Critical factors affecting the high-temperature corrosion performance of iron aluminides
P.F. Tortorelli
a,U, K. Natesan
baOak Ridge National Laboratory, PO Box 2008, Oak Ridge, TN 37831-6156, USA
bArgonne National Laboratory, 9700 S. Cass A¨enue, Argonne, IL 60439, USA
Abstract
Iron aluminides are known to exhibit good-to-excellent corrosion resistance in a number of high-temperature environments.
Under most conditions, this resistance derives from the establishment and maintenance of a sound and adherent alumina layer. Consequently the performance of iron aluminides under different aggressive high-temperature can be related to
Ž .
fundamental factors that affect the development, adhesion, and durability lifetime of protective alumina. Overall corrosion resistance depends not only on thermodynamic stability of the reaction product in a particular environment and its growth kinetics, but also on scale integrity and morphology, the chemical and physical nature of the oxide]metal interface, alloy strength, and the specific composition of the iron aluminide. Despite the multiplicity of factors, a relatively simple lifetime prediction model based on aluminum consumption can be used to evaluate the influences of changes in material parameters and to determine approaches to improving high-temperature corrosion resistance of iron aluminides. Q 1998 Published by Elsevier Science S.A. All rights reserved.
Keywords: High-temperature corrosion; Iron aluminides; Oxidation; Sulfidation; Alumina scales; Scale adhesion; Oxidation lifetimes
1. Introduction
It has been more than 60 years since the potential of iron]aluminum alloys as oxidation-resistant mate- w x rials at high temperatures was first identified 1 . However, it has only been within approximately the last 10 years that studies of their high-temperature corrosion behavior over a wide range of exposure conditions have begun in earnest. As a result, iron aluminides based on Fe Al and FeAl have emerged3 as good-to-exceptional corrosion-resistant alloys for elevated-temperature use in several aggressive envi- ronments. Much of the work on the high-temperature corrosion of iron aluminides has already been ade- w x quately reviewed within the past few years 2]8 . Therefore only the salient features of the corrosion performance of iron aluminides will be summarized with the major emphasis being on recent results not yet included in previous reviews. The remainder of
UCorresponding author.
the paper will then be devoted to a synopsis of what is known about the basic processes and materials properties that underlie the observed corrosion be- havior in these various environments through their influence on the development, integrity, and adhesion of alumina formed on iron aluminides. In view of recent progress in understanding some of these fun- damental factors, it is timely to examine their signifi- cance with respect to the high-temperature corrosion behavior of these alloys vis-a-vis other alumina-` formers and in terms of approaches to improved performance
2. Summary of corrosion performance of iron aluminides in high-temperature environments 2.1. Air or oxygen
Because alloys based on Fe Al and FeAl form3 Al O during exposure to oxidizing gases, they typi-2 3 cally display relatively low oxidation rates when com- pared to iron-based and other alloys that do not form
0921-5093r98r$ - see front matter Q 1998 Published by Elsevier Science S.A. All rights reserved.
Ž .
P I I S 0 9 2 1 - 5 0 9 3 9 8 0 0 9 2 4 - 1
Ž .
alumina or silica in comparable temperature ranges w x5,6 . With respect to other iron- and nickel-based alumina formers, iron aluminides do not offer any particular advantage in terms of high-temperature oxidation rates in air or oxygen at temperatures in excess of ;9008C. Recent extensive studies of Fe Al3
alloys containing 2]5 at.% Cr and various minor additions of oxygen-active elements have shown that their long-term oxidation performance approximately matches that of FeCrAlY alloys and NiAl at 1000 and
w x
11008C, but is inferior at 1200 and 13008C 9,10 . Other work has suggested that oxidation of iron aluminides without these oxygen-active elements is worse in air than oxygen, particularly at 1000 and 11008C, due to internal nitridation below a defective
w x
alumina scale 11 . The Fe Al alloys produced by3 ingot-metallurgy processes tended to have worse oxi-
Ž .
dation behavior that is, exhibit greater spallation
Ž .
than oxide-dispersion-strengthened ODS iron w x
aluminides of similar compositions 9 .
2.2. Sulfur-bearing gases
The reaction of iron aluminides with high-tempera- ture sulfur-bearing gases depends on the nature of the sulfur species and the partial pressure of oxygen in the particular environment. Alumina will form at very low partial pressures of oxygen even in the pres-
w x
ence of a significant sulfur activity 12,13 . In pure sulfur vapor or H S2 rH mixtures, alumina does not2
form. Nevertheless, binary iron]aluminide alloys con- taining G18% Al still show good resistance up to
w x
;7508C 14]18 . At 750 and 9008C in sulfur vapor at 10y3Pa, a Fe Al alloy containing 28% Al only formed3 an Al S external scale, but at sulfur pressures2 3 )133 Pa sulfidation rates increase rapidly with time above
w x 8008C 16 .
In many service environments, sulfur will co-exist with oxygen, carbon, chlorine, andror hydrogen. Ex- amples include flue gases, which contain SO2 and
w Ž .x
have high oxygen partial pressures p O2 and syngas produced from gasification of coal, in which the sulfur
w Ž .x Ž .
partial pressure p S2 is high and p O2 is low. In these cases, the oxygen potential is sufficient to form Al O on Fe Al and FeAl alloys, and corrosion be-2 3 3 havior is controlled by the integrity and adherence of the protective alumina surface product. The impor- tance of the Al O product is shown in Fig. 1, which2 3 indicates that a FeCrAl alloy in H S-bearing mixed2
Ž . Ž .
gases of high p S2 and low p O2 can only approach the corrosion behavior of FeAl if it is preoxidized to form alumina. Because iron aluminides form a pro- tective form of alumina under these conditions, they are uniquely corrosion resistant in mixed-gas environ-
w x
ments containing sulfur and oxygen 5,7,13,19]22 . Figs. 1 and 2 clearly indicate the better sulfidation resistance of FeAl and Fe Al, respectively, with re-3
Fig. 1. Weight change behavior of Fe-40 at.% Al and MA-956 ŽFe]19%Cr]9%Al]Y O2 3. alloys exposed to 2% H S2 ]CO at2
w Ž . y6 Ž . y2 0 x
7008C p S s102 atm, p O2 s10 atm ; from Gesmundo et w x
al. 22 .
Fig. 2. Weight change behavior of Fe]Al alloys isothermally ex-
w Ž . y6 Ž .
posed to H S-H -H O-Ar at 8002 2 2 8C p S s102 atm, p O2 s
y2 0 x
10 atm . Concentrations are in at.%. The ‘q’ indicates the presence of additional minor alloying elements that have little
w x effect on observed behavior; from DeVan 19 .
spect to Fe]Cr]Al alloys similarly exposed to high
Ž . Ž .
p S2 ]low p O2 gases. In Fig. 2, the detrimental effect of Cr on corrosion behavior at 8008C is shown;
resistance to sulfidation is better when the chromium concentration in Fe Al is3 F2 at.% or, alternatively, w x if the aluminum concentration is increased 4,19 . Results from experiments in a similar high-sulfur, low-oxygen environment at 7008C by Schutze and¨
w x
Noth 23 are consistent with these earlier observa-¨ tions regarding the good corrosion resistance of iron aluminides; measured rates were comparable to, or perhaps slightly lower than, those of MoSi and TiAl,2 which, in turn, were superior to standard heat-re- sistant alloys based on iron and nickel. Exposures of a fully-heat treated ODS Fe Al-2 at.% Cr alloy in a3 typical H S-H -H O-Ar gas at 8002 2 2 8C resulted in low corrosion rates that matched those measured earlier for a similar composition produced by conventional
w x ingot metallurgy techniques 24 .
In combustion atmospheres, the corrosion behavior of iron aluminides is very much like that observed in
air or oxygen environments as oxygen partial pres- sures are high while those of sulfur are low. Results
w x for such environments are reviewed in refs. 5,7 . The oxidation rates for Fe Al and FeAl determined for3 purely oxidizing conditions are approximately equal to those for mixed gases containing up to at least 1%
SO .2
2.3. Carbon-bearing gases
Relatively little work regarding the performance of iron aluminides in carburizing environments has been reported. As in mixed gases with sulfur, iron aluminides should be corrosion resistant provided that the relative partial pressures of oxygen and carbon favor alumina stability and the kinetics of Al O2 3 formation are not sluggish. Recent results by Smith et
w x
al. 25 have shown that FeAl alloys have superior high-temperature carburization performance to tradi- tional heat-resistant alloys, presumably because alumina forms and acts as a barrier to carbide forma- tion.
If the carbon fugacity of the environment is suffi- ciently high that mixed carbides of iron and aluminum can form, there is a potential for metal dusting of iron aluminides. Metal dusting of iron-based alloys in- volves the rapid ingress of carbon into the metallic material leading to supersaturation, formation of sur-
w x
face M C carbides, and deposition of carbon 26,27 .3 Ž . Carbon deposition sets the carbon activity ac at the metal surface to unity and the carbides then decom- pose to form metal particles and filamentous carbon in the form of a dust that can be carried away by the flowing gas. Strauß et al. have recently shown that substantial degradation of Fe Al by metal dusting can3 occur under environmental conditions where there is a high carbon activity, but that the addition of
Ž .
chromium 2]4 at.% to the alloy significantly reduces w x
metal loss 28 .
2.4. Chlorine-bearing gases
Natesan has demonstrated that iron aluminides based on Fe Al have adequate corrosion resistance in3 high-temperature gas mixtures containing ;2% HCl w x29 . The presence of HCl in a H S-H -H O environ-2 2 2 ment only slightly degrades the excellent performance
Ž .
of these alloys Fig. 3 . In a combustion gas with HCl, a 5 at.% Cr level significantly improved the corrosion behavior of an Fe Al alloy compared to3 F2 at.% Cr.
This is the reverse of the effect of Cr on the corrosion of Fe Al in the H S-H -H O environments without3 2 2 2
w x
HCl 4,19 . Relatively good performance of Fe Al3 ]Cr
Ž .
alloys 28 at.% Al in a H S-H -H O-HCl gas at 4502 2 2 and 5508C has recently been reported by Saunders et al., but rather substantial corrosion of an Fe-16 at.%
Fig. 3. Weight change rates for Fe Al, FeAl, and heat-resistant3 alloys and for Fe Al and FeAl top coatings on type 316 stainless3
Ž .
steel substrates exposed at 6508C to H S-H -H O -HCl ; from2 2 2
w x Natesan 13 .
w x
Al-5% Cr composition was noted 30 . Bakker has demonstrated that while the high-temperature resis- tance of Fe Al type of aluminides in a H S-H -H O-3 2 2 2 HCl environment is good, there is significant degrada- tion when iron aluminides are cooled in the HCl- bearing gases such that they are exposed to acidic
Ž . w x
condensates downtime corrosion 31 .
2.5. Molten salts and condensed deposits
Iron aluminides are not particularly resistant to hot corrosion by molten sulfates and the presence of such salts significantly increases the corrosion of iron
w x
aluminides in SO -containing mixed gases 5,322 ]34 . It is expected that chromium additions to Fe]Al compositions would decrease corrosion rates in sul-
Ž w x.
fates see, e.g. Gleeson et al. 35 . There is some limited evidence in this regard for iron aluminides w x32 , but the effect of alloying additions has not been fully explored. The corrosion of iron aluminides in contact with solid CaSO surface deposits, as typically4 formed in fluidized coal combustors, is much less severe than that by molten alkali sulfates, and a Fe Al alloy showed good resistance under such condi-3
w x tions 36 .
Laboratory corrosion studies using an ash deposit chemistry typical of coal-fired boilers Fe O , K SO ,Ž 2 3 2 4 Na SO2 4.showed that, after 800 h at 650 and 7008C in a simulated flue gas 14% CO , 10% H O, 4% O ,Ž 2 2 2 balance N2.with 0.25% and 1.0% SO , Fe Al-2 at.%2 3 Cr was more heavily attacked than stainless steels while Fe Al-5 at.% Cr had the same order of weight3 change and thickness loss as type 347 stainless steel w x37 . However, after 16 000 h in the superheater area
Ž .
of a 250-MW coal-fired boiler 620]7308C , an Fe Al-3
5 at.% Cr aluminide showed better relative perfor- w x
mance than other alloys 38 . Preliminary metallurgi- cal examination showed that the iron aluminide had a total wastage 2]10 times lower than type 347 stainless steel and was comparable or better in corrosion per- formance than the most resistant alloys HR3C andŽ Ta-modified type 310 stainless steel ..
Corrosion of iron aluminides by a highly aggressive,
Ž . Ž
oxidizing molten NaNO -KNO -Na O3 3 2 2 used in chemical air separation at 650. 8C proceeds by oxida- tion and a slow release from an aluminum-rich product layer into the salt such that the compositions with higher aluminum concentrations yielded significantly
w x
better resistance 39 . Such Fe]Al alloys are relatively resistant compared to many other metallic materials Žstainless steels, nickel-base alloys, refractory metals ,. particularly when the aluminum concentration ex-
w x
ceeds 30 at.% 40 . Aluminides based on Fe Al would3 be expected to offer excellent resistance to NaNO3 Ž-KNO -Na O at lower temperatures3. 2 2 Ž-6008C or. in nitrate salts that do not contain Na O2 2 Žwhich severely exacerbates the corrosivity of the melt . Pre-. liminary 9008C results from exposure of FeAl to mol- ten NaCl-Na CO2 3 Žused as an oxidizer for waste disposal showed significantly better corrosion resis-. w x tance of this aluminide compared to Inconel 600 41 .
3. High-temperature corrosion behavior of iron- aluminide coatings
The above summary indicates that there are a number of environments for which high-temperature corrosion resistance of iron aluminides containing greater than approx. 20]25% Al extends to tempera- tures at which these alloys have limited or poor me-
w x
chanical strength 2 . Because of this, iron aluminides appear particularly attractive as coatings or claddings on more conventional higher-strength materials which are less corrosion-resistant at high temperatures. Ef- forts at synthesizing iron-aluminide coatings have in-
w x
cluded weld overlay 42,43 , electro-spark deposition w x44 , pack cementation 45 , and thermal spray tech-w x
w x
niques 46 . The fabricability of composite iron- aluminide-clad stainless steel tubing via co-extrusion
w x
has also been investigated 47 . For all these various coatingsrclads, the high-temperature corrosion be- havior in air and mixed-gas environments have been
w x
and are being evaluated 13,24,29,43,48 . An example of the effectiveness of these coatings is shown in Fig.
3. When crack-free iron-aluminide coatings of ap- propriate compositions sufficient aluminum concen-Ž tration, optimized chromium levels, etc. can be syn-. thesized, corrosion results are similar to what are described above for bulk alloys.
4. The nature of protective alumina scales on iron aluminides
The resistance of iron aluminides to chemical degradation in a number of high-temperature envi-
Ž .
ronments see above correlates with the ability of these alloys to quickly form and maintain a protective alumina layer. Therefore studies aimed at under- standing the fundamental and phenomenological as- pects of alumina growth, adherence, and reliability are of importance in developing iron aluminides that are highly corrosion resistant. Much research on alumina scale growth and adherence has been con- ducted, but it only has been within the past few years that such studies have included iron aluminides.
Therefore a significant amount of scientific informa- tion about alumina grown on Fe]Al substrates has recently become available for comparison to what is known about other alumina formers. A summary and assessment of these findings follow.
4.1. Formation of alumina scales
Upon exposure to oxygen-bearing gases, the aluminum in iron aluminides is preferentially oxidized to form a surface layer of almost pure alumina. Ac- cording to the Wagner model of selective oxidation Žsee, e.g. Doychak 3 , Kofstad 49 and Meier et al.w x w x w x.50 , this will occur when the concentration of aluminum exceeds a minimum level in the alloy.
Thermodynamically, this concentration is quite low in most cases and, in reality, a concentration signifi- cantly greater than this is needed to assure that a
Ž .
continuous layer of alumina only is formed. For a cast binary Fe]Al alloy, approx. 16]19 at.% Al is needed to suppress internal oxidation and growth of
w x
iron oxides at 800]9008C 21,51,52 . The addition of chromium reduces this critical aluminum concentra-
w x w x
tion 21,51 , while nickel increases it 51 . Metallurgi- cal factors can also be important: a recent comparison of binary ingot-processed and ODS Fe]Al alloys showed that the latter materials had a lower critical Al concentration for alumina scale formation at 9008C
w x
by approx. 4 at.% 53 . This reduction was attributed to the fine grain size of the alloys stabilized by oxide dispersions.
The aluminum levels present in iron aluminides are well in excess of the critical concentrations discussed above, and, as expected, alumina scales form readily above approximately 5008C upon exposure of these
w x
alloys to an oxidizing environment 19,51,54]56 . Chromium appears to accelerate the formation of
w x protective alumina in the Fe Al system 56 . The3 predominant surface product that forms on iron aluminides between 600 and 8008C has been reported
w x
to be g-Al O 57 , but it is also quite possible for2 3
u-Al O or other forms of alumina to exist in this2 3
w x
temperature range 58 . At higher temperatures, a- Al O is the principal steady-state protective product2 3 w x58 . The temperature at which there is a transition from the other types of alumina to the slower growing
w x
a-Al O appears to be about 9008C 56,59 . This is a2 3
lower transition temperature than what is expected w x
based on results for b-NiAl 60 and may be due to Fe O or Cr O acting as precursor templates for the2 3 2 3
w x
growth of a-Al O on Fe Al alloys 59 . It is also2 3 3
interesting to note that, at least for iron aluminides, the concept of protective alumina can be extended to lower temperature where the non-a forms of Al O2 3
exist. In view of the good corrosion performance of
Ž .
these alloys between 500 and 9008C see above , these lower-temperature Al O -based products can be con-2 3 sidered protective in many of the environments for which results are available.
4.2. Maintenance of alumina layers
For Fe]Al alloys, oxidation lifetime is defined as the time over which an alumina scale provides protec- tive behavior to the underlying substrate. This period is characterized by low oxidation rates associated with the establishment and maintenance of a continuous alumina surface product formed by preferential reac- tion of environmental oxygen with aluminum in the substrate. However, as aluminum is depleted in the substrate as an oxide scale grows or reforms if spalla-Ž tion of the scale occurs eventually the aluminum. concentration in the substrate, C, falls below that needed to maintain an effective surface activity to
Ž .
form alumina C . At this point, non-protectiveb
Ž .
products such as iron oxides , with substantially higher corrosion rates, develop and scaling changes from a process characteristic of a protective regime to one of breakaway oxidation or sulfidation, carburiza-Ž tion, etc., if the alumina was preventing such reac- tions . Therefore oxidation lifetime can be formally. defined as the time-to-breakaway, t , or, equivalently,b the time over which C decreased to C . Beyond thisb point, metal wastage by oxidation or other high-tem- perature reactions is greatly accelerated and struc- tural integrity can be rapidly compromised by loss of load-bearing area. Therefore maximizing tb can ef- fectively extend a component’s high-temperature ser- vice life in a variety of high-temperature environ- ments.
w x
Quadakkers and colleagues 61,62 have studied and modeled the time-to-breakaway for iron-based
w x
alumina formers and Bennett et al. 63 and Wright et w x
al. 64 have applied such an approach to iron aluminides. In these oxidation lifetime models,
Ž . 41rn Ž .
tbs A? C yC ?o b r?hrK 1
where C is the initial aluminum concentration in theo alloy, r is the alloy density, h is the specimen thick-
Ž .
ness, K is the oxidation metal wastage rate, and A is a constant based on the atomic weight ratio of aluminum and oxygen. Eq. Ž .1 clearly shows that oxidation lifetime is directly dependent on the amount of aluminum available for reaction Žthat is, the aluminum reservoir through C and h and inversely. o proportional to the rate at which aluminum is con-
Ž .
sumed K . In this formulation, K accounts not only for the alumina growth rate, typically characterized by a parabolic rate constant, k , where, for a product ofp thickness x after time t,
2 Ž .
x sk ?tp 2
but also for metal wastage caused by scale spallation such that the total mass change by oxidation includ-Ž ing that due to adherent scale as well as what has spalled ,. Dm , can be modeled ast
n Ž .
Dm sK?tt 3
Finally, t increases with decreasing C as this effec-b b tively raises the aluminum reservoir.
Iron aluminides have a potential advantage over other iron-based alumina-forming alloys in that C iso much greater in Fe Al and FeAl. However, despite3 their greater aluminum reservoir, oxidation lifetimes at 1200 and 13008C are not significantly greater for
Ž .
Fe Al alloys 28 at.% Al than those for FeCrAlY3 Žwith approx. 10 at.% Al. w63,64 . Preliminary resultsx for an FeAl with ;37 at.% Al indicated a similar
w x result 65 .
4.3. Al O scale adherence2 3
The less-than-expected oxidation lifetimes of iron aluminides are primarily due to the higher overall aluminum consumption rates of these alloys K, Eq.Ž Ž ..1 which offset the greater aluminum reservoir CŽ o
. w x
yC of these aluminides 63,64 . Much of the dif-b
ference in K ’s for iron aluminides and other iron- based alumina formers relates to greater spallation susceptibility of the former alloys rather than a higher
w x
isothermal scale-growth rate 9,10,24,63,64 . This is illustrated in Fig. 4, which compares the total and
Ž .
specimen weight gains Dm and Dm , respectivelyt s
of a cyclically oxidized Fe]28%Al]5%Cr]0.1Zr to similar data for APM an FeCrAlŽ ]ZrO alloy that is2
the best of its class with respect to scale adherence. w x10 . The difference represented by ŽDm yDm ,t s. which is a direct indication of spallation losses, is negligible for APM, but significant, and increasing with time, for the iron aluminide. Therefore despite their relatively high Co values, iron aluminides will
Fig. 4. Weight change as a function of time for cyclic oxidation of
Ž .
an iron aluminide and a FeCrAl]ZrO alloy APM . Each point2
represents measurement after a 100-h exposure to oxygen at 11008C and rapid cooling to room temperature. Total gain is the weight of specimen and any spalled oxide. Any difference between total gain w x and specimen weight represents spallation loss; from Pint et al. 10 .
have a substantial time-to-breakaway advantage over other iron-based alumina formers only if Al O -scale2 3 adherence can be improved.
Some progress in extending the lifetimes of iron aluminides by decreasing K has recently been re-
w x
ported 66 , but spallation of Al O from these alloys2 3 remains generally worse than for other alumina- formers. In recent years, several factors have been directly or indirectly identified as affecting the adher- ence of alumna on iron aluminides. These include sulfur segregated to the scale]alloy interface, the
Ž .
presence of reactive elements REs in the alloy, interfacial and scale defects, and alloy strength.
4.3.1. Sulfur
The deleterious influence of indigenous sulfur on scale adherence for alumina- and chromia-forming
w x
alloys has been known for more than 10 years 67]71 , but this effect has only recently been documented for iron aluminides. Specifically, Hou et al. have shown that, at 10008C, sulfur segregated to the Al O ]metal2 3
w x
interface on Fe]28%Al]5%Cr 72 , but that inter- facial sulfur was not detected when an Fe]28%
w x Al]5%Cr]0.1%Zr was similarly exposed 73 . Fur- thermore, under the same conditions, the latter alloy showed substantially greater resistance to spallation
w x than the former one 74 .
4.3.2. Reacti¨e elements
The beneficial effect of REs on scale adherence has
w x
been reviewed and analyzed many times 75]78 . One of the first reports of improved scale adherence at- tributed to the presence of a RE in iron aluminides was for small levels of Zr in an Fe]28%Al]5%Cr
w x
alloy 79 and this effect in Fe Al alloys containing Zr3
w x
has now been demonstrated several times 9,10,63,74 . More importantly, the extensive work of Pint et al.
w9,10,80]82 over the past few years has shown that:x Ž .1 there are a number of oxygen-active elements that improve the spallation resistance, and, therefore, de- crease the overall Al consumption rate of Fe Al3 alloys whether these are added as alloying elements
Ž . Ž .
or oxide dispersoids Fig. 5 ; 2 the RE effect in iron aluminides manifests itself, in terms of segregation of the RE and scale composition and morphology, in a similar manner to what is generally observed in other
w x Ž .
alumina-forming systems 58,78,83 ; and 3 REs are generally not as effective in Fe Al. The curves shown3 in Fig. 5 indicate that not all active elements are effective in improving oxidation behavior of Fe Al or3 Ž
w x.
other alumina formers } see Pint 78,83 and that
Ž .
an appropriate optimal concentration is needed to have a RE effect. The presence of a RE in iron
Ž .
Fig. 5. Weight change as a function of time for cyclic oxidation of Fe]28%Al]2%Cr FAS with various reactive element dopants added as
Ž . Ž . Ž .
oxide dispersoids, ingot-processed IrM FAS and Fe]28%Al]5%Cr]0.1%Zr FAL and a FeCrAl]ZrO alloy APM . Each point represents2
w x a measurement after a 100-h exposure to air at 12008C and rapid cooling to room temperature; from Pint et al. 9,10 .
Ž . Ž . Fig. 6. Fracture cross-sections of alumina scales formed on Fe Al alloys oxidized at 12003 8C in O for 96 h. a Fe]28%Al]5%Cr]0.1%Zr; b2
w x Fe]28%Al]5%Cr; from Prußner and Pint 11 .¨
aluminide leads, upon high-temperature oxidation, to development of an alumina scale with a primarily
Ž .
columnar grain structure see Fig. 6a that is quite distinct form the large equiaxed grains of the Al O2 3
Ž .
that grows on a similar alloy without the RE Fig. 6b . This characteristic is typical of alumina formers with alloying additions that are effective in promoting scale
w x
adherence 58,78 and occurs in concert with the transport of RE ions from the bulk to the oxide]metal
w x
interface and through the scale 78 . Such RE segre- gation has been observed for Zr and other additions to iron aluminides, as observed by scanning transmis- sion electron microscopy of thinned cross-sections of
w x oxidized specimens 81 .
4.3.3. Scale and interfacial defects
As shown in Fig. 5, the Fe Al alloy with the best3
Ž .
overall oxidation resistance FAS-Y O2 3 still exhibits a higher Al consumption rate than the benchmark
Ž .
alumina-forming FeCrAl APM due to greater scale spallation. This difference persists despite suppression of sulfur segregation and the related benefits to scale adherence conferred by RE additions. While some improvement may still be gained by further optimiza- tion of the concentration and distribution of the RE, it appears that other factors must help determine the relative spallation tendencies for alumina scales formed on different alloys. One key difference is the greater propensity for interfacial void formation asso-
Ž .
ciated with oxidation of ODS Fe Al alloys Fig. 7 as3
w x
compared with other ODS alumina formers 9,78 . A greater density of such voids would presumably lead to more rapid scale failure and spallation. In addition, there is now some evidence that alumina scales grown on Fe Al alloys have a higher overall density of3 defects Žmeasured using the whole-scale analysis
w x.
method of Hancock and Nicholls 84 when com-
w x
pared to Al O on FeCrAlY alloys 85,86 . The greater2 3 defect density associated with oxide scales on the iron aluminides would then translate into shorter times-
Ž Ž ..
to-breakaway t ’s, Eq. 1b as scale failure will occur earlier for these alloys.
Fig. 7. Transmission electron bright-field image of the metal]scale interface of Y O -dispersed Fe2 3 ]28%Al]2%Cr alloy after 2 h at 12008C in O . The arrows mark some interfacial voids; from Pint et2
w x al. 9 .
4.3.4. Substrate strength
It is generally expected that a weaker alloy will spall less because of stress relaxation and blunting of interfacial cracks by deformation of the substrate.
Scale stresses normally arise during the process of oxide growth and during cooling from the oxidation temperature because of the differences in the respec- tive coefficients of thermal expansion of the alloy and
w x
alumina 87]89 . When scale failure is by a wedge
w x
mechanism 89,90 , the creep rate of the substrate should have the greatest effect. However, a stronger substrate strength cannot be the reason for the rela- tively worse spallation of the Fe Al alloys. Although3 appropriate high-temperature mechanical properties data are not available, these iron aluminides are no stronger, and are probably weaker than, APM and other alumina formers to which they are compared Žsuch as b-NiAl . This is particularly true at 1200 and. 13008C, where the differences in spallation behavior
w x
are greatest 10 . Interestingly, recent results from the oxidation behavior of ODS Fe Al have indicated an3 apparent beneficial effect of substrate strengthening w x9 . As shown in Fig. 8, the presence of an oxide dispersion in a Fe]28%Al]2%Cr produced a flatter, more adherent scale after oxidation at 12008C. This
effect was observed for an alloy with an Al O disper-2 3 sion as well as one that also included Y O . While2 3 flattening the scale and improving scale adhesion in the short-term, the Al O dispersion, unlike a typical2 3 RE addition like Y O , did not change the scale2 3 structure nor confer improved oxidation resistance
w x
over extended times 9 . Nevertheless, the suppression of scale buckling, as shown in Fig. 8, represents an improvement in overall scale adhesion, which ap- peared related to a substrate strengthening effect.
The mechanism underlying this effect appears related to the presence of interfacial voids or any type of deviation from a planar oxide]metal interface. Such perturbations can lead to a redistribution of normally
Ž w x
compressive scale stresses see, e.g. Nicholls et al. 90 w x.
and Hou et al. 91 and thus create tensile forces that lead to convoluted scales by plastic deformation of the substrate and elongation of voids into buckles w9,91 . Presumably, such scale morphologies are morex prone to spallation. Under such conditions, a stronger substrate can better resist such deformation and maintain a flatter scale; it is this type of model that was proposed to explain the results shown in Fig. 8 w x9 .
As mentioned above, the expected inverse relation- ship between creep strength and spallation resistance
w x
is for failure by a wedge mechanism 89,90 and given the highly defective nature of scales on many iron aluminide alloys, and their low high-temperature strength, it is possible that oxide loss can occur in other ways. For example, the interaction of stress and interfacial voids can possibly lead to a buckling failure w x89 , which is normally only associated with very thin
w x
scales 90 . Under these circumstances, substrate creep w x
has little influence on failure 90 and an effect of alloy strength on spallation susceptibility can manifest itself in a different manner.
5. Improvements in the high-temperature corrosion resistance of iron aluminides
5.1. Promoting alumina growth by material modifications
Because the high-temperature corrosion resistance of iron aluminides depends on the growth and stabil- ity of surface alumina in different environments, ma- terial modifications that promote these factors will lead to improvements in the performance of this class of materials. As mentioned above, chromium addi- tions to Fe Al alloys increase resistance to metal3
w x
dusting 28 and may be beneficial for certain molten salt applications. In a broader sense, chromium may
Ž .
promote more rapid growth of alumina see above and this can be a benefit in a number of environ- ments, particularly at lower temperatures. However, the addition of more than 2% Cr to Fe Al is detri-3 mental to corrosion resistance in H S-bearing mixed2
Ž . Ž .
gases with high p S2 and low p O2 because the formation of chromium sulfides makes it more dif- ficult to maintain a continuous protective alumina
w x
layer 19 . For iron aluminides with higher aluminum
Ž .
contents FeAl , higher chromium levels can be used with compromising the excellent sulfidation resistance
w x
in such environments 19 . In addition, there is some evidence that a finer alloy grain size may promote the
w x
formation of Al O on iron aluminides 53 and other2 3 w x
alumina formers 92 . This may prove to be a particu- larly effective approach to achieving continuous alumina layers at lower temperatures. Alloying or processing to achieve a fine grain size may improve corrosion resistance if other factors affecting perfor- mance are not significantly compromised.
5.2. Maximizing time-to-breakaway
Based on the premise that protective alumina con-
Ž . Ž .
Fig. 8. Secondary electron images of the alumina scale surface after 2 h at 12008C in O a unmilled, extruded Fe]28%Al]2%Cr powder; b2
milled, extruded powder which contains an Al O dispersion; and c milled, extruded powder which contains a Y O dispersion. The addition2 3 Ž . 2 3 w x
of an oxide dispersion flattens thea-Al O scale on Fe Al; from Pint et al. 9 .2 3 3
trols the corrosion performance and lifetime in many high-temperature environments, increasing the aluminum concentration within the range represented by iron aluminides appears, by virtue of Eq. 1 , to beŽ . a straightforward method of increasing the time over which alumina will continue to form. However, such an approach may be inconsistent with optimization of other material properties or with processing fac- tors. This is particularly of concern for weld-overlay coatings of iron aluminides, where cold cracking of deposits can occur at aluminum concentrations
w x
)20]25% 42,43 . A better way to maximize tb
Žoxidation lifetime of iron aluminides would be to. decrease the rate of aluminum consumption ŽK in Eq. 1 . As discussed above, this would involve im-Ž ..
proving the scale adherence of alumina on iron aluminides as spallation susceptibility for this class of alloys is worse than for many other alumina formers.
Improvement in spallation resistance can be obtained by addition of small amounts of REs; alloying addi-
Ž .
tions such as Y or Zr that are effective in this regard
w x
have been identified 9,10 and appear to suppress the deleterious effect of sulfur segregation to the oxide]metal interface. Nevertheless, RE-doped iron aluminides still do not match the oxidation perfor- mance of the best alumina formers, particularly at temperatures in excess of 11008C. The addition of RE-oxide dispersions to iron aluminides is more ef- fective in improving spallation resistance, possibly by increasing substrate strength, which promotes less convoluted scales at high temperatures. However, for these ODS alloys, more voids are observed at the
Ž .
scale]metal interface Fig. 7 . At 900]10008C and below, RE-doped iron aluminides show fairly low K
Ž Ž ..
values Eq. 1 and oxidation lifetimes are long. Little work on the development of alumina on iron aluminides below 9008C has been done, but as shown by the above review of corrosion performance, Al O -based reaction products can provide extended2 3 corrosion protection at these temperatures.
6. Conclusions
Based on a review of the corrosion performance of iron aluminides in various high-temperature environ- ments, it was concluded that the ability to form and maintain a protective Al O surface product is the2 3 common factor linking reports of the good corrosion resistance of iron aluminides in many elevated-tem- perature environments. While protective alumina is normally considered to be a-Al O , which is typically2 3
not observed below 9008C, other forms of alumina that develop in the various mixed-gas environments at lower temperatures do provide corrosion protection.
Consideration of a simple oxidation lifetime model based on the ability to maintain a protective alumina
scale provides a way to relate Al levels and funda- mental factors controlling Al consumption to long- term corrosion performance. The high aluminum con- centrations of the iron aluminides increase oxidation lifetime, but, overall, their performance in this regard is not better than FeCrAl alloys with -20% Al, particularly at temperatures in excess of 11008C. This is due to generally greater spallation susceptibilities of the iron aluminides. Therefore improving alumina scale adherence is the key to extended lifetimes com- mensurate with the high aluminum levels in these alloys.
Important factors controlling the development and adherence of protective Al O2 3 on iron aluminides include: 1 appropriate alloying elements and mi-Ž . crostructures to promote alumina growth; 2 the levelŽ . of sulfur in the alloy and its ability to segregate to the scale]alloy interface as controlled by reactive ele-Ž
. Ž .
ments ; 3 the type and concentration of reactive element additions; 4 the nature of scale and interfa-Ž . cial defects the concentrations of which tend to beŽ
. Ž .
higher for oxide grown on Fe Al alloys ; and 5 alloy3 strength.
Acknowledgements
The nature of this paper is such that many of the ideas, explanations, and implications expressed in it reflect not only the products of the authors’ own work and analyses but the understanding brought about by the findings and insights of many others. In this
Ž .
regard, one of the authors PFT is particularly indebted to the influence and input of several of his colleagues at Oak Ridge National Laboratory: B.A.
Pint, I.G. Wright, J.H. DeVan, K.B. Alexander, and K Prußner. This paper has also specifically benefited¨ from collaborations of the authors with B.W. Veal ŽArgonne National Laboratory , P.Y. Hou and R.M..
Ž .
Cannon Lawrence Berkeley National Laboratory ,
Ž .
and M.J. Bennett Harwell Laboratory, retired . Spe- cial thanks go to J.R. DiStefano, B.A. Pint, and I.G.
Wright for review of this manuscript. The research was sponsored by the US Department of Energy’s Division of Materials Sciences and its Fossil Energy AR and TD Materials Program under contract DE- AC05-96OR22464 with Lockheed Martin Energy Re- search Corporation and contract W-31-109-Eng-38 with the University of Chicago.
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