[1962] the oxidation of iron-chromium alloys and stainless steels at high temperatures

Corrosion Science, 1961, Vol. 2, pp. 173-196. Pergamon Press Ltd. Printed in Great Britain.

R E V I E W A R T I C L E

THE OXIDATION OF IRON-CHROMIUM ALLOYS AND STAINLESS STEELS AT HIGH TEMPERATURES*

G. C. WOOD" I"

Department of Chemical Engineering, Faculty of Technology, University of Manchester, England

Abstract--The present state of knowledge of the oxidation of iron--chromium alloys and stainless steels is critically reviewed. A fundamental approach is adopted and special reference is made to recent researches in the field. An indication is given of how the application of certain techniques in carefully controlled experiments would lead to a greater understanding of these widely studied systems. This information should, in turn, make the development of steels for specific industrial purposes considerably easier.

R6sum~----L'auteur fait l'examen critique de l'6tat actuel des connaissances sur l'oxydation des alliages fer-chrome et des aciers inoxydables. II approche son sujet par des consid6rations fondamentales et donne des ref6rences sp6ciales sur les r6cherches r6centes darts ce domaine. I1 donne des indications sur la faqon dont l'application de certaines techniques dans des exp6riences soigneusement contr616es pourrait conduire h une meilleure compr6hension de ces syst6mes largement 6tudi6s. Cette information devrait rendre beaucoup plus facile le d6veloppement de la fabrication d'aciers pour des besoins industriels sp6cifiques.

Zusammenfassung--Der gegenw~irte Stand der Kenntnisse ~iber die Oxydation von Eisen-Chrom- Legierungen wird kritisch betrachtet. Unter besonderer Bedicksichtigung der letzten Untersuchungen auf diesem Gebiet wurde ein grundlegendes Bild des Vorganges entworfen. Es wird ein Hinweis darauf gegeben, wie die Anwendung bestimmter Untersuchungsverfahren in sorgf~iltig iiberwachten Versuchen zu einem besseren Verst/indnis dieses h/ifig bearbeiteten Gebietes fiihrt. Dieser Beitrag sollte deshalb die Entwicklung yon Spezialst~ihlen betdichtlich erleichtern.

I. GENERAL I N T R O D U C T I O N

THE oxidation of stainless steels and related alloys has received considerable at tention dur ing the past thirty-five years. Growth kinetics have been determined by thermo-

gravimetric techniques and the oxidation products have been examined and analysed by a number of methods including optical and electron microscopy, X-ray and electron diffraction and chemical, spectrographic, X-ray-fluorescence and electron-probe analysis. In addit ion, much practical experience has accumulated.

If a steel is exposed unde r mild condit ions a thin protective film is developed but in a severe envi ronment a thick stratified scale is produced. It is desirable in practice to operate in the thin-film region where this is possible. The condit ions leading to

the t ransi t ion from 'protective ' to 'non-protect ive ' behaviour depend on the alloy

* Manuscript received 7 May 1962. t Formerly of the Department of Metallurgy, University of Cambridge, England.

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174 G.C. WOOD

and the environment but are often unpredictable. Indeed there has been considerable disagreement about the structure and thickness of the layers and their mechanism of growth. Lack of standardization of the oxidizing variables has been one cause of discrepancy. It is also evident that the presence of small quantities of alloying or ' tramp' elements in the steels may have a considerable influence in some instances. These factors, together with the observation that protective films are generally very rich in chromium, have led some research workers to turn to the relatively simple pure iron-chromium binary system. However, even in very carefully controlled conditions, reproducibility is still sometimes difficult to attain.

From the practical viewpoint, it is possible in many cases to specify a suitable material, but this becomes increasingly difficult where cost and severe conditions require operation near the borderline between 'protective' and 'non-protective' behaviour. In this situation the only satisfactory procedure at present seems to be to test the materials with the most suitable mechanical properties under conditions very similar to those likely to be experienced in practice.

Fresh impetus has been brought to the problem recently because expensive tech- nological developments have been envisaged or initiated involving the operation of these alloys in different environments. For example, if steam superheaters are to be operated at considerably higher temperatures than at present, a material resistant to steam and heating gases must be provided. The operation of stainless steel in contact with carbon dioxide in the nuclear power industry is another important development.

It appears that there is considerable similarity in the nature of the oxidation of the various alloys in air, oxygen, water vapour, carbon dioxide and certain other gas mixtures, although the rates for a given set of conditions may be very different. Catastrophic oxidation and other cases where the protective film tends to be destroyed with the aid of small proportions of impurities in the environment are somewhat different and will not be discussed in great detail.

This paper has been writtert with the object of presenting the views of the author on the fundamental aspects of this complex subject, reference being made to the original papers that influenced these views. Obviously in matters of controversy it is always advisable for readers to turn to these papers for more detail.

Brief reference is made to practical schemes only where it is felt appropriate. Much of the early work on compatibility has already been reviewed elsewhere. 1

2. PRESENT KNOWLEDGE 2.1 Growth kinetics

(a) General theory. Most of the common growth relations have been observed, which is perhaps hardly surprising considering the wide range of alloys and conditions that have been used. However, in order to attempt some form of generalization, the schematic weight-gain/time curve shown in Fig. 1 may be used.

Under mild conditions a thin 'protective' film is built up at a declining rate according to the curve OAD. When the atmosphere is somewhat more severe, however, after an initial protective period OA (sometimes referred to as an induction period), there is a sudden increase in rate AB (break-through). This stage is often followed by a further reduction in rate BC (self-healing) but in some cases where the oxide stays non-protective BE may be followed instead. In a very aggressive atmo-

The oxidation of iron-chromium alloys and stainless steels at high temperatures 175

i

E 1

i I ii

ii C

Time (i')

FIG. 1. Typical growth curves for the oxidation of iron-chromium alloys and stainless steels

sphere, no induction period is observed and AB is followed directly. The impression gained from the curves often depends upon the scales to which they are plotted and this should always be taken into account when comparing results.

The lengths and shapes of the various portions of the resultant growth curve depend critically on the alloy and environment in a complex way. There may be several breaks of the type shown in a single curve. This preliminary summary should be borne in mind when considering the detailed cases in the following sections. These examples have been loosely classified but it will be appreciated that this is principally for convenience as there are many variations in conditions within the sections.

(b) Pure iron-chromium alloys. A recent study ~, z of the oxidation of pure iron and iron-chromium alloys containing 0"2-10"07o chromium (mainly abraded surfaces) in oxygen at 750--1025°C shows that whereas the former follows a parabolic 'law', the latter obey no simple or reproducible growth relation. At very low chromium contents, the alloys oxidize more rapidly than iron initially but for longer periods the rate is less. As the chromium content is increased the oxidation rate falls at any given temperature. The accelerations in growth rate do not occur at reproducible times or average thicknesses. Similar experiments a on an iron-12 70 chromium alloy (abraded and vacuum annealed) show no increase in rate at 825°C, but a break appears rapidly at 855°C. Abraded specimens of 25-37.5 ~ chromium content follow a parabolic curve when oxidized in oxygen (100 mm partial pressure) at 700°-1200°C. I ron- 37"570 chromium oxidizes more rapidly than iron-2570 chromium below 1000°C. Abraded dure chromium in 76 mm partial pressure of oxygen obeys a parabolic 'law' below 900°C but above 900°C and a film thickness of 4800A the rate increases sud-

176 G.C. Wood

denly, only to fall offagain on further oxidation. 5 Large increases in rate above 1050°C are attributed to lack of adhesion and the volatility of chromium. The oxidation rate is often slower in air. 6

It has also been shown that whereas abraded i ron-14 .4~ chromium gives an induction period in air at 1000°C, none is readily detectable in steam. 7 Abraded, annealed and chemically cleaned alloys containing 5-25 ~ chromium obey a parabolic curve at 700°C and 900°C but a linear one at l l00°C in a 90 °//o a~gon:10~ water vapour atmosphere. 8 The results appear relatively independent of the chromium content for 5-15 ~ chromium alloys within the limits of experimental error or reproducibility. It should perhaps be noted here that this change from parabolic to linear behaviour may be due to the surface adsorption of water vapour and subsequent chemisorption of the oxide ion becoming rate-determining.

(c) h71pure iron-chromium alloys and chromium steels. Russian studies 9, lo on i ron- chromium steels containing 12-13 ~o chromium with other small additions in oxygen and air/steam mixtures show an induction period over a considerable temperature range. Oxygen attacks the steels three times as fast as air and shorter induction periods occur at higher temperatures and higher water vapour pressures. During the induction period a logarithmic ' law' is obeyed but this is succeeded by a parabolic relation. Similar results are obtained for 5.8 ~ chromium-0.38 ~ silicon 11 and 3.2 ~ chromium- 0.5~(, molybdenum '~- steels. In the latter case steam causes a higher oxidation rate than relatively dry air and as the temperature is raised the behaviour approaches that of pure iron.

Other investigators 13 studied pickled 12'5~o and 16~ chromium steels and austenitic chromium-nickel steels in dry oxygen/nitrogen mixtures. The equation y" =- kt holds in the temperature range 815:-950°C (with ,7 between 1 and 2) in the initial stages but the rate increases rapidly after the induction period. The rate may decrease again at a later stage. The length of the induction period decreases as the chromium and nickel contents decrease and as the oxygen partial pressure and temperature increase.

Abraded or electropolished 12 ~ and 16 ~ chromium steels containing a variety of other elements including silicon show an increase in growth rate after an induction period at 870°C? a Both of these alloys and certain austenitic chromium-nickel steels oxidize more rapidly in moist than dry air? a, ~ However, a 26.5 ~o chromium-0.44 0/~ silicon steel gives several breaks of the type illustrated in Fig. 1 and also oxidizes more rapidly in dry than in moist air at 980°C and 1100°C. The austenitic steels also show more than one break in some cases. The silicon present may appreciably affect the growth laws obtained.

Abraded iron-chromium alloys containing 25 ~ and 37.5 ~ chromium and small additions of yttrium (0.5-2.0 ~ ) or certain other rare-earth elements obey a parabolic "law'. 18 The mechanism does, however, seem somewhat different in the early stages of oxidation. Similarly prepared iron-25 ~ chromium-5 ~o aluminium alloys containing other trace elements oxidize rapidly initially but soon follow a parabolic relation in 76 mm of oxygen at 700-1100°C. x7 In the temperature range 900-1050°C the rate is nearly constant. This appears to be related to the type of oxide predominating in the scale changing in this temperature range.

(d) Austenitie nickel-chromium steels. When abraded and electropolished austenitic


steels of the 187o chromium-87o nickel series are oxidized in air at I050°C, a curve of the type OABC in Fig. I is obtained, as Although the length of the induction period is not reproducible, the first points always lie on the same parabolic curve. In the course of the following rapid scaling period (which does not exceed several hours) the mass of oxygen absorbed is a linear function of time. It appears that the increase in weight during the course of this rapid oxidation is approximately a parabolic function of the duration of the induction period. In the third region the oxidation rate falls off rapidly but in an irreproducible manner. Other work on austenitic steels has already been discussed, x3,15

A reasonably comprehensive sample of the thermo-gravimetric results has been given above. Reference may also be made to other studies made on various alloys. 19-82

Some attempts have been made to compare the behaviour of austenitic chromium- nickel steels of various compositions with chromium steels of somewhat similar performance, z~,z3 However, these are of a somewhat 'ad hoc' nature, cannot hold under all conditions of oxidation and will not be further considered here.

(e) The influence of surface finish. As has already been intimated, the initial surface finish and the manner in which the specimen is brought to the reaction temperature may have an appreciable influence on the reaction rate. Abrasion is the most com- monly adopted finish but, as is well known, a surface consisting of internally stressed, disarrayed alloy containing entrained oxide and polishing material, and covered by an oxide formed by preferential oxidation, is obtained? 4 Etching may leave a rough surface enriched in the more noble constituent, whereas electropolishing, although giving a flat surface, often leaves a surface film. If the specimens are raised to temperature in the test atmosphere, in vacuum, or in an inert gas, preferential oxidation will occur. I f hydrogen is used and is not extremely dry the surface effects discovered by French workers, to be described later, may influence the results? 5.38

A few investigators have studied the influence of surface finish on growth kinetics. It has been suggested z4, 35 that the observed induction period is due to the presence of a protective film developed during the preliminary treatment and that this phenomenon is not truly characteristic of the behaviour of the alloy at temperature. Pre-oxidation treatments at elevated temperatures affect the subsequent scaling rate? 3 A recent study has shown that an i r o n - 2 6 ~ ch romium-0 .55~ silicon alloy oxidizes more rapidly in oxygen at 980°C after electropolishing than after abras ionY It is suggested that the most appropriate treatment is abrasion, electropolishing to remove con- taminated and worked alloy, annealing in argon at 1100°C, electropolishing to remove the preferentially oxidized layer and finally a controlled minimum etch. Hydrogen pretreatment of dilute iron-chromium alloys still leads to irreproducible results and irregular growth curves, z.z The longer the hydrogen pretreatment the faster the initial oxidation rate and the slower the long-term rate. It is also interesting to note here that the rates of oxidation of certain stainless steels decrease on the application of a cathodic current across the oxide and increase with a similarly applied anodic current. as

2.2. Morphology and structure (a) Films obtahled by low-temperature pretreatment. It is likely that the thin film

protecting the alloys at ambient temperatures is principally Cr20a, at least for the

178 G.C. WOOD

higher chromium contents. Surface films stripped from polished austenitic nickel- chromium steels and analysed chemically show marked chromium enrichment. 34 Films from brightly polished specimens may contain as much as 90~o CrzO3, the rest being mainly Fe2Oz with no nickel enrichment. In some cases, however, the polishing material is found in the oxide. The film thicknesses are apparently in the range 30-160A depending on the alloy, the nature of the polishing and the figure adopted for the surface roughness factor. The suggestion has been made 34 that the chromium enrichment is in some way associated with surface flow, the rate of oxidation being partly controlled by the free energies of formation of the respective oxides and to a considerable extent by selective oxidation, arising from the restriction of oxygen at the interface between the alloy and the polishing material.

It is of some interest to consider the nature of the films present on various alloys during anodic passivation. Some workers consider that the film on alloys containing more than 16~o chromium is the spinel Fe Fe(~_x) CrxO4 a9 but it seems more likely to be Cr2Oz in most cases? ° It may also be calculated that the passive film on iron- chromium alloys containing up to 12~ chromium is 6 atom layers thick, on allgys up to 18",~ chromium 3 atom layers thick and on pure chromium a monolayer. 41 These calculations are, however, based on a large roughness factor and are open to dispute. Similar protective films on stainless steels have received some discussion 34, 4,,-48 and seem to depend for their resistance on layers enriched in chromium, silicon or molybdenum.

Films left after electropolishing in highly oxidizing solutions probably consist principally of Cr.,O3 containing a greater proportion of Cr 6+ ions than is usual in this oxideY Estimates of film thickness on austenitic nickel-chromium steels anodically treated in dilute hydrochloric acid have been made by electron microscopy. Values of 6.5 A and 30/1 are given for the 'plateau' and 'gas evolution' regions respectively, a9

(b) The morphology of thin films. Having considered the initial film, we can now discuss what happens on heating in various atmospheres.

When iron-chromium alloys containing 18-30~ chromium are electropotished, pretreated in hydrogen in the range 900-1300°C and then preferentially oxidized at the same temperature in a hydrogen atmosphere containing small traces of water

vapour, three distinct areas on log ~l-120/temperature diagrams are observed? 5,36 ,uH~

Very dry hydrogen gives a plane specular surface. As i~tH20 is increased, no oxide is visible but a heavily serrated surface is observed. These striations are due to the surface of the alloy assuming a topography of minimum free energy when a thin chemisorbed layer of oxygen or oxide is formed. At high I~H20 a surface partly covered by numerous Cr203 crystals is observed. The sequence in producing an oxide film may therefore be firstly the growth of a very thin layer followed by the formation of Cr203 nuclei. The number of these depends on the crystal orientation and the gas composition, temperature and pressure. At present the location of the nuclei is not clear but it may be related to dislocations or other defects in the underlying alloy. 5° Once formed, the nuclei grow laterally until the surface is completely covered. At higher temperatures the number of nuclei falls, the surface diffusion rate increases and each nucleus has a larger zone of influence. The film thickness may vary from grain to grain (possibly


due to the dependence of work function on orientation 5x) and in some instances irregular growths such as whiskers and plates develop.

These have certainly been seen under more severe conditions. Thus austenitic chromium-nickel steels oxidized for 24 h at 500°C in 76 mm of oxygen produce these platelets, particularly in the presence of stress, strain and water vapour. 62 They are about 150 A. thick, 30,000/~. high and 60,000 A long and often exist above the main film or scale, forming only a small percentage of the total oxide volume. Similar growths have also been observed in the oxidation of pure iron-chromium alloys in a 90°~ argon:10~o water vapour atmosphere 8 and in oxygen ~ at high temperatures. Again, theories relating these growths to defects in the oxide or alloy have not been satisfactorily developed.

When chromium steels and austenitic chromium-nickel steels are heated in air at 300-700°C, a thin, nearly uniform film develops in the first few minutes and then grows slowly with a [111] fibre axis normal to the surface. 53. 54 The crystal size increases somewhat with time and the preferred orientation becomes more complete. There is some variation from grain to grain. Other similar work 55 shows that the film consists of a mat of fine crystals of 100-1600 A in size. The crystals have an irregular shape and increase in size with the temperature; the thickness of the boundary layer between crystals in the mosaic structure is about 50 A or within the resolving power of the microscopical technique used. It is suggested that there may be a correlation between the size of the oxide crystals and the protective quality of the oxide film. A rather similar mat of larger crystals is formed on pure iron-chromium alloys in air at 1000°C. 5

(c) The structure of thin films. An important contribution has shown that pure iron-chromium alloys containing over 13 ~o chromium are in equilibrium with Cr203 containing a little dissolved Fe203, at least at 1000°C and 1300°C. 56 Alloys of lower chromium content are in equilibrium with Fe Fec2_x ) CrxO~ (O ~ × ~2) . Isothermal phase diagrams have also been constructed and these are used in making further predictions39,56-~9 Oxygen pressure/composition diagrams are presented for the entire system. 29,56 The Cr2Os-Fe2Os system 6°-e4 and the spinels Fe Fet2-x) CrxO4

(O ~ × --<__2) 65-e7 have been further studied or considered by other workers. As is usual, however, thermodynamic considerations are of limited usefulness in systems which are often kinetically controlled.

Under normal high temperature conditions the initial stages of oxidation are extremely rapid until a protective layer has been built up. °'7 It seems clear, however, with most pure iron-chromium alloys, chromium steels and austenitic nickel-chromium steels that the thin protective films (growing according to OA in Fig. 1) are very rich in chromium. In the very first stages of film thickening the layer may actually become richer in chromium, is, 5s probably by the reduction of iron oxide, trapped in the initial film by chromium in the alloy at the alloy/scale interface. Diffraction data.O-15.18,20, 21,27.2.0.30, 42,53-55,68-75 and recent work with the electron-probe micro- analyser 75 appear to confirm that the films consist principally of Cr~Oa together with a little Fe2Os. It is not entirely clear whether these oxides are completely in solid solution or whether the composition of the film is uniform in any dimension. 57, 68 Indeed it has been suggested that stratification occurs at a very early stage. 6s,55,~7 Under some circumstances the spinel Fe Fec.,_x~ CrxO4 (O <= × _<__2) with x nearly 2 may be

180 G.W. WOOD

formed in thin films 13, ~2, 53 and this would still be more protective than the oxides of iron. ~a With extremely low chromium contents, where 'Cr2Oa' or 'Fe Fe(2_x) CrxO4' rich in chromium would not be expected even under conditions of extreme preferential oxidation the protective properties of oxides of iron modified by chromium addition need to be considered.

By analogy with the films obtained on passivation under low-temperature aqueous conditions, it appears that the layer is relatively protective because, although it is a reasonable electronic conductor, its ionic conductivity is low. s° The exact defect structure of CrzO3 which, as has been described, is the principal oxide present in most cases is still incompletely understood. It is probably a p-type semi-conductor r7 under these conditions and grows by cation movement, TM at least below 1300°C. Other workers believe that it operates as an intrinsic or amphoteric semi-conductor above 1300°C but that it behaves as an electron-defect semi-conductor below this temperature due to impurities. 7a, 80 Indeed, as the composition is so nearly stoichiometric, it seems likely that its defect structure depends on its method of formation and on impurities incorporated from the alloy or the atmosphere.

Indirect evidence supports the argument for cation movement at relevant temperatures. Thus, inert marker experiments with pure iron-chromium alloys containing 20 ~o and 30 ~o chromium in oxygen at 1000°C indicate growth by cation movement. 56 Although marker experiments are in some disrepute at the present time, the formation of voids at the alloy/oxide interface are in agreement with the marker positions. Electropolished iron-chromium alloys sometimes oxidize faster than abraded or etched specimensY This could be due to a higher percentage of Cr 6+ (or Cr 4+) ions in a Cr203 lattice developed under very severe oxidizing conditions. Further, work on cobalt-chromium alloy, 81, as and on nickel-chromium alloys 83, 84 indicates a maximum oxidation resistance at about 25 ~o chromium. There is some evidence of an optimum chromium content for the iron-chromium system) 6 This could be explained by the reduction of the number of cation vacancies in the Cr2Oa by the entry of Fe ~+ ions, although the contribution from the volatility of chromium and lack of adhesion at higher chromium contents must not be forgotten. There is some evidence of the volatilization of chromium in the oxidation of high-chromium steels, possibly due to the reduction of Cr20 3 by carbon.14 Oxidation rates of the alloys and of pure chromium in various environments are generally in accordance with the concept of Cr.,O3 being a cation-deficient oxide. (See Section 2.2. (k).)

When the protection is due to Fe Fe(2_x) CrxOa spinel, x should be as high as possible so that the more protective 'normal' spinel is present. Fe CrzO4 is a p-type semi- conductorT~,85 but Fe804 may grow partly by anion movement. It is also probably worth noting at this stage that FeO 86,a7 and NiO are p-type semi-conductors and Fe20386,87, SiO2 and AI20 3 are n-type semi-conductors.

(d) The morphology of break-through. When the conditions are borderline between giving 'protective' and 'non-protective' behaviour, the induction period is often followed by extremely rapid scaling (BC in Fig. 1). This process, sometimes referred to as 'break-through' or in certain cases 'break-away', occurs by the formation of stratified nodules or warts of oxide of varying shape and thickness at apparently randomly distributed sites on the surface. 7,1,, 15,18, 53, 54 The film thickness at which this occurs depends on a number of variables including, of course, the alloy and the


atmosphere. In fact, the length of the induction period and the size and distribution of the nodules vary considerably, even for specimens prepared under nominally identical conditions. Eventually the nodules grow or fresh ones are formed until the surface is completely covered by stratified scale. Growth of the individual crystals of oxide may also occur at this time. As scale growth continues, all the layers thicken, although the total scale may grow at a somewhat reduced rate (BC in Fig. 1). If little protection is afforded, however, or if the oxide flakes off, BE may be followed.

(e) The structure o f scales. Diffraction techniques sa, 54 and electron-probe microanalysis 76 prove that the inner layers of the nodules contain principally a mixture of chromium oxide and iron oxides and the outer regions almost pure oxides of iron. In the completely stratified scales various combinations of 'Cr2Oa', 'Fe Fe(2-x) CrxO4' rich in chromium, 'FeO', 'Fe8Oa' and 'Fe2Os' may appear, usually in this order from the alloy to the outer surface, depending on the conditions. ~,a,7-15'18-2°,29-al, 5e, 57, e7-7o. 88.89 This is, as expected, in order of increasing dissociation pressures as the outer surface is approached. The phases given above are put in quotation marks to indicate that they are not stoichiometric and not necessarily pure. For example, 'Fe2Oa' may contain some chromium and, in systems more complex than iron- ¢hronium, other impurities. 'FeO' tends to decompose on cooling to Fe and 'Fe304' and may also contain some chromium. 57,~9, 70, 90 Other specific phases found in complex alloys are dealt with in Section 2.2. (j).

Electron-probe microanalysis of scales produced on iron-14.4 70 chromium alloys in steam at 950-1000°C clearly reveal two major scale layers, the outer one consisting largely of 'FeO' and the inner one containing a rather complex distribution of oxides of iron and chromium. 91 Sometimes a marked chromium maximum and iron minimum is found near the inner scale/outer scale interface but oxidation for a longer period tends to destroy this relatively simple stratification. Indeed the apparent lack of homogeneity is a factor which has, perhaps, not been entirely appreciated in the past. This means that theories of alloy oxidation in which simplifying assumptions are made must be applied with great care. a2, 9a Further electron-probe studies of iron-14-4~o chromium specimens oxidized in air at 1000°C show that continued oxidation leads to dilution of the chromium maximum in the inner layer, although it sometimes persists in an attenuated form at the inner scale/outer scale interface. 7e Studies by other workers, however, appear to show increased chromium content near the alloy/oxide interface with different alloys. 57

(f) The underlying alloy. During oxidation the underlying alloy is depleted with respect to chromium and with respect to certain other alloying elements in the case of complex alloys. Although it is obviously difficult to get a precise measure at the alloy/scale interface, values within 1 micron of the interface have been obtained, gl For iron-14.4 70 chromium oxidized in steam for 10-30 min, there is a 5-micron wide layer showing depletion, the minimum recorded being 11.570 chromium. The actual minimum at the interface may be rather less. Somewhat similar results have been recorded by other workers using various alloys, a, 11, 31

The appreciable depletion is related to the large activation energy of interdiffusion of iron and chromium. 94 It appears that currently accepted theories on the vacancy mechanism of diffusion do not easily explain diffusion rates in the iron-chromium

182 G.C. WOOD

system. Ring diffusion or some other mechanism may be involved at some compositions.gS. ~s

In another important investigation, 97 the influence of adding 1 ~ of third element on the diffusion of chromium in iron was studied at 950-1000°C. Tin increases the rate, tungsten and nickel have little effect and titanium, silicon, niobium and beryllium reduce the rate by a half to a tenth. It is difficult to accept the conclusion 97 on this basis that tin is undesirable and titanium and silicon helpful in oxidation-resistant materials. Neglecting possible influences on oxide defect structure and adhesion, it might be anticipated that the most suitable alloy would be one where chromium is readily replenished at the interface. Another worker claims that for alloys of transition metals based on iron, the addition of < I ~ of a metal of ionic radius > or < that of iron increases or decreases respectively the rate of diffusion of the alloy and its rate of oxidation. 98 This change of rate is proportional to the difference in ionic radius and the concentration of the added metal.

(g) hTternal oxidation. Under certain conditions the alloys become internally oxidized. When pure iron-chromium alloys are oxidized in oxygen at 1000 ° and 1300°C for short periods and subsequently equilibrated in argon, there is a pronounced change in behaviour at 13 ~ chromium. 56 Above this figure, where 'Cr~Os' is formed in contact with the alloy, no internal oxidation is observed. For lower chromium contents (5~o and 10~ chromium), where Fe Fe(2-x) CrxO4 spinel is formed in contact with the alloy, pronounced internal oxidation occurs, apparently at a parabolic rate. The internal oxide is also spinel. Specimens of iron-14.4 ~o chromium, oxidized in air at 1000°C until break-through had occurred in some places, show another interesting result. 7 Pronounced internal oxidation occurs behind the nodules but not behind the thin protective film. As stratification continues over the whole surface, the internal oxidation becomes general and increases in depth. Evidently it is related to the chromium depletion and to the ability of oxygen to reach the alloy through the inner scale. These results appear consistent with the conclusion from another research that the higher the chromium content the higher the temperature before internal oxidation begins. 75

For pure iron-chromium alloys oxidized in a 9 0 ~ a rgon :10~ water vapour atmosphere at 700-1000cC to give scales, the number of particles increases but the depth of penetration decreases as the alloy chromium content is increased. 8 At 700°C the subscale is thin, the particles are small and are formed preferentially at grain boundaries. At 900°C the subscale is thicker and there is a greater variation in particle size and shape than at 700°C. There is still evidence of preferential formation at grain boundaries in some alloys. At 1100°C the particles are far more massive and irregular in shape but have a lower population even though the penetration is deeper. In an iron-I 0 ~ chromium alloy the particles appear in the microstructure as rods with a Widmanst~itten-type distribution. There is no evidence of preferential grain-boundary precipitation. The incorporation of the subscale into the main scale on continued oxidation may occur. 8

Other workers have also noted internal oxidation, 18. 31, 70, 75, 99 the particles being ,Cr~O3,18, 98 or 'Fe Fe(2_x) CrxO4' spinel. 7°, 98

(h) The mechanism of break-through. Evidently a question of vital importance is the reason for the break-through or transition from a thin 'protective' film to a thick,


stratified, 'non-protective' scale. The arguments presented here are for an initial 'Cr203' film but, as should become apparent, the same principles could hold for 'Fe Fe~2_x) CrxO4' rich in chromium. There are two main theories. Either the 'Cr2Os'-rich film is penetrated by the preferential diffusion of iron ions which form iron oxides at the outer surface or it is rendered 'non-protective' by some form of cracking. On the evidence at present available the author favours the former argument, although it may prove that both factors are active to varying extents under different conditions. The theories are given briefly below but this is obviously a case where reference to the original papers for fuller details is advisable.

The 'chemical' mechanism has been suggested by several authors14.15,1°° but a theory developed in some detail in one paper will be mentioned here. z9 After the 'Cr_~Os'-rich layer is formed it may be attacked from within by the chromium- depleted steel to give initially a Cr~O3-Fe2Oz solid solution. The oxygen potential of this phase increases on the dissolution of iron until it becomes thermodynamically more favourable for a spinel to be in contact with the steel. Thus a phase transformation occurs, resulting in Fe Fe~_x) Cr~O4, which contains a greater proportion of cation vacancies, appearing. This should be an autocatalytic effect and once the complete film, or a complete path through the film, is in this form, the transport of iron ions through the film to form an outer layer of iron oxides would be facilitated. Presumably the spinel can become yet richer in iron and in some circumstances transformation to 'FeO' containing some chromium may occur.

Cracking of the protective film could be brought about by means of several mechanisms. It has been suggested that it may be due to ferrite/austenite phase changes in the underlying alloy, 4 but this could only apply to certain ahoy compositions. Sigma phase formation in certain other alloys may need to be considered in some cases. TM Whilst discussing the various phases that may be developed it is worth noting that austenitic and ferritic steels of similar compositions oxidize at different rates, is, 21, 51 In addition, the abrupt change in initial oxidation rate of iron-chromium alloys at the Curie temperatures may be related to the ease with which electrons can leave the alloy to enter the film.51 Another explanation of cracking would be that stress develops 2,s,ls due to volume changes during oxide growth. These could be super- imposed on any internal stresses in the alloy, and inclusions in the alloy or recrystallization of the oxide 14,15, 53, 57 might play a part. In the case of steels containing silicon, break-through may be caused as a result of silica and voids accumulating at the alloy/oxide interface, la, 15. s; Even if some of these effects are not important in explaining the break-through of the initial film they may play a part in scale growth (see Section 2.2. (i)).

Neither theory has yet been satisfactorily worked out or demonstrated. In the case of the 'chemical' theory, it is perhaps difficult to visualize the exact mechanism whereby the iron ions pass through the film to form outer layers. The iron enters the film initially as Fe s+ but after the phase transformation much of it becomes FEZ+; it is questionable whether this reduction of the iron within the film is reasonable. Another point to be considered is that when the Fe2Os content of Cr2Os-Fe~O 3 solid solution reaches a certain value the solution might be expected to become oxygen- deficient. The irreproducibility of the induction period and the randomness of the nodules are not necessarily incompatible with the theory. The preferential iron

184 G . C . Wood

diffusion could perfectly well occur randomly over the surface, possibly promoted by local defects in the oxide such as dislocations and grain boundaries 5~ and by recrystallization.l~, 15, 5s, 57 The lack of homogeneity in many scale layers would tend to support this view. °1

The main, and considerable, objection to the cracking theory is that it does not explain the formation of virtually pure iron oxides at the outer surface. Any alloy revealed by cracking of the 'Cr203' layer would not be sufficiently chromium- depleted to allow iron oxides to form directly. However, it is important to note that on re-oxidation, in many cases 'Fe Fe(2_,) Cr~O4' spinel would be formed in the cracks rather than 'Cr203' and this would then promote the preferential diffusion of iron.

Evidently overheating of the specimen as oxidation proceeds is not vitally important, at least in the case of i ron -26~ chromium at 1170°C. 1°2

(i) The growth nlechanism of scales. When the chromium content of an iron- chromium alloy is very low, the alloy should scale in much the same way as pure iron except that the thickness, composition and mechanical properties of the 'FeO', 'FezO4' and 'Fe2Oa' layers will be affected by the chromium. The observation that in the initial stages iron-0.2 ~o chromium oxidizes more rapidly than iron in oxygen 2 could be explained by chromium present in the 'FeO' making it even more cation deficient. It has been pointed out, however, that the presence of spinels should decrease the composition range of 'FeO' or even eliminate it as a stable phase in the presence of the alloy in some instances3 °3 There is some evidence for this in the present system.14, 57.9o It seems then that the reduced rates of oxidation of low iron-chromium alloys (including perhaps iron-0.2 ~o chromium for longer times) as compared with iron are due to the presence of 'Fe Fe(2_~) CrxO4' spinel and to the decreased amounts of 'FeO' formed."

Somewhat similar behaviour will occur after break-through with alloys of higher chromium content. It is necessary to explain the continued thickening of all the layers. In some cases this may be done by considering normal diffusion mechanisms, proper regard being given to the dominant diffusing ion in each layer and to boundary conditions. Presumably chromium-rich layers usually control the overall oxidation rate and these gradually become diluted and iron-rich layers grow. However, in some instances there is considerable evidence that cracking and/or pore formation occurs during this stage. 2,a The inner layer becomes porous 8 and cracks and cavities are visible at the inner scale/outer scale interface. 8, :~x These may be formed due to some of the factors discussed in Section 2.2. (h) and are possibly accentuated by mechanical interaction between the various scale layers. It has tentatively been argued that the inner layer then grows principally by the inward passage of oxygen gas (or oxygen ions) and the outer layer by the outward diffusion of electrons and iron ions (or iron ions), provided by the dissociation (or 'ionization') of an oxide of iron at the inner scale/outer scale interface, a, ~1 Some Fe 2+ and Cr 3+ diffusion may still occur in the inner layer. Marker experiments support the above mechanism, 8 and it is interesting to note that a chromium maximum is observed just inside the inner layer at the inner scale/outer scale interface of oxides formed under somewhat similar conditions. ~1 This may represent the position of the original 'Cr203' film, now in a somewhat depleted form. 91 In other cases, the relatively simple stratification is destroyed and


'particles' rather than layers are observed. The similar oxidation rates of various iron-chromium alloys in 9 0 ~ argon: 10~o water vapour atmospheres may indicate that the outer layer, free from chromium, is rate-determining in this particular case. s

(j) The #~uence of alloying and 'tramp' elements. It should be evident from the previous discussion that there is considerable similarity between the oxidation of pure iron-chromium alloys and of various commercial steels. Nevertheless, the results are affected to a greater or lesser extent by the presence of other alloying and ' t ramp ' elements. Oxides of these elements, distributed in the bulk oxide in solid solution or as discrete layers or particles, may affect the defect structure and conductivity and hence the growth rate of the scale and its mechanical properties, in particular its adhesion to the alloy and its plasticity. There is also the possibility of the elements influencing the diffusion rates of iron and chromium in the alloy.

In most cases, the examination of the influence of trace elements has been sub- sidiary to other investigations and the results aresomewhat contradictory. Forexample, no systematic investigation on the effect of the entry of small amounts of relevant elements into the 'Cr~Os' defect structure has been conducted. Some observations worthy of note are described below but no systematic discussion is possible. Many of these elements are, of course, introduced to improve the mechanical properties of the alloy.

Nickel appears to increase the oxidation resistance in many cases and there are factors involved besides the possibility of the alloy being oxidized existing in the austenitic condition. It does not always enter the bulk scale in great quantities and is presumably retained largely in the alloy or in the inner scale at the alloy/scale interface. This effect has been observed with austenitic nickel-chromium steels 18,104 and even in a commercial iron-chromium alloy containing only a small proportion of nickel. 91 The enrichment may restrict the rate of diffusion of iron or chromium to the alloy/scale interface. In other instances, however, nickel enters thin films, 53.55 nodules13,15 or scale is, 15.18, 6s possibly in the form of a spinel. 29, e8 In the case where the film owes its resistance to 'Cr2Os' the presence of small proportions of Ni 2+ might be expected to cut down the oxidation rate.

Manganese has been detected in "protective" films ~. 68. lo5 and in the outer layers of scales, 9,14 often in the form of appreciable quantities of MnCr204. 57,n8 In other cases no segregation is apparent.91

Considerable importance has been attributed to the formation of SiO2 films under a variety of conditions. It appears, as mentioned earlier, that passive films produced under room-temperature, aqueous conditions may be rich in silicon. 43-4e The protective nature and subsequent breakdown of a 26.5 ~ chromium-0.44 ~o silicon steel, and possibly of certain other steels, during oxidation in air at high temperature have been attributed to the accumulation of SiO2 at the alloy/scale interface? 4,1s,31,sT, 71 Several breaks of the kind shown in Fig. 1 are often observed and correspond roughly to the number of SiO2 layers in the scale? 4 Other investigators claim that a layer of Fe2SiO4 is formed at the alloy/oxide interface when i ron-(5 .5-6 .2~ o) chromium- (1-4-3.3 %) silicon is oxidized in air at 800-1200°C. T M However, no SiO2 segregation has been observed under other conditions? 4, 7~, 91,100

Considerable attention has been given to iron-chromium-aluminium alloys. 17,24, 3x,100,107-1~5 A ternary diagram showing the concentration ranges in which A1203,

186 G.C. WOOD

Fe~O3 and Cr2Os respectively predominate in the scale has been constructed. 24 Solid solutions and spinels may also be formed under some circumstances. The upper layers of the alloy are, of course, impoverished in the preferentially oxidized elements and some internal oxidation may occur, which helps to key the oxide to the alloy. It is sometimes held that the best oxidation resistance is observed when the film is AlzO3- rich but superior behaviour under thermal cycling conditions is obtained with Cr203- rich oxide. 112 Small additions of thorium, 114 cerium u3 and calcium 11~ appear helpful, probably by improving the scale adhesion.

The keying of the scale to pure iron-25 Yo chromium and iron-37.5 ~o chromium alloys is also improved by adding small quantities of yttrium, lanthanum, dysprosium, gadolinium and erbium. 16 The external scale is primarily Cr208 but the internal oxide, often found at grain boundaries, is Y208 at 1000°C and YCrO2 at 1200°C. Small thorium additions are said to improve the properties of austenitic chromium-nickel steels. 112

Molybdenum is enriched in passive films formed at room temperatures. ~3-46 Small quantities of this element may be tolerated in stainless steels at high temperatures, but larger quantities can lead to catastrophic destruction of the oxide. 11s. 117 The situation appears to depend on the molybdenum and chromium contents and on the temperature. The location of the molybdenum within the scale may be important.x16,119 Catastrophic oxidation is discussed further in Section 2.2. (k).

Carbon is important where CO could be evolved and cause failure of the film. This situation is presumably related to the relative affinities of carbon and the alloying elements for oxygen. There is also the influence of carbide inclusions in the alloy on the protective qualities of the film to be considered.

(k) The influence of the atmosphere. It was indicated earlier that the mechanism of oxidation is probably essentially the same in oxygen, air, water vapour, carbon dioxide and certain other oxidizing media, namely the break-through of an initially protective 'Cr2Os'-rich layer followed by scale growth. However, the actual rate may show considerable variation and the purpose of this section is to suggest possible reasons for this behaviour.

Most studies have been conducted in dry oxygerd -4, 9,10, la, 1~, 30, 37, 51, 33-57, 68-71, lO2,130 or air that is substantially dry. 7, 9-15,13, 53, 57, 63, 73, 74,11e, 117,13o Somewhat similar behaviour is observed in the two environments but the growth rate in the initial stages and the eventual break-through is appreciably faster as the partial pressure of oxygen is increased. 9,13 Care should be taken in applying results obtained in oxygen to air 5e because the differences may not be entirely due to variation in oxygen partial pressure. Compounds such as Cr2N may be formed under some conditions e, 82 and nitrogen may be trapped in the 'Cr203' lattice. Two N 3- ions replacing three 0 2- ions would tend to make the 'CrzO3' more nearly stoichiometric, so cutting down the oxidation rate. A few results are available for ammonia or cracked ammonia atmospheres, m, 122 Aluminium and chromium nitrides are sometimes formed as subscale where the previously formed oxide layer is not sufficiently protective. 121

In almost all cases, water vapour in small or large quantities increases the scaling rate of pure iron-chromium alloys, 7,a,33 chromium steels 9,1°,1~,14, 33,12°,123-128 and austenitic nickel-chromium steels. 15, as, 31, es, 120,123-127,129,130 The exception is a commercial chromium steel containing 26 .5~ chromium and 0 .44~ o silicon oxidized in

The oxidation of iron--chromium alloys and stainless steels at high temperatures 187

moist air at 980°C and 1100°C. a4 Probably the main reasons for the rate differences are again connected with the influence of atmosphere on the 'CrsOs' defect structure. In a water vapour atmosphere, the film is effectively being produced at a lower oxygen potenta l and may also contain H+ and O H - ions. &dnoting the idea of a slightly cation-deficient 'Cr2Oa' lattice, the presence of O H - ions in the place of some 0 2- ions would effectively increase the number of cation vacancies and hence the growth rate, which is the common result, s, 15.131 In the case of the exception mentioned above, the decreased rate in the presence of water vapour could be explained by protons entering cation vacancies and hence reducing the Cr s+ diffusion, s, 14,15 Evidently the exact behaviour is not yet clear because presumably the possibility of growth by means of O H - ion migration and of cracking due to trapped hydrogen must also be considered. It is also significant that in water vapour oxidation, appreciably more 'FeO' is formed and 'Fe2Os' is generally absent. The absence or reduced quantities of the less ductile 'FeaO4' and 'FesOa' may make some contribution to the increased rates.13s

Interest in carbon dioxide oxidation has been revived in recent years. 1°, 7s, lso, lsa-~as It appears that protection is again due principally to a 'CrsOa'-rich layer. Oxidation may be faster than in oxygen in some cases 1° but small traces of water vapour in the carbon dioxide do not appear to have much influence. 1as

Essentially the same type of process occurs with certain combustion gas mixtures except in the presence of sulphur in some reducing atmospheres, s9 Environments containing sulphur-bearing gases such as SOs and H2S can produce various sulphides, sulphates and other compounds containing oxygen and sulphur. 2°, sg. 1~7-139 Sulphides are less protective than oxides and also non-protective low melting-point eutectics may be formed. High nickel contents are dangerous in certain non-oxidizing atmospheres but increased chromium and aluminium contents improve the resistance of some steels considerably. 14°-14s

Steels containing more than about 20 ~o chromium resist SOs quite well. 14°, 1~4-14e Investigations of oxidation in fuel gas, furnace and similar atmospheres containing small quantities of SOs have been conducted by several workers. 29,1~7-1s9 In the case of 18 ~ c h r o m i u m - 8 ~ nickel stainless steels at 1050°C, in the presence of 0-7-2.1 CO and in the absence of SO2, there is evidence of the dilution of the initial 'CrsOs'- rich layer by iron and some carburization but no break through occurs as in the presence of excess O~. 29 With excess O~, SO~ up to 0.6 ~ does not alter the general behaviour and no sulphide is formed. When CO is present in excess, however, SO2 considerably stimulates attack and a substantial sulphide subscale is formed below the oxide scale. With a large excess of CO, carburization of the steel occurs, the outer scale is thin and little subscale is formed. The influence of SO2 is therefore greatest at moderate CO contents. The relative proportions of the various oxides naturally depend on the atmosphere.

A detailed discussion of the catastrophic oxidation of stainless steels is beyond the scope of this paper but the most important features will be noted. For details, reference should be made to the original papers. 116-11s,laT,147-159 The essential feature is the destruction of the protective 'Cr2Os'-rich oxide by volatile oxides or molten oxides or eutectic mixtures, originating in some instances entirely f rom the alloy but more often as a result of substances in the fuel-oil ash. The oxides MoOs, V205, WOs, BisO3 and

188 G . C . WOOD

PbO are the most serious offenders but the behaviour is accentuated by the presence of Na2So4 and NaCI, probably due to the formation of low melting-point vanadates, or other compounds. At temperatures below possible melting points of the vanadates, vanadium or molybdenum may increase the number of cation vacancies in the oxide and hence increase the oxidation rate. But, as the temperature is raised, extremely catastrophic behaviour begins around likely melting points of eutectics, so it appears that liquid is necessary for the more severe attack. The reaction occurs substantially at the alloy/oxide interface, possibly accompanied by considerable overheating. The ability of the molten phase to absorb and transmit oxygen may be an important factor. Possible remedies include altering the steel composition, purifying or making additions to the fuel and applying protective coatings. The suitable properties of metallic and ceramic coatings cannot be discussed here.

3. FUTURE STUDIES 3.1. General

From the previous discussion it is evident that although this subject has received considerable attention, much remains to be done. It is equally apparent that to make further notable advances, studies in several fields, carefully integrated where possible, are essential. This would undoubtedly be a long and costly procedure but the benefits, not only in the theory of alloy oxidation but also in the industrial field where those alloys based on the iron-chromium, nickel-chromium and iron-nickel-chromium systems are so widely used, should be considerable. It is suggested that selected researches along the following lines would make the postulation of a more general theory a better proposition than at present.

Although the compatibility of commercial steels with industrial atmospheres must still be measured because of the urgency of practical problems, fundamental studies on simpler systems, as recently conducted by some workers, are essential. -0-4, 7.8,16, 3~- a7,~6.76.~1 Carefully controlled experiments with suitable specimens of binary and ternary alloys in simple atmospheres should provide the starting point but these would then need extending to cases where trace elements in alloy or gas appear to play a vital role. In atomic-energy applications it may be necessary to investigate the influence (if any) of such features as atomic particles and radicals.

Thermo-gravimetric techniques must form the basis of many investigations because of their usefulness in detecting rates of formation and discontinuities in growth. It is probable that the irreproducible results reported by many workers will still be detected even under the most stringently controlled conditions, certainly in cases where pores, cracks and blisters are likely to form. However, these are much easier to discuss in simple systems, particularly where the latest physical techniques are available to enable detailed examination of the scale.

3.2. Thermodynamic Considerations The extension of present knowledge of phase diagrams for systems such as

iron-chromium-oxygen, iron-nickel-oxygen and iron-nickel-sulphur should receive attention. It is still uncertain how useful thermodynamic data of this nature are in predicting kinetic behaviour, the problem being somewhat similar to the application of potential/pH diagrams in liquid environments. For example, although Cr203 and


Fe203 should be completely miscible under equilibrium conditions, these compounds do not always give homogeneous solid solutions even after prolonged oxidation at high temperature and scales are often irregular in composition over short distances. Another point which needs considering is to what extent the restriction of a phase area in the phase diagram by alloying is reflected in the proportion of this oxide in the scale.

3.3. The Influence of Surface Finish The question of surface finish should receive considerable attention because, as

described earlier, this appreciably influences the results. Data obtained from specimens prepared by abrasion, electropolishing, microtoming, etching and more refined techniques 37 are of great value because they give information about the mechanism of oxidation as well as the growth rate. It is possible to question whether any of these techniques gives results really characteristic of the true resistance of the 'bare alloy' to oxidation at temperature except under very severe conditions because of the initial film present. Presumably, from the practical viewpoint, the finish giving the slowest initial rate of oxidation, the longest and most reproducible induction period and the least disastrous behaviour after break through is the most desirable. Results for likely industrial finishes are obviously the most valuable in this respect.

3.4. Growth Morphology The nucleation and growth of the initial oxide layer is a problem which should

prove very attractive. By slowing down the rate of this process by means of suitably controlled atmospheres it is possible to observe the striations and nuclei on the metal surface.3~, 3e An investigation of the influence of the atmosphere and of alloy defect structure on the number and distribution of nuclei should be valuable. Perhaps experiments on annealed and cold-worked single crystal faces would be suitable. 5° Careful optical and electron microscopy is also useful in observing the thickening of the initial fine mat of crystals, the nucleation and growth of nodules at break-through and the subsequent development of stratified scales. The size, shape, distribution and regularity of the small crystallites can all be studied as can the development of needles and plates. The elucidation of the mechanism of growth of these last features, produced above films or scales, is important, particularly if they can be related to the defect structure of alloy or oxide, to the atmosphere, or to internal stress.

Optical microscopy of metallographically prepared cross-sections is a necessary technique in observing nodules, layer structures, precipitation of other phases on cooling, internal oxidation and the presence of cracks, cavities and other discontinuities. The distribution of these features across the scales represents a vital part of many studies.

3.5. Structure and Analysis The recently developed technique of electron-probe microanalysis should prove

a powerful tool in following the distribution of elements in the alloy and in the oxide at various stages of oxidation. 78, 91 When the probe is made to scan a square raster of a cross-section through scale and alloy, the qualitative distribution of the con- stituents over a field of interest between 500~ and 50~z square may be obtained. B

190 G . C . Wooo

This is of great importance in revealing the general distribution and any local segregation of elements in particles or layers. The technique of expanded contrast may prove useful in providing greater discrimination. By traversing along any line of interest on the specimen or by stationing the probe at a point, semi-quantitative or quantitative analyses (accuracy -I-, 1 ~o) are obtained. The present limiting thickness for a single valid quantitative reading is about 2,~ but this may be improved in time. Signs of enrichment in a particular constituent may be obtained with slightly thinner layers.

However, as protective films are often much thinner than this a different approach must be adopted in examining them. Instead of preparing cross-sections, filmed alloys are examined directly in plan without further surface treatment. Now, the depth of penetration of the electron beam is a function of the accelerating voltage, so by making measurements at different voltages the distribution of elements through the film may be obtained. Provided the beam can penetrate to the underlying alloy estimates of film thickness may be made. The quantitative theory of this last technique is not yet fully established but it appears that films down to at least 1000/k in thickness may be studied. The usefulness of this technique on a quantitative basis requires checking by comparing the results with those obtained by more conventional if less convenient methods of determining film thickness and composition. It is also possible that by comparing the shapes of curves of percentage alloying element/accelerating voltage for different stages of oxidation, the progress of the initial film formation may be tracked. In some instances taper sections or detached films might be used but for obvious reasons care should be taken in interpreting the results. To take full advantage of the electron-probe technique improved methods of preparing the surface, capable of producing smooth surfaces without relief polishing, and of avoiding surface contamination, are desirable.

The structure, orientation and grain size can be determined by the well-established techniques of electron and X-ray diffraction. However, on occasion these may fail, for example in being unable to differentiate between oxides of very similar lattice spacing such as Cr203, Fe2Os and their solid solutions. There is also a need for the development and application of micro-diffraction methods, capable either of providing results obtained by controlled sampling le°, ~61 or from suitable small areas accurately located on the specimen. If such experiments could be carried out in situ in a hot stage so much the better, le2-164 Depletions in the underlying alloy can also be determined by X-ray techniques 165,166 but electron-probe microanalysis is probably preferable. Radiotracer techniques may also be employed. 167 If, as has recently been suggested, techniques incorporating electron microscopy, electron-probe microanalysis and electron diffraction in a single instrument prove possible, they should prove extremely useful in oxidation studies. Spectrographic analysis should still prove helpful in examining scales for small quantities of trace elements. While on the subject of oxide identification it is worth noting that under polarized light 'CrzOa' is said to appear green whereas 'Fe Fer,_x) CrxQ' spinel is yellow or red. 56

With the techniques briefly described above, it is easy to see how vital data could be readily obtained. For example, the oxidation of pure iron-chromium alloys of different chromium contents could be followed from the thin-film stage to the thick- scale stage. Of particular interest would be the chromium depletion of the alloy, the distribution of elements in the scale including any chromium maxima and the


nature of internal oxidation. Experiments carried out on specimens alternately oxidized for very short periods and then equilibrated 58 would also be instructive because there would be a tendency to keep the equilibrium oxide 'CrzO3' (for alloys containing >13 70 chromium) in contact with undepleted alloy. Provided no side effects of importance occurred during the equilibration, a considerably prolonged induction period would be anticipated. The behaviour could be compared with that obtained during continuous preferential oxidation. This work can of course be extended to more complex systems and in an ad hoc manner to scales produced on commercial steels. Investigations along the above lines have been initiated by the author and are to be continued. 7. 7~9~

Detailed application of this type of study should allow the classification of the various layers and particles that appear as oxidation proceeds under controlled conditions. It is then necessary to investigate the defect structures and mechanical properties of the vital oxides involved. Production of these layers in a form suitable for further examination probably involves growing films oll chosen alloys in suitable controlled atmospheres. 5e

3.6. Defect Structures Appreciable information about the growth mechanism can probably be obtained

by standard inert marker measurements and by radiotracer techniques, care being given to interpretation. Nevertheless, further detailed studies of the conductivity and thermo-electric properties of 'Cr203' and 'Fe Fe~2_x)CrxO4' are essential to clarify their defect structures. As pointed out earlier, in the case of the more important 'Cr~Oz' the composition is nearly stoichiometric and its defect structure may depend on its method of growth. Aaly variation across the film is also of importance.

It therefore seems essential that oxides formed on the alloys and not prepared 'chemically' should be employed. Care must be taken to ensure that any changes due to ageing after oxidation and prior to measurement are kept to a minimum. Perhaps it is too much to hope that quantitative relations between ionic and electronic con- ductivities and the forming conditions can be established. Nevertheless, one of the main hopes of improving the protective value of the 'CrzO3'-rich films is by systematic alteration of the defect structure to suit the alloy and the atmosphere. 5~ At least semi-quantitative data might be accumulated for controlled additions of iron. nickel and other alloying elements and of H +, O H - and N 3- ions to the Cr,,Oz lattice. It would also be interesting to know whether any possible stress or strain in the oxide or its recrystallization alters the rate of ion transport. Defects such as grain boundaries and dislocations in the oxide film could control diffusion locally.

3.7. Mechanical Properties Any advances in the theory of the plastic deformation, creep and fracture of oxides

could prove a boon in the development of oxidation theory. The phase and volume changes accompanying the heating up of the specimen, the oxidation of the alloy and subsequent cooling must be accommodated by the relaxation of the oxide onto the alloy, by the formation of voids or blisters at the metal/oxide interface or in the oxide, by peeling or by cracking. The situation depends on the defect structure of the oxide as well as its mechanical properties and the specimen geometry. Stresses are

192 G.C. WOOD

more likely to be developed if the oxide grows by the inward movement of oxygen rather than by the outward movement of metal and may lead to spontaneous cracking. The outward movement of cations is more likely to lead to cavity formation. The effects of edges and of specimen geometry have been studied by some workers 169-m but require further examination. When a scale becomes detached from the alloy, the diffusion of ions is interfered with and the oxides are often converted to higher oxides. In some instances it appears that vacancies can migrate from one side of an oxidizing specimen to the other through the metal, certain voids growing at the expense of others. Often voids appear in the oxide on one side whereas on the other they are absent and alloy and oxide remain firmly attached, m, 17z

A few workers have considered or investigated the mechanical properties of oxides. Observations of cavity formation in the oxidation of iron and creep experiments on the oxides of iron have been carried out. 1~2 It is concluded that 'FeO' is far more plastic than 'Fe304' or 'FezO3' which may be significant in explaining variations in oxidation behaviour where these phases are formed in different proportions. The influence of alloying on the mechanical properties of oxides is also apparently little understood. However, it appears that the alloying of magnetite should make it more brittle in several ways if a recently proposed mechanism of strain-rate dependent ductility applies to spinels, as it appears to do. z,1°3, z32.173 The theory considers the restoration of order in a deformed lattice by the movement of ions and electrons and is thus affected by their respective diffusion rates. These isolated researches are obviously inadequate to form a general theory and further work is badly needed. Experiments could include the determination of the mechanical properties of detached scales or completely oxidized wires. Preliminary stress measurements could be made by oxidizing one side of thin metal sheets, possibly supported by mica, 174 but more sophisticated techniques are probably required. Conclusions drawn from the direct study of ceramic materials should be applied with caution.

The adherence of oxides to alloys is obviously very important and the subject has recently been reviewed. 175 This property has been related in some instances to the part played by trace elements, 14-z~,n3.zz4 and requires more systematic examination. In this connection, the keying-on of scale by the formation of grain-boundary or other forms of internal oxide should be mentioned. The internal oxidation of some alloys is also worth studying in its own right as well as for any information it may provide about scaling.

3.8. Transfer of matter through cavities In principle the formation of cavities, pores and internal cracks provides a barrier

to continued growth. 178 In certain systems, however, they appear to provide little hindrance, le8,177 Cavities formed instantaneously out of contact with the external atmosphere would begin as vacua and should be filled by gas provided by the dissociation of the oxides with greatest dissociation pressures forming the walls. Little is known of the transfer of matter across cavities by vapour phase transfer or surface diffusion.

3.9. Further practical considerations This paper is not directly concerned with the application of fundamental data to


p rac t i ca l s i tuat ions . H o w e v e r , p rac t i ca l exper ience can c o n t r i b u t e to f u n d a m e n t a l

knowledge . T h e r m a l cycl ing 123, xva a n d effects o f impac t , c r ev i ce s a n d edges , c o n c a v e

a n d c o n v e x surfaces a2s a n d s imi la r fea tures are l ikely to be exper ienced . F u r t h e r , the

inf luence o f h igh gas pressures a n d of cu r ious i m p u r i t y effects a re p r e s u m a b l y m o r e

f r equen t l y m e t in an indus t r ia l rig. I n t e r m i t t e n t life tests 75 a n d s t r a i n - o x i d a t i o n tests a7

m a y p r o v i d e su i tab le da t a in s o m e instances .

Acknowledgements--My thanks are due to the Central Electricity Generating Board for financial support from October 1959-September 1961 during which period this work was started. I am also indebted to Professor A. H. Cottrell and Dr. T. P. Hoar of the Department of Metallurgy at Cam- bridge for their encouragement and interest during this period.

R E F E R E N C E S

1. H. H. UHLtG, The Corrosion Handbook, J. Wiley, New York (1948). 2. D. LAI, R. J. BORG, M. J. BRABERS, J. D. MACKENZIE and C. E. BIRCHENALL, Metallurgy Rept.

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on Metals at Elevated Temperatures, 1954, p. 162.

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