Micromorphological Studies of theCorrosion of Gold AlloysA. J. FortyDepartment of Physics, University of Warwiek, Coventry, U.K.

Considerable insight into the detailed mechanisms by which metals arecorroded can be derived from direct microscopic observations. Thisarticle describes those micromorphological changes occurring near thesurface that can be observed by transmission electron microscopy aftergold alloys have been subjected to anodic dissolution in strong acids.These observations are used to discuss the important problem of corro-sion by selective dissolution.

Direct observation of the surface of a metal after ex-posure to a corrosive environment has frequentlybeen used as a basis for deductions concerning thestructural processes by which chemical reactionsinodify the surface and near-surface layers. Althoughconsiderable information can be obtained in this wayusing simple optical microscopy, as, for example, inthe case of the aqueous corrosion of gold-copperalloys studied by Graf (1), Bakish and Robertson (2)and Pickering (3), a much more detailed picture canbe obtained with the electron microscope, particular-ly when this is coupled with the powerful analyticaltechniques of selected area electron diffraction andX-ray microanalysis. The availability of very highresolution microscopes now makes it possible tofollow changes in the interaal microstructure, the sur-face morphology and the composition of metals andalloys on a scale approaching atomic dimensions.

The use of the transmission electron microscopenecessitates the study of very thin specimens. Greatcare must be taken to ensure that the micromor-phological changes arising from corrosion can bedistinguished from those produced during thepreparation of such thin films from bulk samples. Forthis reason, there may be doubts concerning, thevalidity of some of the observations by Pickeringand Swann (4) and by others who have studied thecorrosion of alloys prepared as thin foils by elec-tropolishing techniques. Ion-sputtering methods forthinning alloys from bulk specimens may also besuspect, because different sputtering rates for thevarious constituents can lead to compositionalchanges. These difficulties have been largely over-come recently by Durkin and Forty (5) who havedeveloped techniques for preparing thin films byvapour deposition of an alloy from its individualcomponents. It. will be shown later how this hascontributed to a very detailed understanding of the

corrosion micromorphology in the special case ofgold-silver alloys.

Selective DissolutionThe most widely studied and possibly the most im-

portant phenomenon involved in the aqueous cor-rosion of góld alloys is that of selective dissolution,whereby the less noble element is preferentially re-moved from the alloy, leaving a gold-rich residue (6).This is the basis of various practical methods for theparting gold from its alloys. It is also thought to be animportant step in the stress corrosion of gold alloyssince rupture of the gold-rich surface layer by an ap-plied stress can lead to the initiation of localized,deeper corrosion and subsequently of a stress corro-sion crack (1, 2, 3). As we shall discuss later, selectivedissolution might also be an important precursor ofother corrosion reactions, such as oxidation.

Fig. 1 Schematic representation on an atomic scaleof the surface of an alloy composed of dissolvable Aatoms and noble B atoms.

K is a kink site on a surface stepN is a non-kink site on a stepT is a terrace site

Gold Bull., 1981, 14, (1) 25

The fundamental question to be answered, as faras the understanding of selective dissolution is con-cerned, is why a gold-based alloy should continue todissolve in this way beyond the stage where the sur-face should be passivated by the gold residue. Suchpassivation might be expected to develop at a veryearly stage, as can be seen from a consideration of theatomic processes that might be occurring on the metalsurface during dissolution. These are depicted intheir simplest form in Figure 1, where we ignoremolecular adsorption, oxidation and complexing ef-fects associated with the electrolyte, and assume thatdissolution involves only ionization and solvation ofthe metal atoms. Dissolution is expected to occurpreferentially from kink sites (K) in the surface stepswhere the atoms are least firmly bound and, at suffi-ciently low potentials, the dissolution current will in-volve predominantly A atoms — the less noblespecies. As dissolution proceeds, however, this cur-rent will be diminished as more and more kink sitesbecome occupied by more noble B atoms. Thereafter,dissolution can proceed only by the removal of Aatoms from non-kink sites (N) on steps or from terracesites (T), which requires a greater activation energy oroverpotential. Eventually, the alloy becomes com-pletely passivated when all the surface sites are oc-cupied by B atoms only. For most alloys, and Cu 3Auin particular, this passivation stage should be reachedafter the removal of A atoms from only a few atomiclayers. This is clearly contrary to practical experience,from which we must infer that this simple model ofselective dissolution is incorrect.

The continuing dissolution of the alloy must meanthat some form of mass transport accompanies thecorrosion reaction in such a way as to expose more ofthe less noble, dissolvable atomic species at the sur-face. There are three possible mechanisms for this.One is an electrochemical transfer of gold from moreanodic to cathodic sites on the surface, so that freshalloy is continuously exposed to the environment inthe anodic regions. However, on thermodynamicgrounds, it is doubtful whether this process isgenerally applicable (7), particularly in the case ofgold alloys where a highly noble species is involved.Nevertheless, it has been suggested (8) that such aprocess might be involved in the formation of corro-sion tunnels. However, as we shall see later, there isan alternative explanation for this. Another possibili-ty is that the less noble component is continuouslyreplenished at the surface by volume diffusion fromthe underlying layers of alloy. The relatively high dif-fusion flux required to support the dissolution cur-rent can be accounted for by an inward flow of latticevacancies generated at the surface by selective dissolu-tion from individual atomic sites. There is con-siderable support for this model from measurements

of partial dissolution currents and changes of surfacepotential, particularly in the case of gold-topperalloys which Pickering and other workers (9, 10) havestudied extensively. Pickering (11) has also shown,using a reflection X-ray diffraction technique, thatchanges of lattice parameter detected in the near-surface layers of gold-topper alloys are consistentwith an accumulation of vacancies injected duringselective dissolution. An additional possible mech-anism by which fresh alloy may be continuously ex-posed to the corrosive medium is by surface diffusionof the residual gold atoms so that they form islands ofnoble metal on the alloy surface. This has been shownto be an important factor in the development of corro-sion micromorphology by the recent work on gold-silver alloys by Forty and Durkin (12). In a latersection of this article it will be suggested that bothvolume diffusion and surface diffusion are of impor-tance in the corrosion reaction. There is aninitial step in which vancancies diffuse to form adisordered zone near the surface, this process beingpossibly accompanied by an oxidation reaction. Thismetastable region then decomposes into gold-richislands which grow by surface diffusion. It will beshown that the growth and coalescence of theseislands lead to the formation of pits and eventuallycorrosion tunnels.

Corrosion TunnellingThe first observations of corrosion tunnels were

made by Nielsen (13) on stainless steels exposed tohot chloride solutions. By examining oxide replicas ofthe surface of such specimens in the transmissionelectron microscope, he was able to reveal the exis-tence of long tunnels of sub-microscopic dimensionswhich penetrated the metal along slip bands andother structural singularities. Nielsen interpretedthese features as localized corrosion tunnels formedby the dislocation lines. We shall see later that thislocalization of the attack along dislocation lines is nota necessary feature of tunnelling, which is a moregeneral form of corrosion associated with imperfec-tions in passivating surface films or with any otherpreferentially anodic site.

The occurrence of corrosion tunnels in gold alloyswas first reported by Pickering and Swann (4). Theseauthors used transmission electron microscopy toexamine the micromorphology of thin foils of variousgold-topper alloys directly after exposure to variouscorrosive environments. Since the foils were preparedby electropolishing of thicker specimens, great carewas taken to avoid introducing corrosion damageartificially prior to exposure. It was found that thedominant form of corrosion morphology consisted ofclosely spaced pits or tunnels, depending on the com-position of the alloy and the nature of the corroding

26 Gold Bull., 1981, 14, (1)

Fig. 2 A colony of corrosion tunnels in a gold-copper alloyfilni after corrosion in ferric chloride solution. Thistransmission electron micrograph was published byPiskering and Swann (4)

reagent. Figure 2 is a typical example of a colony oftunnels formed in the 25 gold/75 copper atomic percent alloy by immersion in aqueous ferric chloridesolution. It is particularly interesting that environ-ments such as ferric chloride solution, known topromote stress corrosion cracking, produce deepcorrosion tunnels, whereas other environments whichdo not induce such cracking, potassium cyanide solu-tion for example, give only shallow pits. It was foundthat when tunnelling occurred, the tunnel radiusvaried significantly with alloy composition, rangingfrom approximately 4 nm for gold/95 copper atomicper cent to 30 nm for the 40 gold/60 copper alloy.

Swann and various co-workers (14) have subse-quently extended these observations to other alloysystems, including stainless steels in hot magnesiumchloride solution and gold-nickel in ferric chloride.Once again, they demonstrated that corrosion tunnel-ling is commonly found under conditions known tolead to stress corrosion cracking in these alloys. Thishas led to a model for transgranular stress corrosion,in which the porous material formed by tunnellingcorrosion along active slip bands is ruptured by theapplied stress to initiate a propagating crack.

Swann (15) has examined the conditions underwhich a corrosion tunnel might form. He suggeststhat a small depression in the surface of an alloy issufficient to establish a stable tunnel if selectivedissolution occurs, so that the side-walls become fullyprotected by the accumulation of atoms of the morenoble species. It is implicit in this suggestion that thenoble component is deposited on the side-walls,probably by surface diffusion or by ionization andre-deposition. It can easily be shown that this leads toa simple geometrical criterion for tunnelling to occur:

R, = 2d(1-c)/c

where R, is the critical tunnel radius, d is thethickness of an atomic layer and c is the atomic frac-tion of noble metal atoms in the alloy. If the tunnelradius is smaller than R, the side-walls are not suffi-ciently protected by the deposited gold atoms and thetunnel will widen by anodic dissolution. At largerradii, sufficient gold atoms will be deposited to causethe tunnel to shrink back to the critical size.

Despite its attractiveness in providing a simplequalitative description of tunnelling, Swann's modelfails to account quantitatively for the sizes of tunnels

observed under the electron microscope, as can beseen from the comparison of measured and calculatedradii given in Table I. Swann (15) has argued that thislarge discrepancy is the result of a tendency for the re-deposited gold to aggregate into small particles on thetunnel walls rather than form a uniform protectivelayer, and he has produced electron microscope im-ages which suggest that such particles do exist. Theformation of particles means, of course, that the tun-nels must grow to radii greater than R, before theside-walls can be fully passivated by a continuousgold deposit. Recent observations by Forty andDurkin (12) on gold-silver alloys (described in thenext section) have shown that selective dissolution ofsilver leads to the growth of gold islands on the sur-face of the corroded alloy and that the tunnelling mor-phology develops as a result of the coalescence ofneighbouring islands. We shall show later how thisleads to an alternative model which gives calculatedvalues of R, which are in much better agreement withmicroscopic observations.

Table 1Observed and Calculated Values of the Radiiof Stable Corrosion Tunnels in Various Gold

Alloys. After (15) and (17)

Alloy, I Exp. R., I Calc. R„ (17), I Calc. R„ (15),

at. per cent nm nm nm

50 Au/50 Ag 28 3033 Au/67 Ag 26 255 Au/95 Ag 4 5

25 Au/75 Cu 10 15 140 Au/60 Cu 30 40 0.516 Au/84 Ni 4.5 5

Gold Bull., 1981, 14, (1) 27

GOLD VAPOUR 1fijIiJI ILVER VAPOUR

I^III,.nm

nm

Fig. 3 Method used by the author and his co-workersfor preparing thin single crystal films of gold-silveralloys for micro-corrosion experiments

Surface Island GrowthSwann's observations described in the preceding

section were made on thin foils prepared from bulkalloys by electropolishing. This technique producessurfaces which may be rough at the atomic scale andcontaminated by films of oxide or some other residuesof electrochemical reactions. These surfaces aretherefore not ideal for fine studies of the microscopicstructural changes that might occur during a corro-sion reaction. For this purpose, smooth, crystallogra-phically well-defined surfaces, free from an initialoxide layer are required. Surfaces of this high qualityhave been produced on gold-silver alloys by Durkinand Forty (5), using the technique of vapour deposi-tion. The procedure is illustrated schematically inFigure 3. First, a layer of silver, about 100 nm inthickness, is deposited at a rate of about 0.5 nm/s ontoa freshly cleaved mica surface held at 285°C. Thesilver deposits epitaxially on the mica, so that a cleansingle crystal surface in a (111) plane orientation isobtained. Next, an intermediate 30 nm thick layer ofpure gold is deposited epitaxially on the silver at therate of 0.1 nm/s, and this is followed immediately by afurther layer of 30 nm of alloy, formed by evaporatingfrom the silver and the gold sources simultaneously atappropriate rates. Thus, the 50 gold/50 silver atomicper cent alloy is formed if both silver and gold aredeposited at 0.1 nm/s; the 25 gold/75 silver per centalloy requires deposition rates of 0.1 nm/s for goldand 0.3 nm/s for silver. The pure gold intermediatelayer, together with its coating of alloy, is thendetached from the mica by floating onto a bath ofdilute nitric acid so that the initial layer of silver isslowly dissolved. The intermediate layer of pure goldfunctions as a buffer to prevent corrosion of the alloyduring this operation.

The resulting film is a single crystal of alloy on agold base with its surface in a well-defined (111)crystallographic orientation. The specimen can be ex-posed to a corrosive environment and then examineddirectly in the electron microscope. All the observa-tions reported in this and the following sections weremade on samples prepared in this way. We describehere our work on gold-silver alloys, but a similartechnique is currently also being used to examine thecorrosion morphology of gold-copper and gold-nickelalloys (16).

The gold-silver system in nitric acid is particularlyuseful as a model for micromorphological studies ofcorrosion by selective dissolution. Silver and goldalloy substitutionally as a single face-centred cubic(fcc) phase across the complete compositional range.The pure metals and all the alloys have almost iden-tical crystal lattice parameters, and this means thatthere is very little lattice strain, and consequently nosignificant microstructural effect, arising from latticeparameter changes during selective dissolution. Fur-thermore, because of the high solubility of silvernitrate, there should be no insoluble corrosion pro-duct. (However, as will be shown later, there is someevidence that gold oxide forms during the selectivedissolution of silver.) Thus, changes in micromor-phology observed after corrosion may be interpretedsimply on the basis of the various forms of masstransport by which the alloys respond to selectivedissolution. It is not possible to draw such unam-biguous conclusions from microscopic studies ofgold-copper alloys where lattice parameters changeextensively and several structural ordering processesmay occur when the composition varies.

The most striking conclusion from microscopicstudies of a wide range of gold-silver alloys preparedas described above and corroded in nitric acid ofvarious strengths, is that there is an extensive rear-rangement of the initially smooth surface leading tothe formation and growth of small, gold-rich islands(12). An example of this is to be seen in Figure 4 inelectron micrographs of a 50 gold/50 silver atomic percent alloy film recorded after exposure for two suc-cessive periods to 50 per cent aqueous nitric acid.Figure 4(a) shows a group of islands formed after aninitial exposure to the acid of 30 s and Figure 4(b)shows the same field of view after a further exposureof another 30 s. Considerable growth of the islands asa result of the more prolonged exposure is evidentfrom a comparison of these two micrographs. Even-tually the islands coalesce and form a continuous goldnetwork extending across the surface of the specimen.The corrosion is then reduced to localized activitywithin the channels and pits, formed as a result ofincomplete coalescence of the gold islands. Thesechannels and pits begin to deepen as their diameter

28 Gold Bull., 1981, 14, (1)

ji

s^.:.

tel ^` .fit í^ ^.l ^C :.J \} ,r wf

yJ ^.^_ ^^ ) É^ ^^ f ^T ^ ^ i r M ^

-H

^^

Fig. 4 Illustrating the formation and growth of Islands during the corrosion of a 50 gold/50 silver atomic per cent alloy filmin 50 per cent nitrie acid solution, (a) is a transmission electron micrograph of the specimen after immersion in acid for 30 sand (h) is the same field after immersion for a further 30 s. After (12)

continues to shrink until a form of tunnellingmorphology is reached. The transmission electronmicrograph in Figure 5 shows this last stage in thecorrosion of a 30 gold(70 silver atomic per cent alloyin concentrated nitric acid. The tunnelling mor-phology so produced is very similar to that found forgold-topper alloys corroded in ferric chloride solutionby Pickering and Swann, and illustrated in Figure 2.

The process of island growth leading to pit and tun-nel formation is accelerated if stronger acid is used orif the silver content of the alloy is increased. This canbe seen by comparing the examples of Figure 4 andFigure 5. In the Jatter case, the increased amount ofsilver and the use of concentrated nitric acid have ledto the final stage of tunnelling, while in the formeronly the coalescence stage of island growth has beenreached. Further growth of the islands and pit forma-tion can also be produced by annealing the corrodedspecimens at approximately 450°C as shown inFigure 6. This observation suggests that the corrosionreaction, which is almost certainly accompanied bythe selective dissolution of silver, leaves the alloy sur-face in a highly disordered state which then re-orders

Fig. 5 Tunnelling corrosion in a 30 gold/70 silver atomicper cent alloy film after exposure for 30 s to concentratednitric acid. After (12)

by thermally-activated surface diffusion of theresidual gold atoms into island structures. The extentof the disorder, and hence the rate of re-ordering, aredependent on the silver concentration and thestrength of the acid. It is thought that the corrosion-induced disordering and the diffusion-controlled re-ordering occur continuously, so that the islands havean increasing passivating effect as they spread acrossthe surface, and the corrosion reaction thereforebecomes confined to the channels between theislands, and ultimately is concentrated at the pits re-maining after coalescence. The expected sequence of

Gold Bull., 1981, 14, (1) 29

w, .

tf.

Fig. 6 Enhanced growth of gold islands on the surface of asample of corroded 50 gold/50 silver atomie per cent alloyafter annealing for 30 s at 450°C. After (12)

GOLD-RICH L GOLD-RICH%/ ISLAND

—f1 ISLAND

^^^^^^^^^^^\^^^\\\\\\\\\\\\\\\\\\\\V\

,o memIVf//4

1

YiiAJ7 am I

TiiiA

Fig. 7 The corrosion-disordering/diffusion-reordering model of corrosion by selective dissolu-tion. After (17)

events is represented schematically in Figure 7. It isinteresting that the island-channel-pit morphology isfound only for alloys containing between 20 and 50atomic per cent gold. This can be accounted for onthe basis of our model because, in the silver-richalloys, island nucleation is difficult due to therelatively small concentration of residual gold atoms,whereas, in the gold-rich alloys, the residual atoms,being more densely packed, migrate across the surface,less easily. However, as we shall show in the last twosections of this article, the occurrence of the islandmorphology might be dependent on an intermediatestage of oxidation of gold during selective dissolutionof silver, and this might form only for the alloys con-taining 20 to 50 atomic per cent gold.

This simple model of corrosion, based on selectivedissolution and surface diffusion, has been examinedmore quantitatively by Forty and Rowlands (17). Itcan be seen from Figure 7 that gold accumulates inever increasing thicknesses as the islands spreadacross the surface. This, together with the restrictionof the occurrence of further corrosion to the channelsbetween islands, means that these channels shrink at acontinuously decreasing rate. As the islands coalesce,the channels break up into isolated pits which thencontinue to deepen at an accelerated rate while theyshrink inwards more and more slowly. The results ofa detailed analysis of the kinetics of pit shrinkagebased on this model are given in Figure 8. We see thatfor alloys having a low gold concentration the pitsshrink at progressively lower rates and never com-pletely fill in — that is, the surface is never complete-ly passivated. For gold-rich alloys (over 50 atomic percent gold), pits are filled in after a finite time, whichdecreases as the gold content of the alloys increases.There is a critical alloy composition with 50 per centgold below which the pits shrink at an exponentiallydecreasing rate. Thus, we find from this simplemodel that there is a 'parting limit' such that onlythose alloys containing less than 50 atomic per centgold will continue to undergo selective dissolution ifsurface diffusion is the controlling mass transportprocess. This is in reasonable agreement with the par-ting limit of between 30 and 50 atomic per cent goldobserved in practice (18).

A New Model for Corrosion TunnellingThe calculations of Forty and Rowlands (17) also

show very clearly how the quasi-stable pits formed byisland coalescence on dilute gold alloys (of less than50 atomic per cent gold) lead eventually to a tunnel-ling morphology. Figure 9 shows, for example, thecalculated profile for a pit in a 33 gold/67 silveratomic per cent alloy formed by the corrosion-disordering/diffusion-reordering process. As the pitshrinks, it reaches a critical radius when it begins to

30 Gold Bull., 1981, 14, (1)

c=0.10

6 8 10TIME, ARBITRARY UNITS

c=0.25

c=0.33

c. 0.40

c 0.50L

12 14 16

Fig. 8 Shrinkage of corro-sion pits on a range ofgold/silver alloys cal-culated on the basis of thecorrosion-disordering/diffusion-reorderingmodel. c is the atomicfraction of gold in thealloys. After (17)

1.0

0.8

0o.s

W^

0.

z0.4 ,

LL c=0.65% c=0.60%

0.2 ``; %`

00 2

deepen faster than its radius contracts and eventuallyit becomes a deep tunnel the radius of which thenchanges only very slowly with depth. Since the depthof such a pit is essentially a linear function of cor-rosion time, the profile shown in Figure 9 can also beinterpreted as describing the development of a pitwith time. .

The transmission electron micrograph shown inFigure 10 gives an 'example where the transition fromthe island-channel-pit morphology to a tunnellingmorphology is taking place for the 33 gold/67 silveratomic per cent alloy corroded in 30 per cent aqueousnitric acid. Where isolated pits have been formed, thepenetration of the alloy has occurred at an acceleratedrate, as can be seen from the enhanced transparencyof these regions.

The critical radius, R, at which the rate of penetra-tion of a pit begins to exceed the rate of shrinkage ofits diameter, corresponds to the stage at which theamount of gold released by selective dissolution ofsilver from a monolayer of alloy at the bottom of thepit is only just sufficient to give monolayer coverageof the side-walls after diffusion. This leads to thegeometrical condition:

Rc = 2N,d/c

where N, is the pit depth, measured in number ofalloy monolayers of thickness d, and c is the atomicfraction of gold in the alloy. N, is dependent on thecomposition of the alloy and on the initial distribu-tion of the gold islands, and thus R, is a complexfunction of c.

Values of R, calculated on the basis of this newmodel are given in Table I where they can be com-

pared with the observed radii of tunnels for variousgold alloys together with the critical radii calculatedfrom Swann's model. Thus, we see that the island-channel-pit model, evolved directly from microscopicobservations of corroded surfaces, is not only more

RADIUS OF PIT, nm

E

a

0rw0

Fig. 9 Expected profile of a corrosion pit on a33 gold/67 silver atomic per cent alloy calculated onthe basis of the corrosion-disordering/diffusion-reordering model

Gold Bull., 1981, 14, (1) 31

realistic than Swann's model, but is also more suc-cessful in accounting quantitatively for the measuredradii of corrosion tunnels. Although island growthhas been observed for gold-silver alloys and, undercertain circumstances, has also been found on gold-copper alloys, it has not so far been found on gold-nickel alloys. However, the close correspondence ofcalculated and observed tunnel radii found for all thegold alloys which are listed in Table I suggests thatisland growth might be a common occurrence for allsuch alloys.

A corrosion tunnel can be expected to penetrate thealloy indefinitely. Branching into secondary tunnelsmight occur at structural inhomogeneities and conse-quently the final morphological state will consist of alabyrinth or interconnected porous structure. Tun-nels can occur preferentially at grain boundaries, slipbands and other microstructural features such as twinboundaries. Thus, corrosion tunnelling leads even-tually to the disintegration of the solid alloy. This canoccur even in the highly perfect single crystal filmsused for the study of island growth, because thesefilms appear to be damaged by the corrosion reaction.The electron micrograph in Figure 11 shows, for ex-ample, the progressive disintegration of a singlecrystal of the 25 gold/75 silver atomic per cent alloyafter a prolonged immersion in dilute nitric acid. Theinitial corrosion by selective dissolution producesislands on the surface and at the same time an exten-sive network of narrow twins is developed within thealloy. The islands grow preferentially along the sur-face traces of these twins, so that narrow corrosionchannels form along these directions. Once the chan-nels reach the critical width for tunnelling, which isclose to the critical radius for a cylindrical tunnel,they penetrate the alloy very rapidly, so that it

Ka I

crumbles into small particles, as seen in Figure 1 l(a).The electron diffraction pattern for the alloy in thisstate is given in Figure 11(b), which shows how theresidual granules have become separated andmisorientated into a polycrystalline powder. It is pro-bable that further selective dissolution of silver fromthese particles leads to the so-called gold-rich 'sludge'commonly found after alloy parting (6).

The Micro-Mechanisms of CorrosionDisordering

From their electron microscope observations ofcorroded gold-silver alloys, Forty and Durkin (12)have deduced that the selective dissolution of silverleaves the surface -in a highly disordered state whichsubsequently recovers by surface diffusion to formgold-rich islands. Very recent work by the author(unpublished) has shown that corrosion disorderingoccurs in a near-surface zone, approximately 5 nmthick, by a complex sequence of processes involvingselective dissolution, oxidation of the residual goldand finally the decomposition of oxide back into gold.

The evidence for this comes from a detailedanalysis of the electron diffraction patterns and highresolution electron micrographs obtained from cor-roded alloys. Figure 12 shows the electron diffractionpatterns obtained for a 25 gold/75 silver atomic percent alloy before and after corrosion in dilute nitricacid. We see that the corrosion reaction has produceda new phase which gives rise to the extra diffractionspots apparent in Figure 12(b). This new phase canbe revealed by dark field electron microscopy usingthe extra spots to form the image that is shown inFigure 13(b). Figure 13(a) is the normai bright fieldimage of this area. The dark field image features astrongly diffracting phase, corresponding to the extraspots in the diffraction pattern, which covers the sur-face of the corroded alloy in well-defined domains.From dark field images of this kind, using all theextra diffraction spots in turn, we deduce that thereare three distinct sets of such domains correspondingto the three-fold symmetry in the diffraction pattern.The crystallographic character of the domains is ap-parent in Figure 14 which is a high resolution imageformed by the interference of beams doubly diffractedby the uncorroded underlying alloy and the corrosionphase. The moiré pattern thus produced, consists of

.rT).

1t f^.

Fig. 10 Transmission electron micrograph of a 33 gold/67 silver atomic per cent alloy film corroded in dilutenitric acid, showing the transition from the island-channelmorphology (I) to a deep pit morphology (P)

32 Gold Bull., 1981, 14, (1)

fringes spaced 1.8 nm apart and changes orientationby approximately 120 0 from one domgin to another.The relative weakness of the extra diffraction spotsand the low visibility of the interference fringes in-dicate that the corrosion phase is thin compared withthe uncorroded alloy and probably extends not morethan 5 nm below the corroded surface.

The structure of the corrosion phase cannot beidentified unambiguously from the electron diffrac-tion patterns, but it is thought to be an oxide of gold

formed during the selective dissolution of silver fromthe alloy. Thus, we can expect the oxide — probablyAu203 — to form within the near-surface disorderedzone during corrosion of silver-rich alloys by strongacids. Gold(III) oxide (Au2O3) has an orthorhombicunit cell (19) with dimensions such that an epitaxialfit can be found with the fcc substrate of uncorrodedalloy in three distinct crystallographic orientations,which is consistent with the tri-domain structure thatis observed.

Fig. 12 Electron diffraction patterns of a 25 gold/75 silver atomie per cent alloy film, (a) before corrosion and (b) after cor-rosion in dilute nitric acid. The additional diffraction spots are due to the formation of a new phase

Gold Bull., 1981, 14, (1) 33

1

1

: ;t

F Z f

J ^. - Z °rify,• = C` .

.t f

)' b i1 }

r' .. f• t t,, T ff ..V t ,^ .tt r

Y.p

} 4 td li ak

fi`.J ."• ._ (mei l t }

.4 5 ,^ "s. +^rffk ,?t"

r!LFig. 13 Surface structure of a 25 gold/75 silver atomic per cent alloy film after corrosion in dilute nitric acid. (a) is atransmission electron micrograph using bright field illumination, showing island structures and (b) is the same areaobserved in dark field illumination using `extra' spots in the diffraction pattern of Figure 12(b), thereby revealing surfacedomains of the `corrosion phase'

The oxide formed under these corrosion conditionsshould be metastable, forming during the rapid disso-lution of silver at the corrosion front, but tending todecompose when the silver dissolution is complete.Thus, a dynamic balance of the following processes:

Selective Dissolution-*Gold Oxidation-*Decomposition

might be established, with either the oxidation and/orthe decomposition step predominating, depending onalloy composition and strength of acid. The decom-position of the oxide might be expected to occurpreferentially at the oxide-acid interface, thereby

yielding free gold atoms which can migrate across thesurface, possibly assisted by partial ionization andsolvation in the acid environment, to form the islandmorphology described earlier. In other words, islandgrowth might be only a secondary result of the moreimportant primary step of oxidation. This idea is sup-ported by the observation that the rapid island growthfound when the alloy is heated to 450°C after cor-rosion is accompanied by the simultaneous disap-pearance of the extra spots in the diffraction pattern.It seems possible, therefore, that island growth arisesonly as a result of oxidation followed by decomposi-tion, as represented schematically in Figure 15.

Fig. 14 High resolution transmission elec-tron micrograph of a corroded 25 gold/75silver atomic per cent alloy film, showingmoiré-type interference fringes within thesurface domains of the `corrosion phase'

34 Gold Bull. 1981, 14, (1)

Fig. 15 Schematic representation of the corrosion ofgold-silver alloys by selective dissolution accom-panied by oxidation followed by decomposition ofthe oxide and island growth. Island growth even-tually leads to pit formation and to tunnellingcorrosion

Generalized DiscussionIn this concluding section, we attempt to present a

unifying review of the observations detailed in theforegoing, by means of a brief generalized discussionof the structural aspects of alloy corrosion by selectivedissolution. The essential steps in the overall corro-sion reaction are:(1) Selective dissolution of the less noble species in

the alloy(2) Formation of surface vacancies and their inward

migration(3) Volume and surface diffusion(4) Oxide formation(5) Oxide decomposition.Of these, some will be dominant and others may beabsent, depending on the specific alloy and acidenvironment.

As we discussed at length at the beginning of thisarticle, selective dissolution of the less noble compo-nent of an alloy should lead to the creation of surfacevacancies. These can migrate across the surface toform pits, steps and other surface rougheningfeatures, or they can assist the migration of theresidual noble metal atoms which leads to islandgrowth; they can also diffuse into the underlying alloyto assist volume diffusion of the less noble metal tothe corroding surface. According to Harrison andWagner (20), a combination of surface and volumediffusion leads to a geometrical instability of the sur-face during selective dissolution. This may be impor-tant in establishing corrosion tunnels if a mechanismlike that proposed by Swann is involved. On the otherhand, if island growth resulting from surface diffu-sion is dominant, then the model for tunnelling pro-posed by Forty and Rowlands appears more realistic.Tunnelling is almost certainly always involved in thefinal stages of corrosive disintegration of an alloy.

The possibility of oxidation of the residual noblemetal as a result of selective dissolution of the lessnoble metal has not previously been considered forgold alloys, although it has long been known that theselective dissolution of zinc from alpha-brass inammoniacal solutions of suitable pH can be accom-panied by a tarnishing reaction involving the oxida-tion of topper (21). It should be emphasized that theinterpretation presented here of the electron diffrac-tion patterns for the corrosion phase formed on gold-silver alloys by selective dissolution as evidente foroxidation is only tentative at this time. However, ifoxidation of gold does occur under conditions ofselective dissolution, it is likely to be a highly specificphenomenon, depending on the particular alloyingcomponent, rather than a general behaviour for allgold alloys. This probably accounts for the fact thatthe island morphology has so far been found as aregular feature only on gold-silver alloys. Clearly, thekinetics of the overall corrosion reaction leading even-tually to dissolution by tunnelling must be dependenton whether or not the intermediate oxidation stateoccurs and, in turn, on whether this leads to islandgrowth and, subsequently, tunnel formation.

AcknowledgementsFigures 2, 4, 5, 6, 7 and 8 have been published previously, as in-

dicated in their captions. The writer is grateful to the authors andpublishers of these articles for permission to include thern here.

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