Journal Archive

Johnson Matthey Technol. Rev., 2021, 65, (4), 535
doi: 10.1595/205651321X16221908118376

Research Progress of Platinum-Based Superalloys for High Temperature Applications

Platinum-aluminium ternary, quaternary and multiple alloys with excellent prospects for future use

  • Changyi Hu, Yan Wei*, Hongzhong Cai, Li Chen, Xian Wang, Xuxiang Zhang, Guixue Zhang, Xingqiang Wang
  • State Key Laboratory of Advanced Technologies for Comprehensive Utilization of Platinum Metals, Kunming Institute of Precious Metals, Kunming 650106, China
  • *Email:

Article Synopsis

Platinum-based alloys are being developed for high-temperature applications with the aim of replacing some of the currently used nickel-based superalloys (NBSAs) and benchmark alloy, PM2000. The platinum-based superalloys have a similar structure to the NBSAs and can potentially be used at higher temperatures and in more aggressive environments because platinum is more chemically inert and has a higher melting point. In this paper, the recent progress in research and development of platinum-based superalloys is overviewed. Firstly, the composition optimisation and structural design of platinum-base superalloys are introduced. The structural characteristics, mechanical properties, oxidation resistance and corrosion behaviour of platinum-aluminium ternary, quaternary and multiple superalloys are summarised. Finally, directions for further research and application of platinum-based superalloys are analysed and prospected.

1. Introduction

Components in many different applications at high temperature and in corrosive environments require materials with excellent high temperature mechanical properties and chemical resistance. Aerospace applications represent an extremely challenging field, for the development of new materials, and for the improvement of the existing ones (1). NBSAs with the Ni3Al intermetallic compounds as the main strengthening phase have been widely used in high-temperature components such as aeroplane engines, industrial gas turbine blades and modern industry fields. After decades of development, the working temperature of NBSAs is about 1100°C (2) and can reach up to 1150°C (3, 4). Recently, a new platinum-modified nickel-base alloy with exceptional high temperature stability has been identified (5, 6). Coarsening studies conducted reveal unusually high volume fractions of morphologically stable γ’ precipitates up to 1200°C, which suggests that the alloy would have excellent performance as a bond coat or single crystal blade. A further increase in the operating temperature of the gas turbine will improve the combustion efficiency, reduce fuel consumption and CO2 emissions, and leads to higher thrust values (7, 8). NBSAs operating temperature is approximately 85% of their melting point. An increasing interest has been shown in developing new alloys based on materials with higher melting points with similar structure to that of NBSAs and capable of being used at 1300°C (9, 10).

Potential candidate materials that may replace NBSAs or iron-based superalloys such as PM2000 mainly include intermetallic compounds, ceramics and ceramic matrix composites, refractory metals and high melting point platinum group metals (11). Some high melting point intermetallic compounds have high-temperature strength, lower diffusion rate and excellent corrosion resistance. However they lack plasticity or fracture toughness at room temperature (12, 13). Components based on ceramics and ceramic matrix composites have high-temperature strength, creep resistance, oxidation resistance and corrosion resistance which are attractive for their potential use in gas turbine engines. Unfortunately, low fracture toughness and brittle behaviour usually associated with ceramics are problems for high-temperature applications. Refractory metals and their alloys have high melting temperatures and for this reason researchers are considering the possibility of using them in the hot parts of gas turbines to replace NBSAs. These refractory metals and alloys lack sufficient oxidation resistance which limits their practical application (14). Superalloys based on platinum group metals (platinum, iridium, rhodium) show extremely strong chemical stability and two-phase structure (face-centred cubic (fcc)/L12), which makes this group of alloys a potential candidate to be developed as high-temperature materials for next-generation gas turbines (1522). However, the main weaknesses of most iridium-based and rhodium-based refractory superalloys are brittleness, high cost and high density.

Unlike iridium and rhodium, platinum has become an essential high-temperature material in special applications. It has a higher melting point than nickel (platinum = 1769°C, nickel = 1453°C), better oxidation resistance, corrosion resistance, chemical stability and does not require coating protection when used at high temperatures (23, 24). Platinum-based alloys have excellent mechanical properties such as high creep strength and ductility, which make them have application potential in the fields of chemical engineering, space technology and glass industry (24, 25). At this time, the research on platinum-based superalloys includes solid solution strengthened, dispersion strengthened and precipitation strengthened alloys as well as platinum group metal compounds (26). Current usage is restricted to solid solution strengthened alloys and dispersion strengthened alloys, the latter being classified as part of the group of composite materials. Solid solution strengthened platinum-based alloys are a family of alloys which have been researched and developed for some time. The compositions and preparation process can be said to be more mature.

All transition group elements have considerable solid solubility in platinum. The elements near platinum in the periodic table form a continuous solid solution with platinum and have different degrees of solid solution strengthening effect on the platinum matrix. At high temperatures, ruthenium, iridium and rhodium have higher tensile strength than platinum and palladium. The high temperature durability and creep rupture strength of iridium are also much higher than that of Pt-Rh alloys. Ruthenium, iridium, rhodium and palladium have become the main solid solution strengthening elements. According to the relationship between stacking fault energy and creep rate, ruthenium and iridium have the largest solid solution strengthening effect on platinum, followed by rhodium, and palladium with the smallest effect (27).

Currently, the research and development on platinum-based solid solution alloys mainly include binary alloys such as Pt-Rh, Pt-Ir, Pt-Ru, Pt-Ni, Pt-W and ternary alloys such as Pt-Pd-Rh and Pt‐Rh-Ru (28). The properties of Pt-Rh alloys are the most stable: an increase in rhodium content leads to higher temperature durability, extended creep life and decreased creep rate (29). However, the improvement of mechanical properties decreases when rhodium content exceeds 30 wt%. In addition, machinability is significantly worse at these levels of rhodium content. The high temperature durability, creep life and creep rate of Pt-Ir alloy are better than the Pt-Rh alloy. On the other hand, Pt-Ir alloys tend to have higher weight losses in oxidising environments above 1100°C due to the selective oxidation of iridium after prolonged exposure to such atmospheres. In addition, for solid-solution strengthened Pt-Rh alloys, the coarsening of crystal grains at high temperatures will lead to reduced alloy strength and premature failure of components.

In order to improve the high-temperature mechanical properties of platinum-based alloys, oxide dispersion strengthened (ODS) alloys with ZrO2 or Y2O3 as reinforcing phase have been researched and developed (30, 31). The fine oxide particles dispersed in the platinum matrix can stabilise the grain boundaries, prevent the movement of dislocations and improve the high temperature fracture strength. However, these ODS alloys have great brittleness, crack sensitivity and cannot withstand severe temperature changes. A particular type of ODS alloys has been developed, namely dispersion hardened platinum (DPH) alloys (32). These alloys are reinforced by dispersion of oxides formed inside the alloy via an internal oxidation process of pure oxygen-reactive elements (i.e., elements with high affinity to oxygen such as cerium, yttrium, scandium and zirconium). These elements are added to the melt and their oxidation is subsequently obtained by appropriate treatments. The stress-fracture strength of DPH alloys is further improved with respect to solution strengthened alloys. For instance, the stress-fracture strength of DPH platinum is higher than that of Pt-10Rh alloy (33), with good plasticity. However, the production process of all the dispersion strengthened alloy families is complicated and problems may arise during welding operations. The strength of the welded joint tends to drop due to excessive formation of oxide particles during welding (24).

Precipitation strengthening is the main strengthening mechanism of platinum-based superalloys currently under research and development. These alloys show higher temperature strength than solution strengthened and dispersion strengthened alloys at 1300°C (3439). However, high density and cost are the major drawbacks to the use of platinum, but it is likely that the platinum-based alloys can be used for the highest application temperature components (10, 40). Due to the excellent properties of oxidation and corrosion resistance, the Pt-Al system superalloys could have potential as coatings on NBSAs or other substrates (10, 22, 41). Their density can be slightly reduced by adding suitable light alloying elements, while the high performance and recyclability of platinum-based alloys make up for the high price (42). This paper summarises the research status and progress of platinum-based superalloy materials. Firstly, we introduce the composition and structural optimisation design of platinum-based superalloys, the structural characteristics and evolution of Pt‐Al-based ternary, quaternary and multi-element superalloys, and their mechanical properties, oxidation and corrosion resistance behaviours. The strengthening mechanisms, the relationship between oxidation, corrosion and alloy composition have been analysed and the results will be presented and discussed. Finally, further research and application prospects of platinum-based superalloys are analysed and discussed.

2. Structure and Composition Design of Platinum-Based Superalloys

A large number of new alloys have been researched and developed based on strong demand for higher working temperature and high temperature resistant structural materials in the aerospace field. The goal of the research is to seek a material with better high-temperature mechanical properties (tensile, fracture, creep and thermomechanical fatigue properties) and environmental stability (resistance to high-temperature oxidation and hot corrosion) than nickel-based superalloys (43). Inspired by the successful experience of precipitation strengthening obtained in the γ matrix (fcc structure) of nickel-based superalloys, much effort is now put into the research and development of platinum-based superalloys with similar structures to the γ/γ’ system found in nickel-based superalloys (7). Research on platinum-based superalloys began at the end of the 20th century. The initial work was mainly on structure and composition design, including the formation of a Pt3X (γ’) precipitation strengthening phase and the selection of solid solution strengthening elements.

2.1 Second Phases and Possible Reinforcement

Platinum can form Pt3X and Pt5X high melting point intermetallic compounds on the platinum-rich side of the phase diagram with transition metals and rare earth metals (44), such as Pt3Al (1550°C), Pt3Hf (2250°C), Pt3Sc (1850°C), Pt3Y (2020°C) and Pt3Zr (2250°C). Their melting points are higher than Ni3Al (1390°C). Most platinum compound precipitation phases with γ’ structure have a high melting point, high thermal conductivity, low thermal expansion coefficient, high strength and a large number of possible slip systems. It can be expected that γ/γ’ platinum alloys will have high thermal strength and precipitation strengthening effects, leading to the possible development of a new generation of precipitation strengthened platinum-based superalloys (45). Figure 1 represents the binary Pt-Al phase diagram (44). The diagram shows that the maximum solubility of aluminium in platinum is about 10 at% at relatively low temperatures, whereas at the eutectic temperature (1507°C) the value is slightly higher. The Pt3Al phase forms at the eutectic temperature and the eutectic reaction can be found in the top right of the diagram. At high temperatures, the Pt3Al intermetallic compound shows a L12 structure. It transforms into a tetragonal structure (D0’c) at lower temperatures. Pt3Al is the most important intermetallic phase.

Fig. 1.

Binary Pt-Al equilibrium phase diagram. Reproduced with permission of ASM International from (44), Copyright 1990; permission conveyed through Copyright Clearance Center, Inc

Table I lists the common precipitation phases and main performance evaluations in platinum-based alloys (46). The Pt3X phase mainly appears in platinum-transition metals and platinum-simple metal alloy systems, while Pt5X mainly appears in platinum-rare earth and platinum-alkaline earth metal alloy systems. The Pt3X phase formed by chromium, vanadium and platinum decomposes at 1130°C and 1015°C respectively, and was excluded from the design study of platinum-based superalloys. Although tin, lead, gallium and other elements can form a L12 structure phase with platinum, they are also excluded due to the low melting point. Initial research on binary systems such as Pt-Zr, Pt-Hf and ternary system alloys such as Pt-Rh-Zr, Pt-Rh-Hf showed that γ’ phases Pt3Zr and Pt3Hf are formed in the alloy which are coherent with the matrix (47). The presence of zirconium and hafnium as solid solution strengthening elements in combination with γ’ precipitation strengthening give these alloys high-temperature mechanical properties, but they all have the problem of poor oxidation resistance. The formation of brittle zirconium and hafnium oxides leads to embrittlement of the material (48).

Table I

Candidate Elements to Form Pt3X Precipitates (46)

Element Structure of Pt3X Melting range, °C Environmental resistance Density, g cm–3 Attributes
Aluminium Low temperature: tetragonal (D024)High temperature: fcc (L12) 1500–1769 Good oxidation resistance. Forms stable external Al2O3 scale which protects the metal from internal oxidation 2.7 Low density. Transformation of Pt3Al at 1290°C. Good oxidation resistance. Lowers solidus temperature
Titanium Tetragonal (D024) at stoichiometric compositions; fcc (L12) in platinum-rich alloys 1769–1800 Prone to oxidation even at low temperatures 4.5 L12 structure. Favourable density. Increases solidus temperature
Vanadium Fcc (L12) 1769–1805 Vanadium absorbs relatively large amounts of oxygen 5.8 The Pt3V phase is only stable to 1015°C
Chromium Fcc (L12) 1769–1785 Chromium has a beneficial effect on hot corrosion and oxidation resistance 7.19 The Pt3Cr phase is only stable to 1130°C
Gallium Fcc (L12) 1373–1769 The effect of gallium additions on high temperature oxidation of alloys is uncertain 5.91 Low melting temperature
Yttrium Fcc (L12). No two-phase (Pt)+Pt3Y region because of the Pt5Y phase 1615–1769 The effect of major yttrium additions on the environmental behaviour of alloys has not been documented 4.5 Highly reactive and difficult to process
Zirconium Tetragonal (D024) at stoichiometric compositions. Fcc (L12) in platinum-rich alloys 1769–1963 Exposure to oxygen causes embrittlement due to the formation of brittle oxides 6.4 L12 structure. Generally embrittling in alloys. Susceptible to oxidation
Niobium <~1100°C: orthorhombic 1769–2000 Oxidises substantially at T >500°C 8.55 Raises solidus temperature. Probably only partially coherent Pt3Nb
>~1100°C: tetragonal (D024)
Tin Fcc (L12) 1365–1769 The high-temperature oxidation behaviour of tin-based alloys is unknown 7.29 Low melting point and is not suitable for high-temperature use
Hafnium Not reported. Pt-rich Pt3Hf with L12 structure has been reported 1769–2000 Exposure to oxygen causes embrittlement due to the formation of brittle oxides 13.1 L12 structure. Highly reactive and difficulty to process
Tantalum Monoclinic (L60) 1769–1970 Prone to rapid oxidation at T >500°C 16.6 Raises solidus temperature. Good high-temperature mechanical properties
Lead Fcc (L12) 915–1769 The high temperature oxidation behaviour of lead has not been extensively studied 11.3 Low melting temperature

The refractory metals tantalum and niobium can increase the solidus temperature of platinum-based superalloys and are worthy of further study. Alloying with aluminium leads to the formation of Pt3Al, a strengthening intermetallic of the Pt3X type that has two different crystal structures: the cubic structure (L12) at high temperatures and allotrope tetragonal structure (D024) at low temperatures. The high-temperature allotrope can be stabilised at lower temperatures by adjusting the composition of the alloy matrix.

It is necessary to study the structure and properties of the Pt3X second phase particles in order to develop precipitation strengthened platinum-based superalloys. However, there are few reports on the experimental measurement of the mechanical properties of Pt3X intermetallic compounds. Adjal (49), Pan (50) and Li (51) investigated the electronic structure, thermodynamic properties, oxidation resistance and mechanical properties of Pt3X intermetallic compounds by simulation calculation methods. The mechanical properties and anisotropy of Pt3M alloy were determined from these studies. The anisotropy is mainly derived from the d electronic state of platinum and the d electronic state (or p electronic state) of M. Pt3Hf has the highest modulus (bulk modulus, shear modulus, Young’s modulus) and hardness, while Pt3Y has the lowest values (51). Liu (52) used first-principles calculations to study the effect of pressure on the structure and mechanical properties of Pt3Al. The study found the elastic modulus, bulk modulus and shear modulus of Pt3Al increase linearly with increasing pressure. Pt3Al changes from tetragonal to cubic structure when the pressure reaches 60 MPa, indicating that the cubic structure of Pt3Al has higher resistance to volume deformation.

2.2 Alloy Composition Design

Table II lists the candidate solid solution strengthening elements and main performance evaluations of platinum-based superalloys (46). Platinum group metals ruthenium, iridium, rhodium, palladium, nickel and rhenium can all be used as solid solution strengthening elements. The alloys Pt-Al-Ni, Pt-Al-Ru, Pt-Al-Cr, Pt-Al-Ti, Pt‐Al‐Re, Pt-Ti-Ru, Pt-Ti-Re, Pt-Ta-Ru, Pt-Ta-Re and Pt‐Nb‐Ru were selected through the comprehensive evaluation of precipitation strengthening, solid solution strengthening and alloy properties. These alloys were optimised and screened and are all characterised by a fcc (platinum) solid solution matrix and fcc (L12) Pt3X precipitation phase, thus a two-phase microstructure. Investigations on structural composition, mechanical properties and oxidation resistance were performed. The results showed that the two-phase structure of Pt-Al-X and Pt-Ti-X alloys confers significant precipitation strengthening, with hardness exceeding 400 Vickers hardness (HV1) and strong resistance to crack initiation and propagation. Aluminium-containing alloys also have better oxidation resistance than the other alloys due to the formation of an aluminium oxide-based protective film on the surface. Internal oxidation was observed in titanium-containing alloys. Therefore, aluminium is considered to be a necessary alloying element for the development of oxidation-resistant platinum-based superalloys (46, 53). For these reasons, the subsequent research mainly focuses on the Pt-Al-X alloy system.

Table II

Candidate Solid-Solution Strengthening Elements (46)

Element Melting range, °C Environmental resistance Density, g cm–3 Attributes
Nickel 1455–2447 A dense, almost pore-free layer of NiO is formed on exposure to oxygen. The diffusion rate of nickel atoms through this layer is low and this limits oxidation 8.9 Similar electronic structure to platinum. Reportedly one of the strongest solid-solution straighteners in platinum. Low cost. Low melting temperature
Ruthenium 1769–2100 Forms volatile oxides above ~1100°C resulting in pgms grain boundary embrittlement 12.2 Greater solid-solution strengthening effect than others because of its hexagonal structure. Reasonable cost
Rhodium 1769–1963 Comparable to platinum. Intercrystalline oxidation does occur but can be controlled by alloying 12.4 Extremely high cost (about four times that of platinum). Limited resources would not be able to sustain a large increase in demand
Palladium 1555–1769 Highest vaporisation rate of all the pgms in the presence of oxygen 12.0 Low melting point. High cost
Rhenium 1769–2450 Pure rhenium is resistant to oxidation up to 1000°C. Above this temperature volatile oxides are formed. Rhenium had been reported to improve the hot corrosion resistance of nickel-based superalloys 21.0 High cost
Iridium 1769–2447 Forms volatile oxides >1196°C. Above 1100°C iridium has a superior oxidation resistance to ruthenium 22.5 High cost. Limited resources would not be able to sustain large increase in demand. Known to have excellent high-temperature mechanical properties

3. Pt-Al-X Ternary Alloy

3.1 Structural Characteristics of Ternary Alloys

In order to obtain an effective strengthening effect on platinum-based ternary alloys it is necessary to determine the low-temperature structure of the Pt3Al phase in the representative binary alloy. Wolff (22) used an electric arc furnace (EAF) to smelt the Pt-12Al (at%) alloy and the material was then subjected to 96 h solution annealing and subsequent ageing at 1350°C. Figure 2 shows some structural features of the Pt-12Al alloy. It can be seen that Pt3Al precipitation phase exhibits strong directional cubic alignment and the dispersed sub-micron nature of the precipitation phase exhibits a typical bimodal size distribution. This special-orientation geometrical configuration is a highly coherent phase which has low mismatch strain. The Pt3Al-γ’ phase additionally has a low-temperature variant and a lath or twin structure.

Fig. 2.

TEM bright-field image of cuboidal Pt3Al precipitates in Pt-12A1 alloy after solution annealing and subsequent ageing at 1350°C (22). The magnification is 10,000 times

The maximum volume fraction of the Pt3Al phase is limited to about 30% due to the maximum solubility of aluminium in platinum. The volume ratio of the precipitation phase in nickel-based superalloys can be as high as 70–80%. Additional information about the microstructural features of Pt3Al can be inferred from Figure 3 where a transmission electron microscopy (TEM) bright-field image microstructure characteristic of the precipitate phase in an alloy with slightly higher aluminium content, namely Pt-14Al (at%) alloy (54). It can be seen from Figure 3 that the martensitic transformation has occurred in the platinum matrix leading to the formation of a clear stacked sheet or plate-like structure. Electron diffraction analysis shows that the stacked plates are twinned with each other, the [001] directions of adjacent tetragonal structure D0’c single packages are perpendicular to each other, and the twin planes are (112) planes. The [001] direction of D0’c is parallel to the <001> direction of the cubic matrix.

Fig. 3.

Bright-field TEM image of a Pt3Al precipitate in a Pt-14Al alloy. Stacked plates or laths e.g. P1 and P2; M = platinum matrix. The arrows indicate individual platelets between stacked plates. Reprinted from (54), Copyright 2007, with permission from Elsevier

Adding transition metals such as nickel, titanium, chromium, ruthenium, iridium, rhenium and tantalum to Pt-Al alloys can improve the strengthening effect of γ’(Pt3Al) precipitates, thermal stability, solid solution strengthening of the matrix and overall properties. Several typical ternary alloy microstructures are shown in Figure 4. The two-phase structure of Pt3Al and platinum solid solution of all alloys has been confirmed by X-ray diffraction (XRD) experiments (55). The alloys with nominal compositions (at%) of Pt-14Al-3Re, Pt-14Al-4Ti, Pt-14Al-4Ta and Pt‐14Al‐4Cr were all characterised by a microstructure consisting of primary Pt3Al surrounded by a fine two-phase eutectic-like mixture of a (platinum) matrix and fine particles of Pt3Al. The proportion of primary Pt3Al in Pt-14Al-4Ti, Pt-14Al-4Ta and Pt-14Al‐4Cr alloys is between 40% and 50% and in Pt‐14Al‐3Re alloy it is about 25% according to optical microstructure analysis. The fine martensite-like lamellar structure was observed in Pt-14Al-3Re (Figure 4(a)) and Pt-22Al-2Ru (Figure 4(e)) alloys. This means the Pt3Al phase has transformed from the high-temperature cubic structure (L12) to the low-temperature tetragonal structure (D0’c). In the Pt-14Al-4Ti and Pt-14Al-4Ta alloys, it seems that the high-temperature Pt3Al phase is formed and maintained. This demonstrates that the third metal elements titanium and tantalum can stabilise the L12 polymorph. It is important to observe that an alloying element must enter the Pt3Al phase in order to prevent the low-temperature transformation of the Pt3Al phase (55). Biggs (56) and Hill (57) have also shown the possibility for other alloy third components (nickel, titanium or chromium) in Pt-Al alloys to stabilise the cubic (L12) structure of Pt3Al.

Fig. 4.

Scanning electron microscopy-backscattered electron micrographs (SEM-BSE) showing the microstructures of the Pt-Al-X alloys: (a) Pt-14Al-3Re; (b) Pt-14Al-4Ti; (c) Pt-14Al-4Ta; (d) Pt-14Al-4Cr; (e) Pt-22Al-2Ru. Reprinted from (55), Copyright 2001, with permission from Elsevier

Hill et al. (58) studied the microstructure and lattice mismatch of Pt-Al-X alloy systems with stabilising elements titanium, chromium, tantalum, ruthenium and iridium. Figure 5 shows TEM images of the typical two-phase microstructure. All the precipitated phases show a bimodal or even trimodal size distribution. Titanium, chromium and tantalum elements enter into the Pt3Al phase to stabilise the cubic L12 structure, the precipitation phase has a cubic appearance with no clearly discernible internal structure (Figure 5(a)). On the contrary, when ruthenium and iridium enter the platinum matrix, the Pt3Al precipitation phase transforms into a D0’c structure, and presents a band-like structure with alternating light and dark distributions (Figure 5(b)). Under higher magnification, it was confirmed that the lamellar structure in the precipitation phase of the D0’c tetragonal structure should belong to the twin structure (54).

Fig. 5.

Typical TEM images of the ~Pt3Al precipitates in Pt-Al-X alloys (X = chromium, iridium, ruthenium, tantalum or titanium), with letters indicating the different size ranges (where P = primary; I, T = intermediate; S = secondary): (a) L12 precipitates stabilised by chromium, tantalum and titanium additions. The inset shows the selected area diffraction (SAD) pattern, confirming the L12 structure; (b) D0’c precipitates stabilised by iridium and ruthenium additions. Reused from (58), Copyright © 2001 by The Minerals, Metals & Materials Society. Used with permission

The lattice misfits between the matrix and precipitate phase in the Pt-Al-X alloy system at room temperature and 800°C were measured respectively by XRD (58). The (220), (211) and (112) diffraction peaks are used to obtain the lattice constants of the platinum solid solution matrix, L12-Pt3Al and D0’c-Pt3Al precipitate phases, respectively (amatrix and appt represent the lattice constants of the platinum solid solution matrix and Pt3Al precipitation phase, respectively). Then, the Lattice misfits δ between the precipitation phase and the matrix are calculated by Equation (i), and the results are listed in Table III. It can be seen that the degrees of mismatch for all alloys are negative and the difference in mismatch degrees at different temperatures is very small. Besides alloys containing ruthenium, the lattice misfit increases at high temperature and the cubic L12 structure (alloys containing chromium, tantalum or titanium) has a lower degree of mismatch than the D0’c structure (alloys containing iridium, ruthenium).



Table III

Lattice Misfits Between Precipitates and Matrix for Selected Pt-Al-X Ternary Alloys (58)

Room temperature
Alloy Pt3Al type amatrix, nm appt, nm δ amatrix, nm appt, nm δ
Pt-10Al-4Cr L12 3.9022 3.8741 –0.0072 3.9390 3.9103 –0.0073
Pt-10Al-4Ir D0’c 3.8983 3.8507 –0.0123 3.9246 3.8747 –0.0128
Pt-10Al-4Ru D0’c 3.9001 3.8530 –0.0121 3.9349 3.8967 –0.0098
Pt-10Al-4Ta L12 3.8941 3.8682 –0.0067 3.9246 3.8961 –0.0073
Pt-10Al-4Ti L12 3.8921 3.8642 –0.0072 3.9246 3.8961 –0.0073

3.2 Mechanical Properties of Ternary Alloys

Figure 6 shows the high-temperature compression strength of Pt-10Al-4Ru compared to Mar-M247 (a nickel-based superalloy) and the tensile strength of PM2000 (an iron-based superalloy) (22). It can be seen that the Pt-10Al-4Ru alloy based on γ/γ〉 precipitation strengthening has higher compressive strength and the ability of withstand higher temperatures than the traditional nickel-based and iron-based superalloys at 1200°C.

Fig. 6.

High-temperature compression strength of Pt-10Al-4Ru alloy compared to Mar-M247 (a nickel-based superalloy) and the tensile strength of PM2000 (an iron-based superalloy) (22)

Süss (59) studied the stress-rupture strength and high-temperature creep properties of the Pt-Al-X (X = chromium, ruthenium, iridium) ternary alloy system. Figure 7 shows the stress-rupture curves of Pt-10Al-4Ru and Pt-10Al-4Cr alloys at 1300°C. The alloys’ interpolated strength levels for a rupture time of 10 h are summarised in Table IV. PM2000 shows the highest high temperature rupture strength among all the tested alloys, but the lower slope of the stress-rupture curve indicates that the alloy has high stress sensitivity and brittle creep behaviour. This means PM2000 alloy is likely to be damaged by stress concentration or short-term overload in practical use. In contrast, the stress-rupture curve of platinum-based alloys has a steeper slope. Pt-10Al-4Cr alloy has the highest strength at 1300°C. The high-temperature durability of this precipitation strengthened alloy at 1300°C is higher than the ODS and DPH alloys (53) and the solid solution-strengthened Pt‐20wt%Rh alloys, and is close to the strength of Pt-30wt%Rh alloys (10). However, due to wide price fluctuations of rhodium as well as processing difficulties, the practical application of Pt-30wt%Rh alloy is restricted (10).

Fig. 7.

Stress-rupture curves of PM2000 and Pt-10Al-4X (X = chromium, ruthenium, iridium) alloys at 1300°C in air. Reprinted from (59), Copyright 2002, with permission from Elsevier

Table IV

Stress-Rupture Strength Rm for PM2000 and Pt-10Al-4X Alloys (59)

Alloy Stress rupture strength (Rm /10 h/1300°C), MPa
PM2000 25
Pt-10Al-4Cr 17
Pt-10Al-4Ru 15
Pt-10Al-4Ir 13

Figure 8 shows the creep curves of the above alloys under the test conditions of 1300°C and 30 MPa (59). The initial stage of creep was not observed in the three platinum-based alloys. After the second-stage creep, the platinum-based alloys undergo the third stage of creep and subsequent rupture. The creep rupture strain was as high as approximately 10–30%. For PM2000 alloy, it is impossible to divide the creep curve into different stages because of the very low creep rate and fracture strain of less than 1%.

Fig. 8.

Creep curves of PM2000 and Pt-10Al-4X (X = chromium, ruthenium, iridium) alloys tested at 1300°C and 30 MPa. Reprinted from (59), Copyright 2002, with permission from Elsevier

3.3 Oxidation Behaviour of Ternary Alloy

The service environment for components such as aeroplane engines, industrial gas turbine blades and aerospace engine thrusters is very harsh: high temperature, strong oxidation and corrosion. Although platinum-based superalloys with potential for high-temperature applications have been confirmed from the perspective of microstructure features and high-temperature mechanical properties, further attention must be given to assessing their oxidation and corrosion behaviour (10). Hill et al. (46) conducted oxidation tests on Pt-Al-Ni, Pt-Al-Ru, Pt-Al-Re, Pt-Nb-Ru and Pt-Ti-Ru in flowing air at 900°C, 1100°C, 1300°C and 1400°C, respectively. The study found that all aluminium-containing alloys show negligible weight loss on oxidation, while the Pt-Ti-Ru and Pt-Nb‐Ru alloy systems have significant mass increases, indicating lower resistance to oxidation. Figure 9 shows optical micrographs of the transverse-sectional morphology of several platinum-based alloys (46). It can be seen that the Pt-Ti-Ru alloy has grain boundary oxidation, while severe internal oxidation has occurred in the Pt-Nb-Ru and Pt‐Ta‐Re alloys. The aluminium-containing alloy has formed a protective aluminium oxide film on its outer surface. This layer prevents oxidation of the underlying metal, conferring better oxidation resistance on this alloy.

Fig. 9.

Optical micrographs for comparison of the transverse sections of the oxidised samples: (a) Pt-14Al-8Ru; (b) Pt-23Ti-7Ru; (c) Pt-24Ta-4Re; (d) Pt-24Nb-3Ru. Reprinted from (46), Copyright 2001, with permission from Elsevier

Figure 10 shows the isothermal oxidation curves of aluminium-containing platinum-based ternary alloys Pt-10Al-4X (X = chromium, iridium, ruthenium, titanium) and PM2000 iron-based superalloys at 1350°C (10). It can be seen that Pt-10Al-4Ti and Pt-10Al-4Ru exhibit a parabolic oxidation law similar to the PM2000 alloy. The Pt‐10Al-4Ir and Pt-10Al-4Cr alloys exhibit a parabolic change in the initial stage of oxidation, and the oxidation rate is relatively high. After that, the oxide film is grown at a logarithmic rate. After 800 h of oxidation a continuous oxide layer is obtained, which has better oxidation resistance than the PM2000 alloy. Experimental research on the microstructure, mechanical properties and oxidation resistance of the Pt-Al-X series of superalloys and comparison with the PM2000 benchmark alloy leads to the conclusion that the highest performing platinum-based ternary superalloys are Pt-10Al-4Cr and Pt-10Al-4Ru (10).

Fig. 10.

Results of the isothermal oxidation tests conducted on Pt-Al-X (X = chromium, iridium, ruthenium, titanium) alloys at 1350°C (10)

4. Pt-Al-Cr-X Quaternary and Multi-Element Alloys

Early studies on platinum-based superalloys mainly focused on the addition of alloying elements to improve oxidation resistance and ensure the γ’ phase has a stable L12 structure. Research shows that the volume fraction of the precipitation phase reaches only about 30% no matter how the heat treatment process is optimised. Hence it is difficult to obtain the desired strength of the alloy (60). Subsequent work mainly focused on the Pt-Al-Cr ternary alloy system. Adding nickel, ruthenium and other alloying elements to form quaternary and multi-element platinum-based superalloys is expected to further improve the microstructure and to enhance the mechanical properties and oxidation resistance. Compared with steel, nickel-based alloys and aluminium alloys, the experimental data and phase diagrams of platinum-based alloys are relatively lacking. In order to develop quaternary and multi-element platinum-based alloys, relevant research institutions in South Africa, Germany and the UK have collaborated to establish Pt-Al‐Ru (61) and Pt-Cr-Ru (62) ternary system and Pt‐Al‐Cr‐Ni (63) quaternary system alloy databases by experiment and first-principles thermodynamic calculations.

4.1 Pt-Al-Cr-Ni Quaternary Alloy

Nickel has a good solid solution strengthening effect on the platinum matrix (64) and its addition can stabilise the L12 structure of the Pt3Al phase. Researchers have added nickel to the Pt-Al-Cr alloy to form a Pt-Al-Cr-Ni quaternary alloy (36). In order to reflect the results of previous research on Pt-10Al-4Cr, the nominal composition ratio of Pt:Al:Cr in the quaternary alloy is designed to be approximately 86:11:3, and the maximum content of nickel is 10 at%. The Pt-11Al-3Cr-(0-10)Ni alloy system has a single-phase structure after solution treatment at 1450°C. The microstructure obtained is similar to that of nickel-based superalloys after ageing treatment at 1000°C, but alloys with nickel content below 6 at% seem to have a lower coherence of precipitation. After ageing treatment, the Pt-11Al-3Cr-6Ni alloy has a maximum γ’ phase volume fraction of about 23%. The cubic precipitate phase is arranged in a straight line with a side length of 200–500 nm, and the degree of mismatch between the precipitate phase and the matrix is about –0.1% (similar to nickel-based superalloys) (Figure 11). Spherical particles are observed in alloys with a nickel content of more than 6 at% and it is believed that the change in the γ’ phase morphology is due to the increase in the nickel concentration and a decrease in the degree of mismatch. Ageing at 1100°C will cause coarsening of γ’ phase and reduce the volume fraction of the γ’ phase. However, as ageing temperature increases, the volume fraction of the γ’ phase of the nickel-containing alloy decreases less than that of the nickel-free alloy.

Fig. 11.

Scanning electron (SE) image of Pt-11Al-3Cr-6Ni after solution heat treatment at 1450°C for 24 h and ageing at 1000°C for 120 h. Reprinted by permission from Springer Nature Customer Service Centre GmbH: Springer Nature, (36), Copyright (2005)

A Pt-(12-15)Al-3Cr-(4-8)Ni alloy with aluminium content near the solubility limit (15 at%) was selected to obtain a platinum-based alloy with high volume fraction of the γ’ phase (36, 37). For alloys with aluminium content less than 13 at% it is possible to obtain a homogenised single-phase structure after heat treatment at 1500°C. Alloys with a higher aluminium content will form a eutectic two-phase dendritic structure even after heat treatment at 1530°C. For alloys with aluminium content less than 13 at%, ageing treatment at 120 h at 1000°C ageing produces a uniformly distributed precipitation of Pt3Al. Figure 12 shows cubic Pt3Al particles with an average side length of 520 nm in the Pt-14Al-3Cr-6Ni alloy. The absolute mismatch between γ and γ’ phases decreases with the increase of nickel content. For alloys with nickel content higher than 5 at%, the slightly negative mismatch value (less than –0.5%) at room temperature and the cubic or spherical particle morphology of the γ’ phase indicate that the γ and γ’ phases are in a coherent state. Increasing the main γ’ phase forming element (aluminium) to 13 at% can increase the volume fraction of γ’ up to 30% (37).

Fig. 12.

SE image of Pt-14Al-3Cr-6Ni after solution heat treatment at 1500°C for 12 h and ageing at 1000°C for 120 h. Reprinted by permission from Springer Nature Customer Service Centre GmbH: Springer Nature, (37), Copyright (2005)

The alloy system Pt-Al-Cr-Ni was studied with an aluminium content limited to 12.5 at% and a chromium content up to 6 at%. The effect of chromium content and heat treatment on the volume fraction of the γ’ reinforcing phase is reported (34). Analysis of the microstructure showed that the dendritic cast structure of Pt‐12.5Al‐3Cr‐6Ni, Pt‐12Al-6Ni and Pt-12Al‐6Cr‐6Ni can be homogenised by heat treatment at 1500– 1510°C. After homogenisation treatment (12 h at 1500°C), the Pt3Al precipitation in the Pt-12.5Al-3Cr-6Ni alloy was almost completely suppressed after water quenching (Figure 13(a)). Air cooling causes the Pt3Al particles (average size 200 nm) to be uniformly distributed with a volume fraction of about 30% (Figure 13(b)). Furnace cooling from 1500°C resulted in cubic and coarse particles distributed in the alloy with a volume fraction of 34% (Figure 13(c)). Increasing the chromium content to 6 at% resulted in Pt3Al with an average particle size of 500 nm and a volume fraction that reached 50% after solution heat treatment for 6 h at 1500–1510°C and air cooling.

Fig. 13.

Secondary electron SEM micrographs of Pt-12.5Al-3Cr-6Ni after homogenisation for 12 h at 1500°C and different cooling regimes: (a) water quenched; (b) air cooled; (c) furnace cooled. Reprinted by permission from Springer Nature Customer Service Centre GmbH: Springer Nature, (34), Copyright (2004)

The γ’ phase dissolves at very high temperatures, therefore the designed platinum-based superalloys have a maximum operating temperature of 1300°C. Figure 14 shows the stress-rupture strength curves of Pt-10wt%Rh, Pt-10wt%Rh DPH and Pt-12Al-6Cr-5Ni alloys at 1300°C (65). The latter has the highest fracture strength. The minimum creep rate of Pt-12Al‐6Cr‐5Ni alloy is almost three orders of magnitude lower than that of the initial Pt‐10Al‐4Cr alloy at 1300°C. Under the stress of 30 MPa, the creep performance of Pt-12Al-6Cr‐5Ni alloy is better than the PM2000 benchmark alloy (Figure 15). Adding a small amount of boron (0.3 at% or 0.7 at%) can significantly improve the creep strength and ductility of the Pt-12Al‐6Cr‐5Ni alloy (66). To investigate the oxidation resistance of Pt-12Al-6Cr-5Ni, Wenderoth et al. (67) studied its isothermal oxidation behaviour of after 400 h exposure to a temperature range of 1100–1300°C in air. A layer of Al2O3 was observed on the surface below which a free γ’-free layer was detected. It was also observed that the size of the γ’-free layer continuously increases with time and temperature. Furthermore, the local concentration of aluminium in the γ’-free layer increases with elevated ageing temperatures. After 400 h exposure at 1300°C a thick polycrystalline Al2O3 scale with large oxide grains developed on the surface. This is in good agreement with the typical behaviour of alloy systems forming protective Al2O3 scales.

Fig. 14.

Stress-rupture strength curve of different platinum-based alloys at 1300°C. Reprinted by permission from Springer Nature Customer Service Centre GmbH: Springer Nature, (65), Copyright (2004)

Fig. 15.

Minimal tension creep rates at 1300°C of Pt-12Al-6Cr-5Ni, Pt-10Al-4Cr and PM2000 alloys. Reprinted by permission from Springer Nature Customer Service Centre GmbH: Springer Nature, (65), Copyright (2004)

4.2 Pt-Al-Cr-Ru Quaternary Alloy

The strong solid solution strengthening element ruthenium can be added to the Pt-Al-Cr ternary alloy to form a Pt-Al-Cr-Ru quaternary alloy. This increases the volume fraction of the γ’ phase and further improves the mechanical performance and oxidation resistance of the alloy (68). Figure 16 shows the typical TEM two-phase microstructure of Pt-12Al-4Cr-2Ru (at%) alloy with the Pt3Al precipitation phase and the (platinum) matrix (69). The volume fraction of the Pt3Al precipitate has increased significantly, with the highest volume fraction of the precipitation phase in platinum-based alloys found so far. The morphology of the precipitation phase is mainly cubic with a side length of about 200 nm and a small amount being irregular. XRD analysis confirmed that the Pt3Al precipitation phase is cubic.

Fig. 16.

TEM micrograph for Pt-12Al-4Cr-2Ru, showing Pt3Al precipitates, an example is marked A (69) CC-BY

The quaternary alloy was prepared by arc melting and then aged in air for 100 h at 1250°C in a muffle furnace. It was then water quenched. Test samples for investigation of mechanical properties were obtained from the bulk material by machining. Table V lists the mechanical properties of several Pt-Al alloys, iron-based (PM2000) and nickel-based (CMSX-4) superalloys (68). From the performance comparison of the three platinum-based alloys, it can be found that the quaternary alloy has the highest hardness value, but its room temperature tensile strength and elongation are the lowest. This is inconsistent with the expected result of ruthenium as a strong solid solution strengthening element. The tensile fracture morphology of the alloy was analysed. It was found that only the Pt‐11Al‐3Cr‐2Ru alloy showed intergranular fracture, while all the ternary alloys had cleavage fracture with some localised signs of dimpling. It is likely that the lower ultimate tensile strength value of the quaternary alloy is related to the intergranular failure mode, which also correlates to the lower plasticity expressed by the lower elongation value. Compared with iron-based and nickel-based superalloys, the tensile strength of platinum-based superalloys is within the ultimate tensile strength range of high-temperature alloys at room temperature.

Table V

Mechanical Properties for Platinum-Based Alloys and Selected High-Temperature Alloys (68)

Alloy Hardness, HV Ultimate tensile strength at room temperature, MPa Elongation, %
Pt-10Al-4Cr 317 ± 13 836 ~4
Pt-10Al-4Ru 278 ± 14 814 ~9
Pt-11Al-3Cr-2Ru 361 ± 10 722 ~1
PM2000 700 14
CMSX-4 870

Odusote (70, 71) studied the isothermal oxidation behaviour of Pt-11Al-3Cr-2Ru (at%) at 1350°C in air. It is found that the oxide layer of the alloy is mainly composed of α-Al2O3 and the thickness of the oxide layer increases with the oxidation time according to a parabolic law (similar to Pt-Al-Cr or Pt-Al-Ru alloys). The growth mechanism of the oxide layer is mainly the diffusion of oxygen atoms into the inner layer along the oxide grain boundary, accompanied by the outward diffusion of a small amount of aluminium atoms. The oxide protective layer is dense, has good adhesion to the substrate and no local discontinuities or detrimental internal oxidation phenomena have been found (Figure 17) (71). These characteristics of the oxide layer indicate that the Pt-11Al-3Cr-2Ru alloy has good oxidation resistance and has the potential for high-temperature applications.

Fig. 17.

SEM-SE cross-section image of Pt-11Al-3Cr-2Ru (at%) specimen after 100 h oxidation in air at 1350°C. Reprinted by permission from Springer Nature Customer Service Centre GmbH: Springer Nature, (71), Copyright (2012)

Alloys belonging to the Pt-Al-Cr-Ru series are currently being developed since they show the best overall performance. However platinum-based alloys have the disadvantages of high price and high density (10). The addition of a cheap and low-density alloy element to replace part of the platinum while maintaining the high-temperature properties and desired microstructure is a subject of current research. Wenderoth et al. (72) added the refractory metal niobium to the platinum-based alloy to improve the high-temperature strength through precipitation hardening. Vanadium belongs to the same group of the Periodic Table as niobium but has a smaller atomic radius and higher solid solubility in the platinum matrix. As well as its effect on precipitation hardening, vanadium may also have a solid solution hardening effect. Odera et al. (73) prepared and analysed four Pt-Al-Cr-Ru-V five-component alloys and two Pt-Al-Cr-Ru-V-Nb six-component alloys by adding vanadium and niobium on the basis of a Pt-12Al-4Cr-2Ru quaternary alloy with excellent properties. Vanadium was added in the range 5.2–19.0 at%, niobium was added in smaller amounts. The content of platinum was reduced to the range 63.2–78.7 at%. The study found that the expected Pt3Al precipitate appeared and that a two-phase structure (matrix with precipitates) was obtained in the four cast alloys, while two alloys had a single-phase structure with vanadium preferentially distributed in the platinum matrix. The hardness of multi-element platinum-based alloys is higher than the quaternary alloys. The optimal vanadium content is about 15 at% to obtain the desired microstructure. A higher vanadium content leads to the formation of a brittle Pt-V mesophase. Niobium loss was too high for its alloying effect to be determined. According to the preliminary experimental results on microstructure and hardness of the multi-element systems based on platinum alloys, these materials have promise for high-temperature applications although further optimisation of their composition is necessary.

5. Research on Hot Corrosion Performance of Platinum-Based Superalloys

For materials used in high-temperature environments, the focus is usually on high-temperature strength (74). However, the creep, oxidation and corrosion resistance of alloys are also important (75). Increasing operating temperature will lead to continuous corrosion. Therefore, it is necessary to evaluate the corrosion resistance of high-temperature materials during the selection process (1, 76). For nickel-based superalloys, the high-temperature strength of the alloy is improved by increasing the content of aluminium and reducing the content of chromium, but the alloy is more sensitive to high-temperature corrosion, and it is necessary to develop and introduce protective coatings (77).

Since platinum-based alloys show excellent performance in various high-temperature applications such as glass manufacturing and corrosive substance processing, platinum-based alloys can be used to solve problems encountered in the aerospace industry (29, 78). Platinum-based superalloys are relatively new high-temperature materials, and there are very few literature reports on their corrosion properties. Fuel or intake air usually contains impurity elements such as sodium, sulfur and vanadium, which may form molten salt corrosion products such as Na2SO4, NaCl and V2O5, which in turn may lead to high temperature hot corrosion (HTHC) of materials (22, 79). Hot corrosion caused by molten salt or corrosive gas accelerates the oxidation degradation of high temperature materials, adversely affects the mechanical properties of the alloy and shortens the service life of high temperature components. There are two types of hot corrosion in nickel-based superalloys: Type I and Type II. Type I hot corrosion is also known as HTHC and usually occurs in the temperature range 850–950°C. Type II hot corrosion is also known as low temperature hot corrosion (LTHC) and generally occurs in the temperature range 650–800°C (79). Type I or HTHC is the main corrosive process in aircraft gas turbine engines.

Maledi et al. (80) studied the accelerated corrosion behaviour of five platinum-based superalloys in analytically pure anhydrous Na2SO4 at 950°C and compared them with NBSA with 1.25 μm thick Pt2Al coating protection or no coating. The experimental results are listed in Table VI (76). The corrosion kinetics of nickel-based and platinum-based superalloys for the first 50 h are shown in Figure 18 and Figure 19, respectively. Due to the protective oxide layer formed on the surface of the platinum-based alloy, the weight gain associated with corrosion is very small. On the other hand, uncoated NBSA form oxides in the initial stage of corrosion, leading to increased mass. After further exposure to the corrosive environment, non-protective oxides form and cause catastrophic corrosion damage. Although the coated NBSA has better corrosion resistance than the uncoated alloy, it still degrades prematurely compared to the platinum-based superalloy. Experiments show that in the molten Na2SO4 environment, platinum-based alloys show superior corrosion resistance compared to both coated and uncoated NBSA. Figure 20 shows the protective layer with strong adhesion formed on the surface of the Pt-10Al-4Cr alloy. The surface protective layer morphology of the Pt-10Al‐4Ru and Pt-11Al-3Cr-2Ru alloys is similar (80). Although the morphology of the surface oxide layer of platinum-based superalloys containing chromium or ruthenium appears porous, their resistance to hot corrosion is higher than that of the platinum-based superalloys containing cobalt. NBSA was subjected to penetration of sulfur under the surface oxide layer, leading to formation of chromium and nickel sulfides. This alloy showed the worst resistance to sulfidation and hot corrosion. Resistance to sulfidation was also the subject of studies in the experimental work of Potgieter et al. (81). The results of the study are shown in Table VI. The tests were performed in a 0.2% SO2-N2 mixed atmosphere and research conclusions were similar to those obtained by Maledi et al. (80).

Table VI

Platinum-Based Superalloys and Nickel-Based Superalloys Together with their Corrosion Kinetics After Treatment in Na2SO4 at 950°C for 540 h (76)

Alloy name Nominal composition, at % Cumulative weight gain during corrosion, mg cm–2
RS-1 Pt-10Al-4Cr 0.00004
RS-2 Pt-10Al-4Ru 0.00008
RS-3 Pt-11Al-3Cr-2Ru 0.0001
P420 Pt-15Al-6Co 0.0001
P421 Pt-15Al-12Co 0.004
CMSX-4 (uncoated) Ni-6.5Cr-11Co-0.3Mo-1.7W-1.8Ta-11.3Al-0.9Ti 0.470
CMSX-4 (coated) Ni-6.5Cr-11Co-0.3Mo-1.7W-1.8Ta-11.3Al-0.9Ti 0.038
Fig. 18.

Corrosion kinetics of the coated CMSX-4 NBSAs during exposure to Na2SO4 at 950°C for the first 50 h. Reprinted from (80) with permission

Fig. 19.

Corrosion kinetics of five platinum-based alloys of various compositions for the first 50 h. Reprinted from (80) with permission

Fig. 20.

SEM secondary electron image showing the thin protective scale on the surface of the platinum-based superalloy RS-1(Pt-10Al-4Cr). Reprinted from (80) with permission

6. Conclusions

More than 20 years of research and development have yielded interesting results for platinum-based superalloys. The results obtained so far are only the tip of the iceberg of a very interesting and topical subject characterised by excellent prospects for future use. Optimised design of alloy composition, microstructure characteristics, mechanical properties and oxidation corrosion behaviour have been achieved. Among the systems studied so far, the Pt-Al-Cr-Ru alloy system has been selected and optimised for excellent performance. The ultimate goal is to develop platinum-based superalloys for application in industrial fields such as gas turbine engines. However, this is a competitive market that is difficult to penetrate with new materials. One possibility would be to exploit the possibility for platinum-based alloys to be used at temperatures 200°C higher than for NBSA (82). During the transition period, platinum-based superalloys could be used in other fields, such as castings, powder metallurgy products and coatings to accelerate their final use in gas turbine engines.

At present, the composition and structure design of platinum-based superalloys mainly follows the successful experience of NBSA development. The melting point of platinum is 316°C higher than that of nickel, but the difference in melting point between platinum-based superalloys and nickel-based superalloys is less than 150°C. The main reason for this discrepancy is the addition of low melting point alloying elements (such as aluminium) which reduces the melting point of the resulting alloy to about 1500°C. The advantage of platinum’s high melting point has thus not yet been fully exploited. The development cycle of a new generation of superalloys is very long and the material cost is relatively high for precious metals. The optimisation process for the design of new platinum-based superalloys could be accelerated with the help of material genome research concepts to further increase the alloy’s temperature tolerance and reduce research and development costs.

Platinum-based superalloys can still be considered a brand new alloy system when compared to, say, NBSA or stainless steels. The accumulation of fundamental data such as phase precipitation mechanisms and alloy properties is far from being complete. There has been little research on the influence of manufacturing processes (for example, precision casting, directional solidification and single crystal preparation) on the formability and mechanical properties of the alloys. There has similarly been insufficient verification and assessment of performance under actual use in typical environmental conditions. To accelerate bringing these materials to market, the level of research and development on platinum-based superalloys needs to be improved urgently.


  1. 1.
    S. Gialanella and A. Malandruccolo, “Aerospace Alloys”, Topics in Mining, Metallurgy and Materials Engineering Series, Springer Nature, Cham, Switzerland, 2020, 570 pp
  2. 2.
    R. C. Reed, “The Superalloys: Fundamentals and Applications”, Cambridge University Press, Cambridge, UK, 2006, 372 pp
  3. 3.
    J. Zhang, L. Wang, D. Wang, G. Xie, Y. Lu, J. Shen and L. Lou, Acta Metall. Sin., 2019, 55, (9), 1077 LINK
  4. 4.
    J. Wu, Y. Liu, C. Li, Y. Wu, X. Xia and H. Li, Acta Metall. Sin., 2020, 56, (1), 21 LINK
  5. 5.
    J. S. Van Sluytman, C. J. Moceri and T. M. Pollock, Mater. Sci. Eng.: A, 2015, 639, 747 LINK
  6. 6.
    J. Rame, S. Utada, L. M. Bortoluci Ormastroni, L. Mataveli-Suave, E. Menou, L. Després, P. Kontis, and J. Cormier, ‘Platinum-Containing New Generation Nickel-Based Superalloy for Single Crystalline Applications’, in “Superalloys 2020”, The Minerals, Metals & Materials Series, eds. S. Tin, M. Hardy, J. Clews, J. Cormier, Q. Feng, J. Marcin, C. O’Brien and A. Suzuki, Springer Nature, Cham, Switzerland, 2020, pp. 71– 81 LINK
  7. 7.
    “Superalloys II: High-Temperature Materials for Aerospace and Industrial Power”, eds. C. T. Sims, N. S. Stollof and W. C. Hagel, 2nd Edn., Wiley InterScience, Hoboken, USA, 1987, 640 pp
  8. 8.
    T. Yokokawa, M. Osawa, K. Nishida, T. Kobayashi, Y. Koizumi and H. Harada, Scr. Mater., 2003, 49, (10), 1041 LINK
  9. 9.
    Y. Yamabe-Mitarai, Y. Gu, C. Huang, R. Völkl and H. Harada, J. Miner. Metals Mater. Soc., 2004, 56, (9), 34 LINK
  10. 10.
    L. A. Cornish, R. Süss, A. Douglas, L. H. Chown and L. Glaner, Platinum Metals Rev., 2009, 53, (1), 2 LINK
  11. 11.
    J.-C. Zhao and J. H. Westbrook, MRS Bull., 2003, 28, (9), 622 LINK
  12. 12.
    Y.-L. Jeng, E. J. Lavernia, R. M. Hayes and J. Wolfenstine, Mater. Sci. Eng.: A, 1995, 192–193, (1), 240 LINK
  13. 13.
    K. Sadananda and A. K. Vasudevan, Mater. Sci. Eng.: A, 1995, 192–193, (1), 490 LINK
  14. 14.
    T. E. Tietz and J. W. Wilson, “Behavior and Properties of Refractory Metals”, Edward Arnold Ltd, London, UK, 1965, 419 pp
  15. 15.
    Y. Yamabe, Y. Koizumi, H. Murakami, Y. Ro, T. Maruko and H. Harada, Scr. Mater., 1996, 35, (2), 211 LINK
  16. 16.
    Y. Yamabe-Mitarai, Y. Ro, T. Maruko, T. Yokokawa, and H. Harada, ‘Platinum Group Metals-Base Refractory Superalloys for Ultra-High Temperature Use’, Structural Intermetallics 1997: 2nd International Symposium on Structural Intermetallics, 21st–25th September, 1997, Champion, USA, eds. M. V. Nathal, R. Darolia, C. T. Liu, P. L. Martin, D. B. Miracle, R. Wagner and M. Yamaguchi, The Minerals, Metals & Materials Society, Warrendale, USA, 1997, pp. 805–814
  17. 17.
    Y. Yamabe-Mitarai, Y. Koizumi, H. Murakami, Y. Ro, T. Maruko and H. Harada, Scr. Mater., 1997, 36, (4), 393 LINK
  18. 18.
    Y. Yamabe-Mitarai, Y. Ro, H. Harada and T. Maruko, Metall. Mater. Trans. A, 1998, 29, (2), 537 LINK
  19. 19.
    Y. Yamabe-Mitarai, Y. Ro, T. Maruko and H. Harada, Scr. Mater., 1998, 40, (1), 109 LINK
  20. 20.
    Y. Yamabe-Mitarai, Y. Gu, Y. Ro, S. Nakazawa, T. Maruko and H. Harada, Scr. Mater., 1999, 41, (3), 305 LINK
  21. 21.
    X. Yu, Y. Yamabe-Mitarai, Y. Ro and H. Harada, Intermetallics, 2000, 8, (5–6), 619 LINK
  22. 22.
    I. M. Wolff and P. J. Hill, Platinum Met. Rev., 2000, 44, (4), 158 LINK
  23. 23.
    ‘The Nickel Industry: Occurrence, Recovery, and Consumption: Elemental Nickel’, in “ASM Specialty Handbook: Nickel, Cobalt and Their Alloys”, eds. J. R. Davis, ASM International, Materials Park, USA, 2000, p3
  24. 24.
  25. 25.
    D. F. Lupton, J. Merker, B. Fischer and R. Völkl, Glastech. Ber. Glass Sci. Technol., 2000, 73, (2), 284
  26. 26.
    C. Y. Hu and S. J. Liu, “New Materials of Precious Metals”, Central South University Press, Changsha, China, 2015, 474 pp
  27. 27.
    Y. T. Ning, Precious Met., 2009, 30, (2), 51
  28. 28.
    “Platinum”, eds. Y. T. Ning, Z. F. Yang and F. Wen, Metallurgical Industry Press, Beijing, China, 2010
  29. 29.
    B. Fischer, A. Behrends, D. Freund, D. F. Lupton and J. Merker, Platinum Metals Rev., 1999, 43, (1), 18 LINK
  30. 30.
    R. Völkl, D. Freund, B. Fischer and D. Gohlke, Key Eng. Mater., 1999, 171–174, 77 LINK
  31. 31.
    A. Niwa, Y. Akita, K. Enomoto, R. Aoyama, H. Akebono and A. Sugeta, Int. J. Fatigue, 2020, 132, 105385 LINK
  32. 32.
    K. Teichmann, C. H. Liebscher, R. Völkl, S. Vorberg and U. Glatzel, Platinum Metals Rev., 2011, 55, (4), 217 LINK
  33. 33.
    J. Merker, B. Fischer, R. Völkl and D. F. Lupton, Mater. Sci. Forum, 2003, 426–432, 1979 LINK
  34. 34.
    S. Vorberg, M. Wenderoth, B. Fischer, U. Glatzel and R. Völkl, J. Miner. Metals Mater. Soc., 2004, 56, (9), 40 LINK
  35. 35.
    S. Vorberg, B. Fischer, M. Wenderoth, U. Glatzel and R. Völkl, J. Miner. Metals Mater. Soc., 2005, 57, (3), 49 LINK
  36. 36.
    M. Hüller, M. Wenderoth, U. Glatzel, R. Völkl, S. Vorberg and B. Fischer, Metall. Mater. Trans. A, 2005, 36, (3), 681 LINK
  37. 37.
    M. Wenderoth, U. Glatzel, R. Völkl, L. A. Cornish, R. Süss, S. Vorberg and B. Fischer, Metall. Mater. Trans. A, 2005, 36, (3), 567 LINK
  38. 38.
    R. Völkl, Y. Yamabe-Mitarai, C. Huang and H. Harada, Metall. Mater. Trans. A, 2005, 36, (11), 2881 LINK
  39. 39.
    M. Wenderoth, R. Völkl, T. Yokokawa, Y. Yamabe-Mitarai and H. Harada, Scr. Mater., 2006, 54, (2), 275 LINK
  40. 40.
    M. V. Whalen, Platinum Metals Rev., 1988, 32, (1), 2 LINK
  41. 41.
    P. Audigié, A. Rouaix-Vande Put, A. Malié, C. Thouron and D. Monceau, Corros. Sci., 2019, 150, 1 LINK
  42. 42.
    L. A. Cornish, J. Hohls, P. J. Hill, S. Prins, R. Süss and D. N. Compton, J. Min. Metall. Sect. B: Metall., 2002, 38, (3–4), 197 LINK
  43. 43.
    J. K. Odusote, L. A. Cornish and J. M. Papo, J. Mater. Eng. Perform., 2013, 22, (11), 3466 LINK
  44. 44.
    “Binary Alloy Phase Diagrams”, eds. T. B. Massalski, H. Okamoto, P. R. Subramanian and L. Kacprzak, 2nd Edn., ASM International, Materials Park, USA, 1990, in 3 volumes
  45. 45.
    Y. T. Ning, Precious Met., 2010, 31, (1), 57
  46. 46.
    P. J. Hill, T. Biggs, P. Ellis, J. Hohls, S. Taylor and I. M. Wolff, Mater. Sci. Eng.: A, 2001, 301, (2), 167 LINK
  47. 47.
    G. B. Fairbank, C. J. Humphreys, A. Kelly and C. N. Jones, Intermetallics, 2000, 8, (9–11), 1091 LINK
  48. 48.
    P. J. Hill, L. A. Cornish and G. B. Fairbank, J. Miner. Metals Mater. Soc., 2001, 53, (10), 19 LINK
  49. 49.
    M. Adjal, S. Méçabih, B. Abbar and B. Bouhafs, Comput. Condens. Matter, 2018, 16, e00328 LINK
  50. 50.
    Y. Pan, D. Pu and Y. Jia, Vacuum, 2020, 172, 109067 LINK
  51. 51.
    Z. Li, K. Xiong, Y. Sun, C. Jin, S. Zhang, J. He and Y. Mao, Comput. Condens. Matter, 2020, 23, e00462 LINK
  52. 52.
    Y. Liu, H. Huang, Y. Pan, G. Zhao and Z. Liang, J. Alloys Compd., 2014, 597, 200 LINK
  53. 53.
    L. A. Cornish, B. Fischer and R. Völkl, MRS Bull., 2003, 28, (9), 632 LINK
  54. 54.
    A. Douglas, J. H. Neethling, R. Santamarta, D. Schryvers and L. A. Cornish, J. Alloys Compd., 2007, 432, (1–2), 96 LINK
  55. 55.
    P. J. Hill, Y. Yamabe-Mitarai and I. M. Wolff, Scr. Mater., 2001, 44, (1), 43 LINK
  56. 56.
    T. Biggs, L. A. Cornish, M. J. Witcomb and M. B. Cortie, J. Phys. IV France, 2001, 11, (PR8), 493 LINK
  57. 57.
    P. J. Hill, L. A. Cornish, P. Ellis and M. J. Witcomb, J. Alloys Compd., 2001, 322, (1–2), 166 LINK
  58. 58.
    P. J. Hill, Y. Yamabe-Mitarai, H. Murakami, L. A. Cornish, M. J. Witcomb, I. M. Wolff and H. Harada, ‘The Precipitate Morphology and Lattice Mismatch of Ternary (Pt)/Pt3Al Alloys’, in “Structural Intermetallics, 2001: ISSI: Proceedings of the Third International Symposium on Structural Intermetallics”, TMS, Pittsburgh, USA, 2001, pp. 527–533
  59. 59.
    R. Süss, D. Freund, R. Völkl, B. Fischer, P. J. Hill, P. Ellis and I. M. Wolff, Mater. Sci. Eng.: A, 2002, 338, (1–2), 133 LINK
  60. 60.
    L. A. Cornish, M. B. Shongwe, B. Odera, J. K. Odusote, M. J. Witcomb, L. H. Chown, G. O. Rading and M. J. Papo, ‘Update on the Development of Platinum-Based Alloys for Potential High-Temperature Applications’, 5th Platinum Conference, 19th–21st September, 2012, Sun City, South Africa, The Southern African Institute of Mining and Metallurgy, Marshalltown, South Africa, 2012, pp. 905–923
  61. 61.
    L. A. Cornish, R. Süss, A. Watson and S. N. Prins, Platinum Metals Rev., 2007, 51, (3), 104 LINK
  62. 62.
    B. A. Watson, R. Süss and L. A. Cornish, Platinum Metals Rev., 2007, 51, (4), 189 LINK
  63. 63.
    J. Preußner, S. N. Prins, M. Wenderoth, R. Völkl and U. Glatzel, Platinum Metals Rev., 2008, 52, (1), 48 LINK
  64. 64.
    J.-C. Zhao, M. R. Jackson, L. A. Peluso and L. N. Brewer, MRS Bull., 2002, 27, (4), 324 LINK
  65. 65.
    R. Völkl and B. Fischer, Exp. Mech., 2004, 44, (2), 121 LINK
  66. 66.
    R. Völkl, M. Wenderoth, J. Preussner, S. Vorberg, B. Fischer, Y. Yamabe-Mitarai, H. Harada and U. Glatzel, Mater. Sci. Eng.: A, 2009, 510–511, 328 LINK
  67. 67.
    M. Wenderoth, S. Vorberg, B. Fischer, R. Völkl, and U. Glatzel, Int. J. Mater. Res., 2007, 98, (6), 463 LINK
  68. 68.
    L. A. Cornish, R. Süss, L. H. Chown and L. Glaner, Platinum Metals Rev., 2009, 53, (3), 155 LINK
  69. 69.
    M. B. Shongwe, M. J. Witcomb, L. A. Cornish and M. J. Papo, J. S. Afr. Inst. Min. Metall., 2012, 7A, 551
  70. 70.
    J. K. Odusote, L. A. Cornish, L. H. Chown and R. M. Erasmus, Oxid. Met., 2012, 78, (1–2), 123 LINK
  71. 71.
    J. K. Odusote, L. A. Cornish and J. M. Papo, Metall. Microstruct. Anal., 2012, 1, (3–4), 142 LINK
  72. 72.
    M. Wenderoth, S. Vorberg, B. Fischer, Y. Yamabe-Mitarai, H. Harada, U. Glatzel and R. Völkl, Mater. Sci. Eng.: A, 2008, 483–484, 509 LINK
  73. 73.
    B. O. Odera, M. J. Papo, R. Couperthwaite, G. O. Rading, D. Billing and L. A. Cornish, J. S. Afr. Inst. Min. Metall., 2015, 115, (3), 241
  74. 74.
    B. A. Pint, J. R. DiStefano and I. G. Wright, Mater. Sci. Eng.: A, 2006, 415, (1–2), 255 LINK
  75. 75.
    T. J. Carter, Eng. Fail. Anal., 2005, 12, (2), 237 LINK
  76. 76.
    J. H. Potgieter, N. B. Maledi, M. Sephton and L. A. Cornish, Platinum Metals Rev., 2010, 54, (2), 112 LINK
  77. 77.
    I. Gurrappa, Surf. Coatings Technol., 2001, 139, (2–3), 272 LINK
  78. 78.
    R. Völkl, D. Freund, A. Behrends, B. Fischer, J. Merker, and D. Lupton, ‘Platinum Base Alloys for High Temperature Space Applications’, in “Materials for Transportation Technology”, ed. P. J. Winkler, Vol. 1, Wiley-VCH, Weinheim, Germany, 2000, pp. 257–260 LINK
  79. 79.
    T. S. Sidhu, S. Prakash and R. D. Agrawal, Current Sci., 2006, 90, (1), 41
  80. 80.
    N. B. Maledi, J. H. Potgieter, M. Sephton, L. A. Cornish, L. Chown and R. S. Süss, ‘Hot Corrosion Behaviour of Pt-Alloys for Application in the Next Generation of Gas Turbines’, Second International Platinum Conference: ‘Platinum Surges Ahead’, 8th–12th October, 2006, Sun City, South Africa, Symposium Series S45, Southern African Institute of Mining and Metallurgy, Johannesburg, South Africa, 2006, pp. 81–90 LINK
  81. 81.
    J. H. Potgieter and N. B. Maledi, Open Mater. Sci. J., 2014, 8, 18 LINK
  82. 82.
    L. A. Cornish, R. Süss, R. Völkl, M. Wenderoth, S. Vorberg, B. Fischer, U. Glatzel, A. Douglas, L. H. Chown, T. Murakumo, J. Preussner, D. Lupton, L. Glaner, N. B. Maledi, J. H. Potgieter, M. Sephton and G. Williams, J. S. Afr. Inst. Min. Metall., 2007, 107, 697


This work is supported by Basic Research Key Program of Yunnan, China (No.2019FA048) and The Major Science and Technology Program of Yunnan, China (Nos. 2019ZE001, 202002AB080001-1).

The Authors

Professor Changyi Hu is the Director of the Research and Development Center of Kunming Institute of Precious Metals, China, as well as the Vice Director of State Key Laboratory of Advanced Technologies for Comprehensive Utilization of Platinum Metals, China. His research interests include alloys, films and coatings of platinum group metals, as well as work pieces of refractory metals for high-temperature applications.

Yan Wei is a Professor at Kunming Institute of Precious Metals, as well as the Head of High Temperature Materials Division at State Key Laboratory of Advanced Technologies for Comprehensive Utilization of Platinum Metals, China. Her main fields of research include alloys of platinum group metals, refractory metals and jewellery alloys.

Professor Hongzhong Cai is a Senior Researcher in the High Temperature Materials Division at Kunming Institute of Precious Metals, China. He is working on films or coatings of platinum group metals and high temperature ceramics, as well as structural materials of refractory metals prepared by chemical vapour deposition.

Li Chen is the Deputy Director of Kunming Institute of Precious Metals, China. He is a Principal Engineer of genome project of precious metal alloy materials in the Alloy Material Group, where he is working on first-principles calculations, database building of precious metal alloys and new high-temperature materials.

Xian Wang is a Researcher in the Research and Development Center of Kunming Institute of Precious Metals, China. He carries out modelling and simulation in computation materials science, and is currently working on finite element simulation and analysis, as well as database building of precious metal materials.

Xuxiang Zhang is a Chief Engineer in the High Temperature Materials Division at Kunming Institute of Precious Metals, China. His research experience includes platinum alloys and refractory metals for high-temperature use. He is now responsible for product structure design and quality analysis.

Guixue Zhang is a Principal Technician in the High Temperature Materials Division at Kunming Institute of Precious Metals, where he is in charge of the High Temperature Material Laboratory. He is mainly responsible for testing mechanical properties and analysing microstructures of high-temperature materials.

Xingqiang Wang is a Technician in the High Temperature Materials Division at Kunming Institute of Precious Metals, China. He is working on platinum alloys, refractory metals and titanium alloys, including the preparation of samples, parts fabrication and production.

Related articles

Effect of Ruthenium Targets on the Growth and Electrical Properties of Sputtering Ruthenium Films

“PGM Market Report May 2023”

Find an article