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Johnson Matthey Technol. Rev., 2021, 65, (4), 519


A Conflict of Fineness and Stability: Platinum- and Palladium-Based Bulk Metallic Glasses for Jewellery: Part II

Processing, tarnish resistance and future developments

  • O. S. Houghton*
  • A. L. Greer
  • Department of Materials Science and Metallurgy, University of Cambridge, 27 Charles Babbage Road, Cambridge, CB3 0FS, UK
  • *Email:

Article Synopsis

The properties and glass-forming ability (GFA) of platinum- and palladium-based bulk metallic glasses (BMGs) for jewellery were introduced in Part I of this two-part review (1). Here, we will describe methods for their processing, tarnishing and corrosion resistance and consider their prospects and future developments.

1. Processing: Injection Moulding Meets Metals

1.1 Fluxing and Melt Contamination

The initial report (2) of fluxing of liquid Pd40Ni40P20 with anhydrous liquid B2O3 showed that this treatment facilitates reliable production of fully glassy samples with larger cross-section. Although other agents such as soda-lime glass (3) have been used, B2O3 has remained the flux of choice for palladium-based and, later, platinum-based systems (415). While fluxing is an additional processing step, it is used because the effects can be dramatic: it can improve GFA, characterised as a reduction in Rc, by one or more orders of magnitude (6, 11, 16, 17). In a Pd-(Cu,Ni)-P BMG, fluxing increases the onset time at the nose of the time-temperature-transformation (TTT) curve for crystallisation in the liquid from 130 s to 200 s (17) and has been critical in achieving BMG castings over 50 mm in minimum linear dimension (18, 19).

Fluxing works by reducing the influence of heterogeneous nucleants that would facilitate crystallisation on cooling the melt. These nucleants are generally considered to be dispersed particles, mostly oxides, in the melt, but fluxing may also prevent contact with nucleants at the melt surface. Fluxing works by preferential wetting of the particles, possibly followed by dissolution of the particles in the flux. The removal of oxide particles from the melt would reduce its oxygen content. Direct verification of this is rare but, in one study, an initial oxygen content of ~12 ppm (by weight) was reduced to ~5 ppm by fluxing (8). Following the initial work (2), it is usual to cycle the melt between solid and liquid states several times; crystallisation in early cycles may assist in driving inclusions to the surface of the alloy, facilitating their take-up by the flux.

Ideally, fluxing would have no effect on the composition of the liquid alloy itself. In particular, when using B2O3, none of the elements in the alloy should have a greater ‘oxygen affinity’ (11), i.e. more stable oxide (as quantified by the relevant chemical potentials), than the boron in the flux. Alloys of the systems Pd-(Cu,Ni)-P and Pt-(Cu,Ni)-P meet this condition. But B2O3 has also been used to flux Pd-Cu-Si (8): in this case, reaction with silicon in the alloy can reduce B2O3, leading to dissolution of boron into the alloy. The addition of boron is detrimental to the GFA and the overall effects of fluxing are therefore complex, varying markedly with the time-temperature profile used in the treatment (8, 11). Nevertheless, by optimising the fluxing of Pd-Cu-Si, its dc can be substantially improved to 11 mm (20) and 15 mm (8). Studies of Pd-Ni-Si-P alloys suggest that composition changes inducing by fluxing with B2O3 can in some cases be beneficial for GFA (7, 17, 21) and lead to improvements in plasticity (22).

The remarkable improvements in GFA for ternary BMGs suggest that the fluxing should be considered in greater detail. The high m-fragility of, and the boron content within, Demetriou et al.’s 950Pt based BMG suggests a sensitivity to fluxing and boron pick-up (23), so enhanced fluxing procedures should be sought. Perhaps the effects of fluxing and overheating (to dissolve preexisting structures in the melt) should be revisited for alloys where it has been previously reported that fluxing does not affect GFA. It would be good to explore the use of alternative fluxes. B2O3 has been very widely used, not least because it remains liquid below the Tg of the alloys being processed. It may be possible to find multicomponent fluxes that remain liquid to similar temperatures, yet have greater resistance to chemical reduction by reaction with the alloy being processed.

In the literature, the action of fluxing has been described as ‘purification’ of the liquid alloy. While fluxing may clean an alloy by removing dispersed particles, there is no evidence that it can reduce levels of dissolved impurities. In contrast, as noted above, the alloy can be contaminated by reaction with the flux. In the production of BMGs, impurities can be detrimental to GFA. The most prominent example is the presence of oxygen in zirconium-based alloys, as briefly reviewed elsewhere (11). It is important to start with raw materials of ultra-low oxygen content and care must always be taken to minimise oxygen pick-up during processing. Various approaches have been adopted to reduce the oxygen content; these include oxygen scavenging with low addition levels of elements such as yttrium and scandium (24). In addition to the scavenging effect, these elements with large atomic diameter are likely to have microalloying effects that improve the GFA (24).

For palladium- and platinum-based glass-forming alloys, in contrast, there is little concern about oxygen as a dissolved impurity. Nevertheless, it remains desirable (in some cases essential) to use raw materials of high purity and to minimise pick-up of impurities during processing.

1.2 Direct Casting

In conventional investment casting of platinum alloys, their high TL (around 2000 K) limits the casting size (25) and introduces many issues such as reactions with the crucible, tarnishing and oxidation, plus significant casting shrinkage and porosity (as high as 5%) (26). While palladium jewellery alloys are cast at slightly lower temperatures (27), the further reductions in casting temperatures offered by BMG-forming compositions are still desirable. The low-lying eutectics of platinum-based and palladium-based BMGs offer an exciting opportunity for easier processing of precious-metal jewellery alloys.

The main challenges in casting BMGs arise from the high cooling rates required to avoid crystallisation and the high viscosity of the glass-forming liquid compared with the melts of usual crystalline alloys. In conventional casting techniques, good form-filling requires long casting times (i.e. slow filling), which are longer for more viscous liquids. For glass formation, techniques with low cooling rates (for example, investment casting), are not suitable or need adapting. In any case, the casting size is limited by dc. The undesirability of cold-working BMGs (due to possible fracture or the formation of unsightly surface steps) requires items to be cast near to their final shape.

Nevertheless, for direct conventional casting, BMGs present many advantages when compared with conventional crystalline alloys and this remains a practical processing route for jewellery. The absence of a first-order phase transition on cooling to form a glass means that there is no substantial shrinkage during casting. Even if any crystallisation were to occur, these glass-forming liquids exhibit small crystallisation shrinkages (28).

Casting porosity is a particular problem for many platinum jewellery alloys (26), motivating the introduction of hot isostatic pressing of castings. In palladium-based jewellery, 950Pd alloys suitable for high-quality investment casting are the focus of active research and development (29). For BMGs, low volume shrinkage (< 0.5%) on casting means near-net-shape casting can be achieved with few casting defects such as porosity and with high surface definition. Their lower casting temperatures also lead to lower thermal stresses during cooling to ambient temperature (30). Fewer polishing steps are required after casting due to the absence of crystallinity; ultimately, there can be an atomically smooth surface finish (31).

Unlike zirconium-based BMGs, which are already used in luxury goods markets (32), BMGs based on precious metals (gold, palladium and platinum) for jewellery have the considerable advantage of being processable in air (4, 30, 3336).

The casting of gold-based BMGs for jewellery has been evaluated in detail (3739). Similar information for platinum- and palladium-based BMGs is not as readily available, but their comparably high fragilities and larger dc values (Table I) mean that their castability is similar, if not better. The only significant difference is the requirement to flux with B2O3 before casting. BMGs based on precious metals perform well compared with so-called ‘benchmark’ BMGs which have exceptional processability (32, 34).

Table I

Comparison of Key Processing Parameters of Different Bulk Metallic Glasses

Alloy TL, °C Casting temperature, °C Requires fluxing dc, mm Tg, °C ΔTx, °C S
Zr41.2Ti13.8Cu12.5Ni10Be22.5 (34, 41, 43) 714 831 No 14 349 77 0.21
Pt57.5Cu14.7Ni5.3P22.5 (31, 34, 43) 540 600-800 Yes 20 236 89 0.29
Pt49.95Si6.4B24Cu16.65Ge3 (40) 655 1050 No 5 306 70 0.20
Pd43Ni10Cu27P20 (34, 43) 554 550–650 Yes 30 305 101 0.41
Pd75Si15Ag3Cu7 (42) 756 No 10 348 74 0.18
Au49Ag5.5Pd2.3Cu26.9Si16.3 (31, 34) 371 450–600 No 5 128 58 0.24

Many BMGs can be made using methods such as tilt casting, suction casting and centrifugal casting. While methods such as suction-casting achieve higher cooling rates and higher-quality castings with low porosity (44), die-casting and tilt-casting techniques are scalable to high-volume production (45). High flow rates can induce crystallisation throughout the sample due to shear thinning of the liquid and should be avoided (46). As a result, methods such as centrifugal casting and continuous casting are difficult to apply to BMGs (38, 47, 48). Some heterogeneous nucleation of crystals is inevitable in contact with the copper mould (49).

For BMGs as a jewellery material, tilt-casting appears the most suitable conventional method. Although maximum cooling rates and form-filling are lower than for suction-casting, the method is suitable for industry and has been successfully used to cast samples in massive copper moulds in many of the studies presented here.

For the most intricate designs, a goldsmith uses investment casting. In this method, high form-filling is achieved by slow cooling. While slow cooling gives enough time for the melt to fill the entire mould (37), it is incompatible with the requirements for glass formation. To circumvent this problem, Eisenbart et al. developed a lost-metal mould casting technique, which is an analogue to investment casting for BMGs (39). In their technique, a wax pattern is produced and covered in a thin layer of conducting material. Electroforming is then used to deposit copper onto the wax pattern. The mass of copper deposited is typically five to 10 times the intended mass of the casting to ensure cooling rates high enough for glass formation (38). The wax pattern is then removed either by melting or chemically and the component is centrifugally cast. Finally, the copper mould is etched away to leave the final casting (39).

This technique could have widespread application for jewellery. It allows the formation of highly intricate shapes (Figure 1) not achievable with a dividable copper mould, and high form-filling (near 100%), while satisfying the requirements for glass formation. Furthermore, the electroforming and etching procedures mean that the copper is recyclable, and low-cost wax patterns can be easily shaped by an artisanal goldsmith or manufactured at industrial scale (38).

Fig. 1.

Jewellery products cast using the lost-metal mould casting technique similar to conventional investment casting. Reprinted from (37) with permission. Copyright 2013, Santa Fe Symposium

Jewellery products cast using the lost-metal mould casting technique similar to conventional investment casting. Reprinted from (37) with permission. Copyright 2013, Santa Fe Symposium

1.3 Thermoplastic Forming

Thermoplastic forming (TPF), a method usually reserved for thermoplastic polymers and conventional oxide glasses, can be applied to BMGs. Aside from improved wear resistance, the ability to be so formed is perhaps the most desirable property of BMGs for jewellery. Current jewellery manufacture is laborious, requiring a highly skilled goldsmith. TPF of hallmark-compliant BMGs offers the potential for fast, economical, large-scale production of jewellery items alongside artisanal design innovations.

To date, BMGs have been successfully shaped via a wide range of TPF techniques such as blow-moulding (5053), hot-stamping (5457), extrusion (5861), hot-rolling (62) and injection-moulding (63). Novel processing possibilities, potentially of use to jewellers for setting gems, include welding and joining methods designed to minimise crystallinity (6466), as well as the incorporation of second phases when processing in the supercooled liquid region (SCLR) (67).

TPF has many advantages over the direct casting of BMGs and the casting of crystalline metals. It leads to even fewer casting defects than the direct casting of BMGs, since this isothermal process allows the relaxation of internal stresses and elimination of porosity. Porosity can be reduced to levels well below 0.2% (30). TPF can give a wide range of complex, thin-walled and even hollow shapes with high dimensional accuracy (5153), as well as surface patterning with nanometre accuracy (5557). Recent advances in stretch blow-moulding allow more complex shapes and plastic strains up to several thousand percent (52) compared with a few hundred percent for conventional blow-moulding (50). From a practical perspective, the substantially lower working temperatures and the ability to process platinum- and palladium-based BMGs in air (35) mean that equipment lifetime is extended, costs are lower and manufacture is safer.

Among the many advantages of TPF, the most important is the ability to produce fully glassy samples with dimensions exceeding dc. This is possible since forming of the final shape and cooling to form the glass are decoupled (34). First, a glassy feedstock material (rods (62), pellets and granules (31, 38), plates (45, 67) or discs (50)) can be cast fully glassy. The feedstock can then be used for TPF to produce a larger component. The use of a pellet feedstock is common practice in polymer processing. It allows fast, easy and economical production since large quantities of simple-shape feedstock can be produced by straightforward casting, while complex shapes are introduced only at the final TPF stage. For Pd-Cu-Si based BMGs, as-cast granules have a SiO2 surface layer. This is thicker than would otherwise be expected (68) due to the lower surface tension of silicon which promotes segregation to the liquid surface (69). As for gold-based BMGs, this SiO2 should be removed before TPF to ensure adhesion between the feedstock particles (38, 66).

Since the BMG must be heated above Tg, TPF is possible only within a limited time-temperature processing window before crystallisation occurs (Figure 2). Crystallisation leads to embrittlement, a loss of thermoplastic formability and of many of the desirable properties associated with the glassy structure (36, 55).

Fig. 2.

Schematic TTT diagram of a BMG showing the window for thermoplastic processing: (a) initial cooling; (b) further processing. Temperature–time profiles are shown for: A Cooling curve for a fully glassy sample; B critical cooling rate (Rc) for a fully glassy sample; C thermoplastic processing in the SCLR, with time limited by the onset of crystallisation; D partial crystallisation achieved by annealing in the SCLR

Schematic TTT diagram of a BMG showing the window for thermoplastic processing: (a) initial cooling; (b) further processing. Temperature–time profiles are shown for: A Cooling curve for a fully glassy sample; B critical cooling rate (Rc) for a fully glassy sample; C thermoplastic processing in the SCLR, with time limited by the onset of crystallisation; D partial crystallisation achieved by annealing in the SCLR

While it would seem logical to perform TPF at temperatures just above Tg to maximise the available processing time, sharp-definition forming requires a low viscosity (below 106 Pa s (45, 70)). As can be seen from a simple treatment of deformation in the SCLR using the Hagen-Poiseuille equation, total deformation increases as the temperature increases (35, 71). As the temperature increases by 20°C, the viscosity, η, decreases by one order of magnitude, but the time to crystallisation onset, tc, decreases by a factor of three (36). Consequently, the maximum possible deformation increases, as has been confirmed experimentally (71, 72). The best TPF is achieved by heating the feedstock quickly to the highest temperature possible, shaping the low-viscosity supercooled liquid into the desired final product, then quenching back into the glassy state before the onset of crystallisation (35, 73). Given that a more intricate design takes longer to process, there is a desire to increase thermoplastic formability. Strategies being researched include optimisation of fluxing which can widen the SCLR (8, 11, 19), and multi-step thermoplastic processing (74).

By comparison with other BMGs, gold-, platinum- and palladium-based BMGs exhibit exceptional thermoplastic formability (Table I) due to their low Tg, high liquid m-fragility, large ΔTx and large S (34, 36, 75). Values of Fscan, a parameter that correlates better with deformability during TPF but is more difficult to measure (34, 75), also show that among BMGs, those based on gold, platinum or palladium are the best BMGs for TPF (Figure 3) (75).

Fig. 3.

Thermomechanical analysis (TMA) results, showing penetration depth during continuous heating (5 K min–1) under a load of 0.05 N (through a 4 mm diameter tip) of three BMGs (Au50Cu25.5Ag7.5Si17, Pt60Cu16P22Co2, Pd35Pt15Cu30P20) and images of the test samples at the end of the test. The shaded bars indicate the SCLR for each composition. The wider SCLR for platinum- and palladium-based alloys results in higher deformability during TPF before the onset of crystallisation. Adapted with permission from (36)

Thermomechanical analysis (TMA) results, showing penetration depth during continuous heating (5 K min–1) under a load of 0.05 N (through a 4 mm diameter tip) of three BMGs (Au50Cu25.5Ag7.5Si17, Pt60Cu16P22Co2, Pd35Pt15Cu30P20) and images of the test samples at the end of the test. The shaded bars indicate the SCLR for each composition. The wider SCLR for platinum- and palladium-based alloys results in higher deformability during TPF before the onset of crystallisation. Adapted with permission from (36)

TTT curves for crystallisation of BMGs are helpful in comparing their stability during TPF. platinum-based and palladium-based BMGs show excellent stability against crystallisation at high temperatures within the SCLR (Figure 4). As a result, these alloys, alongside zirconium-based BMGs, have been widely used for many TPF experiments reported in the literature (Figure 5), from hot-embossing on the micro- and even nanometre scale (56), to blow-moulding and injection-moulding (52, 77). As examples, BMG watch casings can be made by TPF and machining; those using zirconium-based BMGs are now commercially available, while those using platinum-based BMGs have been made in research laboratories (40).

Fig. 4.

Experimentally determined sections of TTT curves scaled as a fraction of Tg for: (a) Au49Ag5.5Pd2.3Cu26.9Si16.3 (grey circles) (76); Pt57.5Cu14.7Ni5.3P22.5 (orange squares) (30); and Pd43Cu10Ni27P20 (blue diamonds) (21). These data points indicate the measured onset of crystallisation. For each composition, substantial times are available for TPF before crystallisation in the SCLR; (b) how the crystallisation rate sharply transitions from limitation by growth (at low T) to limitation by nucleation (at high T) in Pd43Cu10Ni27P20. Blue circles show the onset of crystallisation, while orange triangles show its completion (21)

Experimentally determined sections of TTT curves scaled as a fraction of Tg for: (a) Au49Ag5.5Pd2.3Cu26.9Si16.3 (grey circles) (76); Pt57.5Cu14.7Ni5.3P22.5 (orange squares) (30); and Pd43Cu10Ni27P20 (blue diamonds) (21). These data points indicate the measured onset of crystallisation. For each composition, substantial times are available for TPF before crystallisation in the SCLR; (b) how the crystallisation rate sharply transitions from limitation by growth (at low T) to limitation by nucleation (at high T) in Pd43Cu10Ni27P20. Blue circles show the onset of crystallisation, while orange triangles show its completion (21)

Fig. 5.

TPF from a feedstock using Pt57.5Cu14.7Ni5.3P22.5 (30). Image reproduced with permission

TPF from a feedstock using Pt57.5Cu14.7Ni5.3P22.5 (30). Image reproduced with permission

2. Mechanical Properties: Harder Jewellery Alloys Than Ever Before

The jewellery market standard for hardness is a minimum of 100 HV (78). This is still quite soft and so widely used jewellery alloys, with 150–200 HV, scratch relatively easily (Table II). By comparison, BMGs exhibit hardness exceeding 300 HV in the as-cast state. This high hardness offers the opportunity for jewellery with excellent scratch resistance, as well as for watch components (78).

Table II

Typical Hardness Values and Casting Temperatures of Conventional Platinum and Palladium Jewellery Alloysa

Alloy Hallmark Type TL, °C Hardness, HV
Pd-Ru (27) 950Pd Crystalline 1600 120 (as-cast)
Pd-Ag-Cu (27) 500Pd Crystalline 1220 165 (annealed)
Pt-Ir (27) 950Pt Crystalline 1790 140 (cold-worked)
850Pt 1800 110 (annealed)
Pt-Au (27) 950Pt Crystalline 1750 300 (age-hardened)
Pd79Ag3.5P6Si9.5Ge2 (79) 500Pd (90 wt%) BMG 792 497b (as-cast)
Pt74.7Cu1.5Ag0.3P18B4Si1.5 (23) 950Pt BMG 589 395 (as-cast)

a Data for platinum- and palladium-based BMGs are shown for comparison

b Denotes estimate from yield strength

The non-crystalline structure of BMGs means that they cannot show the dislocation-mediated glide and work-hardening familiar for conventional crystalline alloys. When loaded at temperatures below Tg, the BMGs are much harder than their crystalline counterparts, but eventually the limit of elasticity is indicated by the local rearrangement of atoms in small ‘shear-transformation zones’ (STZs) (80, 81). On further loading, the operation of STZs triggers more STZs and there is a plastic instability in which shear is sharply localised in thin bands. These ‘shear bands’ indicate work-softening.

For most BMGs, a dominant shear band forms under tensile load, leading to macroscopic brittle failure (<1% ductility) (82). Even in the absence of fracture, shear bands lead to readily visible surface marks. In contrast, platinum- and palladium-rich BMGs show remarkably high plasticity in compression and fracture toughness (23, 79, 8284). As noted in Section 3.1 of Part I (1), these properties are associated with high Poisson ratio of the glass (> 0.41) (85) and with high m-fragility of the liquid. In these BMGs, plastic deformation is relatively uniform, occurring through the operation of many shear bands (79). Even for these platinum- and palladium-rich BMGs, however, cold working is not recommended, but jewellery made from them is expected to be particularly durable.

The absence of dislocation-mediated plastic deformation means that BMGs can reach much higher elastic strains than crystalline metals and can approach the theoretical strength (86). The elastic strain limit of BMGs is approximately 2%, six times that of a typical polycrystalline jewellery alloy (approx. 0.3%). BMGs are thus potential hallmark-compliant materials for functional watch components such as springs (23, 30, 40), due to their resulting high resilience (σfɛf/2, where σf and ɛf are the stress and strain at the elastic limit), representing the capacity for the elastic storage of energy. While their resilience is an order of magnitude higher than other hallmark-compliant materials, it is lower than current spring materials (approx. two-thirds of the resilience of specialist spring materials).

The deformation mechanism in metallic glasses (MGs) relates closely to the glass transition; there is a strong correlation between the hardness of BMGs and their respective glass-transition temperature (12). Both deformation and the glass transition occur via collective atomic motion requiring sufficient energy to induce atomic rearrangement. In deformation, this occurs locally within a shear band due to the energy provided by mechanical work, while the glass transition occurs across the entire sample and the energy is provided by heat (9, 12, 81, 87).

Compared with phosphorus-containing BMGs, Pt-Si-B based BMGs show a substantially higher hardness in the as-cast state (40), correlated with higher Tg. This higher hardness comes with an undesirable reduction in plasticity. This could be associated with a liquid of lower m-fragility and, therefore, with a different predominant cause of high GFA.

2.1 Effect of Isothermal Annealing Both Above and Below Tg

The intrinsic metastability of glass means it is vital to consider the effect of ageing and ultimately crystallisation, which can have a dramatic effect on properties (36, 73, 8891). BMGs based on precious metals have low Tg’s compared to other MGs, where Tg can exceed 1000 K (92). At room temperature, BMGs based on precious metals are at a comparatively high fraction of Tg so the structure evolves on a shorter timescale. While this is of most concern for gold-based BMGs (Tg = 130°C (33)), it is also important to consider for platinum- and palladium-based BMGs.

In a broad range of temperature mainly below Tg, ageing of glass occurs via α and β relaxation. β relaxation can be associated with changes in chemical short-range ordering, leading to substantial changes in enthalpy. It is ultimately responsible for sub-Tg embrittlement of BMGs (9395). α relaxation occurs predominantly near Tg or on much longer timescales (9699). It is responsible for densification of the glass as it evolves towards the state of a deeply supercooled liquid. This densification, often described as a reduction in ‘free volume’ (100104), occurs during long sub-Tg annealing, and is associated with hardening (94, 95).

Crystallisation can occur only with a level of atomic mobility equalling or exceeding that necessary for α relaxation. Mostly seen on annealing, crystallisation can also be induced by irradiation, deformation or dealloying. It is classified into polymorphic, eutectic and primary crystallisation (105107). In the polymorphic and eutectic cases, the crystallised regions have the same composition as the original glass; as crystallisation proceeds, the glass composition stays unaltered. In that case, on annealing at a fixed temperature, the crystal growth rate is constant. In contrast, primary crystallisation involves solute partitioning between crystal and glass; the associated diffusion control tends to limit crystallite size and fine dispersions may be obtained (107). Crystallisation often occurs in many stages, via metastable phase combinations, towards the equilibrium phase mixture.

Partial crystallisation can lead to a dramatic increase in hardness (Figure 6). Despite suggestions that this is due to a phase-mixture effect (108), it can be due to solute enrichment of the glassy matrix, implied by rejection of solute from the growing crystalline phase (90, 109). The increase in hardness is predominantly attributed to the evolution of hard, brittle equilibrium phases, but we also note that crystallisation in Pt49.95Si6.4B24Cu16.65Ge3 led to a rise in Tg and hardness of the remaining glassy matrix due to copper enrichment (from 570 HV to 750 HV) (88).

Fig. 6.

Hardness vs. annealing time in the SCLR for three BMGs (Au50Cu25.5Ag7.5Si17, Pt60Cu16P22Co2, Pd35Pt15Cu30P20), showing the effect of crystallisation. Annealing performed for each alloy at a temperature of approximately 1.1 Tg (120°C, 260°C and 305°C for gold-, platinum- and palladium-based BMGs, respectively). Figure redrawn from (36) with permission

Hardness vs. annealing time in the SCLR for three BMGs (Au50Cu25.5Ag7.5Si17, Pt60Cu16P22Co2, Pd35Pt15Cu30P20), showing the effect of crystallisation. Annealing performed for each alloy at a temperature of approximately 1.1 Tg (120°C, 260°C and 305°C for gold-, platinum- and palladium-based BMGs, respectively). Figure redrawn from (36) with permission

The equilibrium crystalline phases that form can have complex crystal structures. Although these complex structures may be regarded as beneficial, in that their slow growth should reduce RC (110112), they are hard and brittle. Partial crystallisation therefore leads to substantial embrittlement and a fall in fracture toughness (36, 55, 90, 91, 94, 95, 104, 108, 109). Cardinal et al. report that fracture toughness already halves as a result of 20% crystallinity (36, 74). Given the initially high fracture toughness of many platinum- and palladium-rich BMGs (23, 79, 82), substantial increases in hardness could be achieved, without the BMGs becoming too brittle. The potential advantages of partially crystallised systems, which can be regarded as ‘metallic-glass-matrix composites’ (MGMCs), should therefore not be ignored.

Surface crystallisation may lead to enhanced wear resistance. Kazemi et al. report that the surface hardness of their Pt-Cu-Si-B-Ge BMGs can increase to 800 HV while maintaining toughness in the rest of the sample (25). Surface treatments of this kind may be desirable, but the effect of surface crystallisation on other properties must also be considered.

From a scientific perspective, the crystallisation of BMGs is complex due to the simultaneous or stepwise crystallisation of many ordered phases. Their GFA, thermal stability and crystallisation rate are highly dependent on the fluxing procedure (17, 113), due to the effect of lower cooling rates and cleaning of the melt and on sample size due to the statistical probability of a heterogeneous nucleation site (114). On a TTT diagram (Figure 4), crystallisation of a glass during annealing is described by a C-curve (4, 17, 21, 113). The nose of the curve is often sharp since the crystallisation rate may transition over a narrow temperature range from being nucleation-limited to growth-limited (21, 113). Understanding the factors that affect crystallisation and crystallisation steps is particularly important for understanding an alloy’s suitability for TPF where crystallisation can readily occur. During use as jewellery, crystallisation is unlikely as the temperature is well below Tg.

3. Tarnishing and Corrosion: No Tarnishing Problem?

Corrosion and tarnishing resistance is key to preserving the lustre of jewellery metals. It is the primary reason why gold and platinum are more highly prized than silver. Resistance to corrosion and tarnishing is also essential to prevent failure due to a deterioration in mechanical properties during use.

Comparing the corrosion resistance of glassy and crystalline alloys is complex and requires consideration of multiple factors. Ideally, a glassy alloy is a defect-free, chemically homogeneous, single-phase material without microstructure. This prevents the formation of local galvanic cells that would accelerate corrosion of less noble phases. Superior corrosion resistance is expected, but glasses are metastable and are therefore in a relatively high-energy state. Thus, they are more chemically active, possibly having corrosion rates higher than their crystalline counterparts (115).

Tarnishing is a well-documented problem for gold-based BMGs; it is the main barrier to their application as a material for jewellery, but to date, there is no industrially viable solution. In general, tarnishing is taken to be a chemical and morphological change of the surface due to leaching or oxide formation (68). For gold-based BMGs, tarnishing is attributed to the high copper and silicon content of glass-forming compositions (116, 117), and to fast diffusion kinetics aided by the glassy state (118, 119). While the issues of tarnish and corrosion resistance have been well documented for gold-based BMGs, studies of platinum-based and palladium-based BMG alloys remain few.

3.1 Corrosion Resistance

Wu et al. compared the corrosion resistance of glassy and crystalline Pd40Ni40P20 (120). The crystalline alloy (formed via vacuum annealing) performed two-to-three times better than the glassy alloy in a range of solutions, due to the presence of several inert phosphides and noble palladium-rich crystalline phases (120). These are the dominant equilibrium phases that form during the crystallisation in many Pt-P and Pd-P-based BMGs (113, 121) so similar corrosion resistance is expected. The formation of an inert surface layer appears critical to corrosion resistance (122). In the corrosion of Pd-Au-Ag-Si BMGs, passivation occurs due to the formation of a palladium-enriched surface layer with underlying silicon-rich glassy phases (123). The resulting anodic passive current density is lower than for SUS 316L stainless steel (123).

Copper-containing BMGs contradict these observations that crystallisation enhances corrosion resistance. Studies of Pd-Cu-Ni-P glasses and their crystalline counterparts in a variety of corrosive solutions show that the heterogeneity of the crystalline material results in an inhomogeneous passive film (124, 125). Subsequent pitting leads to accelerated corrosion (124, 125). During crystallisation, several equilibrium phases, Cu3Pd, Ni2Pd2P, Cu5Pd3P2 and an unknown quaternary phase, form via numerous metastable phases (126). The presence of less noble copper-rich phases amongst noble-metal inert phases leads to galvanic cells and the accelerated corrosion of the copper-rich phases (124). In corrosive chlorinated acidic solutions, copper and nickel dissolution is observed first, followed by the dissolution of palladium and phosphorus. This leaves behind a porous Pd-P network that is gradually dissolved after the outer layers of nickel and copper are removed (125). Dealloying may occur via dissolution of phosphorus followed by the formation of a nanocrystalline Pd-Ni solid solution as ligaments in an initial porous structure (127). Similar pitting was also observed in Pd-Cu-Si alloys (115, 127), suggesting that crystallisation is favourable for corrosion resistance only if the phases that form are all equally inert.

Dealloying in glasses occurs due to preferential leaching out of less noble elements. Their dissolution into the electrolyte allows the remaining noble-metal atoms to rearrange themselves into fine ligaments, resulting in a nanoporous surface (68, 128). Crystallisation due to the deviation of the composition from the glass-forming composition, contributes to the formation of ligaments by nucleation and growth of the crystals in the glassy matrix. Compared with dealloying in crystalline alloys, a finer structure is observed in glassy alloys (128). While the resulting nanoporous surface is desirable for catalysis and related applications (128130), it has a degraded appearance unacceptable for jewellery.

3.2 Tarnishing: A Comparison with Gold-Based Bulk Metallic Glasses

Although few studies have compared tarnish resistance, Rizzi et al. reported that a Pd-Cu‐Si BMG incubated at 37°C in artificial saliva did not measurably tarnish while a gold-based BMG tarnished significantly (68). Tarnishing of palladium-based BMGs may not be as problematic as for gold-based BMGs (68) but, given the similarities in composition, it is not clear why this is so. Rizzi et al. note the existence of a SiO2 layer on the surface of both alloys (68). In gold-based BMGs, this SiO2 layer that forms on casting is an effective diffusion barrier, and prevents tarnishing, but it is quickly worn away and therefore it is not protective (68, 116, 117, 131).

Comparing gold-based BMGs and Pd-Cu-Si-based BMGs, the latter have a Tg that is more than 100 K higher. Atomic diffusion rates, closely linked to β relaxation (118, 119, 132134), when compared at room or body temperature, are therefore likely to be significantly lower in Pd-Cu-Si-based BMGs. This might explain their slower tarnishing. Further studies comparing the behaviour of these alloys are likely to be significant in furthering the understanding of the tarnishing mechanisms of precious-metal-based BMGs.

For Pd/Pt-Ni-P based alloys, there appears to be no tarnishing problem. However, the presence of copper in these alloys impairs corrosion resistance, while replacing it with nickel leads to a lowering of GFA and raises issues of toxicity (36, 123, 135).

Further studies of corrosion and tarnishing of palladium- and platinum-based BMGs at body temperature and in simulated body fluids, such as on gold-based BMGs (38, 68, 116, 117), are required to fully characterise their suitability for jewellery, as well as for other applications such as in dentistry or as biomaterials (68, 123, 136).

4. Outlook and Future Developments

4.1 Suitability for Jewellery

Aside from the central issue of the conflicting requirements for high fineness and high GFA, BMGs for jewellery offer a wide range of desirable properties. In the as-cast state, they have hardness values that are, at the very least, challenging to achieve in conventional crystalline metals. With values exceeding 300 HV, components made from BMGs are suitable for watchmaking, potentially even as functional components, due to their unique combination of high resilience and hardness. BMGs in jewellery have excellent scratch-resistance, while more intricate, thinner, and even hollow designs are possible because of their high strength.

BMGs offer many processing advantages, both for casting and TPF. The lower casting temperatures of BMGs, alongside the absence of volume shrinkage due to crystallisation, result in relatively few casting-related defects (such as shrinkage and residual stresses), and in possibilities for near-net-shape casting and excellent surface finishes in the as-cast state. While the requirement for fast cooling limits the size, characterised as the critical casting diameter, dc, of fully glassy samples, the attainable dimensions do exceed those required for most jewellery. The difficulties of form-filling, due to viscous melts and to high required cooling rates, have been resolved through new casting techniques. After casting, TPF of samples offers a wealth of new shaping opportunities. TPF may, therefore, enable new jewellery designs as well as economical, highly reproducible mass production.

The corrosion and tarnish resistance of platinum- and palladium-based BMGs are relatively unstudied. From the literature available, tarnishing does not appear to be a significant issue, while corrosion resistance is improved by the presence of noble phases on the surface and is superior to some stainless steels in aggressive environments. While surface coatings may appear to be a solution, inevitable wear leads to accelerated corrosion of any exposed underlying metal.

4.2 Further Work

For platinum- and palladium-based BMGs, the main obstacle to their use in jewellery is the need to achieve fully glassy states with over 95 wt% platinum or palladium. Achieving such high fineness while alloying sufficiently to achieve adequate GFA appears, at best, challenging. The development of platinum- and palladium-based BMGs with precious-metal content exceeding 90 wt% and even 95 wt% (23, 79) does, however, give hope. Compositions that would give 950Pt and 950Pd BMGs are expected to show high liquid fragilities and therefore low GFA is expected.

In the search for higher precious metal contents, microalloying appears a key strategy, although its effects are far from fully understood. A deeper understanding of microalloying effects may assist in optimising GFA, but the restrictions imposed by hallmark-compliance mean that extensive microalloying can only be with lighter elements (to maximise the weight percent of platinum or palladium) and so other avenues should also be considered. MGMCs may offer opportunities to raise the precious metal content without sacrificing GFA or leading to a heterogeneous appearance (137).

Fluxing of the melt, to remove oxide particles that may act as heterogeneous nucleation sites, has a dramatic effect on GFA: optimised processing can increase critical casting thicknesses from the millimetre to the centimetre scale (8, 11, 18, 19). But treatment of some palladium- and platinum-based BMG-forming liquids with the usual B2O3 flux can lead to undesirable contamination with boron (23, 40, 79, 138). There is a clear need for the development of more stable fluxes and opportunity for optimisation of fluxing treatments.

Compared to 950Pt BMGs, 950Pd BMGs are even more challenging to achieve. From the compositions reported (Table III), further exploration of Pd-Cu-Si alloys should be considered. While Demetriou et al.’s palladium-based BMG exhibits exceptional mechanical properties, compositions such as Pd79Cu6Si10P5 (138) show higher GFA and higher weight fractions of palladium, due to a smaller degree of required alloying with metalloids. They, therefore, appear a promising area for further study in pursuit of hallmark-compliant BMGs. The Pd-Si binary system (specifically Pd81Si19) is one of the best binary glass-forming systems known (142, 143), due to a large energy barrier to crystallisation (142). As for Pd-P-based BMGs, these compositions have liquids of high m-fragility. With palladium content raised to 95 wt%, the m-fragility is likely to be even higher, resulting in reduced GFA. But, there are still many possibilities for a potential 950Pd hallmark-compliant BMG.

Table III

Summary of Platinum- and Palladium-based BMG Alloy Compositions with dc > 2 mma

Alloy, at% Precious metal content, wt% Contains nickel Tg, K Tx, K ΔTx, K TL, K dc, mm Trg S Hardness, HV
Platinum-phosphorus-based BMGs
Pt60Cu20P20 (30, 121) 86.1 No 522 580 58 850 12 0.61 0.18
Pt60Cu2P22Co2 (30, 36) 86.6 No 506 568 63 881 16 0.58 0.17 402
Pt57.5Cu14.7Ni5.3P22.5 (30, 121) 86.3 Yes 508 606 98 795 16 0.64 0.34 402
Pt42.5Cu27Ni9.5P21 (30, 121) 75.9 Yes 515 589 74 873 20 0.59 0.21 392
Pt60Cu2Ni2P22 (121) 87.0 Yes 506 569 63 881 16 0.58 0.17
Pt74.7Cu1.5Ag0.3P18B4Si1.5 (23) 95.0 No 479 529 50 862 2 0.56 0.13 395
Platinum-silicon-based BMGs
Pt52.5Si10.5B19.5Cu17.5 (30) 86.4 No 581 603 22 951 2–5 0.61 0.06
Pt49.95Si6.4B24Cu2.65Ge3 (30) 85.0 No 579 634 55 928 5 0.61 0.16 570
Palladium-phosphorus-based BMGs
Pd30Ni30P20 (2, 121) 71.0 Yes 590 671 91 970 10 0.61 0.24 538
Pd30Ni10Cu30P20 (19, 121) 60.3 Yes 572 670 98 836 72 0.56 0.37 515
Pd43Ni10Cu27P20 (17, 21) 63.6 Yes 585 716 131 878 30 0.67 0.45
Pd42.5Cu30Ni7.5P20 (18) 62.3 Yes 575 684 109 800 80 0.72 0.48
Pd30Pt17.5Cu32.5P20 (136) 34.4 No 541 622 81 812 50 0.67 0.30
Pd35Pt15Cu30P20 (36, 139, 140) 40.6 No 551 624 73 847 30 0.65 0.25 460
Pd20Pt20Cu20Ni20P20 (121) 25.1 Yes 580 645 65 820 10 0.71 0.27
Pd79Ag3.5P6Si9.5Ge2 (79) 89.6 No 613 644 31 1065 6 0.58 0.07 497b
Pd30Ni30Si4P2 (7) 71.1 Yes 579 688 104 971 >10 0.60 0.27 556b
Palladium-silicon-based BMGs
Pd78Si17Cu5 (42) 91.3 No 635 683 49 1098 3 0.58 0.11 527b
Pd78Si17Ag5 (42) 89.1 No 636 690 54 1107 4 0.58 0.11 537b
Pd23Si17Ag5Cu5 (36) 85.3 No 637 703 66 1175 3 0.54 0.12 542b
Pd75Si17Ag5Cu3 (36) 86.9 No 625 692 67 1063 5 0.59 0.15 532b
Pd75Si17Ag3Cu5 (42) 87.7 No 633 704 71 1148 4 0.55 0.14 538b
Pd77Si15Ag3Cu5 (42) 88.5 No 619 696 77 1050 5 0.59 0.18 517b
Pd75Si15Ag5Cu5 (42) 86.2 No 620 691 71 1034 7 0.60 0.17 513b
Pd75Si15Ag3Cu7 (42) 87.0 No 621 695 74 1029 10 0.60 0.18 536b
Pd79Cu6Si10P5 (138, 141) 91.1 No 609 682 73 995 5 0.61 0.19 492b
Pd79Cu5Ag1Si10P5 (141) 90.7 No 614 684 70 1001 4 0.61 0.18
Pd79Cu4Ag2Si10P5 (141) 90.3 No 613 684 71 1005 5 0.61 0.18
Pd79Cu3Ag3Si10P5 (141) 89.8 No 610 683 73 1005 5 0.61 0.18
Pd79Cu2Ag4Si10P5 (141) 89.4 No 611 676 65 1006 7 0.61 0.16
Pd79Cu2Au2Ag2Si10P5 (141) 87.8 No 617 670 53 1020 3 0.60 0.13
Pd79Cu3Au2Ag1Si10P5 (141) 88.2 No 613 676 63 1021 4 0.60 0.15
Pd79Cu2Au1Ag3Si10P5 (141) 88.6 No 617 683 66 1014 3 0.61 0.17
Pd79Cu3Au1Ag2Si10P5 (141) 89.0 No 618 689 71 1015 5 0.61 0.18
Pd79Au1.5Ag3Si16.5 (140) 88.6 No 636 685 49 3 500b

a ΔTx, Trg, dc and S parameters are given to compare their thermal stability, GFA and thermoplastic formability. Values of dc should be treated with caution, since they are dependent on the processing method as well as on Rc

b Denotes an estimate of hardness from compressive yield stress or yield strength

Lastly, for platinum- and palladium-based BMGs to be applicable in jewellery, they must maintain their lustre. Limited work has been carried out to assess how they may tarnish or corrode over time. Further studies of their performance in simulated body fluids at body temperature are necessary to evaluate their suitability. Such studies may also assist in the development of tarnish-resistant 18 karat gold-based BMGs.

5. Conclusions

Hallmark-compliant BMGs exhibit improved hardness and elasticity over conventional crystalline alloys, leading to much-improved wear resistance for jewellery, watch casings and even functional components in watches. Their lower casting temperatures, minimal shrinkage on casting and thermoplastic formability facilitate high-definition near-net-shape casting with few casting defects and processability via methods usually reserved for polymers and conventional oxide glasses. Consequently, BMGs based on precious metals should be of significant interest in jewellery.

The platinum- and palladium-based BMGs reviewed here show some of the best mechanical properties and highest GFA amongst all known BMGs. Their high plasticity is associated with the high m-fragility of their liquids and thereby also with exceptional thermoplastic formability, superior to any other family of BMGs. Compared with 18 karat gold-based BMGs, the absence of any tarnishing problem is advantageous. However, the strict hallmarking standards for platinum and palladium alloys in jewellery mean that the extent of alloying is restricted. Increasing platinum- and palladium-content to the required high levels appears to lead to a reduction in GFA. This reduction may, in part, be due to such melts having higher m-fragility, but changes in crystallisation behaviour may also play a significant role. Specifically, it is challenging to achieve sufficient GFA to obtain BMGs with the compositions that would be most desirable for jewellery (950Pd or 950Pt). A deeper understanding and subsequent optimisation of microalloying and fluxing procedures are suggested as further areas of study for improvements to GFA in these platinum- and palladium-rich systems.



  1. 1.
    O. S. Houghton and A. L. Greer, Johnson Matthey Technol. Rev., 2021, 65, (4), 506 LINK
  2. 2.
    H. W. Kui, A. L. Greer and D. Turnbull, Appl. Phys. Lett., 1984, 45, (6), 615 LINK
  3. 3.
    X. Wang, Z. Tian, M. Zeng, N. Nollmann, G. Wilde and C. Tang, AIP Adv., 2017, 7, (9), 095308 LINK
  4. 4.
    J. Schroers and W. L. Johnson, Appl. Phys. Lett., 2004, 84, (18), 3666 LINK
  5. 5.
    O. Haruyama, T. Watanabe, K. Yuki, M. Horiuchi, H. Kato and N. Nishiyama, Phys. Rev. B, 2011, 83, (6), 064201 LINK
  6. 6.
    A. Inoue, N. Nishiyama and H. Kimura, Mater. Trans., JIM, 1997, 38, (2), 179 LINK
  7. 7.
    N. Chen, L. Gu, G. Q. Xie, D. V. Louzguine-Luzgin, A. R. Yavari, G. Vaughan, S. D. Imhoff, J. H. Perepezko, T. Abe and A. Inoue, Acta Mater., 2010, 58, (18), 5886 LINK
  8. 8.
    D. Granata, E. Fischer, V. Wessels and J. F. Löffler, Appl. Phys. Lett., 2015, 106, (1), 011902 LINK
  9. 9.
    X. Yang, X. Ma, Q. Li and S. Guo, J. Alloys Compd., 2013, 554, 446 LINK
  10. 10.
    I. Seki, D. V. Louzguine-Luzgin, T. Takahashi, H. Kimura and A. Inoue, Mater. Trans., 2011, 52, (3), 458 LINK
  11. 11.
    D. Granata, E. Fischer, V. Wessels and J. F. Löffler, Acta Mater., 2014, 71, 145 LINK
  12. 12.
    K. F. Yao, F. Ruan, Y. Q. Yang and N. Chen, Appl. Phys. Lett., 2006, 88, (12), 122106 LINK
  13. 13.
    J. S. Yu, Y. Q. Zeng, T. Fujita, T. Hashizume, A. Inoue, T. Sakurai and M. W. Chen, Appl. Phys. Lett., 2010, 96, (14), 141901 LINK
  14. 14.
    A. J. Drehman and A. L. Greer, Acta Metall., 1984, 32, (3), 323 LINK
  15. 15.
    J. J. Wall, C. T. Liu, W.-K. Rhim, J. J. Z. Li, P. K. Liaw, H. Choo and W. L. Johnson, Appl. Phys. Lett., 2008, 92, (24), 244106 LINK
  16. 16.
    A. Inoue, T. Zhang, A. Takeuchi and W. Zhang, Mater. Trans., JIM, 1996, 37, (4), 636 LINK
  17. 17.
    J. Schroers and W. L. Johnson, Appl. Phys. Lett., 2000, 77, (8), 1158 LINK
  18. 18.
    N. Nishiyama, K. Takenaka, H. Miura, N. Saidoh, Y. Zeng and A. Inoue, Intermetallics, 2012, 30, 19 LINK
  19. 19.
    N. Nishiyama and A. Inoue, Mater. Trans., JIM, 1996, 37, (10), 1531 LINK
  20. 20.
    H.-Y. Ding, Y. Li and K.-F. Yao, Chin. Phys. Lett., 2010, 27, (12), 126101 LINK
  21. 21.
    J. Schroers, Y. Wu, R. Busch and W. L. Johnson, Acta Mater., 2001, 49, (14), 2773 LINK
  22. 22.
    N. Chen, D. Pan, D. V. Louzguine-Luzgin, G. Q. Xie, M. W. Chen and A. Inoue, Scr. Mater., 2010, 62, (1), 17 LINK
  23. 23.
    M. D. Demetriou, M. Floyd, C. Crewdson, J. P. Schramm, G. Garrett and W. L. Johnson, Scr. Mater., 2011, 65, (9), 799 LINK
  24. 24.
    A. A. Kündig, D. Lepori, A. J. Perry, S. Rossmann, A. Blatter, A. Dommann and P. J. Uggowitzer, Mater. Trans., 2002, 43, (12), 3206 LINK
  25. 25.
    H. Kazemi, “Alloy Development of a New Platinum-Based Bulk Metallic Glass”, PhD Thesis, École Polytechnique Fédérale de Lausanne, Switzerland, 3rd March, 2017 LINK
  26. 26.
    T. Freyé and J. Fischer-Buehner, ‘Platinum Alloys in the 21st Century: A Comparative Study’, The 25th Santa Fe Symposium, 15th–18th May, 2011, Albuquerque, USA, pp. 201–230 LINK
  27. 27.
    C. W. Corti, ‘Jewellery Alloys – Past, Present and Future’, The Goldsmiths’ Company Jewellery Materials Congress, 8th–9th July, 2019, London, UK, 24 pp LINK
  28. 28.
    S. Mukherjee, J. Schroers, Z. Zhou, W. L. Johnson and W.-K. Rhim, Acta Mater., 2004, 52, (12), 3689 LINK
  29. 29.
    P. Battaini, ‘The Working Properties for Jewelry Fabrication using New Hard 950 Palladium Alloys’, The 20th Santa Fe Symposium, 10th–13th September 2006, Nashville, USA, pp. 19–53 LINK
  30. 30.
    J. Schroers, B. Lohwongwatana, W. L. Johnson and A. Peker, Mater. Sci. Eng.: A, 2007, 449–451, 235 LINK
  31. 31.
    B. Lohwongwatana, J. Schroers and W. L. Johnson, ‘Liquidmetal – Hard 18K and .850Pt Alloys that can be Processed like Plastics or Blown Like Glass’, The 13th Santa Fe Symposium, 20th–23rd May 2007, Albuquerque, USA, pp. 289–303 LINK
  32. 32.
    ‘Swatch Group Signs Exclusive License Agreement With Liquidmetal Technologies’, The Swatch Group Ltd, Switzerland, 10th March, 2011 LINK
  33. 33.
    J. Schroers, B. Lohwongwatana, W. L. Johnson and A. Peker, Appl. Phys. Lett., 2005, 87, (6), 061912 LINK
  34. 34.
    J. Schroers, Adv. Mater., 2010, 22, (14), 1566 LINK
  35. 35.
    J. Schroers, J. Miner. Metals Mater. Soc., 2005, 57, (5), 35 LINK
  36. 36.
    S. Cardinal, J. Qiao, J. M. Pelletier and H. Kato, Intermetallics, 2015, 63, 73 LINK
  37. 37.
    U. E. Klotz and M. Eisenbart, ‘Gold-Based Bulk Metallic Glasses: Hard like Steel, Moldable like Plastics’, The 13th Santa Fe Symposium, 17th–20th May, 2013, Albuquerque, USA, 16 pp LINK
  38. 38.
    M. Eisenbart, “On the Processing and the Tarnishing Mechanism of Gold-Based Bulk Metallic Glasses”, Dissertation, Universität des Saarlandes, Germany, 2015
  39. 39.
    M. Eisenbart, U. E. Klotz, A. Pfund and A. Zielonka, ‘Method for Casting an Object Made of Metallic Glass’, World Patent Appl. WO2014198380
  40. 40.
    H. Kazemi, C. Cattin, M. Blank and L. Weber, J. Alloys Compd., 2017, 695, 3419 LINK
  41. 41.
    A. Peker and W. L. Johnson, Appl. Phys. Lett., 1993, 63, (17) 2342 LINK
  42. 42.
    W. Zhang, H. Guo, Y. Li, Y. Wang, H. Wang, M. Chen and S. Yamaura, J. Alloys Compd., 2014, 617, 310 LINK
  43. 43.
    G. Duan, A. Wiest, M. L. Lind, J. Li, W.-K. Rhim and W. L. Johnson, Adv. Mater., 2007, 19, (23), 4272 LINK
  44. 44.
    A. Inoue, Mater. Trans., JIM, 1995, 36, (7), 866 LINK
  45. 45.
    J. Schroers and N. Paton, Adv. Mater. Proc., 2006, 164, (1), 61
  46. 46.
    B. Lohwongwatana, J. Schroers and W. L. Johnson, Phys. Rev. Lett., 2006, 96, (7), 075503 LINK
  47. 47.
    T. Zhang, X. Zhang, W. Zhang, F. Jia, A. Inoue, H. Hao and Y. Ma, Mater. Lett., 2011, 65, (14), 2257 LINK
  48. 48.
    F. Haag, R. Sauget, G. Kurtuldu, S. Prades-Rödel, J. E. K. Schawe, A. Blatter and J. F. Löffler, J. Non-Cryst. Solids, 2019, 521, 119120 LINK
  49. 49.
    S. Mechler, E. Yahel, P. S. Pershan, M. Meron and B. Lin, Appl. Phys. Lett., 2011, 98, (25), 251915 LINK
  50. 50.
    J. Schroers, Q. Pham, A. Peker, N. Paton and R. V. Curtis, Scr. Mater., 2007, 57, (4), 341 LINK
  51. 51.
    J. Schroers, T. M. Hodges, G. Kumar, H. Raman, A. J. Barnes, Q. Pham and T. A. Waniuk, Mater. Today, 2011, 14, (1–2), 14 LINK
  52. 52.
    R. M. O. Mota, N. Liu, S. A. Kube, J. Chay, H. D. McClintock and J. Schroers, Appl. Mater. Today, 2020, 19, 100567 LINK
  53. 53.
    M. Kanik, P. Bordeenithikasem, J. Schroers, N. Selden, A. Desai, D. Kim and R. M’Closkey, ‘Microscale Three-Dimensional Hemispherical Shell Resonators From Metallic Glass’, International Symposium on Inertial Sensors and Systems (ISISS), Laguna Beach, California, USA, 25th-26th February, 2014, pp. 1–4 LINK
  54. 54.
    P. Sharma, N. Kaushik, H. Kimura, Y. Saotome and A. Inoue, Nanotech., 2007, 18, (3), 035302 LINK
  55. 55.
    G. Kumar, H. X. Tang and J. Schroers, Nature, 2009, 457, 868 LINK
  56. 56.
    J. Schroers, T. Nguyen, S. O’Keeffe and A. Desai, Mater. Sci. Eng. A, 2007, 449–451, 898 LINK
  57. 57.
    Y. C. Chen, J. P. Chu, J. S. C. Jang and C. W. Yu, Mater. Sci. Eng. A, 2012, 556, 488 LINK
  58. 58.
    A. Kato, A. Inoue, H. Horikiri and T. Masumoto, Mater. Trans., JIM, 1994, 35, (2), 125 LINK
  59. 59.
    K. S. Lee and Y. W. Chang, Mater. Sci. Eng. A, 2005, 399, (1–2), 238 LINK
  60. 60.
    L. C. Zhang, J. Xu and E. Ma, Mater. Sci. Eng. A, 2006, 434, (1–2), 280 LINK
  61. 61.
    J. S.-C. Jang, C.-T. Tseng, L.-J. Chang, J. C.-C. Huang, Y.-C. Yeh and J.-L. Lou, Adv. Eng. Mater., 2008, 10, (11), 1048 LINK
  62. 62.
    R. Martinez, G. Kumar and J. Schroers, Scr. Mater., 2008, 59, (2), 187 LINK
  63. 63.
    A. Wiest, J. S. Harmon, M. D. Demetriou, R. D. Conner and W. L. Johnson, Scr. Mater., 2009, 60, (3), 160 LINK
  64. 64.
    L. Shao, A. Datye, J. Huang, J. Ketkaew, S. W. Sohn, S. Zhao, S. Wu, Y. Zhang, U. D. Schwarz and J. Schroers, Sci. Rep., 2017, 7, 7989 LINK
  65. 65.
    Y. Huang, P. Xue, S. Guo, Y. Wu, X. Cheng, H. Fan, Z. Ning, F. Cao, D. Xing, J. Sun and P. K. Liaw, Sci. Rep., 2016, 6, 30674 LINK
  66. 66.
    W. Chen, Z. Liu and J. Schroers, Acta Mater., 2014, 62, 49 LINK
  67. 67.
    J. Schroers, T. Nguyen and G. A. Croopnick, Scr. Mater., 2007, 56, (2), 177 LINK
  68. 68.
    P. Rizzi, I. Corazzari, G. Fiore, I. Fenoglio, B. Fubini, S. Kaciulis and L. Battezzati, Corros. Sci., 2013, 77, 135 LINK
  69. 69.
    O. G. Shpryko, R. Streitel, V. S. K. Balagurusamy, A. Y. Grigoriev, M. Deutsch, B. M. Ocko, M. Meron, B. Lin and P. S. Pershan, Science, 2006, 313, (5783), 77 LINK
  70. 70.
    H. Kato, T. Wada, M. Hasegawa, J. Saida, A. Inoue and H. S. Chen, Scr. Mater., 2006, 54, (12), 2023 LINK
  71. 71.
    J. Schroers, Acta Mater., 2008, 56, (3), 471 LINK
  72. 72.
    E. B. Pitt, G. Kumar and J. Schroers, J. Appl. Phys., 2011, 110, (4), 043518 LINK
  73. 73.
    S. Cardinal, J. M. Pelletier, M. Eisenbart and U. E. Klotz, Mater. Sci. Eng. A, 2016, 660, 158 LINK
  74. 74.
    P. Boordeenithikasem, S. Sohn, Z. Liu and J. Schroers, Scr. Mater., 2015, 104, 56 LINK
  75. 75.
    B. Bochtler, O. Kruse and R. Busch, J. Phys. Cond. Matt., 2020, 32, (24), 244002 LINK
  76. 76.
    G. Fiore, P. Rizzi and L. Battezzati, J. Alloys Compd., 2011, 509, S166 LINK
  77. 77.
    N. Nishiyama and A. Inoue, Mater. Trans., JIM, 1999, 40, (1), 64 LINK
  78. 78.
    C. W. Corti, ‘The Role of Hardness in Jewelry Alloys’, The 22nd Santa Fe Symposium, 18th–21st May 2008, Albuquerque, USA, pp. 103–120 LINK
  79. 79.
    M. D. Demetriou, M. E. Launey, G. Garrett, J. P. Schramm, D. C. Hofmann, W. L. Johnson and R. O. Ritchie, Nat. Mater., 2011, 10, (2), 123 LINK
  80. 80.
    F. Spaepen, Scr. Mater., 2006, 54, (3), 363 LINK
  81. 81.
    A. S. Argon, Acta Metall., 1979, 27, (1), 47 LINK
  82. 82.
    J. Schroers and W. L. Johnson, Phys. Rev. Lett., 2004, 93, (25), 255506 LINK
  83. 83.
    G. Kumar, S. Prades-Rodel, A. Blatter and J. Schroers, Scr. Mater., 2011, 65, (7), 585 LINK
  84. 84.
    G. Kumar, P. Neibecker, Y. H. Liu and J. Schroers, Nat. Commun., 2013, 4, 1536 LINK
  85. 85.
    J. J. Lewandowski, W. H. Wang and A. L. Greer, Philos. Mag. Lett., 2005, 85, (2), 77 LINK
  86. 86.
    L. Tian, Y.-Q. Cheng, Z.-W. Shan, J. Li, C.-C. Wang, X.-D. Han, J. Sun and E. Ma, Nat. Commun., 2012, 3, 609 LINK
  87. 87.
    A. L. Greer, Mater. Today, 2009, 12, (1–2), 14 LINK
  88. 88.
    H. Kazemi, C. Cattin, G. Hodel, T. Pachova and L. Weber, J. Non-Cryst. Solids, 2017, 460, 66 LINK
  89. 89.
    H. W. Bi, A. Inoue, F. F. Han, Y. Han, F. L. Kong, S. L. Zhu, E. Shalaan, F. Al-Marzouki and A. L. Greer, Acta Mater., 2018, 147, 90 LINK
  90. 90.
    Z. C. Zhong, X. Y. Jiang and A. L. Greer, Mater. Sci. Eng. A, 1997, 226–228, 531 LINK
  91. 91.
    J. Qiao, H. Jia and P. K. Liaw, Mater. Sci. Eng. R, 2016, 100, 1 LINK
  92. 92.
    M.-X. Li, S.-F. Zhao, Z. Lu, A. Hirata, P. Wen, H.-Y. Bai, M. Chen, J. Schroers, Y. Liu and W.-H. Wang, Nature, 2019, 569, 99 LINK
  93. 93.
    P. Murali and U. Ramamurty, Acta Mater., 2005, 53, (5), 1467 LINK
  94. 94.
    Z. Chen, A. Datye, P. A. Brooks, M. Sprole, J. Ketkaew, S. Sohn, J. Schroers and U. D. Schwarz, MRS Adv., 2019, 4, 73 LINK
  95. 95.
    A. Datye, J. Ketkaew, J. Schroers and U. D. Schwarz, J. Alloys Compd., 2020, 819, 152979 LINK
  96. 96.
    H. Vogel, Phys. Z., 1921, 22, 645
  97. 97.
    G. S. Fulcher, J. Am. Ceram. Soc., 1925, 8, (6), 339 LINK
  98. 98.
    G. Tammann and W. Hesse, Z. Anorg. Allg. Chem., 1926, 156, (1), 245 LINK
  99. 99.
    C. T. Moynihan, A. J. Easteal, M. A. De Bolt and J. Tucker, J. Am. Ceram. Soc., 1976, 59, (1–2), 12 LINK
  100. 100.
    H. S. Chen, Rep. Prog. Phys., 1980, 43, (4), 353 LINK
  101. 101.
    A. van den Beukel and J. Sietsma, Acta Metall. Mater., 1990, 38, (3), 383 LINK
  102. 102.
    A. Niikura, A. P. Tsai, A. Inoue and T. Masumoto, J. Non-Cryst. Solids, 1993, 159, (3), 229 LINK
  103. 103.
    J. J. Lewandowski, Mater. Trans., 2001, 42, (4), 633 LINK
  104. 104.
    G. Kumar, M. Ohnuma, T. Furubayashi, T. Ohkubo and K. Hono, J. Non-Cryst. Solids, 2008, 354, (10–11), 882 LINK
  105. 105.
    U. Köster and J. Meinhardt, Mater. Sci. Eng. A, 1994, 178, (1–2), 271 LINK
  106. 106.
    M. T. Clavaguera-Mora, N. Clavaguera, D. Crespo and T. Pradell, Prog. Mater. Sci., 2002, 47, (6), 559 LINK
  107. 107.
    R. I. Wu, G. Wilde and J. H. Perepezko, Mater. Sci. Eng. A, 2001, 301, (1), 12 LINK
  108. 108.
    H. S. Kim, P. J. Warren, B. Cantor and H. R. Lee, Nanostruct. Mater., 1999, 11, (2), 241 LINK
  109. 109.
    A. L. Greer, Mater. Sci. Eng. A, 2001, 304–306, 68 LINK
  110. 110.
    A. L. Greer, Nature, 1993, 366, (6453), 303 LINK
  111. 111.
    T. Egami and W. Waseda, J. Non-Cryst. Solids, 1984, 64, (1–2), 113 LINK
  112. 112.
    I.-R. Lu, M. Kolbe, G. P. Görler and R. Willnecker, Mater. Sci. Eng.: A, 2004, 375–377, 754 LINK
  113. 113.
    B. A. Legg, J. Schroers and R. Busch, Acta Mater., 2007, 55, (3), 1109 LINK
  114. 114.
    N. Nishiyama and A. Inoue, J. Non-Cryst. Solids, 2002, 312–314, 575 LINK
  115. 115.
    J.-Y. Yang, J.-L. Gu, S.-Q. Chen, C.-H. Luan, Y. Shao and K.-F. Yao, Mater. Corros., 2018, 69, (11), 1509 LINK
  116. 116.
    M. Eisenbart, U. E. Klotz, R. Busch and I. Gallino, Corros. Sci., 2014, 85, 258 LINK
  117. 117.
    M. Eisenbart, U. E. Klotz, R. Busch and I. Gallino, J. Alloys Compd., 2014, 615, (1), S 118 LINK
  118. 118.
    Z. Evenson, S. E. Naleway, S. Wei, O. Gross, J. J. Kruzic, I. Gallino, W. Possart, M. Stommel and R. Busch, Phys. Rev. B, 2014, 89, (17), 174204 LINK
  119. 119.
    Z. Evenson, T. Koschine, S. Wei, O. Gross, J. Bednarcik, I. Gallino, J. J. Kruzic, K. Rätzke, F. Faupel and R. Busch, Scr. Mater., 2015, 103, 14 LINK
  120. 120.
    Y.-F. Wu, W.-C. Chiang, J. Chu, T.-G. Nieh, Y. Kawamura and J.-K. Wu, Mater. Lett., 2006, 60, (19), 2416 LINK
  121. 121.
    A. Takeuchi, N. Chen, T. Wada, Y. Yokoyama, H. Kato, A. Inoue and J. W. Yeh, Intermetallics, 2011, 19, (10), 1546 LINK
  122. 122.
    C. L. Qin, Y. Q. Zeng, D. V. Louzguine, N. Nishiyama and A. Inoue, J. Alloys Compd., 2010, 504, S 172 LINK
  123. 123.
    N. Chen, C. L. Qin, G. Q. Xie, D. V. Louzguine-Luzgin and A. Inoue, J. Mater. Res., 2010, 25, (10), 1943 LINK
  124. 124.
    F. Qin, G. Xie, T. Wada, S. Zhu and Z. Dan, Mater. Trans., 2013, 54, (8), 1347 LINK
  125. 125.
    L. Y. Watanabe, S. N. Roberts, N. Baca, A. Wiest, S. J. Garrett and R. D. Conner, Mater. Sci. Eng. C, 2013, 33, (7), 4021 LINK
  126. 126.
    C. Ma, N. Nishiyama and A. Inoue, Mater. Trans., 2002, 43, (5), 1161 LINK
  127. 127.
    J. Yu, Y. Ding, C. Xu, A. Inoue, T. Sakurai and M. Chen, Chem. Mater., 2008, 20, (14), 4548 LINK
  128. 128.
    P. Rizzi, F. Scaglione and L. Battezzati, J. Alloys Compd., 2014, 586, (S1), S117 LINK
  129. 129.
    Y. Xue, F. Scaglione, P. Rizzi and L. Battezzati, Corros. Sci., 2017, 127, 141 LINK
  130. 130.
    Y. Xue, F. Scaglione, P. Rizzi, L. Battezzati, P. Denis and H.-J. Fecht, Appl. Surf. Sci., 2019, 476, 412 LINK
  131. 131.
    S. Mozgovoy, J. Heinrich, U. E. Klotz and R. Busch, Intermetallics, 2010, 18, (12), 2289 LINK
  132. 132.
    F. Faupel, W. Frank, M.-P. Macht, H. Mehrer, V. Naundorf, K. Rätzke, H. R. Schober, S. K. Sharma and H. Teichler, Rev. Mod. Phys., 2003, 75, (1), 237 LINK
  133. 133.
    H.-B. Yi, W.-H. Wang and K. Samwer, Mater. Today, 2013, 16, (5), 183 LINK
  134. 134.
    O. Gross, N. Neuber, A. Kuball, B. Bochtler, S. Hechler, M. Frey and R. Busch, Commun. Phys., 2019, 2, 83 LINK
  135. 135.
    C. W. Corti, ‘What is a White Gold? Progress on the Issues!’, Santa Fe Symposium on Jewelry Manufacturing Technology, Albuquerque, New Mexico, USA, May, 2005, pp. 103–119 LINK
  136. 136.
    N. Nishiyama, K. Takenaka and A. Inoue, Appl. Phys. Lett., 2006, 88, (12), 121908 LINK
  137. 137.
    J. H. Na, K. H. Han, G. R. Garrett, M. E. Launey, M. D. Demetriou and W. L. Johnson, Sci. Rep., 2019, 9, 3269 LINK
  138. 138.
    L. Liu, A. Inoue and T. Zhang, Mater. Trans., 2005, 46, (2), 376 LINK
  139. 139.
    N. Nishiyama, K. Takenaka, T. Wada, H. Kimura and A. Inoue, Mater. Trans., 2005, 46, (12), 2807 LINK
  140. 140.
    K. Takenaka, T. Wada, N. Nishiyama, H. Kimura and A. Inoue, Mater. Trans., 2005, 46, (7), 1720 LINK
  141. 141.
    L. Liu, X. Zhao, C. Ma, S. Pang and T. Zhang, J. Non-Cryst. Solids, 2006, 352, (52–54), 5487 LINK
  142. 142.
    N. Chen, Y. Li and K.-F. Yao, J. Alloys Compd., 2010, 504, (S1), S211 LINK
  143. 143.
    D. V. Louzguine-Luzgin, K. Georgarakis, V. Zadorozhnyy, N. Chen, K. Nakayama, G. Vaughan, A. R. Yavari and A. Inoue, Intermetallics, 2012, 20, (1), 135 LINK


OSH thanks the Goldsmiths’ Company Assay Office, London UK, for financial support and Dr Christopher Corti for useful discussions. ALG acknowledges support from the European Research Council under grant ERC-2015-AdG-695487 (ExtendGlass).

The Authors

Owain Houghton is a PhD student at the University of Cambridge, UK. He obtained a BA degree in Natural Sciences and an MSci degree in Materials Science from the same institution. He researches bulk metallic glasses based on precious metals and their potential applications for jewellery as a member of the Microstructural Kinetics Group.

A. Lindsay Greer is a professor of Materials Science at the University of Cambridge. He received MA and PhD degrees from Cambridge and holds Honorary Doctorates from AGH University of Science and Technology, Cracow, Poland, and the University of Sofia, Bulgaria. He was an assistant professor at Harvard University, USA; has held visiting positions in Grenoble, France; St. Louis, USA; Vienna, Austria; and Turin, Italy; and is a foreign principal investigator in the Advanced Institute for Materials Research, Tohoku University, Japan. He leads the Microstructural Kinetics Group, and his current research focuses mainly on metallic glasses. He is a coauthor of some 450 papers.

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