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Casting Alloys
Chapter 19
Casting Alloys
The range of alloys that may be used in the oral environment is limited primarily by the need to avoid corrosion. This leads to the use of two major groups: corrosion-resistant or precious metal alloys, and non-precious but passive alloys. The first is exemplified by gold-based products, the second by cobalt-chromium and similar alloys. Each has advantages and disadvantages as well as special handling considerations that must be taken into account in order to obtain the desired outcome. Corrosion resistance also (usually) requires single (metallic) phase alloys, yet sufficient hardening and strengthening must be obtained for the cast devices to function without permanent deformation under service stresses. Detailed consideration is therefore given to all available mechanisms for hardening, and their appropriateness in the two groups of alloys, in order that their design and handling be understood. In the case of some gold and related alloys the hardening can be reversible, depending on the crystal structure, and controlled by the thermal treatment. Composition is critical in determining mechanical properties, and especially in the cobalt-chromium types. This arises because the composition may lie dangerously close to that at which a brittle structure is obtained. Thermal history is also critical in that carbide grain growth is affected; this can either strengthen or weaken the metal. The selection of alloys for cast devices must be based on a knowledge of the properties of those alloys and their sensitivities to procedural and handling variables. Failure to recognize their individual limitations will result in treatment failure, for it is still a matter of compromise in balancing good and bad aspects of behaviour.
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Whilst the casting process has a long history, and the skills needed for successful casting can be learned by trial and error, it is essentially the properties of the alloys that control the outcome. The market these days is characterized by an enormous and potentially bewildering array of very similar products, even from a single manufacturer. To begin to comprehend the reasons for design choices it is necessary to consider some of the alloy systems in greater detail. However, it is also necessary to point out that in many cases the alloys themselves have often been designed by trial and error, and a comprehensive explanation of their behaviour in suitably scientific terms is simply not possible because the essential data have not been gathered. Nevertheless, it is not too difficult to outline the generic properties of the major classes. When choices of casting alloy are to be made it is primarily the application that provides the guide in that strength, stiffness, density (as it affects the mass of a suitably strong or stiff framework, for example), ductility, suitability for porcelain bonding, and so on must first be considered. A second level of importance is that of the experience of the dental laboratory in handling that alloy. As has already been made plain, the outcome of the entire casting process depends on the combination of investment and alloy, with all the many handling and processing variables that this entails. But, despite the theoretical matching of these factors, it still requires a demonstration that the entire system works to give an acceptable outcome. Part of the system is the skill of the technician in running a process consistently, and being able to calibrate it, as it were, to tune critical steps to ensure success.
§1. Gold Alloys
!1.1
Pure gold Pure gold is the outstanding restorative material. Because of its peculiar combination of properties: tarnish resistance, ductility with work hardening, and the ability to be cold-welded by pressure alone, it is usable as a direct filling material in the form of foil or powder (28§4) (although the skill, effort and expense of this process has led to it being more or less abandoned now). Aided by the deep-rooted and ancient mystical associations for the metal, its colour does not present aesthetic problems for many people, despite the great contrast with tooth material. As has been implied elsewhere (Chap. 24), the patient’s perceptions are at least as important as more fundamental criteria for material selection. However, in its pure state, gold is not very strong for applications other than inlays, i.e. for crowns, bridges and removable partial dentures. The strength advantages of alloys are required, but these require casting rather than direct fabrication techniques. This, of course, precludes work hardening as a means of increasing rigidity because dimensional stability requirements are severe if fit is to be maintained. Recourse to hardening and strengthening by suitable alloying elements must be made. Solid solution hardening The disorder introduced into crystal structures by elements in solid solution has previously been discussed in terms of the ease or otherwise of slip along certain planes of atoms (11§5). The effect is only moderate when there is a close match of atomic radii, as is exemplified by the Ag-Au system (radius ratio 144/143 = -1.01) (Fig. 1.1). But even so, at an atomic ratio of -1:1, when the disorder might be expected to be at its greatest, there is an appreciable peak in the strength. By way of contrast, with Al-Cu "-solid solution alloys the effect is much more
Fig. 1.1 Solid solution hardening in Ag-Au. The peak occurs near an Ag:Au ratio of 1:1.
!1.2
Fig. 1.2 Solid solution hardening in Al-Cu.
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pronounced because the radius ratio (143/128 = 1.12) is very close to the general practical limit for a continuous solid solution, and the distortion is much greater (Fig. 1.2).[1] Indeed, the solubility limit for Cu in Al is 2.5 at% at 548 °C. The ratio for Au-Cu (144/128 = 1.13) is also close to that limit, and while a continuous solid solution is formed above -400 °C, i.e. at all compositions, distinct new phases in fact form below that temperature (Fig. 1.3).[2] Of particular interest are those corresponding to stoichiometries around AuCu3 and AuCu, but it is the latter which is the phase of principal interest in dentistry. A peak in solid-solution hardening is again observed at this atomic ratio (i.e. 1:1).
Fig. 1.3 The Au-Cu system equilibrium diagram and some associated crystal structures. The narrow gap between liquidus and solidus is important for avoiding appreciable coring in these and similar alloys when cast as there will be little time for segregation.
Fig. 1.4 Unit cells for the various phases of the Au-Cu system. Notice that the height (c-axis dimension) of the *-phase is shorter than the a and b dimensions.
The Au-Cu constitutional diagram has a feature known as a solution minimum at about 20 mass% Cu. Here the reaction on heating is " Y Liq that is, the solidus and liquidus coincide at this point. It is therefore a type of point reaction. Both liquidus and solidus have horizontal tangents at this point. This reaction should be distinguished from that of the melting of a pure component (see box, 12§1.2).
The crystal structure for the whole Au-Cu system is essentially f.c.c. (11§3) and, in common with many other solid solutions, the "-phase is one of random substitution (Figs 1.4a, 1.5a). That is, at any point in the lattice of an "-phase alloy of composition AuCu, the probability of finding either kind of atom is exactly 0.5.
!1.3
Ordered phases In *-phase AuCu the structure is ordered. The Au and Cu atoms alternate regularly, which is equivalent to alternating layers of Au and Cu atoms (Figs 1.4b, 1.5b), in a pattern that repeats over long distances. In contrast, a random substitutional solid solution does not have long-range order for the occupancy (element identity) of each lattice site, although the probability of finding a particular kind of atom is fixed at the overall atomic proportion for the phase. Such a structure arises because the layers fit together just a little more compactly, increasing the regularity of the structure, thus lowering its energy. This rearrangement results in a slight distortion from a perfect cubic lattice to a tetragonal one, in which one of the three unit cell dimensions is smaller than the other two (the vertical direction in Fig. 1.4b), the ratio being about 0.935. This structure can be seen to consist in effect of two interlaced lattices, in this case of identical pattern: one of Cu only, the other of Au only (Fig. 1.6). This type of structure is known as a superlattice and is characteristic of ordered solid solutions. Alloys of the composition AuCu3 produce a similar ordered,
Fig. 1.5 Example planes in (a) random and (b) ordered 50:50 solid solution crystal structures.
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low-temperature phase ($, Fig. 1.4c) in which the Au atoms are placed at the corners of the f.c.c. lattice and the Cu at the faces. However, the lattice constants, the unit cell dimensions, are not appreciably different from those of the "-phase in this case, and this phase is of no particular dental interest. It can be seen that the close-packed planes of the f.c.c. structure (albeit with some irregularity because of the random substitutions of the solid solution) are taken further from ideality by the tetragonal distortion. Slip is therefore more difficult because the principal close-packed ‘planes’ (normal to the cube diagonals, Fig. 11§3.23) are not quite planar. This means that for this reason alone * is harder than "-phase at the same composition, an effect that does not operate for $-phase. The significance of the ordered solid solution *-phase, and its importance to dentistry, lies in the fact that it produces a structure markedly harder than just "-phase. While this makes such an alloy much more suitable for use in the dental context, it also allows a special kind of heat-treatment to be used to develop the degree of hardness required for any particular application or, conversely, to permit softening for adjustment which is then followed by rehardening. This process is enhanced by a further source of hardening.
!1.4
Fig. 1.6 A superlattice appears as if there are two separate but intermeshed lattices. This is ordered AuCu, and the two lattices may be described as “ base-centred tetragonal”.
Coherency strain hardening If an alloy with the composition AuCu is cast, rapid cooling or quenching to around room temperature will preserve the high-temperature, random, solid solution structure of the "-phase. There will be little segregation (12§1) because the liquidus and solidus are very close together. But if the temperature is then raised again and held below that of the upper limit of the boundary of the *-phase field (at, say, about 400°C), equilibration will now occur by diffusion, and the superlattice structure will tend to form. The higher the temperature the greater the diffusion rate, because diffusion is a process requiring an activation energy. But because there is a change in unit cell dimensions, notably a reduction in the direction of the c-axis, with concomitant slight increases in the other two directions, the transformation produces a strain in the lattice. The initiation of the " to * phase change will take place at a large number of separate locations within any grain randomly and independently. This will necessarily be with random orientation of the tetragonal c-axis along any of the three axes of the original cubic structure. These regions of differing c-axis orientation within a given grain are called domains. Taken over a large enough region the net strain in each grain will average out to zero, and there will be very little volume change overall. But the variation from one orientation to another from place to place within the grain superimposes on those primary grains a microscopic lamellar or granular structure, and it is this that makes the heat-treated material hard. There are now effectively very many more grain boundaries, i.e. between domains, so dislocations travel much shorter distances before their movement is inhibited. There is less slip. The formation of $-phase does not cause a similar hardening because there is no associated axial dimensional change to induce strain. The *-phase hardening process is entirely reversible: reheating above ~400 °C again produces the "-phase, which may again be preserved by quenching. From another point of view, we can see that it is the periodicity of the structure in the *-phase which is varying. The crystal lattice does not repeat indefinitely in any direction, but every so often skips a step, as it were. The boundaries between areas of different phase (in this wave periodic sense, not the
Fig. 1.7 An atom map showing how domains result in coherency strain in *-AuCu (‘atoms’ roughly to scale).
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structural-constitutional sense) are strained, since the spacing between layers and along rows differs slightly across a domain boundary. This then is known as coherency strain. Slip through such boundaries is therefore more difficult, that is, there is coherency strain hardening. Fig. 1.7 shows what a random section through the junction between four such domains might look like. Since the spacing of the atoms in each kind of domain is different from that in its neighbour, the mismatch generates the strain. Although *-phase is expected to be a little harder because of the tetragonal distortion, on top of what would be expected from the substitional effects, this on its own is not very great. Most of the hardening is due to the coherency strain; the superlattice by itself does not confer the observed change, which is largely due to the domains forming with random orientations. Current gold alloys for dentistry may consist of 6 or more major constituents: Au, Cu, Ag, Pt, Pd, Zn; several other minor additions may also be made. Clearly, the multi-dimensional phase diagram for such a system is going to be extremely complicated, and full details are not at present known for any such system. Detailed explanation of the purpose and effects of any single alloying element in such contexts cannot be given, even though sometimes rather vague statements may be made. However, certain aspects and broad principles can be summarized, and these are applicable no matter which alloy is under consideration.
!1.5
Precipitation hardening When a second phase is allowed to separate from (usually) a high-temperature solid solution (see Fig. 12§3.9), those new precipitated crystals form randomly throughout the original grains of the alloy. They therefore get in the way of dislocation movement, limiting slip. Such alloys are described as being precipitation hardened. The example of the Ag-Cu system has already been discussed (Fig. 12§3.1). There, for example, the separation of the $-phase (solid solution of Ag in Cu) from a high temperature, Cu-rich "-phase (solid solution of Cu in Ag) at a temperature below the solvus results in just such a very fine-grained precipitate. Conversely, a heat treatment above the solvus will redissolve the precipitate to form a solid solution again. Reactions similar to these will play a part in the overall hardenability of such alloys. Time and temperature are the two important aspects of any heat-treatment: to determine the extent of any change, and the direction and rate of that change. A continuous solid solution characterizes Ag-Au alloys (Fig. 12§1.1) as it does Au-Cu at high temperature, but in the ternary system Ag-Au-Cu below ~400 °C this is restricted to the very edges of the diagram (Fig. 1.8). (We shall ignore for the moment the formation of $ and *-phase Au-Cu superlattices.) However, as the temperature is raised the gold-rich boundary of the two-phase field (Au,Ag) + (Au,Cu) moves steadily towards the Ag-Cu boundary; at the same time the Ag- and Cu-rich solvi move away from the corresponding pure metal corners (Fig. 1.9). The two-phase field can be visualized as a dome-shaped volume in the full phase diagram (Fig. 1.10). Note that the faces of the prism are the three, two-component phase diagrams (cf. 14§3.7).
Fig. 1.8 The solid solubility limit in Ag-Au-Cu alloys at 371°C. Compare Fig. 1.3, and Figs 12§1.1, 12§3.1. (The detail of the ordered phases is omitted).
Fig. 1.9 Several solid solubility isotherms of the Ag-Au-Cu system mapped on the same diagram. This is a plan view of the triangular prism of Fig. 1.10; the isotherms are therefore equivalent to the height contours of that surface.
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It should be clear from this that an alloy with a composition corresponding to the two-phase field (i.e. at lowtemperature equilibrium) will, if quenched from the melt, as in a normal casting, maintain the high-temperature, single-phase solid-solution structure. Subsequent ageing at temperatures below the single-phase limit will cause the decomposition of the solid solution into two new solid solution phases.
!1.6
Ternary system phase-composition The compositions of the conjugate phases present in the Ag-Au-Cu system cannot in fact be stated as the necessary analytical work does not yet seem to have been done. As explained in 8§4.2, this cannot be done in ternary systems simply by reference to the phase diagram, using tie-lines as is the case for binary systems; it is necessary to analyse each pair for a well-equilibrated alloy. Thus, unless the position of the plait point is known, and its locus with change in temperature on the single phase field boundary surface (Fig. 1.10), it is impossible even to begin to discuss the compositions of the two phases in an aged Ag-Au-Cu alloy, apart from the simple expectation that one is copper-rich and the other is silver-rich. Even so, we can understand aspects of the behaviour of the system from what we do know.
Fig. 1.10 A view of the three-dimensional reconstruction of the solid solubility limit surface of the Ag-Au-Cu system.
What we can say is that on cooling from the singlephase field, a second phase will precipitate in the matrix of the parent, or more generally a phase-separation (8§4.1) akin to that of eutectoid formation (Fig. 1.11), a minimum diffusion path to establish the appropriate compositions in least time. Of course, grain growth will occur if allowed to remain at a high-enough temperature.
!1.7
Reversible hardening Thus, the hardenability of dental gold alloys is probably mostly due to the decomposition of the hightemperature solid solution into a fine-grain two-phase structure. The effect can be traced in Fig. 1.12, which shows the Vickers Hardness (1§8.2) for alloys in three conditions: (1) alloys that have been quenched after a solution heat treatment. This involves holding the metal above the temperature at which a single solid solution begins to be stable for long enough that diffusion can occur to re-equilibrate the structure, and so followed by quenching, rapid cooling to room temperature such as by sudden immersion in water while hot.
Fig. 1.11 Section through Fig. 1.10 at 60% Au. Note that isothermal tie-lines cannot be drawn in the "1 + "2 field of a pseudo-binary diagram of this kind – in general they lie at an angle to the plane of the page (cf. 8§4.2)
(2) alloys that have been annealed in the two-phase region (aged), similarly to re-equilibrate the structure. This age hardening is perhaps more accurately described as precipitation hardening. (3) those that have been allowed to air-cool after casting, which represents an intermediate cooling rate through the transition temperature.
Fig. 1.12 Variation in indentation hardness as a function of composition for various heat treatments over three constant Au-content sections of the Ag-Au-Cu system. (See Fig. 1.13)
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The effect of varying the gold content is shown. The quenched alloys show only solid solution hardening, while the aged alloys show the effect of the two-phase structure in addition. The air-cooled condition, i.e. which was slow enough to allow some development of the two-phase structure, lies somewhere between. In addition, the hardening effect of the formation of *-phase AuCu is clearly seen at 75%Au (Figs 1.12, 1.13). Available evidence suggests that even though a third element (Ag) is present, superlattice structures of the AuCu-type may still form, and even when the Ag- and Curich two phases are present. The hardenability of the dental Ag-Au-Cu alloys is therefore due to solid solution, precipitation and superlattice effects occurring together. Indeed, this is also known to be the case with Ag-Au-Cu-Pt alloys (cf. Figs 1.15 and 1.21 with Fig. 1.3). Alloy colour The addition of Ag to Au-Cu alloys provides the benefit of precipitation hardening, but it also allows a certain amount of variation in the colour of the alloy (Fig. 1.14). This can be used to offset the reddening effect of the presence of the Cu. If ‘white gold’ is required, it is more likely to be alloyed with Pd, at about 5%, which has quite a strong effect; Pt has only a slight effect. (Ni is a particularly powerful ‘whitener’ of gold alloys for jewellery but because of its known hypersensitizing action is likely to be less used in dentistry). Obviously, such colour adjustments are not concerned with matching tooth tissue, but some people may find one colour more pleasing than another, or there may be a need to match the appearance of two alloys which have different mechanical properties because of different functional requirements.
Fig. 1.13 The positions of the sections of the Ag-Au-Cu system shown in Fig. 1.12.
!1.8
Fig. 1.14 Alloy colour vs. composition in the Ag-Au-Cu system. The names are those conventionally used in the jewellery trade. Silver whitens, copper reddens the alloy.
!1.9
Grain size The Hall-Petch relationship has previously been mentioned (11§4.1) as indicating the effect of decreasing grain size on the strength. Grain size is a function of the cooling rate: faster cooling gives more and smaller crystals. In addition, perfect crystals require time to grow. Rapid cooling does not permit the annealing out of spontaneous stacking faults. Errors will be covered over and locked into the structure by the solidifying metal on top. Thus, the faster the cooling, the more stacking faults there will be, in parallel with the formation of smaller grains. Stacking faults, i.e. dislocations, create strain fields that inhibit slip. There will therefore be some hardening effect simply due to rapid cooling for two reasons: the Hall-Petch effect and stacking faults. Indeed, it could be argued that slow cooling is detrimental. Thus, the investment mould temperature must not be too high to avoid inhibiting solidification too much.
In some countries, “white gold” is used as the name for platinum and platinum alloys (whether or not they actually contain any gold), and even sometimes for rhodium and its alloys. This has to do with the unfamiliarity of platinum as a jewellery metal, and the higher status afforded the word “gold”. Care is therefore necessary in some contexts to be very precise about what is being discussed.
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!1.10 Grain-refining Grain size is a function of the cooling rate, but it does vary from alloy to alloy because the degree of supercooling required for homogeneous nucleation varies, and so the number of nuclei formed under any given set of conditions will also vary. However, this aspect can also be controlled deliberately by the addition of so-called grain-refining alloying elements. In dental gold alloys this is now usually done with iridium (Ir) (although Ru, Re and Rh are also reported). When there is a very large difference between the melting points of two metals, their binary phase diagram tends to show a very large gap between the liquidus and solidus (Au-Pt is one such example which has been studied, Fig. 1.15), in great contrast to the Au-Cu system (Fig. 1.3). In other words, the solubility of the grain-refining element is low. This in fact means that the nucleation of grains occurs very readily. Referring to Fig. 1.15, it can be seen that even at 80% Au at 1350 °C the first solid to form is about 75% Pt.
Fig. 1.15 The equilibrium diagram for the system Au-Pt. Such a diagram indicates a strong tendency to coring - note the great width of the " + Liq field above about 1350/C.
Iridium, which has a melting point of 2454 °C, shows such an effect, only much more exaggeratedly. Its solubility in gold is very low (less than 0.005% by mass) and so can be expected to readily precipitate from the melt. Indeed, it is said that there is always this solid present at the point of casting. It would therefore initiate heterogeneous nucleation (11§2.5) of the Au-alloy, for which process the activation energy is much smaller than for homogeneous nucleation. In other words, the degree of super-cooling required is very small. Grain-refining thus ensures that the distance a dislocation can travel is small, thereby adding to the hardening of the alloy. In passing, we can note that the phase diagram for Au-Pt has a two-phase field of two solid solutions ("1 + "2) enclosed within the general solid solution field which is otherwise continuous (i.e. stretching between the two terminals). Such a system is exactly parallel to the two-liquid systems discussed in 8§4.1, as it shows an upper critical solution point (cf. Fig. 8§4.1), and similar to the two-phase field shown in Fig. 1.11. Thus, this system would also show phase-separation on cooling into that region. Au-Pt is the basis of some alloys used for porcelain-fused-to-metal devices (25§6).
Fig. 1.16 The effect of work-hardening on the Brinell indentation hardness in the system Ag-Pd.
!1.11 Work hardening Although work hardening is important in several metals’ contexts, it is not an option applicable to cast alloys in general – the deformation involved clearly would spoil the dimensional accuracy of any such device. Occasionally, however, work hardening is significant. Thus, in the adjustment of clasps to compensate for residual errors in the casting, the deformation will inevitably harden and therefore embrittle the metal. There is then an increased risk of fracture on any further adjustment. Bear in mind that manual adjustment is in effect a ‘strain-controlled’ action (as opposed to loadcontrolled) (1§3.1) since the target outcome is a certain amount of deformation, but in doing this by hand it is
Fig. 1.17 The equilibrium diagram for the system Au-Pd. The narrow gap between liquidus and solidus suggests little coring in these and similar alloys when cast.
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rather difficult to control the amount. In practice this means that bending should be done in one direction only, gradually approaching the correct shape. Avoidance of work hardening is a good reason for using a reversibly hardenable alloy. But work hardening can also be beneficial when the margins of a crown or inlay have to be burnished, i.e. made to conform very accurately to the adjacent tooth. The hardening limits the possibility of deformation later in service.
!1.12 Related alloys There are, as pointed out above, many other alloying elements possible for dental ‘golds’, and some indeed form alloy systems of their own that have found application in dentistry. We may briefly explore a few examples to illustrate further the general principles.
Fig. 1.18 Variation in Vickers indentation hardness with composition in the system Ag-Au-Pd (annealed).
The Ag-Pd system is one of simple solid solution over the whole composition range (Fig. 12§1.2) (radius ratio: 1.05). No further hardening after the effects of solid-solution is possible except from cold-work, which annealing can then remove. After annealing, such alloys are quite weak and soft (Fig. 1.16), although they still show very clearly the general effect of hardening by solid solution disorder (cf. Ag-Au, Fig. 1.1). Unfortunately, these alloys have a strong tendency when molten to dissolve oxygen, which then appears as disseminated porosity in the cast metal (Fig. 18§4.4). Au-Pd also forms a continuous solid solution (Fig. 1.17) (radius ratio also 1.05) and the ternary system Ag-Au-Pd is similarly one of a continuous solid solution. Although some variation in hardness is apparent in that system (Fig. 1.18), although with a complicated pattern, all of these alloys are still relatively soft and not amenable to age or precipitation hardening.
Fig. 1.19 The equilibrium phase diagram for the Cu-Pd system. The formation of ordered phases is comparable to that in the Au-Cu system.
The Cu-Pd system (radius ratio 1.07) shows superlattice formation (Fig. 1.19), just as in Au-Cu (although not exactly at the expected 50-50 composition). As a result the Ag-Cu-Pd system shows essentially the same features as does Ag-Au-Cu, with age-hardening due to both CuPd superlattice formation and a precipitation of two solid solution phases. Substantial improvement in mechanical properties can be made this way (Fig. 1.20). Notice in particular the peak hardness near the centre of the diagram where most disorder would be expected. The Cu-Pt system (radius ratio 1.09) also shows superlattice formation (Fig. 1.21), and agehardening due to the reversible formation of CuPt is possible.
Fig. 1.20 Variation in Vickers indentation hardness with composition in the system Ag-Cu-Pd (annealed).
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!1.13 General considerations What all this illustrates is that there is considerable scope for the design of alloys, even using metals that are very closely related and have very similar chemical properties. The subtlety and complexity of these systems defeats any attempt at memorizing property values, or even compositions. In fact, this is quite pointless since small changes may cause profound changes in behaviour. It does emphasise the need to consider manufacturers’ data and instructions very carefully when making product selections, or changing product. One such aspect of all of these alloys that must be taken into account in casting them is that the liquidus temperature may vary considerably with composition. Some illustrative examples are given in Figs 1.22 - 1.24.[3] Care must be taken that the alloy is heated sufficiently above the liquidus that casting is successful. Different commercial alloys may show appreciable variation in the casting temperature due to the inclusion of further alloying elements and this must be checked. A second point is the melting range, the vertical (temperature) distance between the liquidus and solidus. The wider this gap the greater the likelihood of coring, the first freezing metal having a grossly different composition to that freezing later (12§1). Homogenization annealing then becomes necessary. This is particularly true when platinum forms an appreciable proportion of gold alloys (cf. Fig. 1.15); it is not normally present to more than about 4 mass% because of the tendency to coring that it conveys. Two-phase alloy systems are likely to be less corrosion-resistant than are single-phase alloys, although both phases may be sufficiently noble or electronegative for this to be quite unimportant in alloys with high enough content of elements such Au and Pt. Even so, the requirements of mechanical, chemical (i.e. corrosion) and handling (i.e. casting) properties need to be carefully balanced for optimum clinical results. But with the emphasis turning more and more to alloys of lower intrinsic cost, the difficulties of maintaining corrosion resistance will increase, and polyphase structures will be more prone to suffer in this way. Gold alloys suffer from sensitivity to the market price of the gold content, but they are also handicapped by limited strength and modulus of elasticity compared with the demands made on removable appliances. Other difficulties include their high density, which thus causes high weights for devices, possibly causing problems in stability and comfort. These kinds of factor justify consideration of other alloy systems.
Fig. 1.21 The equilibrium phase diagram for the system CuPt. Again, the formation of ordered phases is comparable to that in the Au-Cu system.
Fig. 1.22 The liquidus temperature contour map for the system Ag-Au-Pd.
Fig. 1.23 The liquidus temperature contour map for the system Ag-Au-Pd.
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!1.14 Scavengers Because some of the alloying elements of gold alloys are oxidizable in air, copper in particular, and such oxides may form inclusions in the cast metal which have effects such as acting as flaws and embrittling the alloy, there is a need to limit this action with a deoxidizer. Zinc, and sometimes indium (In), may be added to the original formulation to provide a readily-reacting scavenger of dissolved oxygen. The resulting oxides then appear as part of the dross in the crucible. The process and outcome are similar to those in the case of amalgam alloys (14§2.4). Of course, a scavenger is meant to be consumed, and such that little is intended to be present in the actual cast alloy (again, cf. Zn in silver amalgam alloy). Re-use of old buttons is problematic in that the composition must have changed by virtue of such oxidation, and not just by loss of scavengers: Sn, Ga, Si and others are also susceptible to reaction or volatilization.
Fig. 1.24 The liquidus temperature contour map for the system Ag-Cu-Pd.
§2. Chromium Alloys A series of alloys, originally developed for high strength at elevated temperatures, called stellites, have provided an alternative to gold-based alloys for dental castings. Chromium is the nominal principal constituent of the majority of these so-called ‘base metal’ alloys. Like aluminium, it is an extremely reactive metal that forms a tenacious, thin (thin enough to be colourless; Cr2O3 is green in bulk), and impermeable film of oxide which limits further reaction, making it essentially non-tarnishing in air and many aqueous media of near neutral pH. It possesses the further remarkable advantage of conferring these properties on alloys which contain an appreciable proportion of it. This ready production of an oxide film is called passivation, and it has obvious interest in dentistry (13§5). Despite this usefulness, dental alloys are not structurally based on chromium (which is b.c.c., Fig. 11§3.3); Ni, Co and Fe form the greatest proportions of the elements present in commercial products. This is because extensive solid solutions ("-phase) are formed in the Co-Ni and Co-Fe systems at high temperatures; these have the austenitic structure, i.e. f.c.c. Only limited solid-solubility is shown in other systems, such as Cr-Ni and Co-Cr itself. Even so, in the ternary and quaternary systems, the resulting f.c.c. solid solutions have very wide ranges of possible compositions (Figs 2.1, 2.2). That this occurs may be readily understood from the atomic radii of the metals in such alloys (in pm): Fe: 126, Co: 125, Ni: 125, Cr: 129, Mo: 140. The radius ratios for the first four elements do not exceed ~1.03, and Mo tends to be used up to only about 5 mass%. Curiously enough, because these are all f.c.c. structures, the mechanical properties of ductility, softness and low strength are common to all such alloys with relatively little variation – as they stand they are still in fact unsuitable for dentistry. However, the intention is to create, if possible, an alloy with mechanical and chemical properties suitable to the task. The preference is therefore a single-phase alloy, to avoid galvanic effects, passivating, to reduce reaction rates and thus corrosion, while maximizing the disorder in the solid solution to increase the yield point. In order to make such alloys functional in the sense of carrying large stresses over long spans with small cross-sections (in other words, to fit in the mouth), further hardening of such alloys may be achieved in a number of ways. The differences, as well as the similarities, between this set of possibilities and that for gold-based alloys should be noted.
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Fig. 2.1 Parts of four of the ternary isothermal equilibrium phase diagrams of interest in connection with ‘chrome’ dental casting alloys. The features of concern are the extensive " solid solution regions, and the adjacent " + F phase field in each case.
Fig. 2.2 Two of the possible quaternary isothermal equilibrium phase diagrams for the chrome alloy series, showing only the "-solid solution regions (shaded).
!2.1
Other metallic alloying elements The effect of the disorder produced in a solid solution by atoms of different sizes in a lattice has already been discussed. Suitable elements in the present context include Mn (~127 pm) and W (~140 pm). It is important to notice that these alloying elements are required to dissolve in the "-phase solid solution to avoid galvanic corrosion. Manganese can also play a role as a deoxidizer (as does Si in these kinds of alloy). There are in fact many other alloying elements that are used in a bewildering variety of commercial products for dentistry, some of which products are predominantly one metal, some another. There are too many choices available. It is beyond the scope of this book to detail the possibilities and it is probably beyond the ability of any practitioner to understand the resulting plethora of property and behaviour data, usage options and prices. Accordingly, here – as for the similar situation with respect to gold and related alloys – only some broad principles are given.
!2.2
Metallic phase precipitation Aluminium is added to at least one dental alloy which contains much Ni; a fine-grained precipitate of the compound AlNi3 is formed which is believed to contribute greatly to the modulus of elasticity and strength of the alloy. Although this looks as though it may be a serious corrosion risk, the fineness of the precipitate may mean that it is soon dissolved when exposed at the surface, and no great roughness results. However, the
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frequency of occurrence of nickel sensitivity may be enough to preclude its use in biological contexts if only as a prudent precaution. The absence of serious mechanical problems does not adequately offset the risk.
!2.3
Dissolved carbon The solid solutions discussed earlier (§1, §2; 12§1) were substitutional, that is, an atom of one kind was replaced by one of another kind, in other words a substitution at the exact same position in the unit cell. However, there is a further kind of solid solution possible when the solute atom is sufficiently small, and that is interstitial solid solution. This possibility arises because in some lattices the packing of the layers leaves enough space to accommodate, with some distortion of the surroundings, atoms such as B, C, N and O (Fig. 2.3). That is, even in close-packed crystal structures there are necessarily holes – interstices – between the layers of atoms (11§3) (not to be confused with vacancies, 11§6.1). The resulting distortion of the crystal lattice itself is enough to inhibit slip, similar to the effect of the tetragonal distortion of *-AuCu. In fact, it goes much further than this. The movement of dislocations past an interstitial atom is greatly inhibited because it physically gets in the way of the necessary metal atom movement.
Fig. 2.3 Geometry of holes in the common metal crystal structures. The tetrahedral and octahedral holes occur between the close-packed planes of the h.c.p. and f.c.c. structures. The body-centred cubic hole is similar to the tetrahedral, but the tetrahedron is distorted by separating the atoms of two of the possible pairs slightly, making the hole a bit bigger. The triangular prism is not a normal crystal structure but is related to h.c.p. r here is the relative radius of the largest sphere that can fit in the space, compared with the parent atoms. In the right-hand colum, the front-most atoms on the left have been removed to show the fitting interstitial atom..
The maximum radius ratio of an assumed hard, spherical atom that can be fitted into an f.c.c. lattice without distorting it is ~0.41 in an ‘octahedral’ hole (Fig. 2.3). Carbon is the commonest important solute. Its atomic radius, -75 pm, is sufficiently small that it can be squeezed into the octahedral holes of the f.c.c. lattice of Co-Cr alloys, albeit with considerable distortion; the relative radius is about 0.60 that of the metal atoms involved. However, atoms are not hard spheres, they are slightly ‘squashy’ and their shapes also get distorted. Other non-metallic dissolved elements such as boron and silicon may also be used in certain alloys deliberately for similar effects, although these are also effective deoxidizers. Others, such as nitrogen and oxygen, may be unavoidable contaminants under ordinary processing conditions but which nevertheless add to hardness or reduce ductility (cf. amalgam alloy, 14§2.4, and titanium, Fig. 28§1.1) (but see §2.1). Note that substitution by such atoms is not possible because it destabilizes the lattice. Thus, the inclusion of some carbon can provide a substantial increase in the yield point of these Co-Crbased alloys. Even so, the maximum solubility of carbon is only of the order of 0.2 at%. Any excess results in the precipitation of carbides.
!2.4
Precipitated carbides More carbon may be dissolved in the melt than in the solid (see also Steel, Chap. 21), and in the freezing process various carbides may separate. The faster the cooling the smaller the amount, and the finer and more dispersed will be the carbide grains. Slow cooling, or annealing at high temperatures, will cause the precipitation of more carbide and growth of the grains already precipitated (Fig. 2.4).[4] Depending on the amount and distribution of these brittle carbides (e.g. whether intragranular or intergranular), the alloys may be relatively ductile or brittle, high or low strength. Large carbide grains are decidedly detrimental by providing longer continuous brittle paths for crack propagation. The majority of such alloys are therefore two-phase: " f.c.c. solid solution and carbide.
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The important distinction in these systems is indeed where the carbide is to be found. There are two kinds of site. The intragranular carbides harden the alloy because they inhibit dislocation movement. In contrast, the intergranular carbides embrittle the alloy because there are no dislocations to inhibit here. The grain boundary is the path that a crack would normally take, and the brittle and unbonded carbide makes it easier for crack propagation. Carbon is in fact the key element in these alloy systems and, Fig. 2.4 The range of types pattern of carbide precipitate that may be obtained as such, any variation in the amount in the " solid solution matrix according to the thermal history - time and will result in changes in properties temperature - of the casting. Note the variety, and the fact that while these which are almost certainly changes all affect strength and ductility, none is reversible. detrimental. Thus, both oxidizing and carburizing conditions, whereby carbon is removed or added (for example, from melting with an oxyacetylene flame), will clearly be disadvantageous. In fact, alloys intended for oxyacetylene flame melting are deliberately supplied with a lower carbon content than might be expected because of the tendency to pick it up. Conversely, alloys for induction melting tend to have slightly more carbon than finally desired because of the tendency to lose it. In like manner, remelting such alloys from previous sprues and buttons is not recommended because of the likely loss of carbon. Conversely, graphite crucibles or carbon-based investment deoxidizers again are not to be used. While special heat treatment is possible for some alloys to redistribute or modify the carbides (Fig. 2.3), in general the properties are so sensitive to such changes, and to the exact initial conditions, that it is not recommended that it be attempted for dental work. Welding or soldering is particularly risky because of the localized heating and consequent unevenly distributed changes to be expected. Residual carbon in the investment from an incomplete burn-out could also be dissolved in the casting surface, making attention to that particular detail especially important. It should be noted that alloys containing precipitated phases, whether metallic or carbides, have composite structures; their mechanical properties may be understood in part from this fact alone.
!2.5
Sigma phase Extended heat treatment or mixing of these "-phase alloys is not recommended for another reason. It can be seen in Fig. 2.1 that a second phase has been marked for notice: the F-phase. This phase has a complicated tetragonal structure (nominally FeCr, c/a ~ 0.52, 30 atoms to the cell) and nucleates readily at grain boundaries. Its significance, not having close-packed slip planes, and in marked contrast to the parent f.c.c. structure, is that it is in fact extremely brittle and would cause the rapid failure of the appliance if present. Some dental alloys have compositions that lie very close to the " / " + F phase fields common boundary. Under some circumstances, the right conditions for F-phase formation may arise. This risk is especially increased if the composition of the alloy is altered by mixing different products. The problem also arises for repeatedly melted metal, when the effects of the unavoidable differential oxidation will be felt. These effects are in addition to the problems of the loss or gain of carbon. Notice that oxides and so on that might be adhering to old, used metal will segregate on melting – they will not dissolve or get suspended in the melt (18§2.1). In any case, common sense dictates that metal to be reused would be cleaned up by sand-blasting first, although the risk of composition changes remains. It is the same problem that precludes stainless steel (21§2) from use as a casting alloy. There is a very high risk of F-phase forming, and consequent embrittlement, unless the melt is quenched very quickly indeed (working is impossible and further heat treatments are impractical). Clearly, this cannot be done for the lost-wax process as employed in dentistry (Chap. 18), even though this is considered to result in “rapid” cooling.
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The ‘base metal’ alloys are generally understood to be cheaper than the gold-, silver- and platinumcontaining ‘noble’ alloys. This is a major selling point. However, it should be noted that the higher casting temperatures required demand the use of specialized equipment (furnaces and casting machines) as well as investments capable of working under these more demanding conditions. This can make the total cost of working with such alloys comparable with that for the ‘dearer’ alloys. Yield point, modulus of elasticity and density must also be taken into account when selecting an alloy for a given application.
!2.6
Grain refining The chromium-based and related dental casting alloys tend to form rather large grains despite the normally fairly rapid cooling encountered in dental casting. Grain refining is thus essential to attain the best properties in the cast metal, especially in thin sections and narrow clasps (see §1.9, §1.10). Molybdenum (Mo) confers this on an alloy of the present type, and this can be seen to be related to its relatively low solid-solubility, which in turn arises from its large radius ratio. Beryllium also has this desirable effect, as well as tending to lower the liquidus of the alloy by as much as 100 K, and of lowering the surface tension of the melt and thus improving the castability, when present in small quantities. Unfortunately, beryllium is a very toxic metal (28§7) and is present in the fumes produced by the casting process and in the dust arising from the cutting, grinding and polishing of its alloys. The special fume-extraction equipment and other precautions required to prevent harm to laboratory technicians working with it has meant that it is now less-used and alternative alloys have been found. The available sources of hardening for dental casting alloys of the two major classes are summarized in Table 2.1. These approaches are also applicable in principle to any other alloy system, in existence or yet to be invented. Whether any particular one or combination is used depends on a number of factors: alloy composition, device purpose, processing conditions. Yet the primary purpose always has to be the inhibition of dislocation movement – slip, whether by raising the activation energy for movement or by shortening the available path.
Table 2.1
Possible means of hardening dental casting alloys.
Source Solid solution Rapid freezing Grain refining Metallic phase ppt. Tetragonal distortion Coherency strain Interstitials (C) Carbide Cold work **
Gold alloys
Chromium alloys
U U U U U U Y Y
U U U U* Y Y U U
!
!
* risk of corrosion ** not practicable; risk factor
____________________________ References [1]
Wyatt OH & Dew-Hughes D. Metals, Ceramics and Polymers. An Introduction to the Structure and Properties of Engineering Materials. Cambridge UP, 1974.
[2]
Brandes EA & Brook GB (eds). Smithells Metals Reference Book. 7th ed. Oxford, Butterworth-Heinemann, 1992.
[3]
Lyman T (ed). Metals Handbook. Cleveland : American Society for Metals, 1948
[4]
Lyman T (ed). Metals Handbook Vol. 8. Amer. Soc. for Metals, Ohio, 1973.
Casting Alloys
The range of alloys that may be used in the oral environment is limited primarily by the need to avoid corrosion. This leads to the use of two major groups: corrosion-resistant or precious metal alloys, and non-precious but passive alloys. The first is exemplified by gold-based products, the second by cobalt-chromium and similar alloys. Each has advantages and disadvantages as well as special handling considerations that must be taken into account in order to obtain the desired outcome. Corrosion resistance also (usually) requires single (metallic) phase alloys, yet sufficient hardening and strengthening must be obtained for the cast devices to function without permanent deformation under service stresses. Detailed consideration is therefore given to all available mechanisms for hardening, and their appropriateness in the two groups of alloys, in order that their design and handling be understood. In the case of some gold and related alloys the hardening can be reversible, depending on the crystal structure, and controlled by the thermal treatment. Composition is critical in determining mechanical properties, and especially in the cobalt-chromium types. This arises because the composition may lie dangerously close to that at which a brittle structure is obtained. Thermal history is also critical in that carbide grain growth is affected; this can either strengthen or weaken the metal. The selection of alloys for cast devices must be based on a knowledge of the properties of those alloys and their sensitivities to procedural and handling variables. Failure to recognize their individual limitations will result in treatment failure, for it is still a matter of compromise in balancing good and bad aspects of behaviour.
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Whilst the casting process has a long history, and the skills needed for successful casting can be learned by trial and error, it is essentially the properties of the alloys that control the outcome. The market these days is characterized by an enormous and potentially bewildering array of very similar products, even from a single manufacturer. To begin to comprehend the reasons for design choices it is necessary to consider some of the alloy systems in greater detail. However, it is also necessary to point out that in many cases the alloys themselves have often been designed by trial and error, and a comprehensive explanation of their behaviour in suitably scientific terms is simply not possible because the essential data have not been gathered. Nevertheless, it is not too difficult to outline the generic properties of the major classes. When choices of casting alloy are to be made it is primarily the application that provides the guide in that strength, stiffness, density (as it affects the mass of a suitably strong or stiff framework, for example), ductility, suitability for porcelain bonding, and so on must first be considered. A second level of importance is that of the experience of the dental laboratory in handling that alloy. As has already been made plain, the outcome of the entire casting process depends on the combination of investment and alloy, with all the many handling and processing variables that this entails. But, despite the theoretical matching of these factors, it still requires a demonstration that the entire system works to give an acceptable outcome. Part of the system is the skill of the technician in running a process consistently, and being able to calibrate it, as it were, to tune critical steps to ensure success.
§1. Gold Alloys
!1.1
Pure gold Pure gold is the outstanding restorative material. Because of its peculiar combination of properties: tarnish resistance, ductility with work hardening, and the ability to be cold-welded by pressure alone, it is usable as a direct filling material in the form of foil or powder (28§4) (although the skill, effort and expense of this process has led to it being more or less abandoned now). Aided by the deep-rooted and ancient mystical associations for the metal, its colour does not present aesthetic problems for many people, despite the great contrast with tooth material. As has been implied elsewhere (Chap. 24), the patient’s perceptions are at least as important as more fundamental criteria for material selection. However, in its pure state, gold is not very strong for applications other than inlays, i.e. for crowns, bridges and removable partial dentures. The strength advantages of alloys are required, but these require casting rather than direct fabrication techniques. This, of course, precludes work hardening as a means of increasing rigidity because dimensional stability requirements are severe if fit is to be maintained. Recourse to hardening and strengthening by suitable alloying elements must be made. Solid solution hardening The disorder introduced into crystal structures by elements in solid solution has previously been discussed in terms of the ease or otherwise of slip along certain planes of atoms (11§5). The effect is only moderate when there is a close match of atomic radii, as is exemplified by the Ag-Au system (radius ratio 144/143 = -1.01) (Fig. 1.1). But even so, at an atomic ratio of -1:1, when the disorder might be expected to be at its greatest, there is an appreciable peak in the strength. By way of contrast, with Al-Cu "-solid solution alloys the effect is much more
Fig. 1.1 Solid solution hardening in Ag-Au. The peak occurs near an Ag:Au ratio of 1:1.
!1.2
Fig. 1.2 Solid solution hardening in Al-Cu.
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pronounced because the radius ratio (143/128 = 1.12) is very close to the general practical limit for a continuous solid solution, and the distortion is much greater (Fig. 1.2).[1] Indeed, the solubility limit for Cu in Al is 2.5 at% at 548 °C. The ratio for Au-Cu (144/128 = 1.13) is also close to that limit, and while a continuous solid solution is formed above -400 °C, i.e. at all compositions, distinct new phases in fact form below that temperature (Fig. 1.3).[2] Of particular interest are those corresponding to stoichiometries around AuCu3 and AuCu, but it is the latter which is the phase of principal interest in dentistry. A peak in solid-solution hardening is again observed at this atomic ratio (i.e. 1:1).
Fig. 1.3 The Au-Cu system equilibrium diagram and some associated crystal structures. The narrow gap between liquidus and solidus is important for avoiding appreciable coring in these and similar alloys when cast as there will be little time for segregation.
Fig. 1.4 Unit cells for the various phases of the Au-Cu system. Notice that the height (c-axis dimension) of the *-phase is shorter than the a and b dimensions.
The Au-Cu constitutional diagram has a feature known as a solution minimum at about 20 mass% Cu. Here the reaction on heating is " Y Liq that is, the solidus and liquidus coincide at this point. It is therefore a type of point reaction. Both liquidus and solidus have horizontal tangents at this point. This reaction should be distinguished from that of the melting of a pure component (see box, 12§1.2).
The crystal structure for the whole Au-Cu system is essentially f.c.c. (11§3) and, in common with many other solid solutions, the "-phase is one of random substitution (Figs 1.4a, 1.5a). That is, at any point in the lattice of an "-phase alloy of composition AuCu, the probability of finding either kind of atom is exactly 0.5.
!1.3
Ordered phases In *-phase AuCu the structure is ordered. The Au and Cu atoms alternate regularly, which is equivalent to alternating layers of Au and Cu atoms (Figs 1.4b, 1.5b), in a pattern that repeats over long distances. In contrast, a random substitutional solid solution does not have long-range order for the occupancy (element identity) of each lattice site, although the probability of finding a particular kind of atom is fixed at the overall atomic proportion for the phase. Such a structure arises because the layers fit together just a little more compactly, increasing the regularity of the structure, thus lowering its energy. This rearrangement results in a slight distortion from a perfect cubic lattice to a tetragonal one, in which one of the three unit cell dimensions is smaller than the other two (the vertical direction in Fig. 1.4b), the ratio being about 0.935. This structure can be seen to consist in effect of two interlaced lattices, in this case of identical pattern: one of Cu only, the other of Au only (Fig. 1.6). This type of structure is known as a superlattice and is characteristic of ordered solid solutions. Alloys of the composition AuCu3 produce a similar ordered,
Fig. 1.5 Example planes in (a) random and (b) ordered 50:50 solid solution crystal structures.
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low-temperature phase ($, Fig. 1.4c) in which the Au atoms are placed at the corners of the f.c.c. lattice and the Cu at the faces. However, the lattice constants, the unit cell dimensions, are not appreciably different from those of the "-phase in this case, and this phase is of no particular dental interest. It can be seen that the close-packed planes of the f.c.c. structure (albeit with some irregularity because of the random substitutions of the solid solution) are taken further from ideality by the tetragonal distortion. Slip is therefore more difficult because the principal close-packed ‘planes’ (normal to the cube diagonals, Fig. 11§3.23) are not quite planar. This means that for this reason alone * is harder than "-phase at the same composition, an effect that does not operate for $-phase. The significance of the ordered solid solution *-phase, and its importance to dentistry, lies in the fact that it produces a structure markedly harder than just "-phase. While this makes such an alloy much more suitable for use in the dental context, it also allows a special kind of heat-treatment to be used to develop the degree of hardness required for any particular application or, conversely, to permit softening for adjustment which is then followed by rehardening. This process is enhanced by a further source of hardening.
!1.4
Fig. 1.6 A superlattice appears as if there are two separate but intermeshed lattices. This is ordered AuCu, and the two lattices may be described as “ base-centred tetragonal”.
Coherency strain hardening If an alloy with the composition AuCu is cast, rapid cooling or quenching to around room temperature will preserve the high-temperature, random, solid solution structure of the "-phase. There will be little segregation (12§1) because the liquidus and solidus are very close together. But if the temperature is then raised again and held below that of the upper limit of the boundary of the *-phase field (at, say, about 400°C), equilibration will now occur by diffusion, and the superlattice structure will tend to form. The higher the temperature the greater the diffusion rate, because diffusion is a process requiring an activation energy. But because there is a change in unit cell dimensions, notably a reduction in the direction of the c-axis, with concomitant slight increases in the other two directions, the transformation produces a strain in the lattice. The initiation of the " to * phase change will take place at a large number of separate locations within any grain randomly and independently. This will necessarily be with random orientation of the tetragonal c-axis along any of the three axes of the original cubic structure. These regions of differing c-axis orientation within a given grain are called domains. Taken over a large enough region the net strain in each grain will average out to zero, and there will be very little volume change overall. But the variation from one orientation to another from place to place within the grain superimposes on those primary grains a microscopic lamellar or granular structure, and it is this that makes the heat-treated material hard. There are now effectively very many more grain boundaries, i.e. between domains, so dislocations travel much shorter distances before their movement is inhibited. There is less slip. The formation of $-phase does not cause a similar hardening because there is no associated axial dimensional change to induce strain. The *-phase hardening process is entirely reversible: reheating above ~400 °C again produces the "-phase, which may again be preserved by quenching. From another point of view, we can see that it is the periodicity of the structure in the *-phase which is varying. The crystal lattice does not repeat indefinitely in any direction, but every so often skips a step, as it were. The boundaries between areas of different phase (in this wave periodic sense, not the
Fig. 1.7 An atom map showing how domains result in coherency strain in *-AuCu (‘atoms’ roughly to scale).
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structural-constitutional sense) are strained, since the spacing between layers and along rows differs slightly across a domain boundary. This then is known as coherency strain. Slip through such boundaries is therefore more difficult, that is, there is coherency strain hardening. Fig. 1.7 shows what a random section through the junction between four such domains might look like. Since the spacing of the atoms in each kind of domain is different from that in its neighbour, the mismatch generates the strain. Although *-phase is expected to be a little harder because of the tetragonal distortion, on top of what would be expected from the substitional effects, this on its own is not very great. Most of the hardening is due to the coherency strain; the superlattice by itself does not confer the observed change, which is largely due to the domains forming with random orientations. Current gold alloys for dentistry may consist of 6 or more major constituents: Au, Cu, Ag, Pt, Pd, Zn; several other minor additions may also be made. Clearly, the multi-dimensional phase diagram for such a system is going to be extremely complicated, and full details are not at present known for any such system. Detailed explanation of the purpose and effects of any single alloying element in such contexts cannot be given, even though sometimes rather vague statements may be made. However, certain aspects and broad principles can be summarized, and these are applicable no matter which alloy is under consideration.
!1.5
Precipitation hardening When a second phase is allowed to separate from (usually) a high-temperature solid solution (see Fig. 12§3.9), those new precipitated crystals form randomly throughout the original grains of the alloy. They therefore get in the way of dislocation movement, limiting slip. Such alloys are described as being precipitation hardened. The example of the Ag-Cu system has already been discussed (Fig. 12§3.1). There, for example, the separation of the $-phase (solid solution of Ag in Cu) from a high temperature, Cu-rich "-phase (solid solution of Cu in Ag) at a temperature below the solvus results in just such a very fine-grained precipitate. Conversely, a heat treatment above the solvus will redissolve the precipitate to form a solid solution again. Reactions similar to these will play a part in the overall hardenability of such alloys. Time and temperature are the two important aspects of any heat-treatment: to determine the extent of any change, and the direction and rate of that change. A continuous solid solution characterizes Ag-Au alloys (Fig. 12§1.1) as it does Au-Cu at high temperature, but in the ternary system Ag-Au-Cu below ~400 °C this is restricted to the very edges of the diagram (Fig. 1.8). (We shall ignore for the moment the formation of $ and *-phase Au-Cu superlattices.) However, as the temperature is raised the gold-rich boundary of the two-phase field (Au,Ag) + (Au,Cu) moves steadily towards the Ag-Cu boundary; at the same time the Ag- and Cu-rich solvi move away from the corresponding pure metal corners (Fig. 1.9). The two-phase field can be visualized as a dome-shaped volume in the full phase diagram (Fig. 1.10). Note that the faces of the prism are the three, two-component phase diagrams (cf. 14§3.7).
Fig. 1.8 The solid solubility limit in Ag-Au-Cu alloys at 371°C. Compare Fig. 1.3, and Figs 12§1.1, 12§3.1. (The detail of the ordered phases is omitted).
Fig. 1.9 Several solid solubility isotherms of the Ag-Au-Cu system mapped on the same diagram. This is a plan view of the triangular prism of Fig. 1.10; the isotherms are therefore equivalent to the height contours of that surface.
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It should be clear from this that an alloy with a composition corresponding to the two-phase field (i.e. at lowtemperature equilibrium) will, if quenched from the melt, as in a normal casting, maintain the high-temperature, single-phase solid-solution structure. Subsequent ageing at temperatures below the single-phase limit will cause the decomposition of the solid solution into two new solid solution phases.
!1.6
Ternary system phase-composition The compositions of the conjugate phases present in the Ag-Au-Cu system cannot in fact be stated as the necessary analytical work does not yet seem to have been done. As explained in 8§4.2, this cannot be done in ternary systems simply by reference to the phase diagram, using tie-lines as is the case for binary systems; it is necessary to analyse each pair for a well-equilibrated alloy. Thus, unless the position of the plait point is known, and its locus with change in temperature on the single phase field boundary surface (Fig. 1.10), it is impossible even to begin to discuss the compositions of the two phases in an aged Ag-Au-Cu alloy, apart from the simple expectation that one is copper-rich and the other is silver-rich. Even so, we can understand aspects of the behaviour of the system from what we do know.
Fig. 1.10 A view of the three-dimensional reconstruction of the solid solubility limit surface of the Ag-Au-Cu system.
What we can say is that on cooling from the singlephase field, a second phase will precipitate in the matrix of the parent, or more generally a phase-separation (8§4.1) akin to that of eutectoid formation (Fig. 1.11), a minimum diffusion path to establish the appropriate compositions in least time. Of course, grain growth will occur if allowed to remain at a high-enough temperature.
!1.7
Reversible hardening Thus, the hardenability of dental gold alloys is probably mostly due to the decomposition of the hightemperature solid solution into a fine-grain two-phase structure. The effect can be traced in Fig. 1.12, which shows the Vickers Hardness (1§8.2) for alloys in three conditions: (1) alloys that have been quenched after a solution heat treatment. This involves holding the metal above the temperature at which a single solid solution begins to be stable for long enough that diffusion can occur to re-equilibrate the structure, and so followed by quenching, rapid cooling to room temperature such as by sudden immersion in water while hot.
Fig. 1.11 Section through Fig. 1.10 at 60% Au. Note that isothermal tie-lines cannot be drawn in the "1 + "2 field of a pseudo-binary diagram of this kind – in general they lie at an angle to the plane of the page (cf. 8§4.2)
(2) alloys that have been annealed in the two-phase region (aged), similarly to re-equilibrate the structure. This age hardening is perhaps more accurately described as precipitation hardening. (3) those that have been allowed to air-cool after casting, which represents an intermediate cooling rate through the transition temperature.
Fig. 1.12 Variation in indentation hardness as a function of composition for various heat treatments over three constant Au-content sections of the Ag-Au-Cu system. (See Fig. 1.13)
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The effect of varying the gold content is shown. The quenched alloys show only solid solution hardening, while the aged alloys show the effect of the two-phase structure in addition. The air-cooled condition, i.e. which was slow enough to allow some development of the two-phase structure, lies somewhere between. In addition, the hardening effect of the formation of *-phase AuCu is clearly seen at 75%Au (Figs 1.12, 1.13). Available evidence suggests that even though a third element (Ag) is present, superlattice structures of the AuCu-type may still form, and even when the Ag- and Curich two phases are present. The hardenability of the dental Ag-Au-Cu alloys is therefore due to solid solution, precipitation and superlattice effects occurring together. Indeed, this is also known to be the case with Ag-Au-Cu-Pt alloys (cf. Figs 1.15 and 1.21 with Fig. 1.3). Alloy colour The addition of Ag to Au-Cu alloys provides the benefit of precipitation hardening, but it also allows a certain amount of variation in the colour of the alloy (Fig. 1.14). This can be used to offset the reddening effect of the presence of the Cu. If ‘white gold’ is required, it is more likely to be alloyed with Pd, at about 5%, which has quite a strong effect; Pt has only a slight effect. (Ni is a particularly powerful ‘whitener’ of gold alloys for jewellery but because of its known hypersensitizing action is likely to be less used in dentistry). Obviously, such colour adjustments are not concerned with matching tooth tissue, but some people may find one colour more pleasing than another, or there may be a need to match the appearance of two alloys which have different mechanical properties because of different functional requirements.
Fig. 1.13 The positions of the sections of the Ag-Au-Cu system shown in Fig. 1.12.
!1.8
Fig. 1.14 Alloy colour vs. composition in the Ag-Au-Cu system. The names are those conventionally used in the jewellery trade. Silver whitens, copper reddens the alloy.
!1.9
Grain size The Hall-Petch relationship has previously been mentioned (11§4.1) as indicating the effect of decreasing grain size on the strength. Grain size is a function of the cooling rate: faster cooling gives more and smaller crystals. In addition, perfect crystals require time to grow. Rapid cooling does not permit the annealing out of spontaneous stacking faults. Errors will be covered over and locked into the structure by the solidifying metal on top. Thus, the faster the cooling, the more stacking faults there will be, in parallel with the formation of smaller grains. Stacking faults, i.e. dislocations, create strain fields that inhibit slip. There will therefore be some hardening effect simply due to rapid cooling for two reasons: the Hall-Petch effect and stacking faults. Indeed, it could be argued that slow cooling is detrimental. Thus, the investment mould temperature must not be too high to avoid inhibiting solidification too much.
In some countries, “white gold” is used as the name for platinum and platinum alloys (whether or not they actually contain any gold), and even sometimes for rhodium and its alloys. This has to do with the unfamiliarity of platinum as a jewellery metal, and the higher status afforded the word “gold”. Care is therefore necessary in some contexts to be very precise about what is being discussed.
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!1.10 Grain-refining Grain size is a function of the cooling rate, but it does vary from alloy to alloy because the degree of supercooling required for homogeneous nucleation varies, and so the number of nuclei formed under any given set of conditions will also vary. However, this aspect can also be controlled deliberately by the addition of so-called grain-refining alloying elements. In dental gold alloys this is now usually done with iridium (Ir) (although Ru, Re and Rh are also reported). When there is a very large difference between the melting points of two metals, their binary phase diagram tends to show a very large gap between the liquidus and solidus (Au-Pt is one such example which has been studied, Fig. 1.15), in great contrast to the Au-Cu system (Fig. 1.3). In other words, the solubility of the grain-refining element is low. This in fact means that the nucleation of grains occurs very readily. Referring to Fig. 1.15, it can be seen that even at 80% Au at 1350 °C the first solid to form is about 75% Pt.
Fig. 1.15 The equilibrium diagram for the system Au-Pt. Such a diagram indicates a strong tendency to coring - note the great width of the " + Liq field above about 1350/C.
Iridium, which has a melting point of 2454 °C, shows such an effect, only much more exaggeratedly. Its solubility in gold is very low (less than 0.005% by mass) and so can be expected to readily precipitate from the melt. Indeed, it is said that there is always this solid present at the point of casting. It would therefore initiate heterogeneous nucleation (11§2.5) of the Au-alloy, for which process the activation energy is much smaller than for homogeneous nucleation. In other words, the degree of super-cooling required is very small. Grain-refining thus ensures that the distance a dislocation can travel is small, thereby adding to the hardening of the alloy. In passing, we can note that the phase diagram for Au-Pt has a two-phase field of two solid solutions ("1 + "2) enclosed within the general solid solution field which is otherwise continuous (i.e. stretching between the two terminals). Such a system is exactly parallel to the two-liquid systems discussed in 8§4.1, as it shows an upper critical solution point (cf. Fig. 8§4.1), and similar to the two-phase field shown in Fig. 1.11. Thus, this system would also show phase-separation on cooling into that region. Au-Pt is the basis of some alloys used for porcelain-fused-to-metal devices (25§6).
Fig. 1.16 The effect of work-hardening on the Brinell indentation hardness in the system Ag-Pd.
!1.11 Work hardening Although work hardening is important in several metals’ contexts, it is not an option applicable to cast alloys in general – the deformation involved clearly would spoil the dimensional accuracy of any such device. Occasionally, however, work hardening is significant. Thus, in the adjustment of clasps to compensate for residual errors in the casting, the deformation will inevitably harden and therefore embrittle the metal. There is then an increased risk of fracture on any further adjustment. Bear in mind that manual adjustment is in effect a ‘strain-controlled’ action (as opposed to loadcontrolled) (1§3.1) since the target outcome is a certain amount of deformation, but in doing this by hand it is
Fig. 1.17 The equilibrium diagram for the system Au-Pd. The narrow gap between liquidus and solidus suggests little coring in these and similar alloys when cast.
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rather difficult to control the amount. In practice this means that bending should be done in one direction only, gradually approaching the correct shape. Avoidance of work hardening is a good reason for using a reversibly hardenable alloy. But work hardening can also be beneficial when the margins of a crown or inlay have to be burnished, i.e. made to conform very accurately to the adjacent tooth. The hardening limits the possibility of deformation later in service.
!1.12 Related alloys There are, as pointed out above, many other alloying elements possible for dental ‘golds’, and some indeed form alloy systems of their own that have found application in dentistry. We may briefly explore a few examples to illustrate further the general principles.
Fig. 1.18 Variation in Vickers indentation hardness with composition in the system Ag-Au-Pd (annealed).
The Ag-Pd system is one of simple solid solution over the whole composition range (Fig. 12§1.2) (radius ratio: 1.05). No further hardening after the effects of solid-solution is possible except from cold-work, which annealing can then remove. After annealing, such alloys are quite weak and soft (Fig. 1.16), although they still show very clearly the general effect of hardening by solid solution disorder (cf. Ag-Au, Fig. 1.1). Unfortunately, these alloys have a strong tendency when molten to dissolve oxygen, which then appears as disseminated porosity in the cast metal (Fig. 18§4.4). Au-Pd also forms a continuous solid solution (Fig. 1.17) (radius ratio also 1.05) and the ternary system Ag-Au-Pd is similarly one of a continuous solid solution. Although some variation in hardness is apparent in that system (Fig. 1.18), although with a complicated pattern, all of these alloys are still relatively soft and not amenable to age or precipitation hardening.
Fig. 1.19 The equilibrium phase diagram for the Cu-Pd system. The formation of ordered phases is comparable to that in the Au-Cu system.
The Cu-Pd system (radius ratio 1.07) shows superlattice formation (Fig. 1.19), just as in Au-Cu (although not exactly at the expected 50-50 composition). As a result the Ag-Cu-Pd system shows essentially the same features as does Ag-Au-Cu, with age-hardening due to both CuPd superlattice formation and a precipitation of two solid solution phases. Substantial improvement in mechanical properties can be made this way (Fig. 1.20). Notice in particular the peak hardness near the centre of the diagram where most disorder would be expected. The Cu-Pt system (radius ratio 1.09) also shows superlattice formation (Fig. 1.21), and agehardening due to the reversible formation of CuPt is possible.
Fig. 1.20 Variation in Vickers indentation hardness with composition in the system Ag-Cu-Pd (annealed).
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!1.13 General considerations What all this illustrates is that there is considerable scope for the design of alloys, even using metals that are very closely related and have very similar chemical properties. The subtlety and complexity of these systems defeats any attempt at memorizing property values, or even compositions. In fact, this is quite pointless since small changes may cause profound changes in behaviour. It does emphasise the need to consider manufacturers’ data and instructions very carefully when making product selections, or changing product. One such aspect of all of these alloys that must be taken into account in casting them is that the liquidus temperature may vary considerably with composition. Some illustrative examples are given in Figs 1.22 - 1.24.[3] Care must be taken that the alloy is heated sufficiently above the liquidus that casting is successful. Different commercial alloys may show appreciable variation in the casting temperature due to the inclusion of further alloying elements and this must be checked. A second point is the melting range, the vertical (temperature) distance between the liquidus and solidus. The wider this gap the greater the likelihood of coring, the first freezing metal having a grossly different composition to that freezing later (12§1). Homogenization annealing then becomes necessary. This is particularly true when platinum forms an appreciable proportion of gold alloys (cf. Fig. 1.15); it is not normally present to more than about 4 mass% because of the tendency to coring that it conveys. Two-phase alloy systems are likely to be less corrosion-resistant than are single-phase alloys, although both phases may be sufficiently noble or electronegative for this to be quite unimportant in alloys with high enough content of elements such Au and Pt. Even so, the requirements of mechanical, chemical (i.e. corrosion) and handling (i.e. casting) properties need to be carefully balanced for optimum clinical results. But with the emphasis turning more and more to alloys of lower intrinsic cost, the difficulties of maintaining corrosion resistance will increase, and polyphase structures will be more prone to suffer in this way. Gold alloys suffer from sensitivity to the market price of the gold content, but they are also handicapped by limited strength and modulus of elasticity compared with the demands made on removable appliances. Other difficulties include their high density, which thus causes high weights for devices, possibly causing problems in stability and comfort. These kinds of factor justify consideration of other alloy systems.
Fig. 1.21 The equilibrium phase diagram for the system CuPt. Again, the formation of ordered phases is comparable to that in the Au-Cu system.
Fig. 1.22 The liquidus temperature contour map for the system Ag-Au-Pd.
Fig. 1.23 The liquidus temperature contour map for the system Ag-Au-Pd.
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!1.14 Scavengers Because some of the alloying elements of gold alloys are oxidizable in air, copper in particular, and such oxides may form inclusions in the cast metal which have effects such as acting as flaws and embrittling the alloy, there is a need to limit this action with a deoxidizer. Zinc, and sometimes indium (In), may be added to the original formulation to provide a readily-reacting scavenger of dissolved oxygen. The resulting oxides then appear as part of the dross in the crucible. The process and outcome are similar to those in the case of amalgam alloys (14§2.4). Of course, a scavenger is meant to be consumed, and such that little is intended to be present in the actual cast alloy (again, cf. Zn in silver amalgam alloy). Re-use of old buttons is problematic in that the composition must have changed by virtue of such oxidation, and not just by loss of scavengers: Sn, Ga, Si and others are also susceptible to reaction or volatilization.
Fig. 1.24 The liquidus temperature contour map for the system Ag-Cu-Pd.
§2. Chromium Alloys A series of alloys, originally developed for high strength at elevated temperatures, called stellites, have provided an alternative to gold-based alloys for dental castings. Chromium is the nominal principal constituent of the majority of these so-called ‘base metal’ alloys. Like aluminium, it is an extremely reactive metal that forms a tenacious, thin (thin enough to be colourless; Cr2O3 is green in bulk), and impermeable film of oxide which limits further reaction, making it essentially non-tarnishing in air and many aqueous media of near neutral pH. It possesses the further remarkable advantage of conferring these properties on alloys which contain an appreciable proportion of it. This ready production of an oxide film is called passivation, and it has obvious interest in dentistry (13§5). Despite this usefulness, dental alloys are not structurally based on chromium (which is b.c.c., Fig. 11§3.3); Ni, Co and Fe form the greatest proportions of the elements present in commercial products. This is because extensive solid solutions ("-phase) are formed in the Co-Ni and Co-Fe systems at high temperatures; these have the austenitic structure, i.e. f.c.c. Only limited solid-solubility is shown in other systems, such as Cr-Ni and Co-Cr itself. Even so, in the ternary and quaternary systems, the resulting f.c.c. solid solutions have very wide ranges of possible compositions (Figs 2.1, 2.2). That this occurs may be readily understood from the atomic radii of the metals in such alloys (in pm): Fe: 126, Co: 125, Ni: 125, Cr: 129, Mo: 140. The radius ratios for the first four elements do not exceed ~1.03, and Mo tends to be used up to only about 5 mass%. Curiously enough, because these are all f.c.c. structures, the mechanical properties of ductility, softness and low strength are common to all such alloys with relatively little variation – as they stand they are still in fact unsuitable for dentistry. However, the intention is to create, if possible, an alloy with mechanical and chemical properties suitable to the task. The preference is therefore a single-phase alloy, to avoid galvanic effects, passivating, to reduce reaction rates and thus corrosion, while maximizing the disorder in the solid solution to increase the yield point. In order to make such alloys functional in the sense of carrying large stresses over long spans with small cross-sections (in other words, to fit in the mouth), further hardening of such alloys may be achieved in a number of ways. The differences, as well as the similarities, between this set of possibilities and that for gold-based alloys should be noted.
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Fig. 2.1 Parts of four of the ternary isothermal equilibrium phase diagrams of interest in connection with ‘chrome’ dental casting alloys. The features of concern are the extensive " solid solution regions, and the adjacent " + F phase field in each case.
Fig. 2.2 Two of the possible quaternary isothermal equilibrium phase diagrams for the chrome alloy series, showing only the "-solid solution regions (shaded).
!2.1
Other metallic alloying elements The effect of the disorder produced in a solid solution by atoms of different sizes in a lattice has already been discussed. Suitable elements in the present context include Mn (~127 pm) and W (~140 pm). It is important to notice that these alloying elements are required to dissolve in the "-phase solid solution to avoid galvanic corrosion. Manganese can also play a role as a deoxidizer (as does Si in these kinds of alloy). There are in fact many other alloying elements that are used in a bewildering variety of commercial products for dentistry, some of which products are predominantly one metal, some another. There are too many choices available. It is beyond the scope of this book to detail the possibilities and it is probably beyond the ability of any practitioner to understand the resulting plethora of property and behaviour data, usage options and prices. Accordingly, here – as for the similar situation with respect to gold and related alloys – only some broad principles are given.
!2.2
Metallic phase precipitation Aluminium is added to at least one dental alloy which contains much Ni; a fine-grained precipitate of the compound AlNi3 is formed which is believed to contribute greatly to the modulus of elasticity and strength of the alloy. Although this looks as though it may be a serious corrosion risk, the fineness of the precipitate may mean that it is soon dissolved when exposed at the surface, and no great roughness results. However, the
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frequency of occurrence of nickel sensitivity may be enough to preclude its use in biological contexts if only as a prudent precaution. The absence of serious mechanical problems does not adequately offset the risk.
!2.3
Dissolved carbon The solid solutions discussed earlier (§1, §2; 12§1) were substitutional, that is, an atom of one kind was replaced by one of another kind, in other words a substitution at the exact same position in the unit cell. However, there is a further kind of solid solution possible when the solute atom is sufficiently small, and that is interstitial solid solution. This possibility arises because in some lattices the packing of the layers leaves enough space to accommodate, with some distortion of the surroundings, atoms such as B, C, N and O (Fig. 2.3). That is, even in close-packed crystal structures there are necessarily holes – interstices – between the layers of atoms (11§3) (not to be confused with vacancies, 11§6.1). The resulting distortion of the crystal lattice itself is enough to inhibit slip, similar to the effect of the tetragonal distortion of *-AuCu. In fact, it goes much further than this. The movement of dislocations past an interstitial atom is greatly inhibited because it physically gets in the way of the necessary metal atom movement.
Fig. 2.3 Geometry of holes in the common metal crystal structures. The tetrahedral and octahedral holes occur between the close-packed planes of the h.c.p. and f.c.c. structures. The body-centred cubic hole is similar to the tetrahedral, but the tetrahedron is distorted by separating the atoms of two of the possible pairs slightly, making the hole a bit bigger. The triangular prism is not a normal crystal structure but is related to h.c.p. r here is the relative radius of the largest sphere that can fit in the space, compared with the parent atoms. In the right-hand colum, the front-most atoms on the left have been removed to show the fitting interstitial atom..
The maximum radius ratio of an assumed hard, spherical atom that can be fitted into an f.c.c. lattice without distorting it is ~0.41 in an ‘octahedral’ hole (Fig. 2.3). Carbon is the commonest important solute. Its atomic radius, -75 pm, is sufficiently small that it can be squeezed into the octahedral holes of the f.c.c. lattice of Co-Cr alloys, albeit with considerable distortion; the relative radius is about 0.60 that of the metal atoms involved. However, atoms are not hard spheres, they are slightly ‘squashy’ and their shapes also get distorted. Other non-metallic dissolved elements such as boron and silicon may also be used in certain alloys deliberately for similar effects, although these are also effective deoxidizers. Others, such as nitrogen and oxygen, may be unavoidable contaminants under ordinary processing conditions but which nevertheless add to hardness or reduce ductility (cf. amalgam alloy, 14§2.4, and titanium, Fig. 28§1.1) (but see §2.1). Note that substitution by such atoms is not possible because it destabilizes the lattice. Thus, the inclusion of some carbon can provide a substantial increase in the yield point of these Co-Crbased alloys. Even so, the maximum solubility of carbon is only of the order of 0.2 at%. Any excess results in the precipitation of carbides.
!2.4
Precipitated carbides More carbon may be dissolved in the melt than in the solid (see also Steel, Chap. 21), and in the freezing process various carbides may separate. The faster the cooling the smaller the amount, and the finer and more dispersed will be the carbide grains. Slow cooling, or annealing at high temperatures, will cause the precipitation of more carbide and growth of the grains already precipitated (Fig. 2.4).[4] Depending on the amount and distribution of these brittle carbides (e.g. whether intragranular or intergranular), the alloys may be relatively ductile or brittle, high or low strength. Large carbide grains are decidedly detrimental by providing longer continuous brittle paths for crack propagation. The majority of such alloys are therefore two-phase: " f.c.c. solid solution and carbide.
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The important distinction in these systems is indeed where the carbide is to be found. There are two kinds of site. The intragranular carbides harden the alloy because they inhibit dislocation movement. In contrast, the intergranular carbides embrittle the alloy because there are no dislocations to inhibit here. The grain boundary is the path that a crack would normally take, and the brittle and unbonded carbide makes it easier for crack propagation. Carbon is in fact the key element in these alloy systems and, Fig. 2.4 The range of types pattern of carbide precipitate that may be obtained as such, any variation in the amount in the " solid solution matrix according to the thermal history - time and will result in changes in properties temperature - of the casting. Note the variety, and the fact that while these which are almost certainly changes all affect strength and ductility, none is reversible. detrimental. Thus, both oxidizing and carburizing conditions, whereby carbon is removed or added (for example, from melting with an oxyacetylene flame), will clearly be disadvantageous. In fact, alloys intended for oxyacetylene flame melting are deliberately supplied with a lower carbon content than might be expected because of the tendency to pick it up. Conversely, alloys for induction melting tend to have slightly more carbon than finally desired because of the tendency to lose it. In like manner, remelting such alloys from previous sprues and buttons is not recommended because of the likely loss of carbon. Conversely, graphite crucibles or carbon-based investment deoxidizers again are not to be used. While special heat treatment is possible for some alloys to redistribute or modify the carbides (Fig. 2.3), in general the properties are so sensitive to such changes, and to the exact initial conditions, that it is not recommended that it be attempted for dental work. Welding or soldering is particularly risky because of the localized heating and consequent unevenly distributed changes to be expected. Residual carbon in the investment from an incomplete burn-out could also be dissolved in the casting surface, making attention to that particular detail especially important. It should be noted that alloys containing precipitated phases, whether metallic or carbides, have composite structures; their mechanical properties may be understood in part from this fact alone.
!2.5
Sigma phase Extended heat treatment or mixing of these "-phase alloys is not recommended for another reason. It can be seen in Fig. 2.1 that a second phase has been marked for notice: the F-phase. This phase has a complicated tetragonal structure (nominally FeCr, c/a ~ 0.52, 30 atoms to the cell) and nucleates readily at grain boundaries. Its significance, not having close-packed slip planes, and in marked contrast to the parent f.c.c. structure, is that it is in fact extremely brittle and would cause the rapid failure of the appliance if present. Some dental alloys have compositions that lie very close to the " / " + F phase fields common boundary. Under some circumstances, the right conditions for F-phase formation may arise. This risk is especially increased if the composition of the alloy is altered by mixing different products. The problem also arises for repeatedly melted metal, when the effects of the unavoidable differential oxidation will be felt. These effects are in addition to the problems of the loss or gain of carbon. Notice that oxides and so on that might be adhering to old, used metal will segregate on melting – they will not dissolve or get suspended in the melt (18§2.1). In any case, common sense dictates that metal to be reused would be cleaned up by sand-blasting first, although the risk of composition changes remains. It is the same problem that precludes stainless steel (21§2) from use as a casting alloy. There is a very high risk of F-phase forming, and consequent embrittlement, unless the melt is quenched very quickly indeed (working is impossible and further heat treatments are impractical). Clearly, this cannot be done for the lost-wax process as employed in dentistry (Chap. 18), even though this is considered to result in “rapid” cooling.
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The ‘base metal’ alloys are generally understood to be cheaper than the gold-, silver- and platinumcontaining ‘noble’ alloys. This is a major selling point. However, it should be noted that the higher casting temperatures required demand the use of specialized equipment (furnaces and casting machines) as well as investments capable of working under these more demanding conditions. This can make the total cost of working with such alloys comparable with that for the ‘dearer’ alloys. Yield point, modulus of elasticity and density must also be taken into account when selecting an alloy for a given application.
!2.6
Grain refining The chromium-based and related dental casting alloys tend to form rather large grains despite the normally fairly rapid cooling encountered in dental casting. Grain refining is thus essential to attain the best properties in the cast metal, especially in thin sections and narrow clasps (see §1.9, §1.10). Molybdenum (Mo) confers this on an alloy of the present type, and this can be seen to be related to its relatively low solid-solubility, which in turn arises from its large radius ratio. Beryllium also has this desirable effect, as well as tending to lower the liquidus of the alloy by as much as 100 K, and of lowering the surface tension of the melt and thus improving the castability, when present in small quantities. Unfortunately, beryllium is a very toxic metal (28§7) and is present in the fumes produced by the casting process and in the dust arising from the cutting, grinding and polishing of its alloys. The special fume-extraction equipment and other precautions required to prevent harm to laboratory technicians working with it has meant that it is now less-used and alternative alloys have been found. The available sources of hardening for dental casting alloys of the two major classes are summarized in Table 2.1. These approaches are also applicable in principle to any other alloy system, in existence or yet to be invented. Whether any particular one or combination is used depends on a number of factors: alloy composition, device purpose, processing conditions. Yet the primary purpose always has to be the inhibition of dislocation movement – slip, whether by raising the activation energy for movement or by shortening the available path.
Table 2.1
Possible means of hardening dental casting alloys.
Source Solid solution Rapid freezing Grain refining Metallic phase ppt. Tetragonal distortion Coherency strain Interstitials (C) Carbide Cold work **
Gold alloys
Chromium alloys
U U U U U U Y Y
U U U U* Y Y U U
!
!
* risk of corrosion ** not practicable; risk factor
____________________________ References [1]
Wyatt OH & Dew-Hughes D. Metals, Ceramics and Polymers. An Introduction to the Structure and Properties of Engineering Materials. Cambridge UP, 1974.
[2]
Brandes EA & Brook GB (eds). Smithells Metals Reference Book. 7th ed. Oxford, Butterworth-Heinemann, 1992.
[3]
Lyman T (ed). Metals Handbook. Cleveland : American Society for Metals, 1948
[4]
Lyman T (ed). Metals Handbook Vol. 8. Amer. Soc. for Metals, Ohio, 1973.