Platinum Metals Rev., 2009, 53, (3), 138
Plastic Deformation of Polycrystalline Iridium at Room Temperature
UNIQUE PROPERTIES OF IRIDIUM, THE SOLE REFRACTORY FACE CENTRED CUBIC METAL
- Peter Panfilov**
- Alexander Yermakov
- Laboratory of Strength, Ural State University,
- 620083 Ekaterinburg, Russia *E-mail: firstname.lastname@example.org
- Olga V. Antonova
- Vitalii P. Pilyugin
- Institute of Metal Physics, Ural Division of the Russian Academy of Sciences,
- 620219 Ekaterinburg, Russia
Defect structure and its relationship with deformation behaviour at room temperature of iridium, the sole refractory face centred cubic (f.c.c.) metal, are discussed. Small angle boundaries and pile-ups of curvilinear dislocation segments are the main features of dislocation structure in polycrystalline iridium at room temperature, while homogeneously distributed rectilinear dislocation segments were the main element of defect structure of iridium single crystals at the same conditions. Small angle boundaries and pile-ups of curvilinear dislocation segments are formed in iridium single crystals under mechanical treatment at elevated temperatures (≥ 800°C) only. The evolution of defect structure in polycrystalline iridium and other f.c.c. metals under room temperature deformation occurs by the same process: accumulation of dislocations in the matrix leads to the appearance of both new sub-grains and new grains up to the fine grain (nanocrystalline) structure. Neither single straight dislocations nor their pile-ups are observed in iridium at room temperature if small angle boundaries have been formed. This feature may be considered as the reason why polycrystalline iridium demonstrates advanced necking (high localised plasticity) and small total elongation.
Iridium is the only refractory f.c.c. metal. Its melting point is 2446°C, and as one of the platinum group metals (pgms), it exhibits excellent mechanical properties and high resistance to corrosion at elevated temperatures (1, 2). Some features of this metal, such as its poor workability (inclination to brittle fracture under mechanical treatment) and intrinsic fracture mode (brittle transcrystalline fracture) are unlike the deformation behaviour observed in other f.c.c. metals and remain puzzling for the materials science community (3–8). Discussion of the possible mechanisms of deformation and fracture in iridium was begun in Platinum Metals Review more than fifty years ago (9, 10), and continues up to the present. Because refining iridium is a very complicated procedure, segregation of impurities at the grain boundaries is considered the most likely cause of the weak cohesive strength of grain boundaries and, hence, the poor workability of this metal (7, 8, 10, 11, 15). Indeed, it has been shown that high purity iridium can be processed like platinum (16). Polycrystalline iridium free of non-metallic impurities, and its ternary alloy with rhenium and ruthenium, demonstrate both transgranular cleavage and satisfactory plasticity. Traces (∼ 10 ppm) of carbon and oxygen in the matrix induce the appearance of intergranular cleavage on the fracture surfaces and, as a result, workability is catastrophically diminished (7, 11). The extraction of detrimental non-metallic impurities continues to be a significant problem for pgm refiners (7, 8, 15).
Some physical parameters of iridium, such as its strong interatomic bonds and high elastic modulus, give a basis for supposing that brittleness may be an intrinsic property of this metal (5, 6, 20–23). For example, formal substitution of the elastic modulus into cleavage criteria equations leads to the conclusion that iridium is expected to behave like a brittle substance (6, 20, 24). On the other hand, according to empirical theories of plastic deformation, it should be deformed like other f.c.c. metals (13, 20, 25). It is a paradox, but the highest yield stress and strength of iridium single crystals under tension (∼ 100 MPa and ∼ 500 MPa, respectively) become similar to those of other f.c.c. metals when these parameters have been normalised on the elastic modulus (20). This may also be applied to the concept of ‘dislocation mobility’ in iridium, as understood by dislocation theory in the field of continuum mechanics, which describes the motion of single dislocations under external stress (normalised on the elastic modulus). This apparent contradiction with empirical knowledge of the deformation and fracture of f.c.c. metals merits further discussion among the materials science community.
Behaviour of Iridium Single Crystals
Experiments have shown that the intrinsic fracture mode of iridium single crystals is brittle transgranular cleavage. However, monocrystalline iridium is also a highly plastic material, as it fractures after considerable elongation (up to 80%) at room temperature and never fails under compression (19, 20, 26, 27). Hence, the brittleness of iridium in the monocrystalline state is a unique kind of brittle fracture, since cleavage occurs after huge plasticity. Therefore, single crystal iridium may be classified as a plastic but cleavable solid (28). Study of the geometry of deformation tracks on the back surfaces of deformed iridium single crystals has shown that octahedral slip is the sole deformation mechanism active at room temperature. In contrast to other f.c.c. metals, no necking is observed in iridium single crystal samples under tension, in spite of considerable elongation prior to failure. Analysis of the deformation track distribution leads to the conclusion that all plasticity of iridium single crystals is observed during the early stages of plastic deformation (29).
Studies by transmission electron microscopy (TEM) have shown that rectilinear dislocation segments laid along the <110> crystallographic direction are the main element of the defect structure in iridium single crystals at room temperature (11, 26, 29). These straight dislocations are grouped into network-like arrangements, where they are distributed almost equidistantly or homogeneously. The homogeneous distribution of dislocations takes place in thin foils of f.c.c. metals in the early stages of plastic deformation, when dislocation density is expected to be small (30). However, the density of dislocations in iridium single crystals may be so high that the foil is no longer transparent to the electron beam. No small angle boundaries or other types of inhomogeneously distributed dislocations are observed in single crystal iridium deformed at room temperature. Therefore, it may be concluded that the deformation mechanism of iridium single crystals at room temperature is octahedral slip of the perfect dislocation with a <110> Burgers vector (28, 29). This observation distinguishes iridium single crystals from f.c.c. metals with low and intermediate melting points.
The evolution of dislocation structure in iridium single crystals stops at the stage of homogeneous distribution of dislocations in the crystal, which is the first stage of plastic deformation in metallic single crystals (30–32). It should be noted that motion of dislocations never occurs in thin foils of iridium during room temperature TEM experiments, including in situ tension in the column of the microscope. These features, together with the high yield stress of iridium single crystals, lead to the conclusion that <110> dislocations have low ability to move under external stress in the refractory f.c.c. metal at room temperature. As a result, the dominant dislocation arrangements in iridium single crystals cannot transform into small angle boundaries under external stress at room temperature. Such behaviour contrasts with f.c.c. metals having low and intermediate melting points, where this process can occur at room temperature. It may be considered as an additional argument for the statement that all plasticity of iridium single crystals is realised during the early stages of plastic deformation.
High Purity Polycrystalline Iridium
High purity polycrystalline iridium (purity > 99.9%, including < 0.1% metallic impurities and no non-metallic impurities) and iridium alloys also exhibit brittle transcrystalline fracture as an intrinsic fracture mode under tension, but the total elongation of samples having a circular cross-section is less than 10% at room temperature (17–19). This does not mean that iridium wires possess poor plasticity, since advanced necking occurs in the samples even at room temperature, in spite of a considerable traverse rate of 1 mm min−1. The homologous temperature, THomologous, is defined as Texp (K)/Tmelt (K), where Texp = experimental temperature and Tmelt = melting point. For low and intermediate homologous temperatures (THomologous ≤ 0.5), localisation of plasticity induces a visible decrease in the total elongation of wires from 10% down to 5% at THomologous ∼ 0.5 when necking to a point takes place (17). Total elongation begins to increase as soon as the flowing neck of a multi-neck effect appears at THomologous > 0.5. This behaviour at similar traverse rates and homologous temperatures is observed only in gold wires and, hence, high purity polycrystalline iridium may also be considered to be a plastic but cleavable substance.
It is apparent that the deformation behaviour of polycrystalline iridium on the macroscopic scale has many features in common with iridium single crystals and also many differences. For example, advanced necking in polycrystalline iridium wire and the homogeneous distribution of plastic deformation in iridium single crystals under tension both point to the high plasticity of iridium. However, this is due to different effects depending on the microstructure of the material. Currently, no information on the defect structure of polycrystalline iridium is available in the literature. Therefore, the aim of this work is to carry out a TEM study of the dislocation structure of polycrystalline iridium and its evolution during different stages of plastic deformation up to severe deformation. The findings will serve as a basis for further discussion of plastic deformation in polycrystalline iridium at room temperature.
Preparation of Samples
Samples for the research were prepared from high purity polycrystalline iridium, free of non-metallic impurities (7). An electron-beam melted monocrystalline ingot was used for the manufacture of iridium sheets (33). Therefore, the samples contain neither small grain boundaries nor grain boundaries in the initial state. TEM study of iridium single crystals agrees with this conclusion (29). The treatment procedure for work pieces includes forging of the ingot at 1500°C to 2000°C and rolling of the sheet at ≤ 800°C. After that, iridium and its alloys can be deformed like platinum (7).
Recrystallisation annealing of the sheet (for iridium, the recrystallisation temperature is ∼ 1000°C) is not carried out, since it induces a catastrophic drop in workability of iridium due to intergranular brittleness. During processing, grains can be formed under forging only, whereas conditions for the appearance of small angle boundaries exist during both stages of treatment. During the first stage, the iridium matrix should not be contaminated by non-metallic impurities (carbon and oxygen), because iridium oxides are volatile substances at these temperatures. The environment is kept carbon free for this procedure (7). However, iridium does not interact with oxygen at temperatures < 1000°C and carbon lubricant is not used in the rolling mill facility. Hence, segregation of harmful contaminants at the grain boundaries is also avoided during the second stage. The fracture mode of the metal, which has been shown to be a brittle transgranular fracture, may be considered as the proof of this supposition.
Discs with a diameter of 3 mm were stamped from the polycrystalline iridium sheet (thickness of 0.3 mm; grain size ∼ 0.1 mm), which was not recrystallised after processing. Samples were deformed under a ‘shift under pressure’ regime in a Bridgman anvil at room temperature (Figure 1). This procedure allows the highest degrees of plastic deformation to be reached without failure, even for brittle materials (34). The technique of iridium thin foil preparation has been described in References (29, 35). The TEM study was carried out on a conventional 200 kV microscope (JEM-200CX from JEOL).
Deformation and Defect Structure of Iridium
Under compression, high purity iridium behaves like a f.c.c. metal: it never fails under load and exhibits the usual stress-strain curves expected of a f.c.c. metal (20, 27, 36). Therefore, it may be expected that the deformation behaviour of iridium would also be the same as that of other f.c.c. metals, under deformation with reduced tensile (cleaving) stresses. A shift under pressure procedure in a Bridgman anvil allows even brittle and almost unworkable materials to be deformed, due to the suppression of cleaving stresses. In metals, this technique allows severe deformations to be reached, which cannot be obtained under uniaxial (tension or compression) load (34).
Four points were chosen for examination of the microstructure in polycrystalline iridium at different deformation degrees (initial state, 0.5 turn, 1.5 turns and 3 turns), which cover the whole range of structural states from undeformed to severe deformation. The metallographic image of an iridium disc after 1.5 turns is shown in Figure 2. No cracks on the edges or other significant deformation defects are observed on its surface, despite considerable deformation of the material. Radial cracks and separation of the material close to the edges appear after 3 turns (Figure 3). However, even such severe deformation does not lead to failure of the central part of the disc, where the level of cleaving tensile stresses is minimal. Another demonstration of the high plasticity of polycrystalline iridium is the fact that the thin section for TEM observations, which was obtained by electropolishing in the central part of the disc after 0.5 turn, has disappeared or been healed during the next deformation (additional turn). It should be noted that polycrystalline copper and nickel exhibit similar behaviour under the same deformation conditions (37).
TEM study of defect structure has shown that small angle boundaries and dislocation pile-ups, consisting of curvilinear dislocation segments, are the main dislocation arrangements in polycrystalline thin foils in the undeformed state (Figure 4). No rectilinear dislocation segments or their arrangements are observed in the material. It may be proposed that the rectilinear dislocation segments become unstable under thermomechanical treatment of the single crystalline work piece at elevated temperatures, and the initial dislocation structure of iridium single crystals begins to transform into a cellular structure. The same defect structure is observed in polycrystalline f.c.c. metals with low and intermediate melting points at small deformations, and in their single crystals at the third stage of plastic deformation (elongation > 10%) (30, 38–40). Deformation after 0.5 turn does not induce the appearance of new features in the dislocation structure of polycrystalline iridium. Small angle boundaries and pile-ups of curvilinear dislocation segments are the dominant attributes of the defect structure, while the dislocation density in the material naturally increases (Figure 5). The microdiffraction image confirms that local deformation of the material is advanced. However, in contrast to the single crystalline state, rectilinear dislocation segments do not appear in the material under external stress.
A threshold is reached after 1.5 turns, as no cracking or separation of the sample takes place at this stage of deformation, whereas the next stage of plastic deformation leads to the cracking of the disc. The microstructure of the central part of the disc looks the same as after 0.5 turn: its main morphological features are small angle boundaries and pile-ups of curvilinear dislocations close to them, and its microdiffraction pattern points to a severely deformed polycrystalline state (Figure 6). This new kind of defect structure is observed in the vicinity of the disc edge, where deformation of the material is considerably higher than in the centre of the disc. Fine grains, which can reach a few dozen nanometres in diameter, are revealed in the material (Figure 7). Analysis of diffraction patterns confirms the supposition that the nanocrystalline state begins to form in polycrystalline iridium at this stage of plastic deformation. The microstructure of iridium in the middle part of the disc after the last deformation step (3 turns) is shown in Figure 8. This is a fine grain structure, where some grains are approximately 50 nm to 100 nm in diameter, and its microdiffraction pattern is typical of the nanocrystalline state of f.c.c. metallic materials. It should be noted that this nanocrystalline structure is stable in refractory iridium, whereas recovery processes make the same structural state in f.c.c. metals with low and intermediate melting points unstable at room temperature (41). Taking into account the differences between their melting points, it may be supposed that the low ability of <110> dislocations to move under external stress in iridium is the main reason for the stability of the nanocrystalline state in iridium at low homologous temperatures.
Behaviour of High Purity Iridium
The experiment described above has confirmed that polycrystalline iridium, free of non-metallic impurities such as carbon and oxygen, can behave at room temperature like a f.c.c. metal. The ‘shift under pressure’ test in the Bridgman anvil was chosen simply as another deformation scheme (the first being uniaxial compression), where the level of cleaving stresses is minimal. Qualitative analysis of the defect structure of polycrystalline iridium at the different stages of plastic deformation, including examination of its main morphological features and the character of its evolution, does not reveal any difference compared to f.c.c. metals having low and intermediate melting points. Indeed, small angle boundaries and pile-ups of curvilinear dislocation segments are common attributes of the dislocation structure for all f.c.c. metals (30–32). The forming of a fine grain (nanocrystalline) structure under severe deformation also agrees with the supposition that iridium can behave like a f.c.c. metal at room temperature. Therefore, taking into account that octahedral slip of <110> dislocations is the sole deformation mechanism in iridium single crystals at room temperature (28, 29), this mechanism of plasticity may also be considered the dominant one for the polycrystalline metal under the same experimental conditions. In other words, the presence of grain boundaries in the iridium matrix does not provoke additional deformation mechanisms in polycrystalline iridium. This conclusion correlates with the fact that the intrinsic fracture mode of iridium free of non-metallic impurities in both the monocrystalline and polycrystalline states is brittle transgranular cleavage, which does not depend on the presence of grain boundaries in the matrix (19).
In contrast to other f.c.c. metals, perhaps excluding rhodium, which may be considered an analogue of iridium having a melting point of 1963°C (1), the nanocrystalline state in iridium at room temperature is stable. This means that the recovery process has been suppressed at room temperature. This may be explained by the low ability of <110> dislocations to move under external stress, as structural obstacles to dislocation motion, such as second phases or dislocation barriers (sessile dislocations, dislocation ‘forests’ etc.), should be absent in highly plastic f.c.c. metals (30–32). This decrease in the ability of dislocations to move in f.c.c. metals is only seen at low temperatures. Room temperature (∼ 300 K) on the homologous temperature scale for refractory iridium is estimated at about 0.12. For rhodium, nickel, copper and aluminium, this homologous temperature is reached at around 240 K, 150 K, 120 K and 70 K, respectively. Hence, the experiments described in this paper (including thin foil preparation and TEM study) must be carried out at very low temperatures for the majority of f.c.c. metals, which naturally causes great technical problems. Therefore, the existence of a stable nanocrystalline structure at this homologous temperature cannot be considered as a unique property of iridium, at least until this experiment can be carried out for a f.c.c. metal having a low or intermediate melting point. On the microscopic scale (atom level), the low ability of <110> dislocations to move under external stress in iridium may be explained by the fact that it has the strongest interatomic bonds among f.c.c. metals, since it is the sole refractory metal with an f.c.c. lattice. Transition from homogeneously distributed rectilinear dislocations to small angle boundaries is stopped, which leads to the absence of necking in iridium single crystals under room temperature tension. This may also be considered as a consequence of the low ability of <110> dislocations to move under external stress.
According to the well-known empirical scheme of evolution of defect structure in f.c.c. metals, homogeneous distribution of single dislocations leads to homogeneous distribution of plastic deformation in the sample, while the transition from homogeneous distribution to the appearance of small angle boundaries or localised distribution of dislocations correlates with the start of the necking process or the localisation of plasticity in the neck region (31, 32). Indeed, necking always occurs after the stage of homogeneous accumulation of plastic deformation in both single crystal and polycrystalline f.c.c. metals with low and intermediate melting points across a wide temperature range. However, the deformation behaviour of refractory iridium at room temperature depends on the microstructure of samples (single crystal or polycrystalline). In single crystals, necking does not occur: straight <110> dislocations are homogeneously distributed in the material and small angle boundaries are absent. In polycrystalline iridium (iridium wire), necking is the dominant stage of plastic deformation: small angle boundaries and pile-ups of curvilinear dislocations are the main features of the defect structure, while no rectilinear dislocation segments are observed. It should be noted that in both cases, iridium continues to be a highly plastic substance, but this manifests in different ways. This important difference between the deformation behaviour of refractory iridium and other f.c.c. metals needs further experimental study and discussion.
- J. W. Arblaster, Platinum Metals Rev.,1996, 40, (2), 62 LINK https://www.technology.matthey.com/article/40/2/62-63
- R. Weiland, D. F. Lupton, B. Fischer, J. Merker, C. Scheckenbach and J. Witte, Platinum Metals Rev., 2006, 50, (4), 158 LINK https://www.technology.matthey.com/article/50/4/158-170
- G. Reinacher, Metall, 1964, 18, 731
- R. W. Douglass and R. I. Jaffee, Proc. ASTM, 1962, 62, 627
- S. S. Hecker, D. L. Rohr and D. F. Stein, Metall. Trans. A, 1978, 9, (4), 481 LINK http://dx.doi.org/10.1007/BF02646403
- C. Gandhi and M. F. Ashby, Acta Metall., 1979, 27, (10), 1565 LINK http://dx.doi.org/10.1016/0001-6160(79)90042-7
- N. I. Timofeev, A. V. Yermakov, V. A. Dimitriev and P. E. Panfilov, “Metallurgy and Mechanical Behaviour of Iridium”, Urals Branch of Russian Academy of Science, Ekaterinburg, Russia, 1996
- “Iridium”, eds.E. K. Ohriner, R. D. Lanam, P. Panfilov and H. Harada, Proceedings of the international symposium held during the 129th Annual Meeting & Exhibition of The Minerals, Metals and Materials Society (TMS), Nashville, Tennessee, 12th–16th March, 2000, TMS, Warrendale, Pennsylvania, U.S.A., 2000
- F. D. Richardson, Platinum Metals Rev., 1958, 2, (3), 83 LINK https://www.technology.matthey.com/article/2/3/83-85
- B. L. Mordike and C. A. Brookes, Platinum Metals Rev., 1960, 4, (3), 94 LINK https://www.technology.matthey.com/article/4/3/94-99
- C. A. Brookes, J. H. Greenwood and J. L. Routbort, J. Inst. Met., 1970, 98, (1), 27
- L. Heatherly and E. P. George, Acta Mater., 2001, 49, (2), 289 LINK http://dx.doi.org/10.1016/S1359-6454(00)00313-X
- T. J. Balk, K. J. Hemker and L. P. Kubin, Scr. Mater., 2007, 56, (5), 389 LINK http://dx.doi.org/10.1016/j.scriptamat.2006.10.042
- S. P. Lynch, Scr. Mater., 2007, 57, (2), 85 LINK http://dx.doi.org/10.1016/j.scriptamat.2007.03.039
- E. K. Ohriner, Platinum Metals Rev., 2008, 52, (3), 186 LINK https://www.technology.matthey.com/article/52/3/186-197
- J. R. Handley, Platinum Metals Rev., 1986, 30, (1), 12 LINK https://www.technology.matthey.com/article/30/1/12-13
- P. Panfilov, V. Novgorodov and A. Yermakov, J. Mater. Sci. Lett., 1994, 13, (2), 137 LINK http://dx.doi.org/10.1007/BF00416826
- P. Panfilov and A. Yermakov, J. Mater. Sci., 2004, 39, (14), 4543 LINK http://dx.doi.org/10.1023/B:JMSC.0000034148.03387.71
- P. Panfilov, J. Mater. Sci., 2005, 40, (22), 5983 LINK http://dx.doi.org/10.1007/s10853-005-1296-1
- C. N. Reid and J. L. Routbort, Metall. Trans., 1972, 3, (8), 2257 LINK http://dx.doi.org/10.1007/BF02643240
- R. E. MacFarlane, J. A. Rayne and C. K. Jones, Phys. Lett., 1966, 20, (3), 234 LINK http://dx.doi.org/10.1016/0031-9163(66)90340-4
- S. Crampin, K. Hampel, D. D. Vvedensky and J. M. MacLaren, J. Mater. Res., 1990, 5, (10), 2107 LINK http://dx.doi.org/10.1557/JMR.1990.2107
- S. P. Chen, Philos. Mag. A, 1992, 66, (1), 1 LINK http://dx.doi.org/10.1080/01418619208201509
- C. Gandhi and M. F. Ashby, Scr. Metall., 1979, 13, (5), 371 LINK http://dx.doi.org/10.1016/0036-9748(79)90227-8
- P. Haasen, H. Hieber and B. L. Mordike, Z. Metallkd., 1965, 56, (12), 832
- C. A. Brookes, J. H. Greenwood and J. L. Routbort, J. Appl. Phys., 1968, 39, (5), 2391 LINK http://dx.doi.org/10.1063/1.1656565
- A. Yermakov, P. Panfilov and R. Adamesku, J. Mater. Sci. Lett., 1990, 9, (6), 696 LINK http://dx.doi.org/10.1007/BF00721807
- P. Panfilov, ‘Brittle Transcrystalline Fracture in Plastic Face Centered Cubic Metal Iridium’, in“Iridium”, eds.E. K. Ohriner, R. D. Lanam, P. Panfilov and H. Harada, Proceedings of the international symposium held during the 129th Annual Meeting & Exhibition of The Minerals, Metals and Materials Society (TMS), Nashville, Tennessee, 12th–16th March, 2000, TMS, Warrendale, Pennsylvania, U.S.A., 2000, pp. 27–40
- P. Panfilov, J. Mater. Sci., 2007, 42, (19), 8230 LINK http://dx.doi.org/10.1007/s10853-007-1722-7
- P. B. Hirsch, A. Howie, R. B. Nicholson, D. W. Pashley and M. J. Whelan, “Electron Microscopy of Thin Crystals”, Butterworth and Co Ltd, London, 1965, 549 pp
- R. W. K. Honeycombe, “The Plastic Deformation of Metals”, Edward Arnold, London, 1968, 477 pp, Translated into Russian, Mir, Moscow, 1972
- J. P. Hirth and J. Lothe, “Theory of Dislocations”, McGraw-Hill, New York, 1968, 780 pp
- P. Panfilov, A. Yermakov, V. Dmitriev and N. Timofeev, Platinum Metals Rev., 1991, 35, (4), 196 LINK https://www.technology.matthey.com/article/35/4/196-200
- P. W. Bridgman“Studies in Large Plastic Flow and Fracture: With Special Emphasis on the Effects of Hydrostatic Pressure”, McGraw-Hill, New York, 1952, 362 pp
- D. L. Rohr, L. E. Murr and S. S. Hecker, Metall. Trans. A, 1979, 10, (4), 399 LINK http://dx.doi.org/10.1007/BF02697065
- H. Hieber, B. L. Mordike and P. Haasen, Platinum Metals Rev., 1964, 8, (3), 102 LINK https://www.technology.matthey.com/article/8/3/102-106
- Kh. Ya. Mulyukov, S. B. Khahizov and R. Z. Valiev, Phys. Status Solidi A, 1992, 133, (2), 447 LINK http://dx.doi.org/10.1002/pssa.2211330228
- P. B. Hirsch, R. W. Horne and M. J. Whelan, Philos. Mag., 1956, 1, (7), 677 LINK http://dx.doi.org/10.1080/14786435608244003
- J. E. Bailey and P. B. Hirsch, Philos. Mag., 1960, 5, (53), 485 LINK http://dx.doi.org/10.1080/14786436008238300
- J. E. Bailey, Philos. Mag., 1963, 8, (86), 223 LINK http://dx.doi.org/10.1080/14786436308211120
- N. A. Smirnova, V. I. Levit, V. P. Pilyugin, R. I. Kuznetsov, L. S. Davydova and V. A. Sazonova, Fiz. Met. Metalloved., 1986, 61, (6), 1170(in Russian)
The authors would like to thank Professor David Lupton (W. C. Heraeus GmbH, Hanau, Germany) and Professor Easo George (The University of Tennessee, Knoxville, Tennessee, U.S.A.) for helpful discussions. This research was partially supported by the JSC Ekaterinburg Non-Ferrous Metals Processing Plant, the Ministry of Education and Science of the Russian Federation (Grant No. 184.108.40.206/5579) and the U.S. Civilian Research and Development Foundation (CRDF) (Grant No. RUXO-005-EK-06/BG7305).
Peter Panfilov is a Professor of Materials Science at the Ural State University in Ekaterinburg, Russia. His scientific interests are closely connected with the problems of plastic deformation, fracture and processing of the platinum group metals.
Alexander Yermakov is the Technical Director of the platinum group metals manufacturer INTECH at Ekaterinburg, Russia. He has been a visiting researcher at the Ural State University since 1987. His research interest is the study of the properties and creation of industrial technology for the production of noble metals and their alloys.
Olga V. Antonova is a Senior Scientist at the Institute of Metal Physics of Ural Division of the Russian Academy of Sciences in Ekaterinburg. Her area of research activity is TEM study of defect structure and phase transformations in high temperature alloys and intermetallics.
Vitalii P. Pilyugin is the Head of the High Pressure Laboratory at the Institute of Metal Physics of Ural Division of the Russian Academy of Sciences in Ekaterinburg. His research covers deformation behaviour of materials under high pressure including high temperature alloys, structural intermetallics, ceramics and biomaterials. He is also a Professor of Materials Science at the Ural State University in Ekaterinburg, Russia.