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RAPIDLY SOLIDIFIED HIGH- DENTAL ALLOYS

Y . C . DURANDET, B. Eng. (C.E.S.T.I. , France)

A thesis submitted for the degree of Doctor of Philosophy

Department of Chemical Engineering, Materials Engineering Group, The University of Adelaide.

DECEMBER 1990 I

TABLÐ OF CONTENTS

TAEILE OF CONTENTS.

DECIARATION ...... iv

ACKNOWLEDGMENTS V

FOREWORD VI

.lX

CHAPTER T INTRODUCTION... I l.I Copper in dental amalgams 3 l.I.l The advent of high-copper amalgams. 3 L.1.2 Addition of copper and structural changes'... 6 1.2 Constitution of --Copper alloys ro CFIAPTER 2 ALLOY PREPARATION AND METALLOGRAPHIC PROCEDURES... I9

2.1 Alloy preparation I9 2.2 Specimen preparation - Scanning Electron Microscopy .20 2.3 Specimen preparation - Analytical Electron Microscopy 20 2.4 Preparation and testing of dental amalgams...... ' . " . .', 23

CHAPTER 3 THE SILVER-TIN-COPPER PFIASE DIAGRAM 24 3.1krtroduction 24 3.2 Construction of the relevant sections of the Ag-Sn-Cu phase diagfarn. 25 3.3 Solidification under equilibrium conditions of Ag-Sn-Cu alloys .-...34 3.4 Interpretation of the microstructures of high-copper amalgam alloys ..... 3I 3.4.1 Equitibrium microstructures...... '39 3.4.2 Solidification microstructures of furnace cooled alloys 42 4B

CHAPTER 4 OBSERVATIONS ON COMMERCIAL AMALGAM ALLOY POWDERS ...... 50 4.I Microstructure of commercial alloys 50 4. l. I Introduction .. 50 4.1.2 Microstructures of commercial alloys Lojic, 1)rtin andVatiant ....5I 4.f .3 Identification of phases in commercial amalgam alloys 59 ll

4.l.4Discussion ...... 72 4.2 lnfluence of alloy microstructure on some properties of amalgam ...... 7 5 4.2.1Results 76 4.2. L. I Fractographic observations ...... 7 7 4.2.L 2 Electron metallography of amalgams...... 82 4.2.2Discussion .....'."....86 4.3 Influence of alloy microstructure on the formation of the n' (Cu6Sn5) phase during the amalgamaton reaction...... 92 4.3.1 Background...... 93 4.3.1. I The dissolution-precipitation model...... 93 4.3.I .2 T}re tin-diffusion model...... 94 4.3. 1.3 The solid-state transformation. 94 4.3.2 Formation of the r1'-phase in high-copper single composition amalgams 96 4.3.2.I Reaction of with the e copper-tin phase...... 98 4.3.2.2 Reaction of mercury with a ternarSr Ag-Sn- Cu alloy...... f OO 4.3.3 Discussion r06 4.4 Conclusions r09 CHAPTER 5 ELECTRON METALLOGRAPHY OF RAPIDLY SOLIDIFIED E)PERIMENT}\L AG-SN-CU ALI.OYS tt2

5.1krtnoduction tL2 5.2 Validation of experimental approach. r r3 5.2. I Backscattered electron metallograp hy...... I L 4 5.2.2 Calculations of cooling rates during rapid solidification ... r22 5.3 Analytical Electron Microscopy of splat quenched alloy T .. r28 5.3.1 CalibrationforEDSmicroanalysis...... 128 5.3.2 Crystallographic procedures...... I3O 5.3.2.I Structure of the Þ and T silver-tin phases r32 5.3.2.2 Structure of the € copper-tin phase...... I33 5.3.3 Results of Analytical Electron Microscopy...... f 35 5.3.3. I Transmission electron microscopy...... I 35 5.3.3.2 Crystallographic analysis ...... 140 5.3.3. 3 EDS microanalysis resuIts ...... I 5 f 5.3. 4 Interpretation of results and discussion ...... 1 5 f 5.4 Effect of cooling rate on the solidifìcation mode of Ag-Sn- Cualloys...... 155 5.4.L Introduction...... 155 5.4.2 Comparison of the microstructures of slowly cooled and rapidly solidified high-copper amalgam alloys r55 5.4.3 Interpretation ...... 157 5.4.4 Discussion. r59 5.5 Conclusions ... t64 llr

CFIAPTER 6 DIMENSIONAL CFIANGE OF AMALGAMS MADE FROM HIGH-CU GAS-ATOIj||{IZED AG-SN ALLOYS . r66

6.1Introduction r66 6.2 Experimental resr:lts. L67 6.2.L Compressive strength of r68 6.2.2 Dimensional change of r68 6.3 Interpretation of results 169 6.4 Discussion t74 CÉIAPTER 7 CONCLUSIONS AND SUGGESTIONS FOR FUTURE IN\¿ESTTGATION...... I7 B

APPENDICES r83 APPENDIX I Commercial high-copper amalgam alloys...'...... " 184 APPENDIX 2 American Dental Association Specification No.l for alloy for dental amalgam r85 APPENDIX 3 IWDS quantitative electron microanalysis of higþ-ocpper amaþaam allcys...... f 86 APPENDIX 4 EDS quantitative electron microanalysis of a high-copper dental amalgam r89 APPENDIX 5 EDS quantitative microanalysis of Ag-Sn thin foil standard in the AEM r92 APPENDIX 6 Kinematic electron structure factors for the disordered hcp p silver-tin phase at 2OO kV...... 194 APPENDIX 7 Kinematic electron structure factors for the ordered orthorhombic y silver-tin phase at 2OO kV...... 195 APPENDIX 8 Kinematic electron structure factors for the ordered orthorhombic metastable e copper-tin phase at 2OO KV r96

APPENDIX 9 Examples of ORTHEX solutions for the indexing of SADP and corresponding ÀL values ...... 197 APPENDIX fO EDS quantitative microanalysis of splat quenched alloy 60Ag- 27 Sn- 13 Cu in the 48M...... 2O3 205 IV

DECLARATION

This thesis contains no material which has been accepted for the award of any other degree or diploma in any university and to the best of my knowledge and belief, contains no material préviously published or written by another person, except where due reference is made in the text.

WONNE C. DURANDET v

ACKNOWLEDGMENTS

I wish to thank my supervisor Professor D. R. Miller for his encouragement and valuable advice during the investigation and his guidance in the preparation of this thesis.

I acknowledge \Mith pleasure the assistance which Mr. Robert J. Finch provided, particularly in preparing spherical particle alloys by gas- atomuation, and Mr. Ian H. Brown's collaboration in the measurement of dimensional change of the amalgams studied in this thesis. Mr. Huw Rosser of the University of Adelaide Electron Optical Centre provided valuable assistance with electron probe microanalysis and analytical electron microscopy.

I would also like to record my appreciation of the help provided by the General staff of the Department, in particular Miss Vicki A. Greenwood and Mr. Bruce H. Ide.

Finally, acknowledgement is given to tl.e University of Adelaide for supporting this project with an Overseas Post-Graduate Research Scholarship. VI

FOREWORD

The literature on structure, properties and clinical applications of dental amalgam goes back to the late fgth century and as the understanding of the behaviour of dental amalgam developed, there emerged a quite specific vocabulary which is now widely accepted to describe the salient issues involved. For the benefit of readers of this thesis who may not be fully conversant with the dental literature, the following brief account of the process of inserting an amalgam restoration may help to clari$r the principal and commonly used terms.

The dental practitioner prepares a cavity in the decayed tooth, taking care to excavate all the carious material and to provide an appropriate undercut to assist in the mechanical retention of the amalgam without, of course, introducing regions of mechanical weakness, stress concentration, etc. in the remaining parts of the tooth. Cavity liners, which constitute thermal and mechanical barriers between the tooth and the restorative material, are then applied to the surfaces of the cavity.

Up until about 20 years ago, tl.e dentist was supplied with the separate components of the amalgam in the form of powdered amalgam alloy and liquid mercury. Carefully weighed proportions of these two components, defined by the target mercury : alloy ratio, were then mixed together or triturated in a mortar and pestle for a recommended period called the trituration time.

The plastic mass of partially reacted amalgam was removed from the pestle after trituration and tl.en packed into the prepared tooth cavity by vll the process of condensation. The condensation pressure (force applied divided by the cross-section area of the condensing tool) was again the subject of specification in order to ensure optimum properties of the hardened amalgam, minimization of the inevitable porosity, removal of excess mercury and a high level of adaption of the margins of the amalgam to the walls of the tooth cavity.

At the present time, most dentists are supplied by the manufacturers with capsules containing preweighed amounts of the amalgam alloy powder and mercury in separate compartments. These capsules can be inserted directly into a high-speed mechanical vibrator which permits the trituration time to be controlled by the dentist. The mixed amalgam is then frequently transferred, in small increments, to the tooth cavity using an 'amalgam-gun", a syringe-t54>e device which minimizes the possibility of the amalgam being contaminated by contact with moisture prior to condensation. Each increment is condensed in the cavity and the process is repeated until the cavity is filled to the required level.

In practice, the cavities are usually overfilled to permit the process of carving to be undertaken so that the surface of the restorations conforms to ttre anatomy of the original tooth. The finlshing procedure, which is carried out after tl.e amalgam has fully hardened (a few days later), involves polishing tl.e surface with a mild abrasive.

It was recognized, from tJre early work of G.V. Black (1895), that the hardening process of dental amalgam was accompanied by a series of complex changes in volume, tl.e sign and magnitude of which varied with time, composition of the alloy, mercury : alloy ratio, condensation vlu pressure, tritutation time, etc. There is a considerable body of clinical evidence which suggests that a large proportion of amalgam failures frequently involving recurrent caries initiates at tl.e interface between the amalgam and the tooth. The significance of dimensional change in relation to loss of marginal integrity can be easily appreciated. Limits of the permitted dimensional change were therefore specified by the American Dental Association and equivalent bodies in other parts of the world. lX

SUMMARY

Previous studies of the reaction of mercury with binary silver-tin alloys have established the significance of the microstructure, in particular the relative proportions of the silver-tin Þ and y phases, in determining the principal characteristics of the reaction and the properties of the hardened amalgam (Black, f895 ; Gayler, 1935 and f 936 ; Abbott et al., 1986).

What was not clear from those earlier investigations was whether the same principles which were delineated for the binary Ag-Sn system were also applicable to ternar¡r Ag-Sn-Cu alloys which form the basis of the modern commercial dental amalgam alloys. These alloys are now very frequently produced by gas-atomization so that not only are the particles very small (5-60 pm diameter), but the metallurgical structure which develops during rapid solidification is so fine that until recently it has been beyond the limits of resolution of available observational techniques.

In the present project, observations of the microstructure of spherical particle alloys were carried out at the limits of resolution of scanning electron microscopy (SEM) and electron probe microanalysis. However, in order to unequivocally determine the precise nature of the microstructure, it was found necessary to employ the techniques of analytical transmission electron microscopy.

The experimental difficulty of preparing thin foils was surmounted by the use of splat quenched flakes of the alloys which had been solidified at cooling rates comparable to those involved in the production of the spherical particles. Using backscattered electron imaging in the SEM, X the microstructures of the splat quenched flakes were observed to be virtually identical to those of the spherical particles. This is believed, therefore, to substantiate the use of thin foils prepared from the splat quenched flakes as a means of determining the microstructure of the spherical particles and the microstructural changes associated \Mith their heat treatment.

As a conseqlrence of these successful approaches, the very fine distribution of copper in the spherical particle alloys has, for the first time, been determined. It has also been possible to gain a detailed understanding of the processes which occur during the reaction of mercury with Ag-Sn-Cu dental amalgam alloys. I

CFIAPTER T INTRODUCTION

Dental amalgams are produced by mixing liquid mercL¡.ry with silver- tin-copper amalgam alloys. Despite their high cost due to a relatively high silver content, they are the most widely used filling material in restorative (Mclean, 1984). It is a fundamental requirement that they must possess adequate physical properties such as high compressive strength, wear resistance, low creep value, low corrosion susceptibility and dimensional change stability in order that they will resist deformation and masticatory forces in the hostile oral environment.

Unfortunately, in clinical practice, failures of amalgam restorations do occur but the caLrses have not been unequivocally determined (Etderton, 1976). They are frequently charactenzed by such problems as lack of marginal integrity, marginal fracture, protrusion from the tooth cavit5r, excessive tarnish and corrosion, etc.

Fracture of the amalgam around the margin is one of the common causes of failure (Osborne and Gale, I98Oa). This is undoubtedly due to a number of factors which include inadequacies of the inherent properties of the amalgam itself (,Letzel and Vrijhoef, 1984), the patient (including the state of oral hygiene), the technique adopted by the dentist (LeØel et al., 1987), the restoration class size and tooth type (Mahler and Mararttz, 1980), etc.

The present investigation relates essential.ly to the properties of the amalgam itself ; the study of the other factors mentioned above is, of course, in the domain of dental practitioners. 2

Since there is a strong relationshÍp between microstructure, chemical and mechanical properties, it is clear that a thorough understanding of the microstructure of amalgam is essential. As dental amalgam is a composite material consisting of partially reacted particles of the original amalgam alloy bonded together by the products of reaction \Mith mercury, it is evident that the microstructure of the amalgam alloy plays an important part, not only in influencing the nature of the amalgamation process itself, but also in governing the properties of the dispersed components of the composite and hence the mechanical and chemical properties of the amalgam.

Historically, alloys for dental amalgams contained at least 65 wt% silver, not more than 29 wto/o tin and less than 6 vtto/o copper (Appendix 2). Although silver-trn amalgam alloys 'were already used as a filling material in dentistry in Ancient China (Chu FIsi-T'ao, 1958), it was not until the fgth century, with the development of mercury-based compounds and G.V. Black's major contribution to the knowledge of the behaviour of silver-tin dental amalgams in 1895, that the dental profession was supplied with a restorative material of satisfactory clinical performance (Abbott and Miller, f98O). These amalgam alloys are often referred to as "low-copper" or "conventional" alloys. Today, however, most modern silver-tin amalgam alloys contain more than 6,¡vto/o copper. This derives from the development in the early 1970's of the so-called high-copper amalgam alloys containing as much as 30 wtolo copper. Amalgams made from these new alloys have shown not only superior physical properties (Bryant, L979 ; Duke et al., 1982 ; Brockhurst and Beech, f982) but also much better clinical behaviour in terms of corrosion resistance and marginal integrity (DoglÍa et al., 1986 ;Marshall et al., l98O). 3

Research has been undertaken to isolate the reason why high- copper amalgams are superior in service to those made from G.V. Black's composition, but despite many attempts to characterize the metallurgical changes associated with the introduction of copper in amalgam alloys and to study their effects on various properties and clinical performance, the precise role of copper in dental amalgams is still not clearly defined (Sarkar, f979).

t.l Copper in dental amalgams

1.1.I The advent of high-copper amalgams

For many years copper has been a common addition to silver-tin dental amalgam altoys. For example, Ward and Scott (1932) report the composition of five commercial amalgam alloys that were widely used in the United States. These contained up to 5 wto/o copper which was believed to assist in the comminution of the alloy ingot into lathe-cut particles if copper was substituted for silver (Phillips, 1982, p3O6) and to impart improved strength to both the amalgam alloy and the hardened amalgam. The American Dental Association specification (Appendix 2) , ptacing an upper limit of 6 v¡to/o on copper concentration, was firmly supported by investigators of dental materials, citing Gayler's work (1935) as evidence that increasing the concentration of copper resulted in undesirable excessive expansion. Therefore, a copper content of l5-2Oo/o was not to be permitted.

Subsequent to this, however, high-copper ternary alloys containing

between 12 v¡Io/o and 30 wto/o copper were introduced. This followed the work of Schoonover and Souder (f941) who pointed out the potential of 4 silver-tin-copper alloys containing more than 6 tttto/o copper : they found that three of the alloys studied produced amalgams with a resistance to saline corrosion superior to that of conventional amalgams. These alloys contained L5-2O wto/o copper but were somehow neglected in the remainder of their experimental study (Greener, 1979). Consequently, it was not until f963 that an amalgam alloy ("Dispersalloy") containing 12 v,tto/o copper was produced commercialty (Innes and Youdelis, 1963). This was what has come to be known as an "admixed" alloy consisting of lathe- cut silver-tin particles and spherical particles \Mith ttre composition of the silver-copper eutectic (the relative proportions of lathe-cut and spherical particles was such that the average copper content of the amalgam alloy was 12 wto/o) . Amalgam prepared from this alloy had a compressive strength some l5olo higher than previous amalgams and it quickly gained wide acceptance by the Dental Profession in many countries even though it violated, at that time, the requirements of the A.D.A. Specifìcation No.l.

The success of Dispersalloy soon led to tl.e formulation of another new kind of alloy again containing 12 wtolo copper but this time consisting of particles of only one composition (Asgar, 1974) instead of an admix of two different compositions. 1)zpical of this new approach rvere the commercial spherical particle alloys T)rtin, Valiant and later Lojic (Appendix 1) whose compositions were approximately 6OAg-27Sn-ISCu,

SOAg-3OSn-2OCu and 46. 5Ag -29Sln-24. 5Cu respectively.

This was a period of very rapid development of dental amalgam alloys for it coincided with the introduction of atomization techniques to produce amalgam alloys in the form of spherical particles. In 1962, Demaree and Taylor first reported the results obtained from low-copper (less than 6 v¡to/o Cu) spherical particle amalgam alloys and showed that 5 the physical properties of these amalgams were comparable with, if not better than, those of conventional lathe-cut alloys of similar composition and had packing characteristics which were favorably received by dental practitioners.

\Mith the advent of the high-copper alloys, the use of low-copper alloys has diminished markedly in recent years. So, too, has there been an increasing trend towards the use of spherical particle amalgam alloys which, from the manufacturer point of view, offer the advantages of a simplification of the production and quality control processes. In particular, the particles are produced in their desired final form (i.e. a spherical particle powder) directly from the melt where composition can be easily controlled. Further, the particles are of uniform microstructure; sizing and distribution of sizes in the spherical powders are easier.

This new manufacturing technique overcame many of the problems arising in lathe-cut alloy production which required long homogenizing heat treatment of ttre cast ingot, machining to produce fine chips, ball milling, sieving and a further annealing heat treatment to reduce residual stresses introduced in the manufacturing process.

Numerous laboratory and clinical investigations (Eames and

Macnamara, 1976 ; Jørgensen, L976 ; Malhotra and Asgar, 1978 : Osborne et al., I98Ob) showed that many of these high-copper alloys (both admixed and single composition) form amalgams with properties superior to the low-copper amalgams, e.g. greater compressive strength, better corrosion resistance, lower creep and plastic deformation and less marginal breakdown. In response to this evidence, the composition limitations of the American Dental Association Specification No. I were 6 relaxed in 1977 to permit the addition of higher proportions of copper (Appendix 2). This 25- to 30- year delay in producing amalgams resistant to corrosion resulted from a misinterpretation of Gayler's work. Gayler (1935) in fact reported that if copper was substituted for tin so that the concentration of Un dropped below 25 wto/o, expansion would occur i but if copper was substituted for silver so that the tin content was maintained at 27 v'tto/o, no excessive expansion occurred.

1.1.2 Addition of copper and structural changes

Examinations of the microstructures of amalgams made from commerciat high-copper alloys soon revealed very significant differences in structures from those made from low-copper alloys.

The reaction of low-copper amalgam alloy with mercury involves partial dissolution in liquid mercury of the y and/or p silver-tin phases of the original alloy particles (Troiano, f938 ; Otani, l97O ; Schoenfeld and Greener, I97f). Since the solubility of silver in mercury is much lower than that of tin, 0.035 wto/o and 0.6 wto/o respectively (Hansen and Anderko, 1958 ; Fairhurst and Cohen, 1972), silver will reach its saturation concentration in mercury before tin, if the dissolution rates of silver and tin are not too different. Thus, according to Abbott et al. (1986), the silver-mercury B1-phase precipitates first as preferential reaction occurs \Mith the p-phase of the (Ê+y) alloy. This is followed by the heterogeneous nucleation of the tin-mercury compound yz (Snz-SHg). The compound Tr (AgzHgs) subsequently nucleates from the p1-phase and/or when the y-phase of the (Þ+y) alloy eventually reacts \Mith mercury. The reaction products grow as dissolution of alloy particles proceeds and as partially reacted alloy particles become covered \Mith newly formed 7 crystals, mostly Tr, the reaction rate decreases. The alloy is usually mixed with mercury in such a ratio that there is insufficient mercury to completely dissolve the (y + Þ) alloy particles ; conseqLlently, unconsumed particles are present in the hardened amalgam. Some liquid mercury may be entrapped within the reaction products and may diffuse through the matrix of reaction products to react with unconsumed alloy particles, thereby creating voids in the structure.

As a result, the microstructure of conventional amalgams is typically constituted of partially reacted particles of the original alloy surrounded, and bound together, by the reaction products y1 md Yz which form the matrix of the hardened amalgam (Allan et al., 1965 ; Wing, f966 ; Greener et al., 1968). The presence of voids and residual mercury can be minimized by controlling manipulative variables such as mercury : alloy ratio, trituration time and condensation pressure. The setting reaction

occurring in conventional amalgams can be summartzed as follows :

(p+y) alloy + Hg + Þr + Tr +Y2+unconsumed (p+y) atloy

+ yr * ')2 + unconsumed (Þ+y) alloy + voids

where B and y are, respectively, the hexagonal and orthorhombic phases of the silver-tin system, Êr = AgHg phase (hcp), yr = AgzHg3 (ordered bcc) and y2 = Snz-BHg (simple hexa$onal).

This reaction is accompanied by volume changes which may result in an overall expansion or contraction. Excessive expansion can produce pressure on the pulp and postoperative pain, protrusion of the restoration from the tooth cavity and perhaps, fracture of the tooth. Severe contraction can lead to microleakage and recurrent caries. Thus, the I

A.D.A. Specification No.l requires that amalgams neither contract nor expand more than 20 p.rn/crn. Dimensional change of the amalgam depends not only on its manipulation (mercury: alloy ratio, condensation pressure, trituration time and particle size dist¡ibution), but also on the metallurgical structure of the amalgam alloy. This dependence on structure will be discussed in detail in Chapter 6.

Copper can be retained in solid solution up to 3 wtolo in the y-phase and up to 5 wtolo in the p-phase (Chang et a1.,1977) so that it undoubtedly results in the solid solution strengthening of the amalgam alloy. Further additions of copper result in the formation of precipitates of the intermetallic compound Cu3Sn (the e-phase of the copper-tin system) which contribute to the strengthening of the amalgam alloy through precipÍtation hardening. Since a significant proportion of partially reacted amalgam alloy is present in the hardened amalgam, copper additions to the amalgam alloy will clearly contribute to the hardness of the amalgam. Further, Mahler et al. (1975) and Cruickshanks-Boyd (fg83a) reported the presence of copper in solution in the amalgam matrix where a further hardening contribution would be expected.

The dissolution-precipÍtation mechanism which governs the reaction of low-copper amalgam alloys with mercury also applies to that of high-copper amalgam alloys (Okabe et al. ,L978a and b ; Boswell,1979). The reaction of high-copper gas-atomized alloys with mercury can be explained in terms of the dissolution of the Ag-Sn-Cu spherical particles in mercury and then, nucleation and growth of the reaction products. Interestingly, the main reaction product is again the y1 silver-mercury phase. I

However, whereas the matrix of conventional amalgams contains the

T2 tin-mercury phase, that of high-copper amalgams shows very little, or no, y2-phase at all. The microstructure of high-copper amalgams can be described as a composite material made up of unconsumed particles of the original alloy, surrounded and bound together by a matrix of reaction products which consists essentially of the Yr-phase and, this time, of the copper-tin q'-phase (Mahler et aI., L975; Takatsu et aI., L977 ; Okabe et a1.,L977; Marshall et al., 1976; Marshall and Marshall, l98l : Sarkar and Greener, 1972). Voids and residual mercury can be minimized by controlling manipulative factors as in the case of low-copper amalgams. The setting reaction of high-copper amalgams can, therefore, be written as follows :

(Ag-Sn-Cu) alloy + Hg + Þr+ Tr t î' + unconsumed (Ag-Sn-Cu) alloy (+ yz) + yr +q'+ unconsumed (Ag-Sn-Cu) alloy + voids (+Tz)

where t]'= Cu6SD5 , Þ1, Yt and y2 aLre as previously defined'

Here again, the setting reaction is accompanied by volume changes, the sign and magnitude of which depend, of course, on manipulative variables (Brown and Miller, f989). However, the influence of the metallurgical structure of the amalgam alloy on the dimensional change of the resulting amalgam still needs to be delineated. This will be discussed in detail in Chapter 6.

Atthough the literature contains a number of reports of metallurgical studies of commercial high-copper amalgams (Cruickshanks-Boyd, 1983a; Bryant, f 985 ; Mahler and Adey, f984 ; Marshall and Marshall, 1981), only a few have involved the structure of the alloys from which those - 10- amalgams are produced (Mahler andAdey, L977 and 1984;Malhotra and

Asgar, 1977 ; Okabe et al., 1978a and b ; Cruickshanks-Boyd, 1982 ; Bryant, f984). This would seem to result from the fact that previous researchers had to rely on the use of commercially produced alloy particles which'were suitable only for examination using ttre techniques of X-Ray diffraction, light optical and/or scanning electron microscopy and electron-probe microanalysis which do not have the capability of resolving the very fìne structures present in the spherical particle amalgam alloys.

Nevertheless, the presence of the Þ artd/ or y phases of the silver-tin system, and of the e and r¡' phases of the copper-tin system was detected in the single composition type alloys ; the crr (Ad and cr (Cu) phases of the silver-copper system were detected in the spherical particles of admixed- type alloys (e.g. Dispersalloy). It is interesting to note that the phases detected were those suggested by the phase diagram, although the microstructures of the spherical particles produced by gas-atomization might be expected to depart from those predicted by the equilibrium diagram. This was further supported by observations which will be reported in Chapters 4 and 5 which suggest that the ternary Ag-Sn-Cu phase diagram provides a most valuable guide to the structure of as- atomized particles and the changes associated with the approach to thermodlmamic equilibrium which occurs during their subsequent heat treatment.

I.2 Constitution of Silver-Tin-Copper alloys

Binary equilibrium diagrams of the silver-copper, silver-tin and copper-tin systems have been thoroughly investigated (Hansen and Anderko, f958). The nomenclature used by Hansen for the Ag-Sn phase -lr- diagram was modified in the present work to conform with that proposed by Murphy (f926) which has now been generally adopted in the dental literature. (Hansen's e and ( silver-tin phases were redesi$nated as Y and P respectively ; the y and p phases of the copper-tin system were renamed y'and p3 respectively, after Gebhardt and Petzow (1959), to distinguish them from the silver-tin T and B phases). It is suggested that the value of the upper limit of the y-phase field in the silver-tin system should be taken as 26.7 wto/o Sn and not 26.85 wto/o Sn which corresponds to the stoichÍometric formula Ag3Sn. This was believed to represent the boundary between the y and the (y + Sn) phase fields, but an alloy of that composition contains in fact both the T and (Sn) phases (Abbott et al., 1982).

The latest silver-tin-copper phase diagram available in the literature is the one established in f959 by Gebhardt and Petzow who took into account some of the results from Guertler and Bonsack (L927). The latter authors originally investigated the ternary system, but this was before the copper-tin binary was accurately determined. Therefore, most data on the Ag-Sn-Cu phase diagram shown in the present project are from Gebhardt and Petzow, which are believed to be more precise.

No terna4r phases have been reported. Intermediate phases of the Cu-Sn and Ag-Sn binary systems also appear in the Ag-Sn-Cu ternary system. Again, it was found necessaС to change the nomenclature used by Gebhardt and Petzow and to adapt it to that defined previously for the binary systems. In order to help the reader of this thesis to understand the modifications of the various nomenclatures used in the literature, the phase designation adopted in the present work has been summarized ín Table l.l, together with that adopted by other workers for comparison. -t2-

Table l.l : Phase Designation Present Hansen and Gebhardt and Composition Symmetry work Anderko Petzow "Ag3Sn" orthor. v Ê e' o/o 9 Ag-Sn (-I5 Sn) hex. p 5 p' Cu-Sn (-23 o/o Sn) b.c.c Þs B Þs Cu3Sn (HT*) b.c.c v' T Y' cu ô ô v CusrSns t e el Cu3Sn orthor. 7 9 CuzoSno hex. S b s q fl Íì" Cu6Sn5 (HT*) hex. Cu6Sn5 &T-) hex. n' q' 1.1' f.c.c c[l crAg c[,1 Ag-rich solid sol. f.c.c c[ c[cu ct, Cu-rich solid sol. (Sn) (sn) Sn Sn-rich solid sol. tetra. * : HT= High Temperature phase ; LT= Low Temperature phase

From microstructural observations, thermal analysis and X-Ray measurement, Gebhardt and Petzow determined the surfaces of primary crystallzation, the course of the liquid valleys and the position of the melting equilibria (Figure I.f) . In addition, isothermal sections at 5O0 oC and 600 'C (Figure L.2), as well as isopleths (vertical temperature- concentration sections for a constant element content) at 5, 10, 15, 20 and 25 wtolo Sn were established. Only the 20 and 25 wto/o Sn sections which are relevant to the present study are shown here in Figure 1.4. A total of sixteen four-phase reactions \Mere found to take place in the ternary system, seven of which involved the liquid-phase and the other nine were solid-state transformations.

More recently, Fedorov et al. (f98f) investigated the Ag-Sn-Cu system in the tin-rich range (more than 30 wto/o Sn) for solders application and confirmed the existence of seven four-phase liquid equilibria. However, the temperatures of four out of these seven reactions Sn c2 cJ

U1 U7

3 00,

J50 20 60

t 1AO. u3 60 10 ò ñ 2 ^ U 2 t J 5N. 60 40 650' 700' 600'\ 550', 600' o 650 o 550. 650' 700' 750' U6 ß: U¡

80 20

60 0, 6 50' 6 S0' 900' 950' t 000' t0 50.

Ag 20 ?, 10 60 ô0 Cu Copper in'L

Figure 7.7 Liquidus projection in ttre ternary Silver-Tin-Copper system (Gebhardt and Petzow, 1959) 10 60

S't' 't 6000

60 ç 10 \ç çt 7,.t a ?". \' ( s.n a ,"\"ó ç .l! ¡ S . o t.lrJ BO ons'ot 20

o o d+d, Ag 20 10 60 80 Cu Copper tn I

10 60 5000

60 10) s ) .ó lr,t,t ar\

80 20

o

A9 20 10 60 00 Copper ¡n 'l

Fþure 7.2 Isothermal sections in the ternary Silver-Tn-Copper system at 600 oC and 5OO 'C (Gebhardt and Petzow, 1959)

n 80 e.ltu) ñ Jv E no ¿0 60 ô '\e q .c AE I L¡

¿0

I'

80 20 \

Sn A9 n lSnl ¿0 Yo BO q. p Silver in wto/o I'v

oC) Fþure 1.3 Solidus projection (37 in the ternary Silver-Tin-Copper system (Kraft, f979) 20"/o Sn 25 "/. Sn S. þt Srcr 8oo +p t' goo S. q.Pr S .qt ùt'lJt S, atr 0 S.c dt.t) ç 700 S.þ 7oo S.st) I a S+ß rdt t' qtrÞl ,0t 0t $+ dt'?t S, çgl ,Pt S.dt 60o 6oo S.a,'p arþt. S.É q+ srl,.7' o /' d,.dl .þt dÇ soo .g Soo 17, a ¡ l a+dt+7t rz'.e ) lr a dt I cf"d ,' / SrP +c L400 I a¡rp tt-arf, ç/' 4'0s 7t 400 l, çqr+ € o,.p Ç7' q.E d F p r3+7 a¡rã orIrt | 0.8 a¡. [. q+¿a¡.1 .J. Ct qeî c7' 3oo 300 .f þ.t.¡t a.0 t I gep eY 'el arl D.l' cl.f 0'Ú grþ þ rf'.c .l .c qrî ?c 200 2oo ll 5.0.c d.p çÍ d.e p Þ.e açp cc lt ,l n sC loo 100 0 20 10 60 60 0 20 10 æ Copper in L Copper in '/"

Fígure 7.4 Vertical temperature-concentration sections for constant Tin content in the ternary Silver-Tin-Copper system : 20 vrto/o Sn and 25 trrto/o Sn isopleths (Gebhardt and Petzow, lg59) - r6- were found to be different from those reported by Gebhardt and Petzow. Further, whereas the latter authors determined all of the seven quaternary liquid equilibria as invariant peritectic transformations, Fedorov identified six invariant quaternary peritectic reactions and one invariant quaternary eutectic reaction. Differential thermal analysis results from Marcon (1988) confirmed some of the previous findings. According to Marcon, reaction U5 (L+Þ+y+e) occurs at 477 oC and reaction Uz is a eutectic reaction (L-+y+(Sn)+n) occurring at 2I7 oC, in agreement \Mith Fedorov et al. ; reaction UO (L+e-+y+q) occurs at 353 oC, in agreement with Gebhardt and Petzow.

The sixteen four-phase reactions, temperatures and compositions of the coexisting phases are listed in Table 1.2 which includes the results of both GebharCt and Petzow, and Fedorov et al. The compositions of the phases involved in the solid state reactions are from Chang et al. (1977), and should be regarded as only approximate, as they are taken from Gebhardt and Petzow's estÍmated drawings.

The equilibrium phases present at room temperature can be determined from the solidus projection or isothermal section at 37 "C reproduced in Figure 1.3 from Kraft (1979).

In the present project, account will be given in general to the data on the ternary diagram reported by Gebhardt and Petzow, except for the values of the temperatures of the liquid equilibria U3, U4, U5 and U7, which will be those reported by Fedorov et al. (located by the * sign in

Table L.2 ). -L7 -

Table 1.2 : Four-phase equilibria in the Ag-Sn-Cu system

Reaction Temp. Coexisting Composition of phases in "C phases wto/o&g wtoloCu wtoloSn U1: L+cr + o1+B3 605 y@ 48 33.5 t8.5 c[ 6.5 80.5 r3 ct1 87 4 I Þs t4 63 23 U2: L+B3 -t crr+Y' 560 L 44.5 29.5 26 Þs t4 62 24 ct' I 86 3.5 ro.5 y' I3 60 27 U3: L+cr1 -+ Ê+y' 550 L 47.5 25.5 27 (540)* c[1 85 2.5 12.5

B 80.5 3.5 r6 v' t2 59 2g Ua:I-+y '+ p+e 540 L 46 24 30 (530) * v II 58 3I p 78 4 r8 e 1t 53 36 U5: L+B -+ y+e 440 L 43.5 t2 44.5 (475)* p 73 4 23 v 7r.5 2.5 26 e to 52 38 U6: L+e + T+q 350 L 20 6 74 e I 54 38 v 7l 2 27 n 4 38 58 U7: L+q -+ y+(Sn) 225 L 4 o.5 95.5 lL + r1+y+(Sn) 2l8l* n 2.5 37.5 60 v 7l r.5 27.5 (sn) 0.5 I 98.5 Ue: ( -+ õ+y'+e 57O ç 3 65 32 ô 2 67 3r v 4 67 29 t 4 62 34 Ug: Þs -l cr*cr,1*]' 545 Þs t4 63 23 (540) * ct 7 8r t2

CT,1 88 3 9 Y' 2T 62 t7 Continued -r8-

Table 1.2: Four-phase equilibria in the Ag-Sn-Cu system - Continued

Reaction Temp Coexisting Composition of phases in oC phases wto/o&g wtoloCu wtoloSn

I Uro: Y'+e --+ ô+B 525 v I 62 29 e I 57 34 õ 5 6r 34 p 78 5 T7 U 11: cr+]' -> cr1+õ 5 I 5 c[ 5 83 L2 T, L4 59 27 c[, I 87 5 I õ 6 66 28 27 Un: ^{ ' -+ cr1+p+õ 5O5 v l8 55 û, I 84 4 t2 B 82 3 I5 ô 7 64 29

U13: cr1+õ + o+p 45O c[ 1 88 3 9 õ I 63 29 c[ 4 83 r3 p 82 4 t4 Ur+: ô -+ ct+p+e 3OO ô 7 63 30 c[ 2 87 ll p 82 3 r5 e 3 62 35 UrS: €+1 + Y+q' L7O e 8 54 38 n 3 38 59 v 7L 3 26 n 2 40 58 Uro: I -+ 1+q'+(Sn) f 5O n 3 37 60 T 70 3 27

1',|' 2 39 59 (sn) o o roo

Data from Gebhardt and Petzow (f959) Phase composition in solid state reactions from Chang et al (L977) *= Data from Fedorov et al. (f 98I) @= L stands for the liquid-phase -r9-

C}IAPTER 2 ALLOY PREPARATION AND METALLOGRAPHIC PROCEDURES

2.L Alloy preparation

In some parts of this project, commercially prepared amalgam alloy powders u¡ere employed. However, in order to have the ability to exercise total control over a number of important variables involved in the manufacture of these alloys, a laboratory-scale gas-atomizing plant capable of producing 25O grams batches of spherical particle alloys was designed in the laboratory by Mr. Robert J. Finch.

In this equipment, appropriate quantities of high purity silver, tin and copper were melted at 12OO "C in a graphite crucible in a nitrogen atmosphere. To atomve the alloys, a valve at the bottom of the crucible was opened so that a fine stream of molten alloy \Mas permitted to enter the atomizing chamber where it was impinged by two high pressure nitrogen gas jets. The spherical particles so formed were collected by a falling curtain of water and subsequently dried and sieved.

For the production of alloys in the form of splat quenched flakes, arrangements were made to cause the stream of molten alloy to fall directly on a polished rectangular copper plate inserted into the atomizing chamber. For this procedure, the nitrogen gas jets \Mere adjusted so that the splat quenched flakes were continuously blown off the surface of the copper plate. By adjusting the velocity of the molten metal, gas velocity etc., splat quenched flakes of the required dimensions could be readily produced. -20-

In order to perform a heat treatment of the commercially-produced and laboratory-produced alloys, the spherical particle powders and/or the splat quenched flakes were sealed in p5rrex capsules under vacuum of less than lO-2 torr after flushing with argon.

2.2 Specimen preparation - Scanning Electron Microscopy

Metallographic specimens were prepared by embedding the spherical alloy particles and/or splat quenched flakes in cold setting Araldite Epoxy Resin (LCI9I) and by grinding with 600 and f 2OO grades of silicon carbide paper, followed by polishing with 6, I and 0. f pm diamond compounds. Specimens for electron metallography were subsequently carbon coated in a vacuum evaporator to prevent charging in the Scanning Electron Microscope (SEM).

Metallographic examination was carried out using Scanning Electron Microscopy (Philips SEM 5O5 with a Tracor Northern EDS system and Etec Autoscan fitted with a Robinson Back-Scattered Electron Detector). Electron probe microanalysis was undertaken using a JEOL Superprobe 733 model fìtted with EDS and WDS systems.

2.3 Specimen preparation - Analytical Electron Microscopy

The preparation of thin foils of individual spherical particles of gas- atomized alloy powders for study in the transmission electron microscope presents a real challenge. Previous techniques reported involved embedding the particles in resin and subsequent slicing by microtome (santner and Omlor, f978) or ion milling (Omlor and Santner, f980) to produce thin foils. Another technique used was to embed the spherical -2t - particles in an electrolytically deposited nickel foil, followed by electropolishing (Field and Fraser, f978) or ion milling (Matson and Omlor, f 98O) of the composite thus formed. The methods of jet acid thinning, ion milling and microtoming have been shown to be applicable in specific cases \Mith only partially satisfactory results. More recently, Shechtman and Gutmanas (f981) reported a technique based on the cold compaction of powders and subsequent electropolishing of the compact for TEM observation.

In the present project, an alternative approach involving the determination of the structure of more experimentally amenable splat quenched flakes (<3OO pm thick) prepared with the same cooling rates and heat treatments as the spherical particles has been investigated.

Atomization and splat cooling are casting processes in which the heat flow is controlled to significant extent by resistance at the mold- metal interface, or "h", the heat transfer coefficient. During gas- quenching of small spherical alloy droplets, h is influenced by both convection at the surface and radiation, and by the type of gas employed : it is also somewhat dependent on particle size. Here, h coefficient can be oc-l expected to be of the order of IO-3 cal cm-2 sec-I for inert $as atmospheres (Flemings, L974). During splat quenching, h is the splat- substrate heat transfer coefficient and depends greatly on the quality of the contact and adhesion between the splat and the substrate. The technique of splat cooling is generally used to produce novel crystalline microstructures or ¿rmorphous alloys by achieving high conduction cooling rates (fOe-1O1o'C/sec) or high values for h, although equivalent cooling rates can also be achieved in gas-quenching of alloy droplets (Patterson, r982). -22-

Since the aim of the present investi$ation was to reproduce the crystalline microstructures of gas-atomized silver-tin-copper alloys for subsequent transmission electron microscopy study, it was not required to achieve an ideal thermal contact between the splat and the solid substrate. In fact, it was necessary to obtain an "indifferent" thermal contact, as defîned by Ruht (1967), so that h was of less than unity and comparable to that involved durin$ inert gas-atomization.

Discs of 3 mm diameter vr'ere spark machined from the as-splat quenched and heat treated flakes. The discs were mechanically polished on l2O0 grade silicon carbide paper to obtain two parallel sided discs. These were tl.en electropolished in a Fischione jet polishing machine in a solution of 70 o/o ethanol, 20 o/o glycerol and LO o/o perchloric acid under the foltowing conditions : T = - 35'C, V = 35 V, I = 4 mA

Following perforation, the thin foils were removed from the polishing bath with the potential still applied and quickly washed in cold high purity ethanol and then dried in a warm air flow. Specimen which were inadequately thinned around the margins of a perforation and/or unsatisfactorily cleaned after electropolishing were ion beam thinned ("Microlap", Ion Tech Ltd) for 15 to 6O minutes accordingly using the fotlowing conditions : O.5 mA ion current, 5kV voltage and l5 o ion beam angle of incidence.

Transmission/ analytical electron microscopic examination of the thin foils was thereafter performed in the High Voltage High Resolution Analytical Electron Microscope JEOL 2OOO FX equipped with a Tracor Northern EDS system. -23-

2.4 Preparation and testing of denta-l amalgams.

Amalgam specimens measuring 4 mm in diameter by approximately 8 mm long were prepared using an " Ultramat " (Southern Dental Industries, Australia) high speed mechanical triturator and the mechanical method of condensation described in the American Dental Association Specification No. I (Guide to Dental Materials and Devices, Chicago : ADA, L974-L975 i pL7).

Manufacturers instructions regarding alloy : mercury ratio and trituration time (Appendix 1) were followed in the preparation of amalgams made from commercial alloy powders and laboratory-annealed commercial alloy powders. Similar procedures were used for the preparation of amalgams made from experimental alloy powders which were normally triturated in cleaned Lojic-type capsules.

Dimensional change during hardening was measured at 37.5 oC, using a microprocessor-equipped low-stress dilatometer (Brown and lde, f 986) over a period of 24 hours. Compressive strengths measurements were carried out in accordance with the ADA Specifications on amalgams stored for the prescribed time at 37 "C. An Instron model Universal testing machine was used urith a crosshead speed of O.5 mm/min. -24-

C}IAPTER 3 THE SILVER-TIN-COPPER PFIASE DIAGRAM

3.I Introduction

It is clear that the constant tin temperature-concentration sections of the phase diagram considered by Gebhardt and Petzow (f 959) are not immediately applicable to the study of the solidification microstructures of ternar¡r alloys containing between 25 and 30 r¡rtolo Sn.

Likewise, the subsequent work of Darvell (1977) which was carried out before the ADA relaxed its restriction on the copper content of ternary alloys, does not consider alloys containing copper in excess of 6 wto/o. The investigation of Fedorov et al. (f981) proved to be only of limited relevance as their work was aimed principally at the microstructures of silver solders.

It was therefore found desirable to reinterpret Gebhardt and Petzow's published results and to construct vertical temperature- concentration sections of the ternary phase diagram of direct relevance to the present investigation of the solidifi.cation microstructures of the range of altoys which encompasses the commercial high-copper amalgam alloys, particularly those containing 25-30 'wto/o Sn and f O-3O wtolo Cu.

White this type of section does not provide information on the compositions of the phases coexisting in equilibrium, it does show the temperatures at which phase changes occur in alloys of given composition and enables the equilibrium solidification path of the alloys to be delineated. The results of this analysis are presented below. -25-

3.2 Construction of the relevant sections of the Ag-Sn-Cu phase diagram

Verticat temperature-concentration sections for constant tin of contents of 27.5, 29 and 3O ',vtolo Sn, and for constant copper contents lO, L2.5, L5, 20 and 25 tvto/o Cu were determined graphically from the data published by Gebhardt and Petzow (f959) and Chang et al. (L977). The temperatures of the liquid equilibria IJg, IJ4, U5 and Uz (Table L.2) determined by Fedorov et al. (f981) were not considered in the construction of the sections since those authors did not report the composition of the phases involved in those invariant reactions. However, their reported temperatures were accounted for when studying solidÍfication sequences. Only the relevant parts of the constructed Sn- sections (O-4O wtolo Cu concentration range) are given here in Figure 3.1 ; the constructed Cu isopleths are presented in Figute 3.2 .

Liquidus lines on the isopleths 'were drawn by joining points determined from the intersection of the straight line representing a constant copper (or tin) content with is