Tarnish Resistance, Corrosion and StressCorrosion Cracking of Gold Alloys

WS Rapson

Tarnishing, corrosion and stress corrosion cracking (SeC) of gold alloys are related phenomena. Theypresent problems when they occur in gold jewellery, dental gold alloys and electronic devices. They areexploited, however, in the depletion gilding and finishing of gold jewellery; and in the extraction, refiningand fire assaying of gold. There is still much that is not known about these phenomena, but a coherentpicture of their mechanisms is emerging as a result of studies of not only gold alloys themselves, but also ofother alloys. Aspects of this development are discussed.

Tammann's extensive studies (1) established that theresistance of gold to tarnish and corrosion is not greatlyreduced by the addition to it of silver and base metals,so long as the gold content of the resulting alloy is notbelow 50 at.%. This corresponds to about 15.6 caratfor Au-Ag alloys and to about 18 carat for Au-Cualloys.

Tammann's observations are of limitedsignificance, however, when considering the resistanceto tarnish and corrosion in the three areas in which thisproperty is of particular significance. These areas arefirstly, the tarnishing, corrosion and stress corrosioncracking (SCC) of gold jewellery; secondly, thetarnishing and corrosion of dental gold alloys; andthirdly, the tarnishing of gold alloy contacts inelectronic circuitry. The circumstances under whichtarnishing and corrosion occur in these areas differwidely. They may occur under humid but notnecessarilywet conditions in the case of gold jewelleryand gold alloy contacts. The products of corrosionunder such conditions tend to accumulate on thesurface of the alloy as tarnish films.

In the case of dental alloys, however, only thosecorrosion products which do not go into solution inthe saliva and which adhere to the alloy surface resultin formation of tarnish films. Where the gold or goldalloy is applied as a coating, as in some decorative andmany contact applications, substrate and other effectsmust also be taken into account.

Tammann's rule also suffers from the limitation thatit takes no account of the variations in themicrostructures of gold alloys which arise from thephase relationships which are applicable to them and

<$9' Gold Bulletin 1996, 29(2)

their metallurgical histories. Thus alloys which arehomogeneous solid solutions and which do notundergo separation into two immiscible solid phases arein general more tarnish and corrosion resistant thanthose which are mixtures of two phases. In the lattercase, as exemplified by some nickel white gold alloys,the one phase may contain too little gold to make itcorrosion resistant. There is also the possibility ofgalvanic effects between the separated phases where thecompositions of these differ significantly; and theformation of ordered Au-Cu phases may also lowercorrosion resistance.

Grain size is a significant factor. Where it is verysmall the grain boundary energy can become asignificant driving force for corrosion, supplementingthe forces arising from electrochemical differencesbetween the alloying elements and from galvaniccoupling between segregated phases.

A further important factor in the case of somealloys is the presence of high internal stresses in thealloy created by working and which have not beenremoved by stress-relieving annealing or in agehardening. These can make some alloys susceptible tostress corrosion cracking. The fabricator of gold alloyproducts can therefore contribute significantly tomaking his products more tarnish and corrosionresistant. In the case of lower caratage products,however, he cannot do so completely.

Published information in this field is extensive inrespect of both dental gold alloys and the gold alloysused in electrical contacts. In respect of carat goldjewellery, however, it is still very scanty indeed and isembodied in a limited number of publications (e.g.

61

2-14). Some of these relate to Au-Ag-Cu or Au-Ag­Cu-based dental alloys, but have direct relevance to thetarnishing and corrosion of gold jewellery alloys.

The deliberate superficial corrosion of gold alloysin electropolishing, in the mass finishing of goldjewellery and as a first step in the depletion gilding orsurface enrichment of lower caratage alloys, as well asthe smudge phenomenon and the interactions whichoccur between the skin and gold jewellery are notconsidered here.

TARNISH RESISTANCE

Assessment oftarnishing and rates oftarnishingUntil relatively recently assessment was done subjectivelyand therefore only qualitatively. It can be assessedquantitatively, however, as a colour change byspectrophotometry in terms of the colour coordinates oftest specimens before, during and after tarnishing (4,8,14). The use of this technique promises torevolutionize studies in this field, since it permitsassessments of tarnishing rates.

Assessment ofresistance to tarnishingThis involves the question of the environment fortesting. This has proved difficult in the testing ofdental gold alloys and of gold alloy contacts, and muchof the available information in these fields relates to thedevelopment of test environments which simulateactual exposure conditions as closely as possible. So far,no generally acceptable procedures for testing goldjewellery have been developed and the difficultiesinvolved make it unlikely that they will be establishedin the near future. Most investigators have thereforehad to be content to employ the most widely usedmethod for testing of dental gold alloys, namely theTuccillo-Nielsen test (2-4). In this, specimens areexposed alternatively in a rotating container to acorrosive liquid and the saturated moist environmentabove it. The corrosive medium is normally anartificial saliva as used for testing of dental alloys, butin other cases more corrosive saline environmentscontaining sulfites, or sulfides at various pH valuesmay be used.

It has been found (5) in the case of a 14 carat alloythat under the conditions of the Tuccillo-Nielsen test,selective sulfidation of both Ag and Cu occurs.

Such 'environmental' testing may be supplementedby electrochemical observations in which potentio­dynamic scanning is used to measure the currentdensity at an electrode of the test alloy in a test

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solution as it changes with both increasing (forward)and decreasing (reverse) potentials. From such data,corrosion currents may be calculated which reflect,with some qualifications, the resistance of the test alloyto corrosion.

Through the application of these two testprocedures, Fioravanti and German (4) were able toconclude that the desirable microstructural features forresistance to tarnish and corrosion in low (ca. 50%)gold content dental alloys are large grains, a minimumof phase separation and an absence of precipitated andordered phases.

No comparable data on low caratage jewellery goldalloys have been found. Since many 14, 10, and 9 caratgold alloys contain significant percentages of zinc, it isuncertain whether these conclusions are applicable tothem also.

CORROSION

BackgroundCorrosion of carat gold jewellery can occur as a resultof immersion in sea water, chlorinated water or othercorrosive media. It normally involves selective attackon the less noble alloy constituents which is a featurenot only of the tarnishing of carat gold alloys but alsoof their surface enrichment or depletion gilding. It isalso accepted as an initiating step in stress corrosioncracking of gold alloys, and is of major importance indetermining the biocompatability of dental gold alloys.

Our understanding of the mechanisms of this typeof corrosion is based largely upon studies ofhomogeneous Au-Cu and Au-Ag alloys. Some aspectsof these studies will therefore be discussed.

Selective dissolution from homogeneous alloysIn the idealized situation as depicted in Figures 1 and 2(11) selective dissolution ofA atoms could be expectedto occur preferentially from kink sites (K) in thesurface steps where the atoms are least firmly bound.At sufficiently low potentials, the dissolution currentshould involve mainly less noble A atoms and shoulddecrease as more and more kink sites become occupiedby the more noble B atoms. Thereafter, dissolution canproceed only by removal of A atoms from non-kinksites (N) or from terrace sites (T), which requires agreater overpotential or activation energy. Finally, thealloy could be expected to become completelypassivated when all the surface sites are occupied by themore noble B atoms only. This passivation stage shouldbe reached after removal of A atoms from only a few

~ GoldBulletin 1996, 29(2)

Figure 1 Schematic representation on an atomicscale ofthesurface ofan alloy composed ofdissolvable A atomsand nobleB atoms(based on reference 11)K isa kink siteon a surface stepN isa non-kink siteon a stepT isa terrace site

Figure 2 Schematic representation on an atomicscale oftheformation ofapit in a lowgoldcontentAu-Ag alloybydissolution ofAgatomsand inward migration ofthe vacanciesformed (based on reference 11)

atomic layers at the alloy surface.That this does not occur implies that mass transfer

must accompany the corrosion reaction so ascontinuously to expose more of the less noble A atomsat the surface. This is currently perceived as occurringthrough the inward migration of lattice vacanciesformed at the surface, which facilitates not only surfacediffusion, but also volume diffusion of the less noblemetal to the alloy surface.

It is well known, however, that although suchpassivation occurs with high (over 50%) gold contentalloys, it does not occur with gold alloys of lesser goldcontent (1). This anomaly has been interpreted interms of a corrosion-disorderingldiffusion-reordering

@ GoldBulletin 1996,29(2)

model based largely upon electrochemical studies byPickering et at(15-17), and upon micromorphologicalstudies by Forty et at(18, 19).

In terms of this model, the lattice vacancies from theinitial stage in the corrosion process migrate both on thealloy surface (Figure 1) and into the body of the alloy(Figure 2). This facilitates diffusion ofgold atoms on thesurface, where they form gold islands, and in the processexpose more alloy to attack. It also promotes diffusion ofalloy atoms from the body of the alloy to the surface,and in the process creates corrosion tunnels and pits.Both these processes can be observed inmicromorphological studies.

In the case of low-gold alloys, the formation oftunnels and pits proceeds unhindered and leads toultimate disintegration of the alloy. In the case of high­gold alloys, however, the higher gold content is seen asimposing restraints on the development of tunnels andpits. The precise nature of these restraints has not yetbeen fully established.

This corrosion-disorderingIdiffusion-reorderingmodel nevertheless deepens our understanding of thewell-known inquartation procedure which is used notonly in the parting of the gold-silver bead in thedetermination of gold by fire assay, but also in therefining of gold alloys by treatment with acids. In thisprocedure, the gold content of the alloy is deliberatelyadjusted to below about 25% in order to facilitate thedissolution of non-gold metals, and the recovery of theresidual gold in pure form.

Selective dissolution from heterogeneous alloysThe pOSItlOn in respect of the corrosion ofheterogeneous (multiple phase) alloys is much morecomplex than that ofhomogeneous (single-phase) alloysdiscussed above. In the former more than one phase isexposed at the alloy surface, and these may differsignificantly in composition and in susceptibility tocorrosion. The state of affairs is well illustrated in thecase of the nickel white gold alloys, in which the nickel­rich phase may contain little gold and may thereforecorrode preferentially when the alloys are exposed toacid. To a lesser extent, this type of effect may also occurwith some low caratage Au-Ag-Cu alloys. No reports ofstudies ofAu-Ag-Cu-Zn alloys have been noted.

STRESS CORROSION CRACKINGIts nature and significanceSuperficial tarnishing and corrosion are not the onlyforms of environmental attack to which gold alloys of

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Table 1 Dependence ofthesensitivity to stress corrosion ofhomogeneous Au-Ag and Au-Cu alloys on theirgoldcontentwhen exposedto 2% aqueous FeCl3 solution (24)

(a)Au-Agalloys

(b)Au-Cu alloys

Goldwt.%

o1.85.010.015.033.340.050.054.960

33.345.150.055.067.4

Gold at. %

o1.02.8

5.758.821.527.035.440.045.2

1421

24.328.340.0

Stress inkg/mm2

151210II101010101010

1818151822

Time to fracture of thespecimen in minutes

>10,000,500-3,000ca. 1,500ca. 300ca. 60ca. Ica. Ica. 60

>3,000>10,000

ca. Ica. 1ca. Ica. I

>10,000

low caratages of about 14 or less may be susceptible.Such alloys, depending on their compositions, theirmetallurgical histories and the environments to whichthey are exposed, also undergo stress corrosion (orseason) cracking. This phenomenon, which is alsoobserved in other alloys, involves the local rupture ofthe alloy under the combined effects of corrosion andstress at levels well below those at which failure mightbe expected to occur if these agencies were to operateindependently. When it occurs there is little or noevidence of corrosion products or distortion of thealloy, and there may have been no external applicationof stress. Cases of this latter type are caused by internalstresses only. Stress corrosion cracking can develop veryrapidly (Table 1) and in environments which cause nosuperficial attack on the alloy and might therefore beregarded as harmless.

Thus, SCC may be induced not only by exposureto acids during pickling but also a result of contactwith reagents such as ink, traces of hydrochloric acid inthe atmosphere, perspiration, etc. It has frequentlybeen initiated at points of. stress created in annealedlow carat alloys by subsequent stamping. Articles suchas fountain pen nibs, rings, chains, etc, provide well­known examples. Unless the stresses in such articles arerelieved by annealing, they remain as focal points forattack by any reagent capable of inducing corrosioncracking. A case has been quoted (20) in which thecollapse of hard rolled low caratage foil, stored asreceived from the supplier in its original packaging,

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was found to be a result of attack by fumes ofhydrochloric acid from an adjacent workshop. In thesame environment annealed foil remained stable. Greatcare should therefore be taken to relieve stress in lowcaratage gold alloyswherever feasible.

Stress corrosion cracking of gold jewellery is notvery common, because such jewellery is not oftenexposed to corrosive conditions. When it does occur,however, the root causes are usually twofold. First,inadequate relief of internal stresses resulting fromworking of the alloy during manufacture, and,secondly, exposure of the jewellery to corrosiveconditions (e.g, in chlorinated water). It is notnormally encountered, as mentioned above, in alloys ofcaratages above about 14. The subject has beenreviewed in well referenced papers by Dugmore andDes Forges (21) and by Rapson and Groenwald (22).

Its mechanismTwo stages are generally accepted as being involved instress corrosion cracking" namely crack initiation andcrack propagation.

Crack initiation is seen as involving the formationand rapture of embrittling surface film on the alloyresulting from reaction between a base metalconstituent of the alloy and the environment to whichit is exposed. It is essentially an electrochemical step,and the sites at which it occurs depend on the localizeddistribution of stress and the adsorbed species. Thesites have a high energy and occur at such places as the

@ GoldBulletin 1996,29(2)

alloy (23), concluded that:(a) SCC mechanisms, such as anodic dissolution,

hydrogen embrittlement or cleavage, couldneither predict nor explain the recordedobservations;

(b) of the SCC mechanisms put forward so far, thesurface mobility-SCC mechanism is the only onethat explains from both a qualitative andquantitative point of view, the observations madein the investigation.

Finally, it should perhaps be noted that the studiesof SCC in gold alloys have mostly been carried out onbinary Au-Ag and Au-Cu alloys. No systematic studiesof SCC in ternary and quaternary gold jewellery alloyshave been noted.

intersection of the surface with grain boundaries andwith dislocations or faults at the metal surface.

Crack propagation is more difficult to explain andits mechanism has been the subject of differing viewsfor a number of years. One problem is that it mayfollow grain boundaries and result in intergranularstress corrosion cracking (ISCC) or occur acrossgrains and result in transgranular stress corrosioncracking (TSCC). In both cases the initial stage is anelectrochemical one, and there has been a naturaltendency to try to explain the rapid propagation ofthe initial 'crack tip on a similar basis. Thus, thefreshly exposed metal surfaces at the crack tip havebeen postulated to be highly reactive, with the resultthat electrochemical initiation and further crackingproceed continuously at rates which are vastly greaterthan those under normal equilibrium conditions. It Studies ofhomogeneous binary alloyshas not been possible to demonstrate unequivocally Most of the available information pertaining morewhether, and if so to what extent, this actually occurs. specifically to the occurrence of stress corrosionAn alternative explanation, for which increasing cracking in Au-Ag, Au-Cu, Au-Cu-Ag and Au-Cu-Nievidence has been coming available, is that cracking alloys has emerged from the work of Graf and his 00-

(brittle fracture) results from a rapid migration of workers (24-26). Their work was done at a time whenlattice vacancies to the crack tip, where they weaken the only mechanism for crack propagation appeared tothe structure and so promote further cracking. be anodic attack at the crack tip, for which cathodic

This is the basis of a surface-mobility-SCC areas were necessary at other sites.mechanism put forward by Galvele in 1987. In terms Au-Ag alloys are fully solid solution in type, evenof this, environmentally induced crack propagation is at low temperatures, and the same applies to Au-Cudue to the migration of vacancies along the crack alloys which have been heat treated at 700-800aC andsurface under the influence of stress. The role of the quenched in order to prevent phase transformations atenvironment is to change the surface self-diffusivity of lower temperatures. Table 1 shows the manner inthe alloy. A recent study of SCC in a Ag - 20 Au at % which the sensitivity of such alloys to stress corrosion

Table 2 Theaction ofdifferent reagents on a 33.3 Au-66. 7 Ag allryundera tensile loadof10 kg/mm2 (24)

ElectrochemicalReagent Concentration action*

current in mA

HN03 cone 0.8--1.0HCI+HN03 3:1 2.5

(cone)K2Cr207+H2S04 cone 2-3KMn04+H2S04 2-3

Cr03 2%aq 0.3HCI cone 0.2

H2SO4 cone 0FeCI3 2%aq 0.8CuCI2 2%aq 0.5ZnCI2 2%aq 0CrCI 3 2%aq 0NaCI 5%aq 0

CUS04 5% aq 0Fe2(S04h 2%aq 0.1

"In a gold-silver cell**For a 33.3 Au-66.7 Ag alloy undera tensile loadof10 kg/mm2

($S' GoldBulletin 1996,29(2)

Time to fractureof the specimen

in minutes**

60

1340

>15,00035

>15,000>15,000>15,000>15,000

160

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cracking was found by Graf (24) to vary with theirgold content. In the Au-Ag alloys the sensitivity tostress corrosion, which is nil with pure silver, increasesprogressively with gold content to a maximum andthen falls again to zero at a gold content of about 45weight per cent. The Au-Cu alloys, studied over anarrower range of composition, show a similar trend.Clearly, at a certain level of gold content, the alloy isprotected from stress corrosion, just as it is fromsuperficial corrosion.

Au-Ag (and Au-Cu) alloys do not exhibit stresscorrosion in all environments, however, and Table 2indicates why this may be so.

It will be noted that only where the reagent is suchthat a current actually flows between electrodes of Auand Ag placed in it, does corrosion cracking occur. Grafconcluded therefore that only those reagents capable ofpromoting dissolution of silver in an Au/Ag cellcontaining the reagent are capable of producing stresscorrosion cracking in Au-Ag alloys. In other words onlythose reagents produce stress corrosion cracking, whichcan by dissolution of the less noble metal in the alloy,change the surfacestructure of the alloy.

In the case of exposure of Au-Ag and Au-Cu alloysto reagents which dissolve gold, such as aqua regia andpotassium cyanide solutions, Graf and Budke (25)found that with aqua regia susceptibility to stresscorrosion cracking reached a maximum between 20 and35 atomic % Au as it does with other reagants. Up to 40atomic % Au, attack on these alloys clearly occurredboth by superficial and stress corrosion mechanisms,whereas from 40-100% Au stress corrosion is notoperative and failure is by superficial corrosion only.With potassium cyanide solutions, however, Au-Agalloys exhibited no susceptibility to stress corrosioncracking. This was attributed (25) to the fact thataerated cyanide solutions dissolve gold to form the verystableAu(CN); complex, so that any gold cathodic areascreated by initial attack on active sites tend to bedissolved irreversibly. With aqua regia as the corrodingagent, however, it was concluded that gold dissolved atthe corroding surface is reprecipitated on the walls ofcracks or crevices by the action of exposed solute metal,so that the cathodic areas then deemed necessary forstresscorrosion cracking were created.

Graf (26) also concluded that in alloy solidsolutions the reactivity of the grain boundaries and ofdisturbed areas on the grain surfaces is enhanced, andincreases with the concentration of the solute metal upto a maximum of 50%. Thus, they observed that Au­Cu solid solutions which are susceptible to corrosioncracking under stress, also exhibit intercrystalline

corrosion even in the absence of stress. Perhaps ofgreatest significance, however, was the demonstrationthat this enhanced reactivity in the grain boundaryareas of homogeneous solid solutions was furtherincreased whilst they were undergoing flow ordeformation. In the light of this the special role oftensile stresses in stress corrosion cracking ofhomogeneous alloys could be understood. Such stressesproduce highly activated flowing areas at the base ofnotches, so that intercrystalline or transcrystallinecracks developed rapidly. In polycrystalline material thecracks would be predominantly intergranular, whilstwith single crystals the cracks would be transgranular.

Studies of ternary and quaternary gold alloysand the effects ofstructural heterogeneityIn comparison with the susceptibility to stresscorrosion cracking of homogeneous gold alloys, that ofAu-Ag-Cu, Au-Ag-Cu-Zn and Au-Cu-Ni-Zn alloys ismore complex.

In the case of Au-Ag-Cu jewellery, the position isthat at low caratages and approximately equalpercentages of silver and copper, the alloys contain twosolid phases,one silver-richAu-Agand the other copper­rich Au-Cu. Both of these would be expected toundergo stress corrosion in 2% FeCl3 solution.Nevertheless, Graf has stated that alloys containing thesephases are not susceptible to such corrosion, eventhough homogeneous Au-Ag-Cu alloys may be attacked.The less noble phase was seen as protecting the morenoble one, so that an incipient crack was blocked off assoon as it encountered the more noble phase.

In early production of 14 carat fountain pen nibs,for example, Loebich (27) has stated that when ternaryAu-Ag-Cu alloys were used, it was found desirable toage the fabricated nibs. In the aged condition they didnot undergo stress corrosion cracking in use; whereas ifheated to the point where they became homogeneous,cracking by the action of the ink became likely. Whencertain 14 carat quaternary Au-Ag-Cu-Zn alloys areused, however, such ageing is apparently unnecessary.This could be due to the known limiting effect of thezinc on phase separation in these alloys. Thesusceptibility of alloys of this type to stress corrosion isapparently considerably influenced both by their zinccontents and by heat treatment. Analogous anomaliesoccur in the case of the white Au-Cu-Ni-Zn alloys andthese have been discussed by Gra£

In the case of Au-Cu-Ni alloys, the two solidsolution phases which separate from the homogeneousalloy on ageing are a nickel-rich phase containing 3little gold and copper, and a copper-rich phasi

(5i9' GoldBulletin 1996, 29(

containing high gold but little nickel. Graf points outthat the latter phase might be expected to besusceptible to stress corrosion because the base metal(copper) is less noble than the other main component(gold). On the other hand, the former phase shouldshow no tendency to stress corrosion because nickeland copper (as a result of passivation of the nickel)have about the same electrochemical potentials, andthis expectation is confirmed in practice. Thus, agedAu-Cu-Ni alloys contain one phase which should besusceptible, and one phase which should not besusceptible to stress corrosion. The susceptible phase is,however, inhibited and in practice the non-susceptiblenickel-rich phase is strongly attacked and leached out.

These examples illustrate the complex interplay offactors which determine the behaviour of multi­component gold alloys under stress in corrodingenvironments.

Avoidance ofstress corrosion crackingSteps which can be taken to avoid corrosion crackingof low caratage alloys include the following:

Ensure good housekeeping - i.e., do not storealloys in corrosive environmentsDo not pickle excessivelyMinimize residual stresses in the alloy. This hasbeen discussed by Grimwade (9), who emphasizesthe extent to which these stresses arise from non­homogeneous deformation during working (andstamping) and from differential expansion andcontraction during heating and cooling due totemperature gradients within an article. Wroughtor stamped articles for example can be given astress relief anneal at 250°C for 30 minutes,followed by slow cooling. Under these conditionsthey retain the high strength and hardness builtin by cold working but their internal stressesare reduced. Slightly different treatment isrecommended for nickel white gold products.

If such steps as the above do not remedy theproblem, it may be necessary to use a more noble alloycomposition. This can be achieved not only by use of ahigher caratage gold alloy, but also by incorporating alarger percentage of silver in the alloy (21) or byexploring the effects of additions ofzinc.

Stress corrosion cracking problems have beendiscussed by Normandeau (28). Also, Heuberger,Pfund and Raub (29) have reported on the effects ofvarious surface treatments and of galvanic coatings onthe susceptibility to SCC of a low caratage Au-Ag-Cu­Zu alloy. They found that resistance to SCC wasincreased by plating the. alloy with thick (pore-free)

(ei§' GoldBulletin 1996, 29(2)

coatings of pure gold. Surface enrichment of the alloyin gold by oxidation, etching and mechanical polishinggave better results, and a combination of this treatmentwith gold plating was especiallyeffective.

Stress corrosion cracking of carat golds bymercuryClosely related to stress corrosion cracking in aqueousenvironments is the cracking of low carat golds and ofclad or plated golds which occurs as a result of theirexposure under stress to mercury.

Except that cracking results from preferentialdiffusion of the liquid metal at activated boundary sites,rather than from electrochemical attack and migration oflattice vacancies, stress corrosion cracking by mercury isvery similar to stress corrosion cracking by conventionalaqueous agents (30). Thus, SCC apparently does notoccur with pure metals, and with homogeneous Au-Cualloys, the susceptibilityto corrosion cracking by mercuryincreases with copper content, up to a copper content of50%, and then gradually decreases to zero again as thecopper content is increased to 100%. Low meltingsolders in the liquid state can induce stress corrosioncracking in the same way as mercury.

Suiface enrichment or depletion gilding ofcaratgold alloysIn this process, the surface of a low caratage gold alloyis enriched in gold by exposing it to media whichpreferentially dissolve its base metal components. As aresult the surface layers are enriched in gold, and onpolishing a product of more gold-like appearance isobtained. The use of acid media for this purpose has avery long history. Such media, usually in the form ofacidic plant extracts, were used in Mycenae, Troy,Japan and South America. In more recent times,reagents such as hydrochloric acid, sulphuric acid (1:1)and aqua regia have been employed.

It was to be expected that the mechanism of thisprocess would be similar to that discussed above forsimple corrosion of binary Au-Ag and Au-Cu alloys.That this is the case has been demonstrated by Forty(31), who studied the effects of 50% nitric acid onthin single crystal films of Au50-Ag50 alloy, usingtransmission electron microscopy. He found that as thesurface silver atoms were selectively dissolved, theresidual gold atoms underwent surface diffusion toform a maze-like structure of gold islands. The islandsgrew and the channels between them shrank ascorrosion proceeded, so that they finally formed thinepitaxial surface coatings of gold. Although theseobservations were made on single crystals of Au-Ag

67

alloy, it is probable that ternary Au-Ag-Cu alloysbehave similarly, and that the role of burnishing indepletion gilding is to remove imperfections in thegold layer formed in the depletion step. The localizedheat generated during burnishing would assist theprocess.

No analogous studies of the use of alkaline mediasuch as the cyanide-hydrogen peroxide mixtures usedin the finishing of mass produced jewellery items bythe Mulnet (32) or 'chemical bombing' process havebeen noted.

organization, and has specialized in recent years inpromotion of gold and its uses. He is the author ofmany publications, including a recent book (in press),entitled 'The Science and Technology of the IndustrialUses of Gold'. This paper is based partly on material inthis book and partly on a paper presented at the 1995Santa Fe Symposium on Jewelry ManufacturingTechnology.

REFERENCES

G. Tammann, Z anorg. allg. Cbem., 1919,107,12 L.W Laub and J.W Stanford, GoldBull., 1981,

14, 133 N. Sarkar, 'Proc. Santa Fe Symp. Jewelry Manu£

Technol. 1989', ed. D. Schneller, Met-ChernResearch Inc., Boulder, Colorado, USA, 1990, p.107

4 K.J. Fiorovanti and R. M. German, Gold Bull,1988,21,99

5 C. Courty, H.J. Mathieu and D. Landolt, W7erkst.Korr., 1991,42,288

6 J.P. Randin, Surf Coat Tecbnol., 1988,34,2537 J.P. Randin, P. Ramoni and J.P. Renaud, W7erkst.

Korr., 1992,43, 1158

9

42,29519 AJ. Forty and G. Rowlands, Philos. Mag. A,

1981,43, 17120 E.M. Wise, 'Gold: Recovery, Properties and

Applications', van Nostrand Company Inc.,Princeton, New Jersey, USA, 1941

21 J.M.M. Ougmore and C.D. Desforges, GoldBull., 1979, 12, 140

22 WS. Rapson and T. Groenewald, 'Gold Usage',Academic Press, London, UK, 1978

23 G.S. Duffo and J.R Galvele, Metall. Trans. A,1993,24,425

24 1. Graf, in 'Stress Corrosion and Embrirtlement',ed. WO. Robertson, John Wiley and Sons, New

York, USA, 1956, p. 4825 1. Graf and O. Budke, Z Metallk., 1955,46,37826 1. Graf, Z Metallk., 1975,66, 74927 O. Loebich, Z Metallk., 1953,44,28828 G. Normandeau, 'Proc, Santa Fe Symp. Jewelry

Manu£ Technol., 1990 and 1992', ed. D.Schneller, Met-Chern Research Inc., Boulder,Colorado, USA; see Gold Tecbnol., 1991(5),8 and1992(8),2

29 U. Heuberger. A Pfund and Ch.J. Raub,Gawanouchnik,1989,50,2602

30 1. Graf and H. Klatte, Z Metallk., 1955,46,67331 AJ. Forty, Nature, 1979,282,59732 G. Mulnet, French Patent 2, 137,296 (1972)

Deposition of Gold and Silver ColloidMonolayers on Dendrimer-Modified Silicon

Oxide SurfacesThe modification of glass, silicon and ITO glasssurfaces with starburst dendrimers (see GoldBull., 1996, 29, 16), has enabled researchworkers in Freiburg, Germany and Los AlamosNational Laboratory, USA to deposit gold andsilver colloid monolayers onto these surfaces.Georg Bar, Shai Rubin, Russell W Cutts,Thomas N. Taylor and Thomas A Zawodzinski(Langmuir, 1996, 12, 1172) have described apreparative method which is very simple,requiring only two steps, i.e, immersing thesubstrate into the dendrimer solution anddeposition of colloids. The resulting transparentfilms are stable for at least several weeks.

The structure of the resulting films has beeninvestigated using X-ray photoelectron spectro­scopy (XPS), scanning electron microscopy(SEM), and atomic force microscopy (AFM),and this has shown that the size distribution,particle separation, and degree of colloidmonolayer formation can readily be controlled.The optical properties were examined using UV­visible spectroscopy and surface-enhancedRaman scattering (SERS).

Gold and silver colloids ranging from 15 to80 nm in particle diameter have been depositedonto the dendrimer-modified surfaces. XPS datashow that the dendrimers spontaneously adsorb

<00' GoldBulletin 1996,29(2)

on to various.silicon oxide surfaces - it is likelythat some amine groups are protonated underthe conditions of these experiments, thusimparting a net positive charge to the dendrimerlayer. The driving force for adsorption of thedendrimers onto the surface then arises from theinteraction between the positively chargedprotonated amino groups and the hydrophilicsilicon oxide surface.

SEM and AFM showed that the noble metalcolloid particles are well isolated and confinedto a single layer, and that aggregation does notoccur on the surface. Particle size, inter-particlespacing, and surface coverage can be controlledby varying the colloid concentration in solutionand the immersion time of the substrates,resulting in tunable physical and chemicalproperties of the films. UV-visible spectroscopicdata show that the microstructure directlycontrols the optical properties of the layer. TheSERS studies are being used to elucidate therelationship between the nanoscale architectureand the optical properties of the films. Thisnew technology could also have applications incontrol of particle size in the preparation ofsupported gold catalysts.

David Thompson

69