often means
environments that are strongly carburizing, strongly nitriding, or contain some
halogens/halides. The alloy is designed primarily for service above 1000°C (1830°F),
although it has been used at temperatures as low as 600°C (1110°F) in more aggressive
environments.
From the MA family, only S35315 is considered in this study and compared to standard
heat resistant grades, such as 310S, 309S and 800H.
It is well known that the presence of water vapor changes the oxidation behavior of
metals and alloys [1]. Chromia-forming alloys are affected by water vapor in the sense
that the critical amount of chromium to form protective Cr2O3 or (Cr,Fe)3O4 scales
increases [2]. The concept that volatilization of chromium species was responsible for this
effect was first proposed by Ebbinghaus, who reviewed the thermodynamic properties
of volatile chromium-containing species and calculated that CrO3 was the dominant
evaporating species in dry oxygen and CrO2(OH)2 in moist atmospheres [3].
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acom 2 - 2006
3
The influence of alloying elements is observed and discussed, for example, silicon is
assumed to enhance oxidation resistance in water vapor. Moreover, cyclic conditions,
often reflecting service conditions, may influence the behavior of the heat resistant steels
in moist atmospheres. The ability for the steel to maintain its protective oxide scale is
essential in environments where an oxide scale is exposed to high stress, such as erosive/
abrasive conditions and/or when large temperature variations frequently occur.
The present study shows how S35315, compared to other standard heat resistant
steels, is affected by different testing conditions, such as: temperature, water vapor
content in the atmosphere and testing time under isothermal and cyclic conditions.
Effects of the water vapor, alloying elements, and cyclic conditions are discussed.
Experimental
The study is carried out on S35315; some comparison testing is done on 309S, 800H
and 310S. The actual chemical compositions are given in Table 1. All the grades are
commercial and fully austenitic. The rectangular specimens are 2 to 5 mm thick and
have a total surface area ranging from 3 to 6 cm2. The specimens, polished with emery
paper (grade #80 and #180 SiC papers), are washed with de-ionized water in
an ultrasonic bath, then with ethanol and finally dried before testing.
Table 1
Chemical composition of the grades in wt%
C
Si
Mn
Cr
Ni
N
other
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Oxidation tests are carried out in synthetic air with 0 and 40% of water in a tube furnace for
up to 1000 h and in two commercial thermobalances for up to 200 h in the temperature
range 600 – 1200°C. All the tests have a gas flow of 100 ml·min-1 (average net velocities
of 0.25 cm·s-1).
The short time tests are conducted in moist atmosphere
in a thermobalance SETARAM Setsys12, and in dry atmosphere in a thermobalance
SETARAM TG96. The weight change is continuously monitored. The furnace is heated
to the desired temperature, while the specimen is kept at room temperature. The gas is
introduced in the thermobalance. A typical test starts when the specimen is lowered into
the reaction chamber. A cyclic test has a cycle consisting of 2 hours in the hot furnace
and 10 minutes in the cold zone (room temperature in the laboratory air).
Tests in thermobalance:
The long time tests, up to 1000 h, are carried out in a tube
furnace. A quartz twin walled reactor tube has been made for this study. The air + H2O
gas stream is first heated to the thermodynamic equilibrium before it flows through the
specimens. This set-up enables the samples to be exposed to the reaction gas at
equilibrium with minimum contact with the sample holder and with an optimal flow
pattern. Details of the test facility has been described elsewhere [4]. All the grades are
tested at the same time. New specimens are used for each time (20, 168, 336, 672 and
1000 h), which means that the samples do not suffer thermal shocks before the end of
the test. The samples are air-cooled after testing.
Tests in the tube furnace:
Morphology: After visual observation, the specimens are mounted in resin and polished, and
the cross sections are examined and analyzed using light optical microscopy, electron probe
and scanning electron microscopy. Identification of the oxides is performed using SEM/EDX.
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4
Results and discussion
Results at 600°C in air with 10% of water, isothermal oxidation
Results from the isothermal oxidation of S35315 in air + 10% of H2O at 600°C for up
to 668 h is presented in Figure 1.
Fig. 1 Mass change of S35315 in isothermal exposures in air + 10% of H2O at 600°C
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The mass changes are very small for the entire test period, as this grade is high alloyed
and the temperature is low. First, a weight gain is observed, up to about 0.02 mg·cm-2.
Then a weight loss occurs, reaching -0.048 mg·cm-2 at 668 h. As no spallation has occurred
after cooling, this weight loss is due to evaporation. This is supported by a brownish layer
present on the SiO2 furnace tube wall in the cold zone, downstream from the samples.
It is well known that chromia forming steels suffer chromium vaporization in the form of
CrO3(g) above 1100°C [5]. At 600°C, its partial pressure is too low to explain the mass
loss observed. Asteman et al. has studied the oxidation of 304 in O2/H2O atmospheres at
600 – 900°C and showed that a significant amount of chromium evaporates from the
chromium-rich oxide scale, even at 600°C [6]. A brown deposit was observed at the exit
of the furnace tube wall and identified as chromium oxide, as observed in this study.
The predominant chromium-containing vapor species in environments containing
oxygen and water vapor at temperatures below 1100°C has been predicted
thermodynamically to be chromic acid CrO2(OH)2 by Ebbinghaus [3]. More recently
the structure and thermodynamic stability of mixed oxyhydroxides of chromium (g)
were studied in more details [7–8]. Panas investigated the mechanism of evaporation
of chromic acid from Cr2O3(s) and gave the reaction energetics for the formation of
CrO2(OH)2 [9–10]. Their calculated energetic values were in good agreement with their
experiments.
The mass loss can then be attributed to the following reaction (1):
1/2Cr2O3 (s) + 3/4O2(g) + H2O ßà CrO2(OH)2(g)
(1)
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5
Results at 900 ºC in air with 10% of water, early stage of isothermal oxidation
Comparative gravimetric curves in dry air and in air + 10% of H2O continuously recorded
during isothermal oxidation of S35315 exhibit a higher mass gain for the oxidation in
dry air than in air + 10% of H2O, see Figure 2.
Fig. 2
Mass change of S35315 in isothermal exposures
in air with 0 and 10% of H2O at 900°C
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The oxidation kinetic in dry air is slow and fully parabolic to the end of the experiment,
indicative of diffusion-limited oxide growth, while in air + 10% of H2O, the curve shows
a non-parabolic behavior. As it has been previously seen at 600°C with evaporation
of chromium species, this non-parabolic behavior results from the combination of a
parabolic oxidation and linear evaporation.
Results at 900°C in air with 10 and 40% of water, long-term exposure under
isothermal conditions
The long-term behavior of S35315 in air + 10% of H2O and air + 40% of H2O at 900°C
for up to 1000 h in isothermal conditions is illustrated in Figure 3a. Each point
represents the mass change of one single sample. Spallation does not occur during
cooling for any of the samples.
Fig. 3a Mass change of S35315 in long-term isothermal exposures
in air with 10 and 40% of water at 900°C
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6
Fig. 3b Mass change of 309S in isothermal exposures
in air with 0 and 10% of water at 900°C
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Obviously, the simultaneous events of oxidation and evaporation are also observed
here. With 10% of water, the oxidation rate is slow and no global mass loss is observed.
In air with 40% of water, a mass gain occurs first, followed by a slight mass loss, showing
more evaporation than oxidation. Obviously, the more water vapor in the atmosphere,
the more evaporation of chromic acid occurs.
According to Asteman [11], the evaporation rate depends on water vapor
concentration in the atmosphere, gas velocity and temperature. Even with up to 40% of
water in the atmosphere and test duration up to 1000 h, no breakaway oxidation occurs,
as for other heat resistant grade like 309S (Figure 3b).
The parabolic oxidation of 309S in dry air becomes non-parabolic in air with 10%
of water, as for S35315, showing evaporation of chromium species. However, after only
100 h of testing, a catastrophic breakaway oxidation occurs, leading to a very rapid
oxidation rate. An extensive spallation occurred when cooling the sample down to room
temperature at the end of the experiment.
The corresponding micrographs for S35315 tested for 1000 h in air with 40% of water
are shown in Figure 4a with element distribution maps.
Fig. 4a SEM cross-section micrograph of S35315 oxidized in air + 40% H2O at 900 ºC
for 1000 h (a) with corresponding EDS maps for Cr (b), Fe (c), Mn (d) and Si (e)
(a)
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(b) Cr
(c) Fe
(d) Mn
(e) Si
7
Figure 4a shows that the oxide scale consists of a three-layer structure. At the metal/
oxide interface, a thin silicon-rich oxide layer, compact and continuous penetrates
several micrometers down into the alloy along the steel grain boundaries. This is a
typical morphology for this grade. These results are in accordance with Jonsson showing
a continuous silica layer at 900°C on S35315 with and without water vapor in the
atmosphere [12]. Next, the main layer consists of a compact and continuous chromia
layer, followed by an outer, thin manganese-rich oxide layer, covering the entire surface
of the sample. No iron-rich oxide is observed, as breakaway oxidation does not occur on
S35315 at 900°C for up to 1000 h in air with 40% of water.
Fig. 4b SEM cross-section micrograph of 309S oxidized in air + 10% H2O at 900 ºC
for 1000 h (a) with corresponding EDS maps for O (b), Cr (c), Fe (d) and Ni (e)
(a)
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(b) O
(c) Cr
(d) Fe
(e) Ni
8
Figure 4b shows a cross-section of 309S oxidized in air with 10% of water at 900°C for 1000 h.
The thick and porous oxide scale consists of a chromium-nickel oxide inner layer and
a iron-rich oxide layer at the outer part of the scale, presenting many cracks and pores.
The mechanism for breakaway oxidation in air/H2O has already been discussed in
previous papers [13].
Hultquist et al. have suggested that hydrogen, originating from water vapor, is
incorporated in the oxide and thereby a change of its defect-dependent properties,
including metal and oxygen diffusion [14–16]. Khanna and Kofstad proposed that
protons (H+) from water vapor affect the diffusion properties of the grain boundaries
[17]. However, definite experimental evidence is needed to show that the effect of water
vapor on breakaway oxidation is related to hydrogen defects.
Jinian proposed another mechanism for the breakaway oxidation: pores, microchannels
or microcracks formed in the oxide scale during the initial oxidation stage allow ambient
gas molecules to penetrate to the metal/scale interface and react with the chromiumdepleted alloy [18]. Since H2O molecules can react with the chromium-depleted alloy
surface, non-protective iron oxide forms and releases H2. Then, this H2 reduces chromia
and generate more H2O. Therefore, as soon as the microcracks in the Cr2O3 scale
appears, the oxidation is accelerated.
Depletion of chromium has been described by Evans as Intrinsic Chemical Failure
(InCF) when the chromium concentration decreases below the thermodynamic value
for Cr/Cr2O3 equilibrium, and Mechanical Induced Chemical Failure (MICF) when the
depleted surface is exposed to the oxidizing atmosphere [19].
According to Henry et al., the catastrophic oxidation due to water vapor is the result
of an intrinsic effect on the scale growth mechanism [20]. This breakaway oxidation is
suggested to be initiated by the arrival of oxygen- and hydrogen- containing species in
the chromium depleted metal/oxide interface to form non-protective iron oxides.
Asteman et al. showed that in water vapor containing atmospheres the formation of
volatile species, such as CrO2(OH) or CrO2(OH)2, are responsible for the breakdown
of the chromia scale [6, 11]. The evaporation of chromium species results in a depletion
of chromium contained in the oxide scale. For Asteman, this leads to poorer protective
properties of the oxide scale. The resistance of chromia-forming steels is enhanced by the
following substrate factors: high chromium concentration, fast diffusion in the bulk, and
a high density of steel grain boundaries.
9
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Asteman’s work was in accordance with a previous study, additionally suggesting that the
transition protective/non-protective oxide is due to a combination of different events [13].
Evaporation of chromium species leads to an increased chromium diffusion and to
chromium depletion in the zone below the metal/oxide interface. A critical chromium
concentration, below which the chromia does not re-form, leads to the formation of a
non-protective iron-rich oxide scale and the catastrophic breakaway oxidation.
This is also in agreement with what is observed here. S35315 contains 25% of Cr and
does not show any breakaway and iron-rich oxide formation, whereas 309S with 22.5%
of Cr does.
Besides the chromium effect, Pettersson showed recently that nickel alloying plays a
role in the oxidation of chromia forming alloys [21]. Her calculations of predicted phase
equilibria for ternary Fe-20Cr-xNi alloys demonstrated two effects of nickel alloying.
The first is a slight increase in chromium activity, promoting chromia formation and the
second is a doubling of the chromium diffusion coefficient, enhancing the ability of the
steel to repair its protective chromia scale. This was supported by data from Jönsson and
Åkermark [22, 23].
Silicon is known to have a beneficial effect on isothermal oxidation kinetics in both
dry and moist atmospheres [24]. On experimental alloys Fe-20Cr-35Ni + either 0.5% or
2% of Si, Pettersson has shown that the formation of a silicon oxide sublayer suppresses
breakaway oxidation at 1000°C in air with 10% of water [25].
The combination of high Cr, Ni and Si contents plays a preponderant role in the
oxidation behavior of S35315, making the alloy very resistant to moist atmospheres
and thus to more aggressive atmospheres, such as carburizing or nitriding environments.
S35315 has been previously studied, subjected to carburization and nitridation at
temperatures ranging from 850 –1200°C. The results showed that S35315 was the most
corrosion resistant grade in these environments compared to N06601, 309S and 310S [26].
Results at 900°C under cyclic conditions in air with 0 , 10 and 35% of water
The study carried out under cyclic conditions gives us information on the adhesion of
the oxide scales. The curves are recorded for the entire cycle (time of exposure and time
of cooling to room temperature). The weight losses, due to spallation during cooling, are
represented in the curves by a vertical step downward.
For S35315, in cyclic oxidation testing with a 2-hour hot dwell time at 900°C,
Figure 5, the same trend as in the isothermal exposures is observed.
Fig. 5
Mass change of S35315 in cyclic exposures
in air with 0, 10 and 35% of water at 900 ºC
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The addition of 10% and 35% respectively, water vapor in the atmosphere enhanced the
evaporation of chromium species, but did not lead to breakaway oxidation.
10
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Results at 1000°C under cyclic conditions in air with 0 and 20% of water
The oxidation behavior of S35315 at 1000°C in air with 20% of water, Figure 6a, is
characterized by the same trend as for the other temperatures, with mass loss after 120 h due
to the simultaneous events of oxidation and evaporation; evidence of more evaporation
than oxidation. Spallation does not occur during cooling for the entire test period.
Fig. 6a Mass change of S35315 in cyclic exposures
in air with 0 and 20% of water at 1000°C
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Fig. 6b Mass change of 310S in cyclic exposures
in air with 0 and 20% of water at 1000°C
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The presence of water vapor in the atmosphere during cyclic oxidation at 1000°C has
two effects on 310S, Figure 6b. The first is evaporation as discussed previously. The
second is that spallation occurs when cooling down after only 130 h of testing while this
does not happen in dry atmospheres at 1000°C for 150 h. The water vapor decreases the
ability of 310S to maintain its protective oxide scale under cyclic conditions. Its presence
decreases the Critical Mass Gain value (CMG, cf. below) of 310S to about 0.8 mg·cm-2.
A method to describe the adhesion of the oxide growing on different heat resistant
grades has been previously proposed [27]. Cyclic oxidation kinetic curves at different
temperatures show a break point (more mass loss during cooling than mass gain during
oxidation). When the mass gain reaches a critical level (break point in the curve), the
sample loses more weight during cooling (spallation) than it gains during oxidizing.
This Critical Mass Gain (CMG) is a characteristic of the alloy and does not depend on
the temperature. The higher value the better.
11
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Results at temperatures above 1000°C and up to 1200°C
under cyclic conditions in dry air
Previous studies on the MA family have shown that S35315, even at 1100°C for up to
200 h, suffers no spallation during cooling. S35315 develops an oxide scale that is very
resistant to thermal shocks, Figure 7a [28].
Fig. 7a Mass change of S35315 in cyclic exposures in dry air
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Fig. 7b Mass change of 310S in cyclic exposures in dry air [27]
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12
Fig. 7c Mass change of 800H in cyclic exposures in dry air
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CMG (Critical Mass Gain) for
310S, S35315 and 800H in dry
atmospheres [27, 28]
Table 2
The CMG for S35315 was found to be
3.8 mg·cm-2, Table 2 [28].
310S has a low oxidation resistance
-2
under cyclic conditions, Figure 7b. At
CMG (mg·cm )
1100°C, the break point (more mass loss
310S
1.3
during cooling than mass gain during oxidS35315
3.8
ation) is reached very quickly. The CMG in
dry atmosphere of 310S has been found to
800H
1.7
be 1.3 mg·cm-2 [27]. This study shows that
even a high alloyed grade as 800H has a
low CMG value (1.7 mg·cm-2), Figure 7c.
As previously reported in this paper, the oxide scale developed on S35315 is very resistant
to severe thermal shocks. This is to be related to the presence of alloying elements
such as cerium. It is known that the most important effect of reactive elements (RE) is
improving scale adhesion [29]. RE promotes the formation of a stronger metal/oxide
interface [30]. S35315 contains 0.04%Ce added as misch-metal, thereby enhancing its
ability to retain its oxide scale. On the other hand, S35315 contains a non-negligible
amount of silicon. A negative aspect of silicon has been reported as increasing the oxide
scale spallation under thermal cyclic conditions [25, 31–32]. This aspect is not observed
in our study.
As spallation did not occur on S35315 during cyclic oxidation in moist atmosphere,
Si compensate
Cr each other
Ni – siliconNcausing other
the two alloying elementCeffects obviously
spallation
effect is 12.26
stronger. 0.070
309S and cerium promoting
0.056 adhesion.
0.36 Evidently,
1.49 the cerium
22.62
-
310S
0.040
0.45
1.52
25.10
19.10
0.044
-
S35315
0.042
1.29
1.35
24.98
35.03
0.144
Ce
800H
0.016
0.43
0.69
20.35
32.70
0.009
Al
Ti
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13
Conclusions
The present study shows that S35315, as other standard heat resistant steels, is affected
by the water vapor content in the atmosphere. The non-parabolic kinetic from
isothermal oxidation is the result of the combination of a parabolic oxidation with
evaporation of chromic acid CrO2(OH)2.
From the isothermal results, S35315 does neither show any breakaway oxidation
in moist atmosphere up to 40% of water nor iron-rich oxide formation at any test
temperatures, while 309S does so, already after 100 h exposure at 900°C.
For S35315, in cyclic oxidation testing with a 2-hour hot dwell time at 900°C,
the addition of up to 35% water vapor in the atmosphere enhanced the evaporation
of chromium species compared to dry air, but did not lead to breakaway oxidation.
Even at temperatures as high as 1000°C, the oxide scale remains protective while on
310S; spallation during cooling occurs after only 130 h of exposure in air with 20% of water.
The presence of water vapor decreases the Critical Mass Gain value (CMG) of 310S to
about 0.8 mg·cm-2.
The CMG in dry air, characteristic of the alloy, was found to be as high as 3.8 mg·cm-2
for S35315, 1.3 mg·cm-2 for 310S and 1.7 mg·cm-2 for 800H.
It has been shown that the composition combination, high Cr, Ni and Si contents
and presence of cerium, plays a dominating role in the oxidation behavior of S35315.
Adding Cr, Ni and Si makes the alloy very resistant to moist atmospheres and thus to
more aggressive atmospheres. The oxide scale developed on S35315 is very resistant to
severe thermal shocks, thanks to the presence of cerium in the steel.
Acknowledgements
The author wishes to acknowledge B. Ivarsson and C. Lille at Avesta Research Centre of
Outokumpu for their contribution and helpful discussions. Thanks are also expressed
to S. Amy, Faurecia, for his help with the testing. Special thanks are addressed to R.
Lindström for performing the sample preparations and thermogravimetric testing.
Furthermore, the author wants to thank E. Torsner, Outokumpu, for her contribution to
this paper.
References
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2. D.L. Douglass, P. Kofstad, A. Rahmel, G. C. Wood, Oxid. Met. 45 (5/6), pp.
529 – 620, 1996
3. B. B. Ebbinghaus. Thermodynamics of gas phase chromium species: The chromium
oxides, the chromium oxy-hydroxides and volatility calculations in waste incineration
processes. Combustion and flame 93, pp.119–137, 1993
4. S. Amy. Water Vapour Effect on the Heat Resistant Grades. Diploma work,
Outokumpu, 2000
5. C.S.JR Tedmon, Electrochem. Soc. J. 113 (8), pp. 766 –768, Aug. 1966
6. H. Asteman, J-E. Svensson L-G. Johansson and M. Norell, Oxid. Met. 52 (1/2),
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Reproduced with permission from NACE International, Houston, TX.
All rights reserved. Paper No 04678 presented at CORROSION/2006,
Houston, TX. © NACE International 2006.
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15
Passive Films on Stainless
Steel – Recent Nano-Range
Research
Claes Olsson, Outokumpu, Avesta Research Center
Gunilla Herting and Inger Odnevall Wallinder
The Royal Institute of Technology, Department of Corrosion Science
Drottning Kristinas väg 48; SE 100 44 Stockholm, Sweden
Introduction
Stainless steels owe their corrosion resistance to an oxide film, about 10 atom layers
thin. The film forms spontaneously in most environments and adopts its thickness and
composition to the surroundings. By changing the composition of the steel, it is possible
to influence the kinetics of film formation. Using the different alloy elements in a
stainless steel to form protective films stable in aggressive environments is one example
of nano scale engineering that has been around for almost a century.
Passive films on stainless steel are still subject to a lively interest within the research
world. The scientists are currently working on problems such as: How fast are
composition and thickness responding to change in environment? Which mechanisms
are controlling film dissolution and growth? To what extent are the different alloying
elements leaching out into the environment? In recent years, the environmental aspects
have gained in importance, during manufacturing as well as end-use. For many years,
the predominant raw material for stainless steel production has been scrap, there is thus
a long tradition as a fully recyclable material.
A more detailed picture of the chemistry and composition of the passive film can
be found in a recent review by Olsson and Landolt [1]. The experiments on metal
dissolution presented can be found in a paper by Herting et al. [2].
When is a passive film formed ?
At room temperature, pure iron has a cubic face centered structure. This means that
every iron atom has eight nearest neighbors. A simple form of a stainless alloy is formed
by replacing part of the iron atoms by chromium. For a random mixture, the likelihood
that at least 50% of the chromium atoms gets at least one other chromium as nearest
neighbor is 1/8. This corresponds to the empirical limit of 11– 12% for the chromium
concentration where the steel becomes stainless. At this concentration, it is possible to
form a continuous network of chromium that can keep the metal matrix together during
selective dissolution of iron. The corresponding threshold for forming a chromium
network in the oxide is about 18%. The mathematical background to this is known as
percolation theory and treats the construction of networks in a wide sense. It is used
also in other fields, for example for the construction of efficient telecommunications
networks.
The protective oxide film is formed through selective oxidation of the different
alloying elements. Iron and chromium are preferentially oxidized at the metal/oxide
interface. Nickel is remaining enriched in its metallic state under the film. Iron is
migrating almost ten times faster than chromium through the oxide. Thus, the oxide film
will be enriched in chromium and most of the iron will be dissolved. The cation fraction
of chromium in the film can amount to 80% under favorable conditions. When the steel
is passive, the oxide film contains a cation fraction of chromium higher than 50%.
By a proper choice of alloying elements, for example by adding molybdenum, nitrogen
or tungsten, it is possible to direct the passive film towards stability under various
conditions.
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How fast is the passivation?
The protective oxide film grows and stabilizes very fast. After an external perturbation,
the film will respond with a thickness and composition change within seconds. After the
initial fast adjustment, the film continues to develop towards an enhanced crystalline
order for several hours or even days.
The electrochemical potential is an external parameter that can be easily controlled in the
laboratory. Figure 1 shows how a passive film grows during and after such a potential
change. The thickness change curve was recorded in the electrolyte by weighing the
passive film during the experiment with a quartz crystal nanobalance. This equipment
makes it possible to record mass changes in real time with a weight resolution
corresponding to 4% of a metallic monolayer. The thickness change abates already
a couple of seconds after stopping the potential sweep. This type of experiments are
performed to find out which processes limit the growth of the oxide film, or to study
effects of chloride adsorption.
Fig. 1
Passive film growth on a stainless steel (continuous line) during and after
a change in electrochemical potential (solid). The experimental uncertainty
is given by the dashed lines. The majority of film growth, about 6 Ångström,
occurs during the sweep and continues only for a few seconds after the
sweep has stopped [3].
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How is the oxide film growing?
The driving force for film growth is the difference in “generalized potential” between
the oxidizing medium, frequently a water-based solution or the atmosphere, and the
underlying metal. This potential drives ion migration through the film. The passive layer
is normally so thin that the growth is limited by a reaction at either of the film interfaces;
one such reaction is the formation of cation vacancies in the oxide film at the metal/
oxide interface. The alternative is rate control by dissolution of cations or incorporation
of oxygen at the outer film interface. The thickness equilibrium is so sharp that a 10%
decrease in film thickness would lead to a tenfold increase in migration rates. Thus, it is
not necessary to introduce cracks or other defects all the way down to the metal to create
a local instability – it is sufficient to replace a couple of oxygen atoms with chloride to
locally increase the ion migration velocities through the film.
What type and how much material will dissolve?
If one studies a common steel grade: 18Cr-9Ni (EN 1.4301, AISI 304), one finds that
the predominating specie dissolving from the sample is iron. On a freshly ground surface,
the passive film has not yet stabilized. While studying such a surface during its initial
passivation process, one will initially see a somewhat higher rate of dissolution that abates
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17
with time. During this exposure time, the passive film is built up with water and oxygen
and reaches a maximum thickness. This film stabilization leads to lower dissolution.
Figure 2a shows a clear enrichment of chromium in the steel surface during exposure
to rain water and a simultaneous reduction of the amount of dissolved metal. During
continued exposure, the dissolution rates will decrease until they reach very low levels
at steady state conditions, cf. Fig. 2b. A passive film weighs about 520 ng·nm-1·cm-2.
The dissolution rates indicated in Fig 2b. are thus less than 10 % of an oxide monolayer.
Fig. 2a Dissolution of chromium during an initial passivation process on a ground
plate used for kitchen sinks.
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Fig. 2b Dissolution rates of chromium from a ground stainless steel. The dissolution
decreases with time and reaches a very low rate after a couple of hours
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Practical consequences
The dissolution rates of different cations are – even for the worst case – at least one order
of magnitude below today’s hygienic limits. Previously, dissolution rates from stainless
steels have been estimated by using values from the pure constituting alloy elements.
This gives a severe overestimate of the amount of iron and nickel dissolved.
The most common delivery surface is cold rolled and pickled. For this final treatment,
the material has already obtained a passivated surface during the final pickling stage.
Evidently, this results in an increased stability of the material even for the first contact
with a chemical product. For a bright annealed surface, the response is slightly different.
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18
For this case, the passive film will undergo a compositional change, whereas the
thickness is less affected. The final oxide on a bright annealed material is formed during
the cooling phase after the furnace and will thus have a higher iron fraction than the
standard pickled surface. Another type of delivery surface is a ground or brushed finish.
For this case, the oxide film will have a composition that closely corresponds to the
bulk composition of the alloy. This surface remains until the first contact with a liquid,
at which point the film will be reshuffled with dissolution rates corresponding to the
reactivity of the liquid. The result will be the same for all the above surface finishes: a
protective film enriched in chromium with a thickness corresponding to the surrounding
medium. This adjustment process is so fast that highly sensitive measurement equipment
is necessary to unveil it. The amount of dissolved metal is so low that it is only recently
that the analytical methods, e.g. ICP-AES/MS, have reached detection limits where the
presence of the different alloying elements can be quantified with the necessary accuracy.
The dissolved quantities that have been indicated are all at least one order of
magnitude lower than the hygienic limits valid today. This is true for the initially
very short formation phase of the passive film. The dissolution rates will decrease
approximately logarithmically with time. After the short intial passivation phase, the
metal dissolution rates will thus decrease with several orders of magnitude.
Concluding remark
Passive films on stainless steels is one example of nano-engineering that has been around
for almost a century. The films are not at all passive. They constantly adopting to
changes in the environment. To better describe their flexible nature and rapid adjustment
to changes in the environment, they are rather to be described as active films.
References
1. C.-O. A. Olsson and D. Landolt, Electrochim. Acta (2003)
2. G. Herting, I. Odnevall Wallinder and C. Leygraf, J. Electrochem.
Soc. 152 B23(2005)
3. C.-O. A. Olsson and D. Landolt, J. Electrochem. Soc. 147 (11) 4093 (2000)
The Authors
Claes Olsson is associate professor in engineering physics at Uppsala University and
employed at the Outokumpu Avesta Research Center. He has worked extensively with
real time measurements of passive film growth on stainless steels and valve metals.
Gunilla Herting. M.Sc. is working on finishing her PhD, focussed on studies of metal
dissolution from stainless steels and related materials, at the department of corrosion
science at the Royal Institute of Technology.
Inger Odnevall Wallinder is associate professor in Corrosion Science at the Royal
Institute of Technology and has predominantly been working with atmospheric
corrosion. She has identified a number of new corrosion products. Her present research
is focussed on environmental- and health aspects of metal dissolution from different
types of surfaces.
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