Analysis of corrosion and oxidation on metals by spectroscopic means

Lectures Vortr ige

Analysis of corrosion and oxidation on metals by spectroscopic means

N. S. MeIntyre 1 D. Johnston l, W. J. Chauvin 1, , , F. E. Doern 1,**, T. C. Chan 1, E. McAlpine 2, D. Lister 2, and V. S. Shastri 3

1 Surface Science Western, Faculty of Science, University of Western Ontario, London, Ontario N6A 5B7, Canada 2 Systems Materials Branch, Chalk River Nuclear Laboratories, Chalk River, Ontario KOE iJO, Canada 3 Processes and Materials Research Laboratory, CANMET, Canada Department of Energy Mines and Resources,

550 Booth Street, Ottawa, Ontario K1R 7K7, Canada

Untersuehung der Korrosion und Oxidation an Metallen mit Hilfe spektroskopischer Methoden

Zusammenfassung. Diese Arbeit betont drei Gesichtspunkte oberflfiehenanalytischer Untersuchung, die ffir die Charak- terisierung von Korrosionsmechanismen als besonders wesentlich erachtet werden: die chemische Bindung, die mikroskopische Analyse und die verbesserte Steuerung experimenteller Bedingungen. Die R6ntgen-Photoelek- tronenspektroskopie (XPS oder ESCA) ist ffir die Unterscheidung chemischer Spezies besonders erfolgreich gewesen. Die Anwendung von XPS beschr/inkt sich jedoch im wesentlichen auf die/iuBere Oberfl/iche einer Schicht. Die Aufnahme von Tiefenprofilen mit Hilfe von Ionenbeschul3 vermindert hfiufig die chemische Information, die vom Schichtinnern und vom Metall-Oxid-Interface erh/iltlich w/ire; und hier spielen sich oft die Korrosionsvorg/inge ab. Mechanische Methoden des Diinnschichtabtrags sind somit aufgefordert, XPS bis an die Leistungsgrenze zu bringen. Mikroskopische Untersuchungen von Korrosionsschichten mittels Raster-Auger-Mikroskopie (SAM) oder Sekundfir- ionen-Massenspektrometrie k6nnen Effekte wie die Korngrenzensegregation und die Oxidbildung aufdecken. Ein weiterer bedeutender Gesichtspunkt der Analyse von Korrosionsschichten ist die Behandlung der Metallober- fl/ichen selbst. Die oxidierte Oberfl/iche kann durch Behandlungsfehler wie z.B. Kaltbearbeitung oder Ver- unreinigung stark beeinflugt werden. An Hand von drei verschiedenen Beispielen werden diese Aspekte illustriert: (i) die Analyse einer Molybd/in-Schutzschicht auf Stahl; (ii) die Untersuchung von Korrosionsprodukten auf 304-Edelstahl, und (iii) die Effekte des H +-Ionenbeschusses zum Ausheilen mechanischer Sch/iden auf der Oberfl/iche von Ni-Cr- Legierungen, vor der Oxidation.

Summary. This paper addresses three aspects of the surface analytical experiment which are felt to be particularly crucial in characterizing corrosion mechanisms: chemical specia- tion, microscopic analysis and improved control of experimental conditions. X-ray Photoelectron Spectroscopy (XPS or ESCA) has been particularly effective in differ- entiating chemical species. The major chemical use of XPS

* Present address: Tracor Northern Inc., Rexdale, Ontario, Canada

** Present address: Whiteshe, ll Nuclear Research Establishment, Pinawa, Manitoba, ROE 1LO, Canada

Offprint requests to." N. S. McIntyre

is limited however to the outer surface of the film. Depth profiling by ion bombardment frequently degrades the chemical information available in the film interior and at the metal-oxide interface, and this is often where corrosion mechanisms are established. Mechanical methods of film penetration are thus encouraged to use XPS to its maximum capability. Microscopic analysis of corrosion surfaces using Scanning Auger Microscopy (SAM) or Secondary Ion Mass Spectrometry can reveal effects such as grain boundary seg- regation and oxide nucleation. Another important aspect of corrosion film analysis is the preparation of the metal surface itself. The oxidized surface can be strongly influenced by preparation artefacts such as cold work or impurities. Three different studies are used to illustrate these aspects: (i) the analysis of a molybdenum inhibitor layer on a steel surface; (ii) the study of corrosion product release from 304-stainless steel; and (iii) the effects of H + ion bombardment to anneal mechanical work on Ni-Cr alloy surfaces, prior to oxidation.

I. Introduction

The characterization of corrosion films on a metal presents a special challenge for the surface analyst. While the major products of a corrosion reaction on a surface are readily detectable by an ever-increasing array of sensitive analytical techniques, the interpretation of the complex mechanisms underlying most processes is more intractable. The range of surface chemical, physical and mechanical forces related to corrosion and oxidation is indeed broad. In order to make sense of much of the surface analytical data from a corrosion test, it is necessary to impose strict controls on some of the many variables which govern corrosion. While this require- ment holds true of any respectable corrosion test, it is doubly so for those where surface analytical techniques play a part; the atomic layer sensitivity of many of these methods is all the more prey to extraneous information from an ill- controlled experiment.

The capabilities of surface techniques for corrosion film analysis have, in the past few years, been remarkably improved, both by the rapid advances in instrumentation technology, and by the efforts of analytical spectroscopists to develop better qualitative und quantitative routines for interpretation of the data.

This paper will attempt to outline the present capabilities of analytical spectroscopy for characterizing corrosion processes on surface. Firstly, a brief review of the present state

255 Fresenius Z Anal Chem (1986) 324:625-634 Springer-Verlag ] 986

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Table 1. Surface analytical techniques for corrosion studies - present capabilities

Technique Elemental information Chemical information Information depth

Z range best detection instant by pro- limit filing

Spatial resolution

1. Electron-emission (a) XPS (ESCA) a 3-92 0.05% Extensive (oxidation state 2 nm 0.5 gm 150 gm

and structure) (b) AES 3 -92 0.3% Considerable (structure) 2 nm 2 p.m 0.1 pm

2. Particle emission (a) SIMS 1-92 0.0001% Scanty 1 nm 10 gm 1 gm (b) ISS 3-92 0.1% Undeveloped 0.1 nm 2 gm 10 ~tm (c) RBS 2-92 b 0.1%" None 5 nm 2 ~tm 100 pm

3. X-Ray emission (a) SEM-EDX 9-92 0.1% None ~ 2 gm 1 gm

4. UV-IR emission (a) IR Reflectance None - Extensive (molecular vibrations) Variable 100 gm (b) Raman None - Extensive (molecular vibrations) Variable 1 gm

a XPS X:ray photoelectron spectroscopy, AES Auger-electron spectroscopy, S I M S Secondary ion mass spectroscopy, ISS ton scattering spectrometry, RBS Rutherford scattering spectrometry, SEM-EDX Scanning Electron Microscopy-Energy Dispersive X-ray Spectroscopy

b Elemental sensitivity varies strongly with matrix

of the analytical techniques will be given, followed by a discussion of three examples of corrosion studies from the authors' laboratory. These examples illustrate most of the techniques discussed in the review. Moreover, they serve to illuminate the large amount of work left to be done in achieving a thorough understanding of any of the processes.

II. Analytical spectroscopic techniques for corrosion studies

For the purposes of this discussion we will consider that the thickness of the corrosion or oxidation film does not exceed 2 ~tm and is usually much less. Analytical techniques considered are those capable of identifying either specific chemical elements and their chemical interactions or molecular vibrations on surfaces. The capabilities of eight fairly common techniques are summarized in Table 1.

It is worthwhile examining the importance to corrosion studies of the various characteristics of each technique. Many of the newer techniques are sensitive to almost all chemical elements (see Table 1); however, of these, XPS measurements yield detection limits less affected by atomic number and a quantitative response less affected by matrix. Thus, its value as a survey technique cannot be under- estimated. The detection limits achievable with SIMS for some elements under some conditions are certainly in the < 1 ~g/g, and this can be important for detection of elements such as chlorine or sulphur which initiate corrosion even in low concentrations.

Chemical information on corrosion films has to date largely been provided by X-ray photoelectron spectroscopy (XPS or ESCA). In most situations, where oxide mixtures are present, it is essential that chemical interactions be associated with a specific element and this militates against non-specific techniques such as IR reflectance spectroscopy. Moreover, by this time, a solid record of XPS binding energy data has been compiled [1, 2] which is becoming increasingly useful for identification of unknown phases. In recent years, XPS spectral resolution has been substantially improved with the use of the X-ray monochromators and once this

modification is in general use, a general upgrading of spectral data should result. Auger spectra of many elements also contain chemical information and these have also been catalogued in the form of the Auger parameter [1, 2].

There ist no doubt that spatial resolution is of prime importance to any corrosion-related study. Analysis by XPS is limited to areas of ~ 150 gm diameter or larger, because of technological limitations in the X-ray gun. Corrosion phenomena usually propogate in crevices, or grain boundaries, or are affected by the phase composition represented within a particular metallic grain. Consequently, a surface technique with sub-micron spatial resolution is as necessary as chemical characterization. Scanning Auger Microscopy, which uses a sub-micron electron beam as a probe, is thus considered as a corrosion study technique equal in importance to XPS. Recently, it has become clear that SIMS is also an important technique for microscopic surface analysis of some trace elements.

Analysis of concentration profiles as a function of depth into the corrosion layer can reveal gradients which sometimes lead to explanation of film growth mechanisms. The use of ion bombardment techniques, familiar to all surface analysts, sometimes is not the optimum approach to profiling corrosion films. Most corrosion films contain some chemically unstable species, particularly near the outer surface (hydroxides, etc.), which are susceptible to decom- position under most ion bombardment conditions. Thus XPS-related studies of chemical composition must be done with caution. Moreover, corrosion films of any appreciable thickness (>0.1 gin) do not, in general, exhibit very sharp profile gradients because of the variety of solution and solid state transport processes which determine the film growth and the usual roughness of the film. Instead, it is often more effective to separate layers of a film mechanically, either by using a mechanical cratering device or be removing layers selectively with adhesive tape, ultrasound or chemical stripping. The use of several different experimental approaches to film removal often consumes less time than a single ion bombardment profile of a moderately thick film, and reveals more about its composition. Ion bombardment

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depth profiles are still useful for characterizing thin, even films, such as those formed by short-term gas phase oxidation.

The use of these approaches to surface characterization are discussed below in the context of several different applications involving both aqueous corrosion and gas phase oxidation processes.

Ill . Characterization of a corrosion inhibitor surface film

The rate of aqueous corrosion of carbon steel has, under some conditions, been found to be reduced by the addition of molybdate or tungstate species to the solution. In particular, steel pipelines carrying coal slurries have been shown to be passivated under working conditions by such agents. The mechanisms of inhibition are being investigated using XPS and AES using conditions somewhat modified from those used in the original tests.

In contrast to the original tests which used 'as received' pipe surfaces with no other treatment, the present experiments involved surfaces finished to a high polish (0.5 pm A1203 finish) to emphasize grain structure, and analyzed before reaction to ensure that no impurities were present.

The experiments described were intended to determine the effects of solution pH and reaction time on the formation of the molybdenum-containing inhibitions at pH 3.0, 5.0 and 7.2 for periods of 1, 3 and 5 h. During reaction, the surfaces were potentiostatically controlled at a voltage con- sistent with the passivation potential obtained in a cy- clic voltammogram. In the case of pH 3 this potential was - 3 0 0 mV (vs. SCE). In all other pH regimes the potential used was - 100 mV (vs. SCE). The chemical structure of the inhibitor layer was analyzed by XPS and, in Fig. I a - d some detailed changes in the molybdenum chemical composition are found in the Mo(3d) spectra. The ratio of major surface elements, as determined by XPS is also given for each specimen illustrated.

The Mo(3d) spectra shown in Fig. I were obtained under high spectra resolution and clearly discernible changes in peak shape and position are found between surfaces treated at different pH regimes and for different periods. These shapes were mathematically fitted using fixed doublet Gaussian peaks of line width found typical of high oxidation Mo(3d) peaks under the instrumental conditions used.

The conditions for the blue-colored surface film formed at low pH (Fig. i a) are of particular interest, since inhibition by molybdate ion is found to be least effective in this regime. In Fig. l a there are three major molybdenum species identified, with centroids at 230.8 (peak 1), 231.7 (peak 2) and 232.7 (peak 3). The major peak at 231.7 clearly cannot be associated with MoO3 (232.4 ___ 0.1 eV) the highest oxidation state oxide of a molybdenum oxide normally expected. We tentatively ascribe the observed oxide to a mixed oxide/ hydroxide 'blue' oxide species likely to contain both + 5 and + 6 oxidation states [3]. The oxide formed initially at pH 5 (Fig. 1 b) has a distinctly different shape, and this is in- terpreted as an increased amount of MoO3 or some other Mo + 6 oxide species. At a slightly later stage in this same pH regime the amount of Mo § 6 oxide has increased further and iron as an hydroxide identified over the general surface. Finally, at pH 7 the oxide surface has, even during the earliest growth period, a substantial concentration of iron as iron hydroxide as determined from the Fe(2p) peak shape

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[4]. All molybdenum appears bonded in one Mo + 6 oxidation state (232.5 eV). Thus, the molybdenum oxide forming an inhibiting film at low pH is chemically different than that at higher pH regimes.

The surface of the inhibiting oxide at pH 3 was also analyzed by Scanning Auger Microscopy to obtain an analysis of the surface structure under high spatial resolution. Figure 2a shows a scanning secondary electron micrograph of the oxide surface formed at pH 3 after 1 h. The oxide exhibits a cracked surface with the cracks, in general, following the pattern of the underlying grain boundaries. An Auger spectrum of the face of the oxide (Fig. 2 b) shows only molybdenum and oxygen, while the spectrum taken at a boundary intercept (Fig. 2c) shows the presence of considerable iron. Thus, even under relatively passivating conditions, some regions of the grain boundary are not inhibited from corrosion by the mixture molybdenum oxides. It is expected that it is these areas which actively

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Fig. 2. Measurement of microscopic surface concentrations on the pH 3 inhibition surface shown in Fig. la. A scanning secondary electron micrograph (a) shows the surface fissures which are coincident with grain boundaries, b and e show Auger spectra taken on the grain face and at a boundary intersection, respectively. Auger spectra were taken with a Perkin-Elmer Physical Electronics Model 600 SAM using an electron beam current of 10 nA and a voltage of 5 Kv

corrode at higher pH regimes and cause the observed mixing of iron and molybdenum species all across the surface oxide. The use of well formed and characterized surfaces has been important in identifying microscopic anomalies in the inhibition treatment of the steel surface.

IV. Studies of corrosion product release from stainless steel surfaces under boiler conditions

In water-cooled nuclear reactors, radioisotopes of cobalt are primarily responsible for the radiation fields that cause occupational hazards. The source of those isotopes is the corrosion of the materials of construction of the reactor

Fig. 3. Scanning secondary electron micrographs of loop-tested stainless steel - 304 surfaces, a Specimen 7 - exposed for 4 h; b Specimen 9 - exposed for 12 h. The linear trails of corrosion product debris seen in this micrograph appear to follow scratches in the surface

coolant circuits and the release of corrosion products to the reactor core where cobalt-59 is transmuted to cobalt-60.

The study described in this section makes use of a variety of surface analytical techniques to better understand the mechanisms of corrosion and of corrosion products released from polished stainless steel surfaces. The specimens studied were polished to a I pm finish and exposed to coolant (280~ 10MPa) in a circulation loop with sufficient turbulent flow, that corrosion products released never return to the surface. The coolant pH (10.4) was maintained with LiOH and the dissolved oxygen was < 4 gg/kg using a hydrogen overpressure. After certain periods of exposure, coupons were removed and analyzed by techniques such as SEM, Auger Spectroscopy and imaging SIMS.

Surfaces removed from the loop after 4 and 12 h of reaction are illustrated by SEM images in Fig, 3. After 12 h, visible corrosion product particles are detected, mostly in

628 258

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Fig. 4. Auger depth profiles obtained for several 304-stainless steel surface loop tested for different time periods. For profiling, a 1 gA, 4 keV Ar + ion beam was rastered over a 2 x 2 mm area. The sputter rate was determined from a calibrated thickness of iron oxide. A 3 kV, 10 nA electron beam was used to excite Auger spectra

the sub-micron range. After 500 h exposure an even heavier deposit is found with a greater percentage of particles in the 1 - 2 gm range.

Such films can be depth profiled using ion bombardment and surface techniques like Auger Spectroscopy. Figure 4 shows some oxygen depth profiles which given an impression of oxide thickness on the surfaces of several specimens exposed to cooland for different times. While a gradual increase of oxide thickness with exposure is measured, it is product not to attach too much quantitative significance to the thickness measurements. For example, the 4 and 12 h test specimens produce identical profiles despite different surface morphologies. Reproducibility in the removal rate of the oxide is governed by the roughness of the metal surface and, more importantly, by the size and distribution of the oxide particles, many of which are considerably thicker than the depths profiled in Fig. 4. The oxide contribution from such particles will be diminished by the redeposited metal sputtered from these parts of the surface clear of oxide particles. In a sense, then, these profiles represent an average depth of oxide, but even this is only valid for specimens with the same roughness and particle size distribution.

Stainless steel surfaces corroding under the relatively low oxygen activity in this loop might be expected, from previous experience, [5] to contain enriched concentrations of chromium in the surface layers. Auger depth profiling of any specimen could not detect a relative increase in chromium near the metal surface. However, selected area analysis of certain portions of the unsputtered oxide was more successful. Figure 5 shows Auger point spectra (0.1 gA diameter region) of a 500 h loop-tested surface. The surface crystals are determined to be iron-nickel oxides with no detectable chromium. However, the featureless oxide layer underlying these crystals was found to have atom ratio O :Fe :Cr :N i of 60:20:12:9. Clearly, the Cr :Fe ratio is higher than that in the unoxidized metal of this alloy. Thus, there is some evidence for accumulation of chromium relatively near the metal interface.

Secondary Ion Mass Spectrometry has been able to provide supplementary information on these surface films by

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virtue of its high detection sensitivity combined with an ability to produce distributional information. The SIMS instrumentation used to determine spatial distributions of elements in a corrosion film is the Cameca IMS-3f. This instrument, termed an ion microscope, allows images of the spatial distribution of each corrosion element to be obtained rapidly as the surface film is penetrated by the primary ion beam. The sensitivity is sufficiently high that even minor elements in the alloy such as cobalt, can be measured or imaged as a function of depth into the film.

A good example of the sensitivity of imaging SIMS is seen in Fig. 6. The 150 ~tm field images shown represent the spatial distributions of lithium, nickel and cobalt on the outer surface ( 0 - 5 0 n m ) of a 12h loop test specimen. Lithium, from the coolant, accumulates in surface scratches, while the nickel distribution shows no such relationship to the surface topography. Cobalt, however, shows definite enrichment into localized patches usually < 5 pm in diame-

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Fig. 6. Secondary ion mass spectrometric (SIMS) images of lithium, nickel and cobalt contributions on a 304 stainless steel surface loop exposed for 12 h (see text for loop conditions). Ion micrographs were taken with a Cameca IMS-3fion microscope using an O3 8 keV ion beam of 100 nA current rastered over a 250 x 250 gm area. A 150 ~tm area was imaged and photographed using exposures of 1 - 10 s. The darkened area seen in the Ni-58 image is an artefact due to regional insensitivity in the image intensifier, a Li-7; b Ni-58; c Co-59

Fig. 7 SIMS images of a 304 stainless steel surface loop exposed for 48 h. Series I: Primary beam penetration range: 2 0 - 40 nm. The Fe-56 image (not shown), a Li-7; b Cr-52; c Ni-58; d Co-59

ter. Thus, there is a definite cobal t preferential enrichment in the surface oxide, compared to nickel at early stages of the reaction. Surface oxides, so enriched, if released into the coolant would cause activity field higher than those calculated based on bulk cobal t concentrat ion and assuming similar behavior of cobal t and nickel.

Figures 7 - 9 shows ion images taken at three different depths in a corrosion film on a 500 h loop test surface. The images in Fig. 7 represent depths of 2 0 - 4 0 rim. Li thium distr ibut ion differs from all other elements indicating an independent phase. Chromium, from the bright spots, is

associated, to some extent with nickel distr ibution, but the light background in this chromium image also suggest that much of it is not associated with either nickel or cobalt . These lat ter two elements are now rather closely associated, in contras t to the more independent behaviour of cobal t following shorter exposures. The cobal t and nickel appear to be preferentially dis tr ibuted along scratch marks in the surface p robab ly as a result of an accelerated growth mecha- nism spawned by chemical condit ions within the crevice. In Fig. 8, the dis tr ibut ion at a depth of 8 0 - 1 2 0 nm is shown. The l i thium-containing phase has decreased in intensity.

630 260

Fig. 8 SIMS images of same loop test specimen as in Fig. 7. Series II: Primary beam penetration: 80-120 nm. The Fe-56 image, a Li-7; b Cr-52; e Ni-58; d Fe-56

Fig. 9. SIMS images of same loop test specimen as in Figs. 7 and 8. Series III: Primary beam penetration: > 500 nm. No Li-7 image was detected at this stage, a Cr-52; b Fe-56; c Ni-58

Some of the nickel and iron present is agglomerated into particles while a more general distribution is also observed.

As in Fig. 7, chromium distribution is more general than any of the other corrosion elements. In Fig. 9 (depth > 500 nm), the major contribution from nickel comes from the bulk alloy. This can be ascertained by the strong secondary ion channeling pattern caused by metal alloy grains of different crystallographic orientation. The iron image also shows some channeling effects but to a much lesser extent, thus suggesting that some finely-distributed iron oxides are present. Again, the chromium image is the most diffuse,

suggesting an oxide distribution in close proximity to the metal interface. Thus, the sequence of ion images o f the corrosion film has revealed the rather rapid migration of cobalt to the surface in localized distributions at early stages in the oxidation, and its incorporation in nickel-containing oxide phases at later stages. Chromium is identified as hav- ing a very diffuse distribution even near the metal interface - evidence that it is present as an independent oxide phase or reacted with iron oxide. In all the above analyses, no attempt has been made to quantify changes in image intensity and to correlate the corrected intensity data for

261 631

several different elements (three dimensional depth profiling). Such corrections would be time-consuming, but along with digitized image acquisition, they would represent a major improvement in our knowledge of corrosion processes on surfaces containing many irregularities.

As shown above, one of the problems of the corrosion analyst is to recognize the movement or reaction of elements which are usually present in massive quantities in the underlying metal only a few nanometers distant or as corrosion products in the solution. Another means for following such reactions is to "spike" the metal with an enriched or radioactive isotope and to follow its progress with appropriate mass spectrometric or radiocounting pro- cedures. In the study of 304-stainless steel corrosion product release much information was gained by following the movement of radioisotopes implanted in the specimen surface prior to loop exposure. Activity released from the surface was monitored elsewhere in the loop [6] while that re- distributed within the film was monitored by sequential ion bombardment and activity counting experiments. For such experiments all of the 25 mm 2 coupon surface was etched with a 4 keV, 8 mA Ar + beam, except for a small area kept covered for serving as a reference level for profilimetry in determining sputter rates. Four radionuclides Co-60, Co-58, Cr-51 and Fe-59 were implanted to a depth of 15 nm into the specimen surface and activity depth profiles were run on loop-tested and "as received" surfaces using gamma spectrometry to monitor each gamma photo peak simulta- neously during the profile. The sequential measurements determine the quantity of radioactivity removed during each sputter. The differential activity plots shown in Fig. 10, give a measure of the distribution of chromium and iron radiotracers before loop-testing and after a 48 h exposure to the loop. Following corrosion, the chromium, iron and cobalt (not shown) are detected in a bimoda ! distribution; the portion nearest the surface is assumed to be within the oxide film. The surface of the loop-tested specimen was then chemically stripped with an oxalate/citrate solution [6] resulting in the loss of most of the large surface crystals, such as those seen in Fig. 3 b. All surface radionuclides ex- cepting Cr-51 were reduced by 3 - 1 0 times, while Cr-51 activity was virtually unchanged. This suggests strongly that chromium [u the underlying metal migrates to the surface during loop exposure, but while at the surface does not form mixed oxides with iron from that alloy. The chromium oxide layer so formed is resistent to chemical etching, while oxides of other metals are removed relatively effectively. Activity measurements in the loop [6] support this premise. Cobalt and iron are released to the loop, but very little cobalt is detected.

Thus, the corrosion behaviour of stainless steel under reactor conditions has been deduced using a variety of surface techniques. Dependence on only one or two of these would have resulted in considerably less certainty of the mechanisms and particularly of the role of chromium in achieving some measure of surface protection.

V. Initial oxidation of Ni-Cr alloy surfaces

Superalloys such as Alloy 600 are frequently used in high temperature gas phase environments where high strength and low oxidation rate are required. A study of the initial oxidation of this alloy (7% Fe, 16% Cr, balance Ni) and an

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18% Cr binary nickel alloy has been carried out under low pressure high temperature conditions (Po2 < 1 x 10 -4 tort, T = 600 ~ C) [7]. Of particular interest was the effect of alloy surface preparation on the form and composition of the oxide produced.

The previous study on 304 stainless steel corrosion product release showed that the surface oxidation was profoundly affected by surface scratches and polishing blemishes, presumably as a result of mechanical work imparted to the outermost atomic layers of metal. In order to separate such effects from those associated with the chemistry and crystallographic orientation of the alloy it is necessary to anneal stresses from the surfaces prior to beginning the oxidation experiments.

Annealing of Ni-18% Cr alloy surfaces has been carried out using two alternate methods: (i) heating to reaction temperature in hydrogen and (ii) hydrogen ion bombardment of the surface prior to oxidation. Hydrogen ion bombardment of a 1 c m 2 central area of a polished 4 cm x 4 cm alloy surface was done using a 4 keV beam to a dose of I x 1014 at/cm 2. After oxidation at 600~ for 3 miu at 1 x 10 -5 tort O2 bombarded and non-bombarded areas were compared using optical microscopy. Fig. 11 compares optical micrographs of the central H + ion bombarded area and an area just outside the ion bombardment region. The oxide in the ion bombarded area is very even and shows a

632 262

Fig. 11 Effects of surface pre-treatment on oxide growth. Optical micrographs ( x 100) of Ni- 18% Cr alloy surfaces are shown after oxidation for 3 min at 600~ under 1 x t0 5 torr oxygen, a Region pre-treated by H § ion bombardment at 4 keV to 1 x 1014 at/cm2; b region ~ 1 mm outside of pre-treated area

profound dependency on the orientation of the underlying grain (as determined by channeling contrast). By contrast, the region outside the A + area is more influenced by polishing marks, and any grain orientation effects become negligible as one moves further from the central H + ion bombarded region.

Figure 12 shows, using SEM images, better detail of the oxides produced in the central area. Three main oxide patterns (labelled VI, VII and VIII) are identified, each with some degree of recognizable patterned growth. Depth profiles of specific microscopic areas on two different grains (VI and VII) (see Fig. 13) shows a profound difference in surface composition with grain orientation. One grain surface is strongly chromium oxide while a second orientation is nickel-rich. Such findings may provide more repro- ducible conditions for studying metal surface oxidation.

The Auger depth profiles shown in Fig. 13 contain more analytical information than is usual for layers formed by longer, aqueous processes. The interfaces shown here are sharp and, in the case of the enriched chromium surface

oxide, a chromium depletion region can be detected in the metal immediately below the oxide interface. Thus, clearly there are certain types of corrosion oxide films that benefit strongly from traditional ion bombardment depth profiling.

This study has shown the almost overriding surface effects of mechanical polishing and the desirability of examining, in parallel, the effects of corrosion on an annealed surface.

VI. Conclusions

1. There is no one spectroscopic technique which is appropriate for most corrosion requirements. Probably AES has provided to date the greatest single benefit to the science. Frequently, several complementary approaches are needed.

2. Corrosion and oxidation films are profoundly influenced by surface preparation. Thus the surface analyst and the corrosion engineer must design experiments together to drive much benefit.

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Lectures

u~ hi 7-

Z hi F- Z

n n- Ld L9 < r~ Ld F- U DJ

O {)

a GRAIN aFr

50

50

I NICKEL

.ROMIU

I CARBON b GRAIN VlT- dark area

T NICKEL

"CHROMIUM

~ ' ~ J ' - - - C~ARBON

I - ~ OXYGEN I 0 1 2 3

SPUTTER TIME (min)

Fig. 13. Auger depth profiles of grain faces VI and VII as shown in Fig. 12

Acknowledgement. The financial support of the Electric Power Re- search Institute for part of this work is gratefully acknowledged.

Fig. 12. Scanning electron micrographs at different magnifications of the oxide structure on grain faces bombarded with hydrogen ion before reaction

R e f e r e n c e s

1. Wagner CD, Riggs WM, Davis LE, Moulder JF, Muilenberg GE (1978) Handbook of X-ray photoelectron spectroscopy. Perkin- Elmer Corporation, Eden Prairie, MN

2. Wagner CD (1983) Auger and photoelectron energies and the Auger parameter: a data set. In: Briggs D, Seah MP (eds)Practi- cal surface analysis by Auger and X-ray photoelectron spectroscopy. John Wiley and Sons, Chichester

3. Cotton FA, Wilkinson G (1974) Advanced inorganic chemistry. Wiley, New York, p 938

4. McIntyre NS, Zetaruk DG (1977) Anal Chem 49:1321 5. McIntyre NS (1983) Uses of Auger electron and photoelectron

spectroscopies in corrosion science. In: Briggs D, Seah MP (eds) Practical surface analysis by Auger and photoelectron spectroscopy. John Wiley and Sons, Chichester, p 417

6. Lister DH, McAlpine E, Mclntyre NS (1983) Proceedings of the BNES conference on reactor water chemistry (in press)

7. McIntyre NS, Kover L, Chan TC. J Vacuum Sci Technol (accepted)

Received November 25, 1985

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Analysis of corrosion and oxidation on metals by spectroscopic means

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