Advanced Search
RSS LinkedIn Twitter

Journal Archive

Johnson Matthey Technol. Rev., 2021, 65, (4), 506

doi:10.1595/205651321x16045078967011

精细与稳定的冲突:珠宝用铂基和钯基块体非晶合金:第一部分

铂基和钯基块体非晶合金的介绍与性能

对于珠宝用金属,高硬度和相应的抗刮性备受追捧。珠宝用传统晶态合金通过合金和大量的(热力和机械)加工提高硬度,但硬度水平很难超过300 HV。因此,基于贵金属、铸态硬度超过300 HV的块体非晶合金(BMG),一经问世便得到珠宝和制表行业的极大关注。这些材料的非晶态结构不仅具有极高的硬度,而且还可通过热塑性成型(TPF)使金属像塑料一样成形。对于更传统的珠宝制造商,BMG还具有高精度和近净形铸造性能。金基合金长期以来一直主导着主要用BMG领域,因为它们可以满足18克拉纯度检验证明要求。尽管基于铂或钯的BMG具有优良的热塑成形性能,并且不存在任何已知的变色问题,但要在通常用于珠宝的更严格品质证明标准(铂或钯重量纯度≥ 95%)条件下实现良好的非晶成形能力(GFA),绝非易事。这项由两部分组成的评论围绕铂基和钯基BMG展开讨论,重点主要集中于它们在珠宝首饰中的应用潜力以及需要进行的深入研究。

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

Introduction and properties of platinum- and palladium-based bulk metallic glasses

  • O. S. Houghton*
  • A. L. Greer
  • Department of Materials Science and Metallurgy, University of Cambridge, 27 Charles Babbage Road, Cambridge, CB3 0FS, UK
  • *Email: osh24@cam.ac.uk
SHARE THIS PAGE:

Article Synopsis

For the metals used in jewellery, high hardness and the associated scratch resistance are much sought after. Conventional crystalline alloys for jewellery are alloyed and extensively processed (thermally and mechanically) to improve hardness, but it is difficult to reach values beyond 300 HV. The advent of bulk metallic glasses (BMGs), based on precious metals and with hardness exceeding 300 HV in the as-cast state, is therefore of great interest for both jewellery and watchmaking. The non-crystalline structure of these materials not only gives high hardness, but also the opportunity to shape metals like plastics, via thermoplastic forming (TPF). For more traditional jewellery manufacture, BMGs also exhibit high-definition and near-net-shape casting. Gold-based alloys have long dominated the consideration of BMGs for jewellery as they can comply with 18 karat hallmarks. Although BMGs based on platinum or palladium possess excellent thermoplastic formability and are without known tarnishing problems, achieving useful glass-forming ability (GFA) within the more restrictive hallmarking standards typically used for jewellery (≥95 wt% platinum or palladium) is at best challenging. In this two-part review, platinum- and palladium-based BMGs are discussed, focusing on their potential application in jewellery and on the further research that is necessary.

1. Introduction

The metals conventionally used in jewellery are polycrystalline. Within each grain of these structures, the atoms are arranged in a regular periodic lattice. In metallic glasses (MGs), in contrast, there is no such long-range order: the atoms are densely packed in a solid state that has a liquid-like structure (1). Since they were first reported by Klement et al. in 1960 (2), MGs have been widely studied. Glass can be described as a hardened, cooled liquid.

The remarkably different structure of MGs leads to a set of unusual properties (Table I). Their high strength and hardness and a large elastic strain limit result in improved scratch and wear resistance. They also open up opportunities for novel intricate designs, making possible thinner and even hollow sections (6) and the use of hallmarked alloys in functional components of watches such as springs (11). The glassy structure also has many advantages for manufacture: near-net-shape casting with minimal casting defects and thermoplastic formability (1214) are desirable for design innovation and economical mass production.

Table I

A Summary of the Advantageous and Disadvantageous Properties of Bulk Metallic Glasses and their Implications for Jewellery (310)

Properties Implication for jewellery
Advantages High yield stress Better scratch and wear resistance
Thinner, more intricate designs
Minimal shrinkage on casting Good surface definition
Low porosity and residual stresses
Near-net-shape casting
Thermoplastic formability Economical scale of production
New processing and designs
Surface patterning
Lower casting temperature Ease of production
Good castability
Lower density Smaller volume fills the same size mould (cost savings)
Disadvantages Four or more elements required for glass formation Difficulties in achieving high fineness
High viscosity and high cooling rates Issues with form filling
Limited time available for processing
Deformation mode Macroscopic brittle failure
Inability to reshape or cold work

The ability to form a glassy state in metals was first observed on splat cooling an Au-Si eutectic alloy. Klement et al. produced partially glassy samples a few micrometres (<30 μm) thick (2). These samples were unstable but demonstrated that glassy metals form on rapid cooling (>106 K s–1) of the liquid. Further research has led to the development of a range of glasses that can be cast with minimum dimension exceeding 1 mm. These are known as ‘bulk’ metallic glasses (BMGs). The first metallic glass with dimensions exceeding 1 mm was a palladium-based system cast by Chen and Turnbull (15). Further improvements came from melt fluxing that removes oxides (or other melt inclusions) that might promote crystallisation (16) (Section 1.1, Part II (17)). The most significant developments came in the early 1990s from separate research by Inoue et al. (18) and by Peker and Johnson (19).

In the years since, many glass-forming compositions have been identified in a wide range of systems including those based on the precious metals gold, silver, platinum and palladium (Figure 1) (14, 15, 28, 38). The critical cooling rate for glass formation (Rc), required to avoid crystallisation on cooling from the liquid, has generally been reduced, in some cases dramatically even to values below 100 K s–1. Consequently, BMGs can now be cast on the centimetre scale using a more extensive range of casting techniques. The same BMGs can also have excellent thermal stability, so a wide range of TPF techniques usually reserved for polymers and oxide glasses are now applicable (1214).

Fig. 1.

One measure of the ease of glass formation is the critical casting diameter (dc), i.e. the diameter below which rod-shaped samples can be cast fully glassy. The optimum values obtained with different base metals cover a wide range. Data from (2039)

One measure of the ease of glass formation is the critical casting diameter (dc), i.e. the diameter below which rod-shaped samples can be cast fully glassy. The optimum values obtained with different base metals cover a wide range. Data from (20–39)

The interest in BMGs for jewellery was sparked by Schroers et al.’s development of platinum- and gold-based BMGs that are compliant with 850Pt and 18 karat gold UK hallmarking standards (Table II) (4, 14, 28). The 950Pt, 950Pd, Sterling silver and gold hallmarks above 18 karat conflict with the alloying required for high GFA. Interest in intermediate hallmarks for palladium, such as ‘585Pd’ (≥58.5 wt% palladium), exists but is not widespread (40). The compliance of gold-based BMGs with the much used 18 karat hallmark drew most attention (4, 41, 42). The majority of jewellery-related research has focused on their improvement, with a particular need to inhibit their abnormally fast tarnishing (3, 4345). To date, silver-based systems have comparatively low GFA and low silver contents (38); they are not suitable for jewellery.

Table II

Relevant Hallmarking Standards for Precious Metals (40)

Base precious metal Hallmark Minimum weight fraction
Gold 14 karat 58.5%
18 karat 75.0%
Platinum 850Pt 85.0%
900Pt 90.0%
950Pt 95.0%
Palladium 500Pd 50.0%
950Pd 95.0%
Silver Sterling/925 92.5%
Britannia/958 95.8%

The increase in popularity of platinum and palladium for jewellery means that developments in platinum- and palladium-based BMGs should not be ignored. Platinum has been widely used as a jewellery metal with varying popularity since the early 20th century. The metal has been used extensively by jewellers, including Carl Fabergé, for its lustrous appearance (46). A resurgence since the early 1960s and the introduction of platinum hallmarking standards in 1975 means that today around 26% of all platinum consumption is for use in jewellery (47). Palladium is becoming an ever more popular choice since the introduction of hallmarking standards in 2010.

Platinum, palladium and gold are face-centred cubic (fcc) metals that, when pure, are so soft (hardness below 75 HV) as to be practically useless. Consequently, it is standard practice to alloy them with other metals such as copper, silver, ruthenium, rhodium, iridium, cobalt, chromium, gallium and indium (8, 40, 4853). The limited alloying achievable within high fineness hallmarks or stamping regulations outside the UK (Table II) means most alloys are solid solutions. They therefore require extensive thermomechanical treatments to achieve desirable hardness (40, 48). A final hardness of 150–200 HV is typical (8). Minimum requirements for jewellery and watchmaking are 100 HV and 300 HV, respectively (8).

Even extensive thermomechanical treatments of two-phase platinum alloys fail to achieve hardness values surpassing 300 HV (48). There is a clear need for alternative routes to hallmark-compliant platinum and palladium alloys with high hardness and improved scratch resistance. Furthermore, platinum and palladium have high melting points requiring casting temperatures over 2000 K (approx. 1750ºC) (11, 28). Such high temperatures introduce many problems for jewellery casting such as reactions with the crucible, tarnishing and oxidation at high temperature and shrinkage-related defects (28).

As discussed in this review, hallmark-compliant BMGs offer a solution to many of the problems currently facing precious-metal jewellery alloys (Table I), but the alloying requirements for GFA and hallmark compliance are mostly opposing and therefore present a significant challenge. Following a brief introduction to the science of MGs, platinum- and palladium-based BMGs suitable for jewellery are discussed, alongside their desirable mechanical properties, processability and corrosion resistance.

2. Bulk Metallic Glasses: A Brief Explanation

MGs are formed when the molten liquid is cooled fast enough to avoid crystallisation. On cooling, the liquid becomes more and more viscous and below the glass-transition temperature (Tg) the atoms undergo substantial thermal arrest; they are ‘frozen’ into a liquid-like arrangement, so crystallisation cannot readily occur. The glass transition is a kinetic phenomenon, not a thermodynamic one. Compared with conventional oxide glasses and polymers, the viscosity of metallic liquids is low, associated with the isotropic non-directional nature of metallic bonding (54). Correspondingly, MGs have comparatively high Rc. There are several approaches to facilitate glass formation by reducing Rc.

The ‘confusion principle’ proposes that glass formation can be assisted if there are many competing equilibrium crystalline phases such that crystallisation kinetics is sluggish (55). For this reason, the highest GFA is often seen in quaternary or higher-order alloys. Alloying elements are chosen to be of different atomic radii. Radii differing by at least 12% promote dense packing in the liquid and may also lead to topological instability in any given crystal structure, thereby further promoting vitrification (56). High GFA is also favoured by having strongly negative enthalpies of mixing between the elements (57), promoting chemical uniformity of the liquid.

Despite having no long-range order, the structure of MGs shows substantial short-range order (SRO) and medium-range order (MRO). The structure of the liquid and the subsequent MG can be considered in terms of stable atom-centred clusters that reduce mobility and inhibit atomic rearrangement that would lead to crystallisation. The liquid viscosity η decreases rapidly with increasing temperature just above Tg (Figure 2). This decrease is described by the kinetic ‘fragility’, a term that refers to the progressive breakdown of order in the liquid as it is heated and not to a mechanical property. One parameter used to characterise the fragility is m (58), Equation (i):

(i)

Fig. 2.

An Angell plot (58) showing the variation in viscosity with temperature above Tg. Ortho-terphenyl, silica and conventional window glass are shown for comparison with several platinum- and palladium-based BMGs and a so-called ‘benchmark glass-former’ zirconium-based BMG. Data obtained from (59, 60)

An Angell plot (58) showing the variation in viscosity with temperature above Tg. Ortho-terphenyl, silica and conventional window glass are shown for comparison with several platinum- and palladium-based BMGs and a so-called ‘benchmark glass-former’ zirconium-based BMG. Data obtained from (59, 60)

Liquids with low m are termed ‘strong’ and those with high m are termed ‘fragile’. While the liquids from which MGs are formed are all ‘fragile’ compared with the liquids giving conventional oxide glasses (i.e. their viscosity is more sensitive to temperature above Tg), their carefully chosen compositions make them ‘strong’ compared to pure-metal melts or to the liquids from which typical crystalline engineering alloys are cast.

Even so, there is a significant range of m-fragility amongst those liquids that form MGs. Those liquids with high m (‘stronger’ liquids) have viscosities that decrease less on heating above Tg. Compared to other liquids at the same T/Tg, these liquids have a higher viscosity and therefore lower atomic mobility. Atoms in a ‘stronger’ liquid are less able to rearrange and in the absence of other factors, such liquids should be better glass formers (61, 62). Characteristics favouring high GFA have been summarised by Inoue as ‘empirical rules’ (57), which have led to the development of BMGs with Rc well below 100 K s–1 and the casting of glassy products on the centimetre scale (63).

Even when a MG can be made, it is difficult to determine Rc and the associated dc with precision. In any case, in trying to develop new compositions, it would be preferable to have a guide to GFA based on parameters derived only from thermophysical data. Several parameters have been suggested, reflecting the thermodynamics and kinetics involved in glass formation, and these can show a good correlation with Rc (63).

The reduced glass-transition temperature Trg is defined by Equation (ii):

(ii)

Turnbull proposed that a high value of Trg, more specifically > 2/3, would indicate a good glass-forming composition (1). A high Trg means a small interval between TL and Tg (Figure 3). This minimises the temperature interval in which crystallisation is possible both thermodynamically and kinetically. While it is possible to raise Tg, the glass transition remains poorly understood. The compositional variation of Tg is comparatively weak, however, so the variation of TL is critical. Eutectic or near-eutectic compositions have strongly suppressed TL and so remain the preferred, but not the only, choice for high-GFA alloys (38). As will be discussed, the composition of these low-lying eutectics, typically in the range 20–30 at% metalloid, is problematic for high fineness BMGs.

Fig. 3.

Schematic of a simple binary eutectic phase diagram with minimal solid solubility. The variation of the liquidus temperature TL and the glass-transition temperature Tg with composition is such that their ratio, the reduced glass-transition temperature Trg, shows a sharp maximum at the eutectic composition. This maximum correlates with compositions of high GFA

Schematic of a simple binary eutectic phase diagram with minimal solid solubility. The variation of the liquidus temperature TL and the glass-transition temperature Tg with composition is such that their ratio, the reduced glass-transition temperature Trg, shows a sharp maximum at the eutectic composition. This maximum correlates with compositions of high GFA

Once formed, the glassy structure remains metastable. On heating, the atomic mobility increases. Once the temperature exceeds ~0.6 Tg, relaxation occurs (i.e. changes in the glassy structure that do not involve crystallisation or phase separation). First, there are local atomic rearrangements, known as β relaxation (64). These occur in a chain-like manner and lead to short-range chemical ordering (9, 65). Near Tg, global and coordinated rearrangements may occur, known as α relaxation (64), which allow the transition from glass into the supercooled liquid state. Ultimately, the supercooled liquid will crystallise into the equilibrium crystalline phases.

On heating in calorimetry, crystallisation shows its onset at Tx (Figure 4). Like Tg, Tx is kinetically controlled and is, therefore, a function of the heating rate (66). As a sample is heated at a higher rate, the transitions to the liquid and to the crystalline states must occur faster, therefore at higher temperature: the measured values of Tg and Tx are higher. The region bounded by Tg and Tx is the supercooled liquid region (SCLR) and is characterised by ΔTx, Equation (iii):

(iii)

Fig. 4.

An annotated calorimetry trace of Pt60Cu16P22Co2 metallic glass on heating, showing the heat-transfer rate (q) as a function of temperature (T). The trace shows the glass-transition temperature as measured on heating (Tg,h), the onset crystallisation temperature (Tx), the SCLR (ΔTx), the melting temperature (Tm) and the liquidus temperature (TL). Adapted from (28)

An annotated calorimetry trace of Pt60Cu16P22Co2 metallic glass on heating, showing the heat-transfer rate (q) as a function of temperature (T). The trace shows the glass-transition temperature as measured on heating (Tg,h), the onset crystallisation temperature (Tx), the SCLR (ΔTx), the melting temperature (Tm) and the liquidus temperature (TL). Adapted from (28)

The SCLR is important as the region relevant for TPF (12, 13). The width of the SCLR, ΔTx, also indicates the stability of the supercooled liquid. Inoue proposes that an excellent glass-forming system has a high Trg (for high GFA) and a wide ΔTx (for excellent thermal stability) (67).

In the SCLR, the metastable supercooled liquid can undergo homogeneous viscous flow, reaching strains of several hundred percent (13, 6870). After some time of holding in the SCLR, crystallisation begins and ultimately proceeds to completion. A wider SCLR (greater ΔTx) indicates greater resistance to crystallisation and improves the thermoplastic formability of the MG (14). For best correlation with thermoplastic formability, Schroers argues that ΔTx should to be normalised by (TLTg) (12, 14). A large ΔTx/(TLTg) means a larger viscosity change within the SCLR, allowing greater deformation during TPF. The S parameter, Equation (iv):

(iv)

although omitting the effect of liquid fragility (Figure 2), correlates remarkably well with thermoplastic formability, making it an effective tool for comparison of different glass compositions (12, 71). While thermoplastic formability shows even better correlation with other parameters (for example, constant heating formability Fscan (13, 71)), S has the advantage that it can be calculated from a single calorimetric experiment (Figure 4), so is more readily adopted as a simple guide. In this review, Trg, ΔTx and S are used to compare GFA and thermoplastic formability.

3. Glass-Forming Ability of Platinum- and Palladium-Based Bulk Metallic Glasses

The remarkable properties of platinum-based and palladium-based BMGs and the particularly high GFA of palladium-based compositions (Figure 1) means they have been the subject of substantial research. While being initially developed for scratch-resistant, hallmark-compliant jewellery, it was found that some platinum-based BMGs show a record-breaking combination of excellent strength and toughness (7274).

3.1 Platinum-Phosphorus and Palladium-Phosphorus Based Systems

Metal-metalloid alloy systems can be excellent glass-formers. This is attributed to the small metalloid atoms occupying interstitial sites between the dense-random-packed metal atoms. The strong bonding between metal and metalloid atoms further helps to give a stable dense packing (75). Pt-P and Pd-P binary eutectics are popular starting points for BMGs with further alloying additions, such as nickel and copper, leading to a substantial increase in GFA (7678). From a topological standpoint, these additions help to form efficiently packed clusters. As proposed by Miracle et al., the form of these clusters depends on the relative sizes of the constituent atoms (7982).

Following the discovery of palladium-based bulk glass formers by Chen and Turnbull (15), the Pd-Ni-P ternary eutectic has received particular attention. Partial substitution of nickel with copper leads to a dramatic rise in GFA. The heats of mixing for Pd-Cu, Ni-Cu and Cu-P pairs are more strongly negative than for Pd-Ni (76) and lead to a change in the dense packing of atoms in the viscous liquid (77). Further refinements in composition resulted in an alloy with dc on centimetre scale and wide ΔTx (83). The maximum GFA was later attributed to having the same chemical SRO of nickel and copper around phosphorus (78). These BMGs are particularly stable against crystallisation (8488). When crystallisation eventually occurs in the SCLR, several ordered and complex crystalline phases are formed in a single cooperative transformation. The excellent GFA is attributed to the difficulty of crystallising the ordered Pd3Cu phase. Compared to other BMGs, the nose of the crystallisation curve on the time-temperature-transformation (TTT) diagram (the temperature at which crystallisation is fastest) for Pd40Ni10Cu30P20 lies at a time 10 times longer than for the benchmark glass-former VitreloyTM 1 (85, 89). The time before crystallisation onset is, however, still short: for a Pd43Ni10Cu27P20 sample held just above Tg this is 104 s (84, 89). Below Tg, times to the onset of crystallisation are orders of magnitude longer due to the dramatically lower atomic mobility.

While palladium-based BMGs have been studied since the 1970s, platinum-based BMGs are a comparatively recent development. In 2004, Schroers and Johnson reported two novel platinum-based BMG compositions, Pt57.5Cu14.7Ni5.3P22.5 and Pt60Cu16Co2P22 (28). These show high GFA when their melts are fluxed with B2O3 (dc > 10 mm), good thermal stability (ΔTx > 60 K), exceptional thermoplastic formability (S > 0.20), as well as being processable in air (28) (which is difficult for zirconium-based BMGs, for example).

For jewellery, Pt60Cu16Co2P22 is particularly attractive due to its high platinum content (satisfying the 850Pt hallmark) and the absence of skin-sensitising nickel (28, 41, 42, 45). On the other hand, Pt57.5Cu14.7Ni5.3P22.5 attracted much attention from researchers. Its plastic strain of 20% and fracture toughness of 80 MPa m1/2 far exceed any previously reported values for BMGs (28, 72). Further studies have reported fracture toughness for both palladium- and platinum-rich BMGs as high as 200 MPa m1/2 (73, 74) — a value comparable to many low-carbon steels widely used as structural materials.

This unique combination of high hardness and high plasticity has been linked to their high Poisson ratio (>0.4) in the glassy state, which itself is related to high m-fragility in the liquid state (7274). High Poisson ratio corresponds to a low ratio of shear modulus G to bulk modulus B. The low value of G/B indicates that resistance to shear (proportional to G) is low compared to resistance to cavitation (proportional to B). While a typical BMG would fail by the operation of a single dominant shear band, leading to macroscopic plasticity below 1%, these BMGs flow by the operation of multiple shear bands leading to high macroscopic plasticity (72).

While high m-fragility of the glass-forming liquid is associated with desirable fracture toughness, it is also associated with reduced GFA (62, 87). The exceptional GFA of both Pt-P and Pd-P based BMGs is therefore surprising (7274, 76, 77, 90, 91). As noted above, high GFA is typically associated with strong liquids (62). Fragile liquids do not aid glass formation through high viscosity (low atomic mobility), but high GFA may result also from a low thermodynamic driving force for crystallisation, or a high crystal-liquid interfacial energy (inhibiting crystal nucleation) (87, 92). Studies report that a low driving force for crystallisation stabilises Pd43Cu27Ni10P20 while high interfacial energy stabilises Pt57.5Cu14.7Ni5.3P22.5 (87). The nature of stabilisation helps to explain why fluxing has such a pronounced effect on the GFA of these alloys, since the presence of any heterogeneous nucleation sites, notably oxide inclusions, substantially lowers their resistance to crystallisation.

Measurement of crystallisation kinetics over the full temperature range from TL down to Tg is helpful in understanding the mechanisms. A classical TTT diagram shows the times necessary for the progress of crystallisation in the supercooled liquid upon isothermal holding at each temperature. The times, for example for crystallisation onset, follow a C-curve in which the minimum time (at the nose of the curve) lies between TL and Tg. At higher temperature than the nose, the kinetics is controlled by crystal nucleation, and at lower temperature by crystal growth. In most cases, for example, for a zirconium-based BMG-forming alloy, the C-curve is asymmetric with the temperature of the nose much closer to TL than to Tg (62). In contrast, the C-curves for palladium- and platinum-based alloys are more symmetric (84, 87, 91). For Pd43Ni10Cu27P20 (89), for example, the nose lies roughly halfway between TL and Tg. As the m-fragility of palladium-based BMG-forming liquids is relatively high (m>50 (93)), the kinetics nearer to TL should be relatively accelerated, which, in the absence of other factors would impart greater asymmetry to the C-curve. The special feature of Pd43Ni10Cu27P20, and palladium- and platinum-based BMG-forming compositions in general, is thus identified as difficulty in crystal nucleation (89). Indeed, the nucleation is mostly possible only because of the influence of heterogeneities (89). This explains why fluxing (which can remove heterogeneous nucleation sites, notably oxide inclusions, Section 1.1, Part II (17)) can dramatically improve the GFA of these alloys.

Recent work on the structures of Pd-P and Pt-P based BMGs provides insight into the critical features required for high GFA (92, 9497). Detailed studies show that the crystal-liquid interfacial energy in Pt-Cu-(Ni/Co)-P BMGs is three times that in kinetically stabilised zirconium-based BMGs (94). The dramatic improvement in GFA with overheating above TL, to dissolve all preexisting structures in the liquid, provides further evidence that these glasses are stabilised by the barrier to crystal nucleation (94, 96).

In the search for BMGs with high weight fractions of platinum or palladium, understanding the liquid kinetics is imperative. Studies suggest that Pt-P based liquids with higher platinum content are more fragile (95, 97). This is relevant, given the desire in jewellery for platinum contents exceeding 95 wt% (approx. 70 at%). In Pd-(Cu,Ni)-P BMGs, increasing m-fragility with higher palladium contents is attributed to the less pronounced bifurcation into Pd-Cu-P and Pd-Ni-P clusters with different coordination shells around the central phosphorus atom (77, 95, 97). These two structural units were suggested to develop MRO that stabilises the supercooled liquid and is responsible for the stronger liquid behaviour (95). Topological-based assumptions that Pd-P and Pt-P liquids have similar structures led to the expectation that the same structural changes would occur in Pt-P liquids (95). In practice, the liquids show substantial differences in SRO and MRO (97, 98).

While icosahedral SRO dominates in Pd-P liquids, Pt-P liquids contain many more trigonal prismatic structural units, leading to the observation of pronounced MRO. This MRO is not observed at high temperatures, unlike the icosahedral SRO in Pd-P liquids, resulting in more pronounced ordering during cooling towards Tg (95, 97). As a result, Pt-P liquids have a much broader distribution of cluster connection schemes, comprising the more flexible two- and four-atom connections, whereas the stiffer three-atom connections prevail in the Pd-P-based liquids (97). The presence of SRO at high temperatures and a relative lack of ordering on the medium length-scale in Pd-P-based liquids lead to a low entropy of fusion compared with Pt-P-based liquids. This explains the low driving force for crystallisation in Pd-P-based liquids (62, 83, 87, 98, 99) as well as the more pronounced sensitivity of plasticity to the cooling rate during glass formation (97, 100, 101). Regardless of the differences when compared with Pt-P-based liquids, the high m-fragility of Pd-P-based liquids does mean that there is still substantial ordering on cooling towards Tg; fragility is dependent on the rate of ordering near Tg, not on the type of ordering (92).

3.2 Bulk Metallic Glasses for Jewellery

The primary restriction on BMGs for jewellery is hallmarking standards (as practised in the UK) and stamping regulations (US equivalent). These hallmarking standards place a minimum weight fraction on platinum, palladium, gold and silver to guarantee quality (Table II). Of these hallmarks, 18 karat gold, 950Pt, 900Pt, 950Pd and sterling silver are predominantly used for jewellery.

Alongside the requirement for a minimum weight fraction of precious metal, jewellery alloys must also be without skin-sensitising elements such as nickel (42, 45, 102). Consequently, many of the otherwise attractive glass-forming compositions are unsuitable for jewellery. Although studies suggest that the ion release rate of nickel from glassy alloys is well below the legal requirement (45), there is an industry-wide desire to eliminate such elements (103). Nickel-free BMGs are also of substantial interest for dental and biomedical applications (104, 105).

Schroers and Johnson’s discovery of 850Pt BMGs was, therefore, a breakthrough (28). The high GFA, high m-fragility and low Tg of these liquids are desirable for high-definition TPF (41, 71, 106). Low casting temperatures and the glassy structure itself appear to solve many issues in jewellery manufacture (Section 1, Part II (17)), while the unique properties of BMGs can be exploited for scratch-resistance and other property enhancements (Section 2, Part II (17)). Still, it is the high 950Pt and 950Pd hallmarks that are the most desirable for jewellery. Given the inevitably high capital cost that would be associated with the manufacture of BMGs using techniques that are novel and unfamiliar to goldsmiths, BMG jewellery would, at least initially, be expensive. Hence 950Pt and 950Pd hallmark-compliant BMGs would be desirable, if not, essential. It is conceivable that the majority of consumers would justify the cost of these materials by fineness alone. While the 900Pt hallmark is widely used by jewellers and offers a wider ‘compositional space’ for alloy development, it would not command as high a price.

Attempting to achieve such high-fineness hallmarks heavily restricts the extent of alloying, and yet, as discussed above, alloying is necessary to achieve high GFA. The platinum-rich (95 wt% platinum) and palladium-rich BMGs (90 wt% palladium) developed by Demetriou et al. (73, 74) are, therefore, both scientifically and industrially significant. Through microalloying (i.e. minor additions of elements), they were able to achieve high weight fractions of platinum and palladium in fully glassy samples, albeit with relatively poor GFA. While Demetriou et al. are correct that their 950Pt BMG would be suitable for jewellery (73), this risks ignoring practical issues. Techniques such as tilt casting, suitable for industrial jewellery manufacture, cannot achieve the cooling rates achieved by expensive, small-scale, laboratory techniques such as suction casting. In an industrial environment, it will be challenging to cast these alloys into a fully glassy state. Furthermore, jewellery alloys often contain a higher fraction of precious metal than hallmark requirements stipulate. This tolerance ensures that all items produced are hallmark-compliant. Increasing the weight fraction of platinum or palladium any further in these alloys is likely to result in the loss of useful GFA.

The work of Demetriou et al. (73, 74) shows that, with difficulty, 950Pt and high fineness palladium (90 wt%) hallmark-compliant BMGs are achievable, but the best alloys developed so far are still relatively poor bulk glass formers. The necessity to add more than approximately 15 at% metalloids to achieve high GFA means that achieving high fineness is difficult.

Microalloying may be key in achieving better compositions, although its effects are poorly understood. Minor additions can have a significant impact on GFA: adding 0.3 at% silver to Pt74.7Cu1.5P18B4Si1.5 increases dc by a factor of two, albeit from a low level (73). Additions of up to 2 at% gold or silver increase GFA in Pd79Cu6Si10P5 (86), rationalised using a topological argument (although joint additions yield less impressive results).

The effect of microalloying on properties of BMGs appears specific to the added element (107, 108), suggesting that both chemical and topological effects are important. From a topological viewpoint, microalloying is expected to increase the variety of local atomic configurations (108110). This may affect local atomic rearrangements (β relaxation) linked to spatial heterogeneity (111) and therefore the properties of BMGs below Tg. Faster β relaxation due to microalloying (109) is expected to significantly affect a wide range of properties, from corrosion and tarnishing behaviour (65, 112) to sub-Tg embrittlement (97, 112115). If microalloying is to be employed in platinum- and palladium-based systems to permit an increase in the precious-metal weight fraction, a deeper understanding is required of its effect on glass properties and GFA.

3.3 Phosphorus-Free Bulk Metallic Glasses

Alloys containing phosphorus are challenging to process. Its reactivity means additional processing steps, such as prealloying, are required (11). Pt-P and Pd-P binary eutectics are, therefore, not ideal starting points for glassy jewellery alloys (11).

For platinum-based BMGs, extensive work by Kazemi et al. has identified the Pt-Si-B ternary eutectic as a suitable starting point (11) for alloy development. The eutectic composition lies at a high atomic percentage of platinum, while the weight percentage of platinum is aided by the low atomic masses of silicon and boron. Many other M-Si-B eutectics (M = nickel, cobalt, iron) have high GFA, which is rationalised by topological arguments (67, 116, 117). Prior studies, comparing the GFA of Pd-Si and Pd-Si-B, further support expectations of a high-GFA platinum-based BMG (118). By substituting copper for platinum, and germanium for silicon, at the Pt-Si-B ternary eutectic composition, a Pt-Cu-B-Si-Ge BMG with a dc of 5 mm was developed (11).

In contrast to phosphorus-containing BMGs, Kazemi et al.’s Pt49.95Cu16.65Si6.4Ge3B24 alloy is unaffected by fluxing. Prolonged fluxing with B2O3 did not reduce the oxygen content (119). While oxygen-scavenging by alloying with scandium was successful in reducing oxygen content, the GFA was not improved (119). Instead, all GFA was lost when more than 2 at% scandium was substituted for copper (119). Scandium and holmium additions were found to sharply increase Tg and Tx, with a concomitant rise in hardness. These additions also led to a rise in TL, ultimately leading to a loss of GFA (lower Trg) (119). While oxygen scavenging, minor additions of large-atomic-radius rare-earth metals and prolonged fluxing all improve the GFA of Pt-P (87, 120, 121), their failure to do so for Pt-Si-B suggests that the controlling factors are different. To elucidate these and to optimise alloy compositions, it would be helpful to know the fragility of the liquid: no values have been reported, but the liquid is expected to be relatively fragile.

Apart from Pd-P, only the Pd-Si eutectic has been reported as a suitable starting point for BMGs. Early work by Chen and Turnbull focused on the Pd-Cu-Si ternary eutectic, yielding high-GFA alloys (15). The similarity in weight fraction of palladium between Pd-P and Pd-Si means, as for platinum-based BMGs, that high fineness and high GFA conflict with one another. Many Pd-Si-based glass formers have low weight fractions of palladium and are therefore not suitable for jewellery. They have, however, attracted much interest for their use as biomedical and dental materials due to their high plasticity, absence of skin-sensitising elements and excellent corrosion resistance (104, 105).

Although high fineness is still challenging, the Pd-Si binary eutectic serves as an excellent starting point for BMGs. Pd-Si binary alloys are remarkably good glass formers; the glasses exhibit high activation energies for crystallisation (122), resulting in high thermal stability. However, only after a careful choice of elements with low atomic mass and suitable atomic radii is the suggestion of a 950Pd hallmark-compliant BMG plausible.

Reports of Pd79Au1.5Ag3Si16.5 (dc = 3 mm) (105) and Pd75Si15Ag3Cu7 (dc = 10 mm) (123) BMGs support these claims. These are two of the highest fineness phosphorus-free palladium-based BMGs reported so far, with weight fractions around 90 wt% palladium before further alloying. Pd75Si15Ag3Cu7 is also readily formed using TPF and does not require B2O3 fluxing, making it easier to manufacture (123). While a phosphorus-free BMG would be desirable, small additions of phosphorus can improve GFA and permit an increased weight fraction of palladium (124).

Further research on 950Pt and 950Pd hallmark-compliant BMGs with a GFA suitable for industrial manufacture is challenging. The inevitable high m-fragility of these alloys means that high cooling rates and carefully optimised processing are likely to be requirements. Research has mostly focused on Pt-(Cu,Ni)-P and Pd-(Cu,Ni)-P based BMGs (16, 20, 41, 42, 73, 74, 84, 89, 90, 104, 125) with some interest shown in Pt-Si-B and Pd-Cu-Si based BMGs (85, 86, 88, 105, 123, 124).

Further optimisation through careful alloying and microalloying may help to identify additional 950Pt BMGs and novel 950Pd BMGs with sufficient GFA. While other systems such as Pd-Ni-S BMGs or novel high-entropy BMGs exist, they have low palladium contents and so do not seem suitable for jewellery (5, 126). Li et al. suggest that there may exist up to 106 BMG compositions (127), so the search for new glass-forming compositions should continue. New combinatorial techniques to assess ΔTx and Trg (128130), and computational techniques (131133), allow more efficient sampling of compositional space for undiscovered high GFA compositions. A new search with these techniques may yield results that cannot readily be predicted and rationalised using empirical rules and experimentally determined phase diagrams.

BACK TO TOP

References

  1. 1.
    D. Turnbull, Contemp. Phys., 1969, 10, (5), 473 LINK https://doi.org/10.1080/00107516908204405
  2. 2.
    W. K. Jun, R. H. Willens and P. Duwez, Nature, 1960, 187, (4740), 869 LINK https://doi.org/10.1038/187869b0
  3. 3.
    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 http://www.santafesymposium.org/2013-santa-fe-symposium-papers/2013-gold-based-bulk-metallic-glasses-hard-like-steel-moldable-like-plastics
  4. 4.
    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 http://www.santafesymposium.org/2007-santa-fe-symposium-papers/2007-liquid-metal-hard-18k-and-850pt-alloys-that-can-be-processed-like-plasticsor-blown-like-glass
  5. 5.
    J. Schroers, B. Lohwongwatana, W. L. Johnson and A. Peker, Appl. Phys. Lett., 2005, 87, (6), 061912 LINK https://doi.org/10.1063/1.2008374
  6. 6.
    J. Schroers, Q. Pham, A. Peker, N. Paton and R. V. Curtis, Scr. Mater., 2007, 57, (4), 341 LINK https://doi.org/10.1016/j.scriptamat.2007.04.033
  7. 7.
    R. Martinez, G. Kumar and J. Schroers, Scr. Mater., 2008, 59, (2), 187 LINK https://doi.org/10.1016/j.scriptamat.2008.03.008
  8. 8.
    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 http://www.santafesymposium.org/2008-santa-fe-symposium-papers/2008-the-role-of-hardness-in-jewelry-alloys
  9. 9.
    H.-B. Yi, W.-H. Wang and K. Samwer, Mater. Today, 2013, 16, (5), 183 LINK https://doi.org/10.1016/j.mattod.2013.05.002
  10. 10.
    I. Gallino and R. Busch, J. Miner. Metals Mater. Soc., 2017, 69, (11), 2171 LINK https://doi.org/10.1007/s11837-017-2573-6
  11. 11.
    H. Kazemi, C. Cattin, M. Blank and L. Weber, J. Alloys Compd., 2017, 695, 3419 LINK https://doi.org/10.1016/j.jallcom.2016.12.017
  12. 12.
    J. Schroers, Adv. Mater., 2010, 22, (14), 1566 LINK https://doi.org/10.1002/adma.200902776
  13. 13.
    J. Schroers, Acta Mater., 2008, 56, (3), 471 LINK https://doi.org/10.1016/j.actamat.2007.10.008
  14. 14.
    J. Schroers, J. Miner. Metals Mater. Soc., 2005, 57, (5), 35 LINK https://doi.org/10.1007/s11837-005-0093-2
  15. 15.
    H. S. Chen and D. Turnbull, Acta Metall., 1969, 17, (8), 1021 LINK https://doi.org/10.1016/0001-6160(69)90048-0
  16. 16.
    H. W. Kui, A. L. Greer and D. Turnbull, Appl. Phys. Lett., 1984, 45, (6), 615 LINK https://doi.org/10.1063/1.95330
  17. 17.
    O. S. Houghton and A. L. Greer, Johnson Matthey Technol. Rev., 2021, 65, (4), 519 LINK https://www.technology.matthey.com/article/65/4/519-534/
  18. 18.
    A. Inoue, T. Zhang and T. Masumoto, Mater. Trans., JIM, 1990, 31, (3), 177 LINK https://doi.org/10.2320/matertrans1989.31.177
  19. 19.
    A. Peker and W. L. Johnson, Appl. Phys. Lett., 1993, 63, (17), 2342 LINK https://doi.org/10.1063/1.110520
  20. 20.
    N. Nishiyama, K. Takenaka, H. Miura, N. Saidoh, Y. Zeng and A. Inoue, Intermetallics, 2012, 30, 19 LINK https://doi.org/10.1016/j.intermet.2012.03.020
  21. 21.
    H. B. Lou, X. D. Wang, F. Xu, S. Q. Ding, Q. P. Cao, K. Hono and J. Z. Jiang, Appl. Phys. Lett., 2011, 99, (5), 051910 LINK https://doi.org/10.1063/1.3621862
  22. 22.
    M. Q. Tang, H. F. Zhang, Z. W. Zhu, H. M. Fu, A. M. Wang, H. Li and Z. Q. Hu, J. Mater. Sci., Technol., 2010, 26, (6), 481 LINK https://doi.org/10.1016/S1005-0302(10)60077-1
  23. 23.
    T. Zhang, R. Li and S. Pang, J. Alloys Compd., 2009, 483, (1–2), 60 LINK https://doi.org/10.1016/j.jallcom.2008.07.224
  24. 24.
    Q. Zhang, W. Zhang and A. Inoue, Mater. Trans., JIM, 2007, 48, (11), 3031 LINK https://doi.org/10.2320/matertrans.MEP2007201
  25. 25.
    Q. Zheng and J. Xu, J. Appl. Phys., 2007, 102, (11), 113519 LINK https://doi.org/10.1063/1.2821755
  26. 26.
    F. Guo and S. J. Poon, Appl. Phys. Lett., 2003, 83, (13), 2575 LINK https://doi.org/10.1063/1.1614420
  27. 27.
    Y. Zeng, N. Nishiyama, T. Yamamoto and A. Inoue, Mater. Trans., JIM, 2009, 50, (10), 2441 LINK https://doi.org/10.2320/matertrans.MRA2008453
  28. 28.
    J. Schroers and W. L. Johnson, Appl. Phys. Lett., 2004, 84, (18), 3666 LINK https://doi.org/10.1063/1.1738945
  29. 29.
    Y. Zhou, Y. Zhao, B. Y. Qu, L. Wang, R. L. Zhou, Y. C. Wu and B. Zhang, Intermetallics, 2015, 56, 56 LINK https://doi.org/10.1016/j.intermet.2014.09.003
  30. 30.
    A. Inoue, F. L. Kong, Q. K. Man, B. L. Shen, R. W. Li and F. Al-Marzouki, J. Alloys Compd., 2014, 615, S2 LINK https://doi.org/10.1016/j.jallcom.2013.11.122
  31. 31.
    E. S. Park and D. H. Kim, J. Mater. Res., 2004, 19, (3), 685 LINK https://doi.org/10.1557/jmr.2004.19.3.685
  32. 32.
    O. N. Senkov, D. B. Miracle, V. Keppens and P. K. Liaw, Metall. Mater. Trans. A, 2008, 39, (8), 1888 LINK https://doi.org/10.1007/s11661-007-9334-z
  33. 33.
    T. Zhang, Q. Yang, Y. Ji, R. Li, S. Pang, J. Wang and T. Xu, Chin. Sci. Bull., 2011, 56, (36), 3972 LINK https://doi.org/10.1007/s11434-011-4765-8
  34. 34.
    A. Inoue, T. Zhang, A. Takeuchi and W. Zhang, Mater. Trans., JIM, 1996, 37, (4), 636 LINK https://doi.org/10.2320/matertrans1989.37.636
  35. 35.
    L. Zhang, E. Ma and J. Xu, Intermetallics, 2008, 16, (4), 584 LINK https://doi.org/10.1016/j.intermet.2007.12.016
  36. 36.
    T. Xu, S. Pang, H. Li and T. Zhang, J. Non-Cryst. Solids, 2015, 410, 20 LINK https://doi.org/10.1016/j.jnoncrysol.2014.12.006
  37. 37.
    H. Guo, W. Zhang, C. Qin, J. Qiang, M. Chen and A. Inoue, Mater. Trans., JIM, 2009, 50, (6), 1290 LINK https://doi.org/10.2320/matertrans.ME200809
  38. 38.
    K. J. Laws, K. F. Shamlaye and M. Ferry, J. Alloys Compd., 2012, 513, 10 LINK https://doi.org/10.1016/j.jallcom.2011.10.097
  39. 39.
    N. C. Wu, L. Zuo, J. Q. Wang and E. Ma, Acta Mater., 2016, 108, 143 LINK https://doi.org/10.1016/j.actamat.2016.02.012
  40. 40.
    C. W. Corti, ‘Jewellery Alloys – Past, Present and Future’, The Goldsmiths’ Company Jewellery Materials Congress, 8th–9th July, 2019, London, UK, 24 pp LINK https://www.assayofficelondon.co.uk/media/2560/jewellery-alloys-past-present-future-c-corti.pdf
  41. 41.
    J. Schroers, B. Lohwongwatana, W. L. Johnson and A. Peker, Mater. Sci. Eng.: A, 2007, 449–451, 235 LINK https://doi.org/10.1016/j.msea.2006.02.301
  42. 42.
    S. Cardinal, J. Qiao, J. M. Pelletier and H. Kato, Intermetallics, 2015, 63, 73 LINK https://doi.org/10.1016/j.intermet.2015.04.003
  43. 43.
    M. Eisenbart, U. E. Klotz, R. Busch and I. Gallino, Corros. Sci., 2014, 85, 258 LINK https://doi.org/10.1016/j.corsci.2014.04.024
  44. 44.
    M. Eisenbart, U. E. Klotz, R. Busch and I. Gallino, J. Alloys Compd., 2014, 615, (1), S118 LINK https://doi.org/10.1016/j.jallcom.2013.11.167
  45. 45.
    P. Rizzi, I. Corazzari, G. Fiore, I. Fenoglio, B. Fubini, S. Kaciulis and L. Battezzati, Corros. Sci., 2013, 77, 135 LINK https://doi.org/10.1016/j.corsci.2013.07.036
  46. 46.
    S. R. Dale, Platinum Metals Rev., 1993, 37, (3), 159 LINK https://www.technology.matthey.com/article/37/3/159-164/
  47. 47.
    A. Cowley, ‘Pgm Market Report’, Johnson Matthey, London, UK, May, 2019, 56 pp LINK http://www.platinum.matthey.com/services/market-research/may-2019-pgm-market-report
  48. 48.
    J. Brelle, A. Blatter and R. Ziegenhagen, Platinum Metals Rev., 2009, 53, (4), 189 LINK https://www.technology.matthey.com/article/53/4/189-197/
  49. 49.
    T. Biggs, S. S. Taylor and E. Van der Linger, Platinum Metals Rev., 2005, 49, (1), 2 LINK https://www.technology.matthey.com/article/49/1/2-15/
  50. 50.
    G. Ainsley, A. A. Bourne and R. W. E. Rushforth, Platinum Metals Rev., 1978, 22, (3), 78 LINK https://www.technology.matthey.com/article/22/3/78-87/
  51. 51.
    C. Mshumi and C. Lang, Platinum Metals Rev., 2007, 51, (2), 78 LINK https://www.technology.matthey.com/article/51/2/78-82/
  52. 52.
    T. Fryé 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 http://www.santafesymposium.org/2011-santa-fe-symposium-papers/2011-platinum-alloys-in-the-21st-century-a-comparative-study
  53. 53.
    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 http://www.santafesymposium.org/2006-santa-fe-symposium-papers/2006-the-working-properties-for-jewelry-fabrication-using-new-hard-950-palladium-alloys
  54. 54.
    S. Schneider, J. Phys.: Condens. Matter, 2001, 13, (34), 7723 LINK https://doi.org/10.1088/0953-8984/13/34/316
  55. 55.
    A. L. Greer, Nature, 1993, 366, (6453), 303 LINK https://doi.org/10.1038/366303a0
  56. 56.
    T. Egami and W. Waseda, J. Non-Cryst. Solids, 1984, 64, (1–2), 113 LINK https://doi.org/10.1016/0022-3093(84)90210-2
  57. 57.
    A. Inoue, Acta Mater., 2000, 48, (1), 279 LINK https://doi.org/10.1016/S1359-6454(99)00300-6
  58. 58.
    C. A. Angell, J. Non-Cryst. Solids, 1991, 131-133, (1), 13 LINK https://doi.org/10.1016/0022-3093(91)90266-9
  59. 59.
    A. Takeuchi, H. Kato and A. Inoue, Intermetallics, 2010, 18, (4), 406 LINK https://doi.org/10.1016/j.intermet.2009.08.015
  60. 60.
    C. A. Angell, ‘Strong and Fragile Glass Formers’, in “Relaxations in Complex Systems”, eds. K. L. Ngai and G. B. Wright, National Technical Information Service, US Department of Commerce, Springfield, USA, 1985, pp. 3–11
  61. 61.
    O. N. Senkov, Phys. Rev. B, 2007, 76, (10), 104202 LINK https://doi.org/10.1103/PhysRevB.76.104202
  62. 62.
    R. Busch, J. Schroers and W. H. Wang, MRS Bull., 2007, 32, (8), 620 LINK https://doi.org/10.1557/mrs2007.122
  63. 63.
    O. Haruyama, T. Watanabe, K. Yuki, M. Horiuchi, H. Kato and N. Nishiyama, Phys. Rev. B, 2011, 83, (6), 064201 LINK https://doi.org/10.1103/PhysRevB.83.064201
  64. 64.
    G. P. Johari and M. Goldstein, J. Chem. Phys., 1970, 53, (6), 2372 LINK https://doi.org/10.1063/1.1674335
  65. 65.
    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 https://doi.org/10.1103/PhysRevB.89.174204
  66. 66.
    P. Badrinarayanan, W. Zheng, Q. Li, S. L. Simon, J. Non-Cryst. Solids, 2007, 353, (26), 2603 LINK https://doi.org/10.1016/j.jnoncrysol.2007.04.025
  67. 67.
    A. Inoue, Mater. Trans., JIM, 1995, 36, (7), 866 LINK https://doi.org/10.2320/matertrans1989.36.866
  68. 68.
    H. J. Leamy, T. T. Wang and H. S. Chen, Metall. Mater. Trans. B, 1972, 3, 699 LINK https://doi.org/10.1007/BF02642754
  69. 69.
    J. P. Patterson and D. R. H. Jones, Mater. Res. Bull., 1978, 13, (6), 583 LINK https://doi.org/10.1016/0025-5408(78)90182-4
  70. 70.
    F. Spaepen, Scr. Mater., 2006, 54, (3), 363 LINK https://doi.org/10.1016/j.scriptamat.2005.09.046
  71. 71.
    B. Bochtler, O. Kruse and R. Busch, J. Phys. Cond. Matt., 2020, 32, (24), 244002 LINK https://doi.org/10.1088/1361-648X/ab7ad7
  72. 72.
    J. Schroers and W. L. Johnson, Phys. Rev. Lett., 2004, 93, (25), 255506 LINK https://doi.org/10.1103/PhysRevLett.93.255506
  73. 73.
    M. D. Demetriou, M. Floyd, C. Crewdson, J. P. Schramm, G. Garrett and W. L. Johnson, Scr. Mater., 2011, 65, (9), 799 LINK https://doi.org/10.1016/j.scriptamat.2011.07.035
  74. 74.
    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 https://doi.org/10.1038/nmat2930
  75. 75.
    A. L. Greer, Mater. Today, 2009, 12, (1–2), 14 LINK https://doi.org/10.1016/S1369-7021(09)70037-9
  76. 76.
    A. Inoue, N. Nishiyama and H. Kimura, Mater. Trans., JIM, 1997, 38, (2), 179 LINK https://doi.org/10.2320/matertrans1989.38.179
  77. 77.
    C. Park, M. Saito, Y. Waseda, N. Nishiyama and A. Inoue, Mater. Trans., JIM, 1999, 40, (6), 491 LINK https://doi.org/10.2320/matertrans1989.40.491
  78. 78.
    J. Qin, T. Gu, S. Pan, X. Bian and T. Zhang, Sci. China Technol. Sci., 2013, 56, (2), 376 LINK https://doi.org/10.1007/s11431-012-5083-3
  79. 79.
    D. B. Miracle, W. S. Sanders and O. N. Senkov, Philos. Mag., 2003, 83, (20), 2409 LINK https://doi.org/10.1080/1478643031000098828
  80. 80.
    D. B. Miracle, J. Non-Cryst. Solids, 2004, 342, (1–3), 89 LINK https://doi.org/10.1016/j.jnoncrysol.2004.05.017
  81. 81.
    D. B. Miracle, Acta Mater., 2006, 54, (16), 4317 LINK https://doi.org/10.1016/j.actamat.2006.06.002
  82. 82.
    D. B. Miracle, J. Non-Cryst. Solids, 2003, 317, (1–2), 40 LINK https://doi.org/10.1016/S0022-3093(02)01981-6
  83. 83.
    I.-R. Lu, G. Wilde, G. P. Görler and R. Willnecker, J. Non-Cryst. Solids, 1999, 250–252, (2), 577 LINK https://doi.org/10.1016/S0022-3093(99)00135-0
  84. 84.
    J. Schroers and W. L. Johnson, Appl. Phys. Lett., 2000, 77, (8), 1158 LINK https://doi.org/10.1063/1.1289033
  85. 85.
    K. Takenaka, T. Wada, N. Nishiyama, H. Kimura . and A. Inoue, Mater. Trans., 2005, 46, (7), 1720 LINK https://doi.org/10.2320/matertrans.46.1720
  86. 86.
    L. Liu, X. Zhao, C. Ma, S. Pang and T. Zhang, J. Non-Cryst. Solids, 2006, 352, (52–54), 5487 LINK https://doi.org/10.1016/j.jnoncrysol.2006.09.025
  87. 87.
    B. A. Legg, J. Schroers and R. Busch, Acta Mater., 2007, 55, (3), 1109 LINK https://doi.org/10.1016/j.actamat.2006.09.024
  88. 88.
    H. Kazemi, C. Cattin, G. Hodel, T. Pachova and L. Weber, J. Non-Cryst. Solids, 2017, 460, 66 LINK https://doi.org/10.1016/j.jnoncrysol.2017.01.025
  89. 89.
    J. Schroers, Y. Wu, R. Busch and W. L. Johnson, Acta Mater., 2001, 49, (14), 2773 LINK https://doi.org/10.1016/S1359-6454(01)00159-8
  90. 90.
    N. Nishiyama and A. Inoue, Mater. Trans., JIM, 1996, 37, (10), 1531 LINK https://doi.org/10.2320/matertrans1989.37.1531
  91. 91.
    I.-R. Lu, M. Kolbe, G. P. Görler and R. Willnecker, Mater. Sci. Eng.: A, 2004, 375–377, 754 LINK https://doi.org/10.1016/j.msea.2003.10.260
  92. 92.
    N. A. Mauro, M. Blodgett, M. L. Johnson, A. J. Vogt and K. F. Kelton, Nat. Commun., 2014, 5, 4616 LINK https://doi.org/10.1038/ncomms5616
  93. 93.
    G. Duan, A. Wiest, M. L. Lind, J. Li, W.-K. Rhim and W. L. Johnson, Adv. Mater., 2007, 19, (23), 4272 LINK https://doi.org/10.1002/adma.200700969
  94. 94.
    O. Gross, S. S. Riegler, M. Stolpe, B. Bochtler, A. Kuball, S. Hechler, R. Busch and I. Gallino, Acta Mater., 2017, 141, 109 LINK https://doi.org/10.1016/j.actamat.2017.09.013
  95. 95.
    O. Gross, B. Bochtler, M. Stolpe, S. Hechler, W. Hembree, R. Busch and I. Gallino, Acta Mater., 2017, 132, 118 LINK https://doi.org/10.1016/j.actamat.2017.04.030
  96. 96.
    D. Granata, E. Fischer, V. Wessels and J. F. Löffler, Appl. Phys. Lett., 2015, 106, (1), 011902 LINK https://doi.org/10.1063/1.4905174
  97. 97.
    O. Gross, N. Neuber, A. Kuball, B. Bochtler, S. Hechler, M. Frey and R. Busch, Commun. Phys., 2019, 2, 83 LINK https://doi.org/10.1038/s42005-019-0180-2
  98. 98.
    I. Gallino, O. Gross, G. D. Fontana, Z. Evenson and R. Busch, Alloys J. Compd., 2014, 615, (1), S35 LINK https://doi.org/10.1016/j.jallcom.2013.12.006
  99. 99.
    G. Wilde, G. P. Görler and R. Willnecker, Appl. Phys. Lett., 1994, 65, (4), 397 LINK https://doi.org/10.1063/1.112313
  100. 100.
    G. Kumar, S. Prades-Rodel, A. Blatter and J. Schroers, Scr. Mater., 2011, 65, (7), 585 LINK https://doi.org/10.1016/j.scriptamat.2011.06.029
  101. 101.
    G. Kumar, P. Neibecker, Y. H. Liu and J. Schroers, Nat. Commun., 2013, 4, 1536 LINK https://doi.org/10.1038/ncomms2546
  102. 102.
    M. Calin, A. Gebert, A. C. Ghinea, P. F. Gostin, S. Abdi, C. Mickel and J. Eckert, Mater. Sci. Eng. C, 2013, 33, (2), 875 LINK https://doi.org/10.1016/j.msec.2012.11.015
  103. 103.
    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 http://www.santafesymposium.org/2005-santa-fe-symposium-papers/2005-what-is-a-white-gold-progress-on-the-issues
  104. 104.
    N. Nishiyama, K. Takenaka and A. Inoue, Appl. Phys. Lett., 2006, 88, (12), 121908 LINK https://doi.org/10.1063/1.2186512
  105. 105.
    N. Chen, C. L. Qin, G. Q. Xie, D. V. Louzguine-Luzgin and A. Inoue, J. Mater. Res., 2010, 25, (10), 1943 LINK https://doi.org/10.1557/JMR.2010.0246
  106. 106.
    H. Kato, T. Wada, M. Hasegawa, J. Saida, A. Inoue and H. S. Chen, Scr. Mater., 2006, 54, (12), 2023 LINK https://doi.org/10.1016/j.scriptamat.2006.03.025
  107. 107.
    G. R. Garrett, M. D. Demetriou, J. Chen and W. L. Johnson, Appl. Phys. Lett., 2012, 101, (24), 241913 LINK https://doi.org/10.1063/1.4769997
  108. 108.
    N. Nollmann, I. Binkowski, V. Schmidt, H. Rösner and G. Wilde, Scr. Mater., 2016, 111, 119 LINK https://doi.org/10.1016/j.scriptamat.2015.08.030
  109. 109.
    R. Hubek, M. Seleznev, I. Binkowski, M. Peterlechner, S. V. Divinski and G. Wilde, J. Appl. Phys., 2018, 124, (22), 225103 LINK https://doi.org/10.1063/1.5047846
  110. 110.
    R. Hubek, M. Seleznev, I. Binkowski, M. Peterlechner, S. V. Divinski and G. Wilde, J. Appl. Phys., 2020, 127, (11), 115109 LINK https://doi.org/10.1063/1.5142162
  111. 111.
    F. Zhu, H. K. Nguyen, S. X. Song, D. P. B. Aji, A. Hirata, H. Wang, K. Nakajima and M. W. Chen, Nat. Commun., 2016, 7, 11516 LINK https://doi.org/10.1038/ncomms11516
  112. 112.
    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 https://doi.org/10.1016/j.scriptamat.2015.02.026
  113. 113.
    P. Murali and U. Ramamurty, Acta Mater., 2005, 53, (5), 1467 LINK https://doi.org/10.1016/j.actamat.2004.11.040
  114. 114.
    C. C. Gilbert, R. O. Ritchie and W. L. Johnson, Appl. Phys. Lett., 1997, 71, (4), 476 LINK https://doi.org/10.1063/1.119610
  115. 115.
    J. J. Lewandowski, Mater. Trans., 2001, 42, (4), 633 LINK https://doi.org/10.2320/matertrans.42.633
  116. 116.
    M. Vasquez, E. Ascasibar, A. Hernando and O. V. Nielsen, J. Magn. Magn. Mater., 1987, 66, (1), 37 LINK https://doi.org/10.1016/0304-8853(87)90125-9
  117. 117.
    I. W. Donald, H. A. Davies and T. Kemény, J. Non-Cryst. Solids, 1982, 50, (3), 351 LINK https://doi.org/10.1016/0022-3093(82)90095-3
  118. 118.
    A. Inoue, T. Aoki and H. Kimura, Mater. Trans., JIM, 1997, 38, (2), 175 LINK https://doi.org/10.2320/matertrans1989.38.175
  119. 119.
    H. Kazemi, “Alloy Development of a New Platinum-Based Bulk Metallic Glass”, PhD Thesis, École Polytechnique Fédérale de Lausanne, Switzerland, 3rd March, 2017 LINK https://infoscience.epfl.ch/record/225963?ln=en
  120. 120.
    Z. P. Lu and C. T. Liu, J. Mater. Sci., 2004, 39, 3965 LINK https://doi.org/10.1023/B:JMSC.0000031478.73621.64
  121. 121.
    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 https://doi.org/10.2320/matertrans.43.3206
  122. 122.
    N. Chen, Y. Li and K.-F. Yao, J. Alloys Compd., 2010, 504, (S1), S211 LINK https://doi.org/10.1016/j.jallcom.2010.02.079
  123. 123.
    W. Zhang, H. Guo, Y. Li, Y. Wang, H. Wang, M. Chen and S. Yamaura, J. Alloys Compd., 2014, 617, 310 LINK https://doi.org/10.1016/j.jallcom.2014.07.214
  124. 124.
    L. Liu, A. Inoue and T. Zhang, Mater. Trans., 2005, 46, (2), 376 LINK https://doi.org/10.2320/matertrans.46.376
  125. 125.
    N. Nishiyama, K. Takenaka, T. Wada, H. Kimura and A. Inoue, Mater. Trans., 2005, 46, (12), 2807 LINK https://doi.org/10.2320/matertrans.46.2807
  126. 126.
    A. Kuball, B. Bochtler, O. Gross, V. Pacheco, M. Stolpe, S. Hechler and R. Busch, Acta Mater., 2018, 158, 13 LINK https://doi.org/10.1016/j.actamat.2018.07.039
  127. 127.
    Y. Li, S. Zhao, Y. Liu, P. Gong and J. Schroers, ACS Comb. Sci., 2017, 19, (11), 687 LINK https://doi.org/10.1021/acscombsci.7b00048
  128. 128.
    S. Ding, J. Gregoire, J. J. Vlassak and J. Schroers, J. Appl. Phys., 2012, 111, (11), 114901 LINK https://doi.org/10.1063/1.4722996
  129. 129.
    S. Ding, Y. Liu, Y. Li, Z. Liu, S. Sohn, F. J. Walker and J. Schroers, Nature Mater., 2014, 13, 494 LINK https://doi.org/10.1038/nmat3939
  130. 130.
    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 https://doi.org/10.1038/s41586-019-1145-z
  131. 131.
    E. Perim, D. Lee, Y. Liu, C. Toher, P. Gong, Y. Li, W. N. Simmons, O. Levy, J. J. Vlassak, J. Schroers and S. Curtarolo, Nat. Commun., 2016, 7, 12315 LINK https://doi.org/10.1038/ncomms12315
  132. 132.
    J. J. Han, C. P. Wang, J. Wang, X. J. Liu, Y. Wang and Z. K. Liu, Mater. Des., 2017, 126, 47 LINK https://doi.org/10.1016/j.matdes.2017.04.030
  133. 133.
    D. C. Ford, D. Hicks, C. Oses, C. Toher and S. Curtarolo, Acta Mater., 2019, 176, 297 LINK https://doi.org/10.1016/j.actamat.2019.07.008
 

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.

Read more from this issue »

BACK TO TOP

SHARE THIS PAGE: