Journal Archive

Platinum Metals Rev., 2006, 50, (4), 158
doi: 10.1595/147106706X154198

High-Temperature Mechanical Properties of the Platinum Group Metals


  • R. Weiland
  • D. F. Lupton*
  • Engineered Materials Division,
  • W. C. Heraeus GmbH, Hanau, Germany
  • Email:
  • B. Fischer
  • J. Merker
  • C. Scheckenbach
  • Department SciTec,
  • Precision-Optics-Materials-Environment, University of Applied Sciences Jena, Germany
  • J. Witte
  • Melting Technology,
  • SCHOTT Glas, Mainz, Germany

Article Synopsis

In order to provide reliable data on the high-temperature deformation behaviour of iridium, the high-temperature material properties such as stress-rupture strength, high-temperature tensile strength and creep behaviour are determined for pure iridium in the temperature range 16502300°C. Analyses of the stress-rupture curves and the creep behaviour of pure iridium samples at 1650°C, 1800°C and 2000°C imply that the fracture behaviour is controlled by two different fracture mechanisms depending on test conditions, in particular applied load and test temperature. The existence of the different fracture modes is confirmed by SEM examination of the fracture surface of samples ruptured at high temperatures. Anomalies in the creep curves and the results of high-temperature tensile tests indicate that dynamic recrystallisation plays an important role in the high-temperature deformation behaviour of pure iridium.

Due to their excellent chemical stability, oxidation resistance, and resistance to the action of many molten oxides, the platinum group metals (pgms): iridium, platinum and rhodium are widely used for high-temperature applications involving simultaneous chemical attack and mechanical loading (1). Although iridium is more sensitive to oxidation than platinum or rhodium, it is the most chemically resistant of all metals. Its resistance to attack by stable oxide melts is maintained up to temperatures above 2000°C.

The melting point of iridium (2454°C) (2) and its high strength even at temperatures above 2000°C make it a particularly suitable material for applications under extreme thermal and mechanical conditions which preclude the use of platinum alloys or rhodium. Important applications of iridium and iridium alloys are as crucibles for pulling single crystals (e.g. yttrium-aluminium garnet (YAG)) and components for manufacturing and processing high-melting special glasses.

A knowledge of the high-temperature properties of a material, for instance stress-rupture strength and creep behaviour, is crucially important for the design of components used at high temperatures. The current investigation is part of an extensive test programme focused on the determination of the high-temperature mechanical properties of the pgms, such as the stress-rupture strength, creep behaviour (3) and elastic properties (4). In this work new investigations into the high-temperature properties of iridium are presented for the temperature range between 1650°C and 2300°C. The results are discussed in conjunction with data determined from earlier studies (3).

Methodology for Stress-Rupture and High-Temperature Tensile Tests

The stress-rupture strength and the creep behaviour of pure iridium and iridium alloys were determined with a testing facility developed at the University of Applied Sciences Jena. The testing device, for the measurement of high-temperature material properties up to 3000°C, is shown schematically in Figure 1. It consists of a gas-tight specimen chamber which permits investigations either in air or under a protective gas atmosphere. In the case of iridium and iridium alloys, a gas mixture of argon with 5 vol.% H2 was used to protect the material from oxidation and thus avoid a reduction in cross-section of the sample due to evaporation of volatile oxides (5, 6).

Fig. 1

Schematic diagram of equipment for high-temperature creep measurements

The load can be applied in two different ways. For the constant-load stress-rupture experiments the load is applied via a steel pull-rod by means of calibrated weights. For the high-temperature tensile tests the specimen chamber is mounted in a commercial servomotor-driven test machine and the steel pull-rod is connected to the load cell at the crosshead of the test machine. This allows a controlled variation of the applied load. Non-standard specimens (with typical dimensions of 120 mm × 4 mm × 1 mm) were used for all measurements. The samples were laser cut from hot rolled sheet material. The sample orientation was chosen parallel to the rolling direction.

Direct electrical heating achieves high heating and cooling rates for the samples. The ohmic heating method allows easy access to the sample, and generally straightforward operation.

The temperature is measured by a non-contacting technique using a digital pyrometer (INFRATHERM IS10). The infrared pyrometer has a small measurement spot (approximately 0.5 mm in diameter). Due to the ohmic heating the highest temperatures are found in the central part of the sample. This region is therefore scanned continuously by the pyrometer via a tilting mirror. By storing the maximum value of emitted radiation, the maximum temperature at the surface of the sample may be determined. This value is used to adjust the heating current via a thyristor regulator connected to the primary winding of a 100 kVA transformer. The sample, short-circuited across the secondary winding of the transformer, is heated by alternating current at 50 Hz. Over a zone 30 mm in length around the centre of the sample the temperature usually does not vary by more than ± 5°C. Once “necking” occurs in the sample, the temperature outside the necking region decreases, whereas the temperature within the necking region remains constant at the intended value. The design of the equipment thus guarantees uniformity of temperature throughout the duration of the test, despite the sample deformation.

The strain is measured with a non-contacting video extensometer consisting of a 17 mm charge coupled device (CCD) camera with 1280 × 1024 pixel resolution. A special arrangement of telecentric lenses allows only near-parallel rays to pass the aperture, thus minimising perspective distortions caused by variations in the distance to the object. Both the CCD camera and frame grabber are controlled by “SuperCreep” software, developed at the University of Applied Sciences Jena, which uses digital image analysis.

As mentioned above, ohmic heating causes the highest temperature to be limited to the central part of the sample, to which creep deformation is normally also limited. Strain at this part of the sample is determined by “SuperCreep” from continuous measurements of the distance between two markers. Suitable markers for high-temperature tests on sheet materials are made by laser machining samples of the material with four small shoulders (Figure 2). The distance between the two corresponding markers on the same side of the sample is 10 mm. Since the part of the sample between the markers experiences a uniform temperature, the exactness of strain measurements can be guaranteed, without their being influenced by the temperature gradients near the ends of the sample.

Fig. 2

Image of a creep sample with markers for the video extensometer (7)

A detailed description of the testing facility and the algorithm for strain measurement is given in (7) and (8).

Material Preparation

The iridium raw material was melted inductively at 2550°C in air in a zirconia crucible. After elemental analysis the ingot was forged at temperatures between 1400°C and 1600°C. The forged material was then hot rolled at moderate temperature to 1 mm thick sheet. The iridium sheet was finally subjected to a special annealing procedure so as to recrystallise the deformed material without significant grain growth.

Scanning Secondary Ion Mass Spectrometry (Scanning SIMS)

The microanalytical investigations were performed with a Cameca IMS 4f-E6 scanning secondary ion mass spectrometer. Secondary ion mass spectrometry (SIMS) allows the detection of very small amounts of impurity elements in the matrix. Since both the species of detectable secondary ions and their detection limits differ as between the positive and negative secondary ion spectra, different primary ions were chosen for the excitation of the secondary ions. Oxygen primary ions were used for the investigation of the positive secondary ion spectrum emitted by the iridium samples. The emission of the negative spectrum was induced by caesium primary ions. It could thus be ensured that all possible impurity elements contained in the iridium samples were detected.

Metallographically prepared samples were used for the scanning SIMS investigations. So as to be able to investigate impurity levels both inside the grains and at the grain boundaries, areas of the samples containing grain boundaries were chosen.

It should be mentioned that the intensity of the emitted secondary ion spectrum is dependent on the crystallographic orientation of the grains. Thus, if several grains with different orientations are contained in the area under investigation, this will be indicated by differences in the brightness attributable to the respective grains due to differences in ionic emissivity. Thus grain boundaries may be identified in the secondary ion spectrum, even if they do not contain significant amounts of impurities.

Stress-Rupture Strength Results

The stress-rupture strength of pure iridium was determined in the temperature range 1650–2300°C. The results of these investigations are summarised in Figure 3. The present experiments showed an excellent degree of reproducibility; the results are in good agreement with those of corresponding measurements for shorter testing times reported earlier (3). In contrast to the conclusion in (3), additional data on stress-rupture strength obtained recently, together with a detailed analysis of the rupture behaviour (see below under ‘Fracture Behaviour Results’), led to the conclusion that the stress-rupture data can best be approximated by two intersecting lines. The discontinuity in the slope of the stress-rupture curves correlates very well with a change in the fracture behaviour of the samples examined. Under high loads pure iridium shows a ductile fracture mode, whereas under low loads and long times to rupture, iridium tends towards brittle intercrystalline fracture. In fact, the discontinuity does not usually occur as sharply as indicated in Figure 3.

Fig. 3

Stress-rupture strength of pure iridium in the temperature range between 1650°C and 2300°C

Samples taken from near the discontinuity often show mixed fracture modes, partly intercrystalline and partly transcrystalline. Since insufficient stress-rupture data in the range of the transition are available, the stress-rupture curves for the temperatures between 1650°C and 2000°C are approximated for the sake of simplicity in the manner shown in Figure 3. In particular, the stress-rupture curve at 2000°C shows a very pronounced change in the slope at loads between 4 and 5 MPa. This leads to strongly reduced stress-rupture strength values at testing times longer than 300 h. It is not yet clear whether this steep decrease in slope can be attributed solely to the change in fracture mechanism, or whether the effect of weakening of the grain boundary coherence is enhanced by very small amounts of impurities accumulating at the grain boundaries after long testing times at high temperatures. As reported below, secondary ion mass spectrometric investigations showed that the impurity content in the iridium samples examined is very low. Nevertheless, it cannot be excluded that even very small amounts of impurity elements accumulating at the grain boundaries can have a detrimental effect on grain boundary coherence, thus leading to significantly reduced times to rupture.

Measurements at 1650°C and 1800°C have been performed up to approximately 1000 h duration. Extrapolations to testing times longer than 10,000 h may not be considered meaningful. Data on the stress-rupture behaviour of pure iridium at 2200°C and 2300°C are available up to times to rupture of approximately 500 h and 150 h, respectively, and not as yet for longer times to rupture. Within the available range the stress rupture curves do not show a visible discontinuity in slope. Because of the very distinct decrease in slope in the stress-rupture curve at 2000°C under low loads, no extrapolations beyond the measured times to rupture have been performed for the data at 2200°C and 2300°C.

The interpolated and extrapolated data on the stress-rupture strength of pure iridium at different temperatures are given in Table I.

Table I

Stress-Rupture Strength of Pure Iridium at Various Temperatures

Time to rupture, h Stress-rupture strength, MPa

1650°C 1800°C 2000°C 2200°C 2300°C

1 31.8 24.4 14.1 7.1 5.4
10 27.7 18.4 8.9 4.4 3.3
100 15.6 11.0 4.6 2.7 2.0
1000 8.8 7.0 1.5
10,000 5.0 4.4

Creep Behaviour Results

The investigation of the high-temperature deformation behaviour of iridium revealed that, depending on test temperature and load, its creep behaviour can be described by two types of creep curve which differ significantly in shape. Particularly at the lowest test temperature of 1650°C, under moderate load, the creep behaviour is represented by typical creep curves such as those in Figures 4(a) and 4(b). These figures represent the creep behaviour of pure iridium at 1650°C under a constant load of 13 MPa exhibiting the three well known stages of creep – primary, secondary (or steady-state), and tertiary – frequently reported in the literature.

Fig. 4

Creep curves: (a) of pure iridium at 1650°C under a constant load of 13 MPa in Ar/H2 atmosphere, and (b) corresponding creep rate as a function of time. The mean creep rate for the period 16.95152.54 h is 1.3 × 10−7 s−1

Under higher load, and particularly at higher test temperatures, the creep curves of iridium show significant anomalies. In the range of steady-state creep the creep curve contains different plateaus, as shown in Figure 5. These plateaux are separated by an acceleration of the elongation. This acceleration of creep is clearly visible in Figures 5(b) and 5(d). The phenomenon may be caused by dynamic recrystallisation, as was reported in (9) and (10), whose authors obtained creep curves of similar shape when investigating the creep behaviour of lead and copper, respectively. It can be seen in Figures 5(c) and 5(d) that in some cases more than one discontinuity occurs in the secondary creep stage. This indicates that dynamic recrystallisation takes place successively several times during the secondary stage of creep. This is called periodic or cyclic creep (11).

Fig. 5

Elongation and creep rate of pure iridium as a function of time in Ar/H2 atmosphere at 1800°C for loads of: (a) and (b) 13 MPa. The mean creep rate at a constant load of 13 MPa for the period 5.6851.12 h is 1.0 × 10−6 s−1. (c) and (d) 9.5 MPa. The mean creep rate at a constant load of 9.5 MPa for the period 22.5225 h is 2.6 × 10−7 s−1

These accelerations of creep complicate the determination of a constant creep rate in the secondary creep range. The creep rate has therefore been calculated as an average for a time range from 10% to 90% of the period of measurement. Thus additional contributions to the average creep rate from the accelerated creep in the transient creep stages are included in the values. This average creep rate has been used instead of the minimum creep rate for the calculation of the Norton plot (Figure 6). In the Norton plot, the average creep rate for pure iridium determined in this way for each test temperature is shown as a function of the initial applied stress on a double logarithmic scale.

Fig. 6

Average stationary creep rate of pure iridium as a function of the applied initial stress for temperatures between 1650°C and 2300°C (Norton plots)

For the calculation of approximate trend lines, the stress dependence of the average creep rate has been assumed to obey a power law (Equation (i)):

  (d/dt)av = f(S, Tn    (i)

where is the elongation and f (S, T) is a function of the structure of the material and of the temperature. For the isothermal representation in the Norton diagram, f(S, T) has been assumed constant. The effects of structural changes, for instance, due to recrystallisation, have been addressed to some extent by the use of the average creep rate, determined as described above. The Norton exponent, n, contains information about the nature of the prevalent creep mechanism in the sample. For creep mechanisms that are based only on the diffusional transport of material due to vacancy gradients either inside the grains (12, 13) or along the grain boundaries, a nearly linear stress dependence of the creep rate will be obtained, and n will be close to or equal to unity. For creep processes that are determined mainly by diffusion-controlled dislocation climb, n falls between 3 and 5 for many common pure metals. In rare cases, n values up to 11 have been found.

The distinction between the different creep mechanisms under low and high loads mentioned in the preceding section must also be taken into account in the Norton plots. The data for each test temperature between 1650°C and 2000°C are therefore approximated by two intersecting straight lines. As already explained for the stress-strain diagram, the information available for 2200°C and 2300°C does not allow this distinction to be made for these temperatures.

For test temperatures between 1650°C and 2000°C, the values obtained for the Norton exponent under low loads n1 fall in the range between 4 and 5.5, and are thus in good agreement with the above-mentioned values of the Norton exponent for diffusion-controlled dislocation climb in many other pure metals. For conditions under which iridium exhibited a ductile fracture mode, i.e. under high loads, the Norton exponent n2 took considerably higher values (5.7 < n2 < 13.7), indicating that under these test conditions diffusion-controlled dislocation climb is not the prevalent mechanism of deformation. The values for f(S,T), n1 and n2 determined by approximation of the experimental data in the Norton plot using Equation (i) are listed in Table II.

Table II

Coefficients for the Approximation of the Quasi-Stationary Creep Rate as a Function of Stress According to Equation (i)

Temperature, °C f(S,T) f1(S,T) f2(S,T) Norton exponents

n n1 n2

1650 1.3 × 10−13 4.5 × 10−25 5.35 13.68
1800 1.2 × 10−11 4.4 × 10−16 4.40 8.28
2000 5.1 × 10−10 3.6 × 10−11 4.11 5.70
2200 2.5 × 10−9 5.38
2300 2.1 × 10−8 4.99

The temperature dependence of the stationary creep rate can be expressed by an Arrhenius term. Thus, Equation (i) can be rewritten in the form (Equation (ii)):

  (d/dt)av = ασnexp–(QC/RT)    (ii)

where α is a factor dependent on structure, QC is the activation energy for creep, R is the gas constant and T is the temperature in K.

The activation energy, QC, for the creep mechanism can be obtained by plotting ln(d/dt) versus 1/RT. If QC is independent of temperature, ln(d/dt) will show a linear dependence on 1/RT, with the slope of the straight line equal to the activation energy QC. In the present investigations it was not possible to determine QC in the way described, as the creep tests were performed under constant load, not under constant stress. In this case the true stress is a function of the elongation, thus invalidating this method for determining the activation energy.

Fracture Behaviour Results

The iridium samples showed excellent ductility at the temperatures investigated. Particularly at high loads, rupture strain values up to 100% have been measured. At lower loads, the values for the rupture strain proved to be considerably smaller. This finding is in accordance with the observation that iridium exhibits two different fracture modes in stress-rupture tests at high temperatures. According to the literature (9) this change in fracture mode should be accompanied by a discontinuity in the slope of the log-log plot in the stress-rupture diagram (Figure 3).

Whereas intercrystalline fracture occurs without necking at low loads (Figures 7(a) and 7(b)), together with the appearance of extensive intercrystalline crack formation across large parts of the sample (Figures 7(c) and 7(d)), the fracture at high stresses occurs by the transcrystalline mode, accompanied by extensive necking as shown in Figures 8(a) to 8(c). This is in accordance with textbook reports (11) of creep crack behaviour at high temperatures for many materials which are prone to brittle intercrystalline fracture under low loads, but tend to transcrystalline fracture under high loads. It should be mentioned that the numerous cracks and fissures occurring along the grain boundaries during the creep tests under low loads make an additional contribution to the elongation of the sample. This apparent ductility does not, however, influence the intrinsically brittle mode of fracture. Samples exposed to high stresses exhibited only very few intercrystalline cracks. As a consequence the high apparent ductility of those samples can be assumed to be identical to the true ductility of the material.

Fig. 7

SEM images of the surface area near the fracture for pure iridium after stress-rupture test at 1800°C under an initial load of 6.7 MPa: (a) and (c) ×20; (b) ×100; (d) ×50

Fig. 8

SEM images of the surface area near the fracture for pure iridium after stress-rupture test at 1800°C under an initial load of 18 MPa: (a) ×10; (b) ×50; (c) ×500; (d) ×200

On the surface of the samples exhibiting ductile fracture behaviour, slip bands that have formed in the necked region are clearly visible. These are very distinct, and in some cases extend across grain boundaries, sometimes covering several grains. Overlapping slip bands in different directions were observed in some grains. This indicates that different slip systems have been activated during the creep experiments. It is not yet clear whether the slip bands are formed only in the final stage of creep deformation, when the load-carrying cross-section of the sample is significantly reduced due to progressive strain and the true stress and creep rate increase rapidly, or whether the slip bands are formed throughout the creep experiment.

High-Temperature Tensile Tests

These tests were carried out in order to examine the behaviour of pure iridium at high temperatures under dynamic loading. Figure 9 summarises the results for the temperature range between 1600°C and 2300°C. The shapes of the stress-strain curves, particularly those for temperatures between 1800°C and 2100°C, are typical for materials that undergo dynamic recrystallisation. After a very steep rise in the stress-strain curve at the beginning of the tests, the slope of the curve decreases, probably due to both plastic deformation and softening of the material caused by dynamic recovery. When dynamic recrystallisation commences, the softening becomes more severe and the load decreases. This is followed by a period at a more or less constant load – in some cases this phase exhibits a slight oscillation – until the stress drops rapidly when the sample ruptures. This behaviour is typical for materials that are recrystallising dynamically during the high-temperature tensile tests.

Fig. 9

Stress-strain diagram for pure iridium at different temperatures

A comparison of the stress-strain curves with the results of the SEM investigations reported in the previous section shows that the slip bands, shown in Figures 8 and 10, are formed only in samples exposed to loads close to or greater than the yield stress determined in the tensile test at the corresponding temperature. In the present investigation, the deformation mechanism under high loads at high temperatures appears to be based on an interaction of plastic deformation due to dislocation slip with common creep deformation. The plastic deformation in turn leads to an increase in the internal deformation energy, thus promoting the initiation of dynamic recrystallisation.

Fig. 10

SEM image (×200) of the surface morphology of an iridium sample after creep test at 1800°C under 23 MPa load

Yield strength, Rp0.2, tensile strength, Rm, and tensile elongation, A, as determined in the high-temperature tensile tests are plotted in Figure 11 as a function of temperature. Measurements were made on at least three samples for each test temperature. The values given for Rp0.2 and Rm are averages over all three measurements. The standard deviation is indicated as an error bar for each data point, revealing excellent reproducibility for the measurements. The values for A, however, showed a greater scatter. Moreover, some of the samples did not rupture between the markers, in which cases it was not possible to determine the tensile elongation with the “SuperCreep” software. No error bars are given for the relevant data points in Figure 11.

Fig. 11

Yield strength, Rp0.2, tensile strength, Rm, and tensile elongation, A, of pure iridium as a function of temperature

Metallographic Investigations

The microstructure of pure iridium exposed to the influence of high temperatures and different loads during the creep tests was evaluated metallographically. Comparative investigations were carried out on longitudinal sections of samples before and after creep testing. The microstructure of a sample in the initial state (i.e. before creep testing) is shown in Figure 12. The sample exhibits a uniform microstructure with an average grain size of about 100 μm.

Fig. 12

Longitudinal section of pure iridium in the initial recrystallised state

Comparison with metallographic sections of samples that were exposed to high temperatures under different loads (Figures 13 and 14) shows that during creep tests specimens have undergone the expected severe grain coarsening. Due to the high temperatures and long test times, very coarse grains with grain sizes up to 4 mm have formed. Furthermore, the samples exhibit strong intercrystalline crack formation as mentioned above. Figures 13 and 14 show that several – partly very deep – cracks have formed within the same sample. Nevertheless, the material withstood this damage and cracked at a different position, several hours later. Dynamic recrystallisation is apparent in areas of high stress concentration and strong deformation, for instance at crack tips (Figure 13(a)), and close to the fracture in the necking areas (Figure 15). Particularly in the area around the crack tips, this may have led to a reduction in local stresses. As a consequence the crack propagation may have stopped in this area and continued in another part of the sample.

Fig. 13

Longitudinal sections of pure iridium after creep test at 1800°C, 6.7 MPa, 1403.7 h

Fig. 14

Longitudinal sections of pure iridium after creep test at 1800°C, 13 MPa, 56.8 h

Fig. 15

Longitudinal section through the fracture tip of pure iridium after creep test at 1800°C, 8.3 MPa, 385.9 h

The metallographic images show the existence of individual voids along grain boundaries and their coalescence into large pores. These are typically found in materials that have undergone creep deformation. This phenomenon is most frequently observed on grain boundaries that are oriented perpendicular to the direction of applied stress. This often leads to the formation of intercrystalline creep cracks during the final period of creep deformation.

Microanalytical Investigations

As has been demonstrated in previous investigations (14–19), iridium generally tends to brittle intercrystalline fracture due to trace impurities at grain boundaries. Investigation by SIMS on the present iridium samples has shown a very high purity for the material. Only very small traces of impurities in the ppm range were detected, showing no enrichment at the grain boundaries. All the elements detected were distributed homogeneously in the matrix.


Due to its outstanding properties iridium is particularly suited to applications under extreme thermal, chemical and mechanical conditions. In order to obtain the materials data necessary for the design of high-temperature equipment and the numerical simulation of its service performance, the stress-rupture strength, creep behaviour and tensile properties of iridium have been measured over the temperature range 1650–2300°C.

The investigations were performed on hot rolled iridium sheet. The results showed a very good degree of reproducibility. The iridium samples exhibit very high stress-rupture strength. A discontinuity in the slope of the stress-rupture curves indicates the existence of two different fracture modes, depending on the temperature and the initial load applied to the samples. The existence of different fracture mechanisms was confirmed by the examination of the fracture surfaces. The change in fracture mode is probably caused by different deformation mechanisms prevalent under the various test conditions.

A significant anomaly was observed in the creep behaviour of pure iridium – the creep curves contained plateaus in the range of steady-state creep. Metallographic examination, investigations by SEM, and high-temperature tensile tests indicated that dynamic recrystallisation may be the cause of this phenomenon. A further influence may be the activation of various slip systems which can be deduced from the observed slip bands.

Microanalytical investigations by means of scanning SIMS showed a very high purity of the iridium heats investigated, without any enrichment of trace impurities at the grain boundaries.

Initial results on the stress-rupture strength of an iridium-rhenium alloy doped with hafnium and molybdenum indicate that this alloy exhibits rupture times three to four times longer than for pure iridium. Moreover, this alloy shows outstanding ductility, comparable to that of pure iridium. This alloy is therefore of particular interest for high-temperature applications and is the subject of ongoing research.


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The authors would like to thank Margit Friedrich (SEM investigations), Erik Hartmann and Frank Lehner (creep tests) from the Department SciTec of the University of Applied Sciences Jena for their support for these investigations.

The Authors

After studying Materials Science at the University of Saarbrücken, Germany, Reinhold Weiland worked as a research associate at the Max-Planck-Institute for Metals Research in Stuttgart, and received his doctoral degree from the University of Stuttgart. Since 2002 he has been working at W.C. Heraeus GmbH in Hanau as a development project manager. His main interests are the processing and manufacture of precious metals products and composite materials.

Prof. Dr David Lupton is Development Manager, Engineered Materials Division, W. C. Heraeus. His main interests are the manufacture and applic-ations of precious metal and refractory metal products.

Prof. Dr.-Ing. habil. Bernd Fischer studied mechanical engineering and materials science at the Technical University Chemnitz, Germany. After more than 25 years at the University of Jena, Bernd Fischer was appointed to the Chair of Materials Science at the University of Applied Sciences Jena in 1992. For many years, his research interests have included the properties and applications of noble and refractory metals.

Jörgen Merker studied materials science at the Technical University, Dresden, where he also earned his Ph.D. degree. Between 1995 and 2000, he worked as a development project manager for W. C. Heraeus GmbH, Hanau. After a short time in the R&D department of KM Europa Metal AG, Osnabrück, he was appointed to a Professorship in Materials Engineering and Applied Metals Science at the University of Applied Sciences Giessen-Friedberg in 2002. In 2006 Dr Merker became Professor of Materials Technology and Materials Testing at the University of Applied Sciences Jena.

Carolin Scheckenbach studied materials technology at the University of Applied Sciences in Jena, Germany, where she received her diploma degree. After one year as a research associate, she is now working as a trainee in the Department of Melting Technology and Hot Forming at Schott AG, Mainz.

Jörg Witte studied chemistry at the Technical University of Darmstadt, Germany, from which he received his doctoral degree in Materials Science. Since 1999 he has worked at Schott AG, Mainz, where he is currently head of the Materials Competence Department. His main interests are the properties of refractory materials and failure analysis.

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