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IntroductionIridium is the only refractory f.c.c. metal. Its

melting point is 2446ºC, and as one of the plat-inum group metals (pgms), it exhibits excellentmechanical properties and high resistance to corro-sion at elevated temperatures (1, 2). Some featuresof this metal, such as its poor workability (inclina-tion to brittle fracture under mechanical treatment)and intrinsic fracture mode (brittle transcrystallinefracture) are unlike the deformation behaviourobserved in other f.c.c. metals and remain puzzlingfor the materials science community (3–8).Discussion of the possible mechanisms of defor-mation and fracture in iridium was begun inPlatinum Metals Review more than fifty years ago(9, 10), and continues up to the present. Becauserefining iridium is a very complicated procedure,segregation of impurities at the grain boundaries isconsidered 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 impuri-ties, and its ternary alloy with rhenium andruthenium, demonstrate both transgranular cleav-age and satisfactory plasticity. Traces (~ 10 ppm)of carbon and oxygen in the matrix induce theappearance of intergranular cleavage on the fracture surfaces and, as a result, workability is cat-astrophically diminished (7, 11). The extraction ofdetrimental non-metallic impurities continues tobe a significant problem for pgm refiners (7, 8, 15).

Some physical parameters of iridium, such asits strong interatomic bonds and high elastic mod-ulus, give a basis for supposing that brittlenessmay be an intrinsic property of this metal (5, 6,20–23). For example, formal substitution of the

Platinum Metals Rev., 2009, 53, (3), 138–146

Plastic Deformation of PolycrystallineIridium at Room TemperatureUNIQUE PROPERTIES OF IRIDIUM, THE SOLE REFRACTORY FACE CENTRED CUBIC METAL

By Peter Panfilov* and Alexander YermakovLaboratory of Strength, Ural State University, 620083 Ekaterinburg, Russia; *E-mail: peter.panfilov@usu.ru

and Olga V. Antonova and Vitalii P. PilyuginInstitute 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 ofiridium, the sole refractory face centred cubic (f.c.c.) metal, are discussed. Small angleboundaries and pile-ups of curvilinear dislocation segments are the main features of dislocationstructure in polycrystalline iridium at room temperature, while homogeneously distributedrectilinear dislocation segments were the main element of defect structure of iridium singlecrystals at the same conditions. Small angle boundaries and pile-ups of curvilinear dislocationsegments are formed in iridium single crystals under mechanical treatment at elevatedtemperatures (≥ 800ºC) only. The evolution of defect structure in polycrystalline iridium andother 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-grainsand new grains up to the fine grain (nanocrystalline) structure. Neither single straightdislocations nor their pile-ups are observed in iridium at room temperature if small angleboundaries have been formed. This feature may be considered as the reason why polycrystallineiridium demonstrates advanced necking (high localised plasticity) and small total elongation.

DOI: 10.1595/147106709X463318

elastic modulus into cleavage criteria equationsleads to the conclusion that iridium is expected tobehave like a brittle substance (6, 20, 24). On theother hand, according to empirical theories ofplastic deformation, it should be deformed likeother f.c.c. metals (13, 20, 25). It is a paradox, butthe highest yield stress and strength of iridium sin-gle crystals under tension (~ 100 MPa and ~ 500MPa, respectively) become similar to those ofother f.c.c. metals when these parameters havebeen normalised on the elastic modulus (20). Thismay also be applied to the concept of ‘dislocationmobility’ in iridium, as understood by dislocationtheory in the field of continuum mechanics, whichdescribes the motion of single dislocations underexternal stress (normalised on the elastic modu-lus). This apparent contradiction with empiricalknowledge of the deformation and fracture off.c.c. metals merits further discussion among thematerials science community.

Behaviour of Iridium SingleCrystals

Experiments have shown that the intrinsicfracture mode of iridium single crystals is brittletransgranular cleavage. However, monocrystallineiridium is also a highly plastic material, as it frac-tures after considerable elongation (up to 80%) atroom temperature and never fails under compres-sion (19, 20, 26, 27). Hence, the brittleness ofiridium in the monocrystalline state is a uniquekind of brittle fracture, since cleavage occurs afterhuge plasticity. Therefore, single crystal iridiummay be classified as a plastic but cleavable solid(28). Study of the geometry of deformation trackson the back surfaces of deformed iridium singlecrystals has shown that octahedral slip is the soledeformation mechanism active at room tempera-ture. In contrast to other f.c.c. metals, no neckingis observed in iridium single crystal samples undertension, in spite of considerable elongation priorto failure. Analysis of the deformation track distri-bution leads to the conclusion that all plasticity ofiridium single crystals is observed during the earlystages of plastic deformation (29).

Studies by transmission electron microscopy(TEM) have shown that rectilinear dislocation

segments laid along the <110> crystallographicdirection are the main element of the defect struc-ture in iridium single crystals at room temperature(11, 26, 29). These straight dislocations aregrouped into network-like arrangements, wherethey are distributed almost equidistantly or homo-geneously. The homogeneous distribution ofdislocations takes place in thin foils of f.c.c. met-als in the early stages of plastic deformation, whendislocation density is expected to be small (30).However, the density of dislocations in iridiumsingle crystals may be so high that the foil is nolonger transparent to the electron beam. No smallangle boundaries or other types of inhomoge-neously distributed dislocations are observed insingle crystal iridium deformed at room tempera-ture. Therefore, it may be concluded that thedeformation mechanism of iridium single crystalsat room temperature is octahedral slip of the per-fect dislocation with a <110> Burgers vector (28,29). This observation distinguishes iridium singlecrystals from f.c.c. metals with low and intermedi-ate melting points.

The evolution of dislocation structure in iridium single crystals stops at the stage of homogeneous distribution of dislocations in thecrystal, which is the first stage of plastic deformation in metallic single crystals (30–32). It should be noted that motion of dislocationsnever occurs in thin foils of iridium during room temperature TEM experiments, including in situtension in the column of the microscope. Thesefeatures, together with the high yield stress of iridium single crystals, lead to the conclusion that<110> dislocations have low ability to moveunder external stress in the refractory f.c.c. metalat room temperature. As a result, the dominantdislocation arrangements in iridium single crystalscannot transform into small angle boundariesunder external stress at room temperature. Such behaviour contrasts with f.c.c. metals having lowand intermediate melting points, where thisprocess can occur at room temperature. It may beconsidered as an additional argument for the statement that all plasticity of iridium single crystals is realised during the early stages of plasticdeformation.

Platinum Metals Rev., 2009, 53, (3) 139

High Purity Polycrystalline Iridium High purity polycrystalline iridium (purity

> 99.9%, including < 0.1% metallic impuritiesand no non-metallic impurities) and iridium alloysalso exhibit brittle transcrystalline fracture as anintrinsic fracture mode under tension, but thetotal elongation of samples having a circularcross-section is less than 10% at room tempera-ture (17–19). This does not mean that iridiumwires possess poor plasticity, since advancednecking occurs in the samples even at room tem-perature, in spite of a considerable traverse rateof 1 mm min–1. The homologous temperature,THomologous, is defined as Texp (K)/Tmelt (K),where Texp = experimental temperature andTmelt = melting point. For low and intermediatehomologous temperatures (THomologous ≤ 0.5),localisation of plasticity induces a visible decreasein the total elongation of wires from 10% downto 5% at THomologous ~ 0.5 when necking to apoint takes place (17). Total elongation begins toincrease as soon as the flowing neck of a multi-neck effect appears at THomologous > 0.5. Thisbehaviour at similar traverse rates and homolo-gous temperatures is observed only in gold wiresand, hence, high purity polycrystalline iridiummay also be considered to be a plastic but cleavable substance.

It is apparent that the deformation behaviourof polycrystalline iridium on the macroscopic scalehas many features in common with iridium singlecrystals and also many differences. For example,advanced necking in polycrystalline iridium wireand the homogeneous distribution of plasticdeformation in iridium single crystals under ten-sion both point to the high plasticity of iridium.However, this is due to different effects dependingon the microstructure of the material. Currently,no information on the defect structure of poly-crystalline iridium is available in the literature.Therefore, the aim of this work is to carry out aTEM study of the dislocation structure of poly-crystalline iridium and its evolution duringdifferent stages of plastic deformation up to severedeformation. The findings will serve as a basis forfurther discussion of plastic deformation in poly-crystalline 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 meltedmonocrystalline ingot was used for the manufac-ture of iridium sheets (33). Therefore, the samplescontain neither small grain boundaries nor grainboundaries in the initial state. TEM study of iridi-um single crystals agrees with this conclusion (29).The treatment procedure for work pieces includesforging of the ingot at 1500ºC to 2000ºC androlling of the sheet at ≤ 800ºC. After that, iridiumand its alloys can be deformed like platinum (7).

Recrystallisation annealing of the sheet (foriridium, the recrystallisation temperature is~ 1000ºC) is not carried out, since it induces a cat-astrophic drop in workability of iridium due tointergranular brittleness. During processing, grainscan be formed under forging only, whereas condi-tions for the appearance of small angle boundariesexist during both stages of treatment. During thefirst stage, the iridium matrix should not be conta-minated by non-metallic impurities (carbon andoxygen), because iridium oxides are volatile sub-stances at these temperatures. The environment iskept carbon free for this procedure (7). However,iridium does not interact with oxygen at tempera-tures < 1000ºC and carbon lubricant is not used inthe rolling mill facility. Hence, segregation ofharmful contaminants at the grain boundaries isalso avoided during the second stage. The fracturemode of the metal, which has been shown to be abrittle transgranular fracture, may be considered asthe proof of this supposition.

Discs with a diameter of 3 mm were stampedfrom the polycrystalline iridium sheet (thicknessof 0.3 mm; grain size ~ 0.1 mm), which was notrecrystallised after processing. Samples weredeformed under a ‘shift under pressure’ regime ina Bridgman anvil at room temperature (Figure 1).This procedure allows the highest degrees of plas-tic deformation to be reached without failure,even for brittle materials (34). The technique ofiridium thin foil preparation has been described inReferences (29, 35). The TEM study was carriedout on a conventional 200 kV microscope (JEM-200CX from JEOL).

Platinum Metals Rev., 2009, 53, (3) 140

Deformation and Defect Structureof Iridium

Under compression, high purity iridiumbehaves like a f.c.c. metal: it never fails under loadand exhibits the usual stress-strain curves expectedof a f.c.c. metal (20, 27, 36). Therefore, it may beexpected that the deformation behaviour of iridi-um would also be the same as that of other f.c.c.metals, under deformation with reduced tensile(cleaving) stresses. A shift under pressure proce-dure in a Bridgman anvil allows even brittle andalmost unworkable materials to be deformed, dueto the suppression of cleaving stresses. In metals,this technique allows severe deformations to bereached, which cannot be obtained under uniaxial(tension or compression) load (34).

Four points were chosen for examination of themicrostructure in polycrystalline iridium at differ-ent deformation degrees (initial state, 0.5 turn, 1.5turns and 3 turns), which cover the whole range ofstructural states from undeformed to severe defor-mation. The metallographic image of an iridiumdisc after 1.5 turns is shown in Figure 2. No crackson the edges or other significant deformationdefects are observed on its surface, despite consid-erable deformation of the material. Radial cracksand separation of the material close to the edgesappear after 3 turns (Figure 3). However, evensuch severe deformation does not lead to failure ofthe central part of the disc, where the level ofcleaving tensile stresses is minimal. Anotherdemonstration of the high plasticity of polycrys-

talline iridium is the fact that the thin section forTEM observations, which was obtained by elec-tropolishing in the central part of the disc after 0.5turn, has disappeared or been healed during thenext deformation (additional turn). It should benoted that polycrystalline copper and nickel exhib-it similar behaviour under the same deformationconditions (37).

TEM study of defect structure has shown thatsmall angle boundaries and dislocation pile-ups,consisting of curvilinear dislocation segments, arethe main dislocation arrangements in polycrys-talline thin foils in the undeformed state (Figure 4).No rectilinear dislocation segments or theirarrangements are observed in the material. It may

Platinum Metals Rev., 2009, 53, (3) 141

Sample

Bridgman Anvil

Fig. 1 Bridgman anvil used for the ‘shiftunder pressure’ test on the iridium samples

0.5 mm

Fig. 2 Iridium disc after ‘shift under pressure’ testing ina Bridgman anvil after 1.5 turns

be proposed that the rectilinear dislocation seg-ments become unstable under thermomechanicaltreatment of the single crystalline work piece atelevated temperatures, and the initial dislocationstructure of iridium single crystals begins to trans-form into a cellular structure. The same defectstructure is observed in polycrystalline f.c.c. metalswith low and intermediate melting points at smalldeformations, and in their single crystals at thethird stage of plastic deformation (elongation> 10%) (30, 38–40). Deformation after 0.5 turndoes not induce the appearance of new features inthe dislocation structure of polycrystalline iridium.Small angle boundaries and pile-ups of curvilineardislocation segments are the dominant attributesof the defect structure, while the dislocation

density in the material naturally increases (Figure5). The microdiffraction image confirms that localdeformation of the material is advanced. However,in contrast to the single crystalline state, rectilineardislocation segments do not appear in the materi-al under external stress.

A threshold is reached after 1.5 turns, as nocracking or separation of the sample takes place atthis stage of deformation, whereas the next stageof plastic deformation leads to the cracking of thedisc. The microstructure of the central part of thedisc looks the same as after 0.5 turn: its main mor-phological features are small angle boundaries andpile-ups of curvilinear dislocations close to them,and its microdiffraction pattern points to aseverely deformed polycrystalline state (Figure 6).This new kind of defect structure is observed inthe vicinity of the disc edge, where deformationof the material is considerably higher than in thecentre of the disc. Fine grains, which can reach afew dozen nanometres in diameter, are revealed inthe material (Figure 7). Analysis of diffraction pat-terns confirms the supposition that thenanocrystalline state begins to form in polycrys-talline iridium at this stage of plastic deformation.The microstructure of iridium in the middle partof 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 to100 nm in diameter, and its microdiffraction pat-tern is typical of the nanocrystalline state of f.c.c.metallic materials. It should be noted that this

Platinum Metals Rev., 2009, 53, (3) 142

0.5 mm

Fig. 3 Iridium disc after ‘shift under pressure’ testing ina Bridgman anvil after 3 turns

1 μm

Fig. 4 Microstructure of polycrystallineiridium at room temperature in the initial(undeformed) state (Inset: Electronmicrodiffraction pattern taken from thesame place)

nanocrystalline structure is stable in refractoryiridium, whereas recovery processes make thesame structural state in f.c.c. metals with low andintermediate melting points unstable at room tem-

perature (41). Taking into account the differencesbetween their melting points, it may be supposedthat the low ability of <110> dislocations to moveunder external stress in iridium is the main reason

Platinum Metals Rev., 2009, 53, (3) 143

0.5 μm

Fig. 5 Microstructure of polycrystallineiridium at room temperature after 0.5 turns(Inset: Electron microdiffraction patterntaken from the same place)

0.5 μm

Fig. 6 Microstructure of polycrystallineiridium at room temperature after 1.5 turns(centre of the disc) (Inset: Electronmicrodiffraction pattern taken from thesame place)

0.5 μm

Fig. 7 Microstructure of polycrystallineiridium at room temperature after 1.5 turns(edge of the disc) (Inset: Electronmicrodiffraction pattern taken from thesame place)

for the stability of the nanocrystalline state in irid-ium at low homologous temperatures.

Behaviour of High Purity Iridium The experiment described above has confirmed

that polycrystalline iridium, free of non-metallicimpurities such as carbon and oxygen, can behaveat room temperature like a f.c.c. metal. The ‘shiftunder pressure’ test in the Bridgman anvil waschosen simply as another deformation scheme(the first being uniaxial compression), where thelevel of cleaving stresses is minimal. Qualitativeanalysis of the defect structure of polycrystallineiridium at the different stages of plastic deforma-tion, including examination of its mainmorphological features and the character of itsevolution, does not reveal any difference com-pared to f.c.c. metals having low and intermediatemelting points. Indeed, small angle boundaries andpile-ups of curvilinear dislocation segments arecommon attributes of the dislocation structure forall f.c.c. metals (30–32). The forming of a finegrain (nanocrystalline) structure under severedeformation also agrees with the supposition thatiridium can behave like a f.c.c. metal at room tem-perature. Therefore, taking into account thatoctahedral slip of <110> dislocations is the soledeformation mechanism in iridium single crystalsat room temperature (28, 29), this mechanism ofplasticity may also be considered the dominantone for the polycrystalline metal under the sameexperimental conditions. In other words, the pres-

ence of grain boundaries in the iridium matrixdoes not provoke additional deformation mecha-nisms in polycrystalline iridium. This conclusioncorrelates with the fact that the intrinsic fracturemode of iridium free of non-metallic impurities inboth the monocrystalline and polycrystalline statesis brittle transgranular cleavage, which does notdepend on the presence of grain boundaries in thematrix (19).

In contrast to other f.c.c. metals, perhapsexcluding rhodium, which may be considered ananalogue of iridium having a melting point of1963ºC (1), the nanocrystalline state in iridium atroom temperature is stable. This means that therecovery process has been suppressed at roomtemperature. This may be explained by the lowability of <110> dislocations to move under exter-nal stress, as structural obstacles to dislocationmotion, such as second phases or dislocation bar-riers (sessile dislocations, dislocation ‘forests’ etc.),should be absent in highly plastic f.c.c. metals(30–32). This decrease in the ability of dislocationsto move in f.c.c. metals is only seen at low temper-atures. Room temperature (~ 300 K) on thehomologous temperature scale for refractory iridi-um is estimated at about 0.12. For rhodium, nickel,copper and aluminium, this homologous tempera-ture is reached at around 240 K, 150 K, 120 K and70 K, respectively. Hence, the experimentsdescribed in this paper (including thin foil prepara-tion and TEM study) must be carried out at verylow temperatures for the majority of f.c.c. metals,

Platinum Metals Rev., 2009, 53, (3) 144

0.5 μm

Fig. 8 Microstructure of polycrystallineiridium at room temperature after 3 turns(Inset: Electron microdiffraction patterntaken from the same place)

which naturally causes great technical problems.Therefore, the existence of a stable nanocrys-talline structure at this homologous temperaturecannot be considered as a unique property of irid-ium, at least until this experiment can be carriedout for a f.c.c. metal having a low or intermediatemelting point. On the microscopic scale (atomlevel), the low ability of <110> dislocations tomove under external stress in iridium may beexplained by the fact that it has the strongest inter-atomic bonds among f.c.c. metals, since it is thesole refractory metal with an f.c.c. lattice.Transition from homogeneously distributed recti-linear dislocations to small angle boundaries isstopped, which leads to the absence of necking iniridium single crystals under room temperaturetension. This may also be considered as a conse-quence of the low ability of <110> dislocations tomove under external stress.

Conclusions According to the well-known empirical scheme

of evolution of defect structure in f.c.c. metals,homogeneous distribution of single dislocationsleads to homogeneous distribution of plasticdeformation in the sample, while the transitionfrom homogeneous distribution to the appearanceof small angle boundaries or localised distributionof dislocations correlates with the start of thenecking process or the localisation of plasticity inthe neck region (31, 32). Indeed, necking alwaysoccurs after the stage of homogeneous accumula-tion of plastic deformation in both single crystal

and polycrystalline f.c.c. metals with low and inter-mediate melting points across a wide temperaturerange. However, the deformation behaviour ofrefractory iridium at room temperature depends onthe microstructure of samples (single crystal orpolycrystalline). In single crystals, necking does notoccur: straight <110> dislocations are homoge-neously distributed in the material and small angleboundaries are absent. In polycrystalline iridium(iridium wire), necking is the dominant stage ofplastic deformation: small angle boundaries andpile-ups of curvilinear dislocations are the mainfeatures of the defect structure, while no rectilineardislocation segments are observed. It should benoted that in both cases, iridium continues to be ahighly plastic substance, but this manifests in dif-ferent ways. This important difference between thedeformation behaviour of refractory iridium andother f.c.c. metals needs further experimentalstudy and discussion.

AcknowledgementsThe authors would like to thank Professor

David Lupton (W. C. Heraeus GmbH, Hanau,Germany) and Professor Easo George (TheUniversity of Tennessee, Knoxville, Tennessee,U.S.A.) for helpful discussions. This research waspartially supported by the JSC Ekaterinburg Non-Ferrous Metals Processing Plant, the Ministry ofEducation and Science of the Russian Federation(Grant No. 2.2.2.2/5579) and the U.S. CivilianResearch and Development Foundation (CRDF)(Grant No. RUXO-005-EK-06/BG7305).

Platinum Metals Rev., 2009, 53, (3) 145

1 J. W. Arblaster, Platinum Metals Rev., 1996, 40, (2), 622 R. Weiland, D. F. Lupton, B. Fischer, J. Merker, C.

Scheckenbach and J. Witte, Platinum Metals Rev., 2006,50, (4), 158

3 G. Reinacher, Metall, 1964, 18, 7314 R. W. Douglass and R. I. Jaffee, Proc. ASTM, 1962,

62, 6275 S. S. Hecker, D. L. Rohr and D. F. Stein, Metall. Trans.

A, 1978, 9, (4), 4816 C. Gandhi and M. F. Ashby, Acta Metall., 1979, 27,

(10), 15657 N. I. Timofeev, A. V. Yermakov, V. A. Dimitriev and

P. E. Panfilov, “Metallurgy and Mechanical Behaviourof Iridium”, Urals Branch of Russian Academy ofScience, Ekaterinburg, Russia, 1996

8 “Iridium”, eds. E. K. Ohriner, R. D. Lanam, P. Panfilovand H. Harada, Proceedings of the internationalsymposium held during the 129th Annual Meeting& Exhibition of The Minerals, Metals and MaterialsSociety (TMS), Nashville, Tennessee, 12th–16thMarch, 2000, TMS, Warrendale, Pennsylvania, U.S.A.,2000

9 F. D. Richardson, Platinum Metals Rev., 1958, 2, (3), 8310 B. L. Mordike and C. A. Brookes, Platinum Metals Rev.,

1960, 4, (3), 94

References

11 C. A. Brookes, J. H. Greenwood and J. L. Routbort,J. Inst. Met., 1970, 98, (1), 27

12 L. Heatherly and E. P. George, Acta Mater., 2001,49, (2), 289

13 T. J. Balk, K. J. Hemker and L. P. Kubin, Scr. Mater.,2007, 56, (5), 389

14 S. P. Lynch, Scr. Mater., 2007, 57, (2), 8515 E. K. Ohriner, Platinum Metals Rev., 2008, 52, (3), 18616 J. R. Handley, Platinum Metals Rev., 1986, 30, (1), 1217 P. Panfilov, V. Novgorodov and A. Yermakov, J.

Mater. Sci. Lett., 1994, 13, (2), 13718 P. Panfilov and A. Yermakov, J. Mater. Sci., 2004,

39, (14), 454319 P. Panfilov, J. Mater. Sci., 2005, 40, (22), 598320 C. N. Reid and J. L. Routbort, Metall. Trans., 1972, 3,

(8), 225721 R. E. MacFarlane, J. A. Rayne and C. K. Jones, Phys.

Lett., 1966, 20, (3), 23422 S. Crampin, K. Hampel, D. D. Vvedensky and J. M.

MacLaren, J. Mater. Res., 1990, 5, (10), 210723 S. P. Chen, Philos. Mag. A, 1992, 66, (1), 124 C. Gandhi and M. F. Ashby, Scr. Metall., 1979, 13, (5),

37125 P. Haasen, H. Hieber and B. L. Mordike, Z. Metallkd.,

1965, 56, (12), 83226 C. A. Brookes, J. H. Greenwood and J. L. Routbort,

J. Appl. Phys., 1968, 39, (5), 239127 A. Yermakov, P. Panfilov and R. Adamesku, J. Mater.

Sci. Lett., 1990, 9, (6), 69628 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 & Exhibitionof The Minerals, Metals and Materials Society (TMS),Nashville, Tennessee, 12th–16th March, 2000, TMS,Warrendale, Pennsylvania, U.S.A., 2000, pp. 27–40

29 P. Panfilov, J. Mater. Sci., 2007, 42, (19), 823030 P. B. Hirsch, A. Howie, R. B. Nicholson, D. W. Pashley

and M. J. Whelan, “Electron Microscopy of ThinCrystals”, Butterworth and Co Ltd, London, 1965,549 pp

31 R. W. K. Honeycombe, “The Plastic Deformation ofMetals”, Edward Arnold, London, 1968, 477 pp,Translated into Russian, Mir, Moscow, 1972

32 J. P. Hirth and J. Lothe, “Theory of Dislocations”,McGraw-Hill, New York, 1968, 780 pp

33 P. Panfilov, A. Yermakov, V. Dmitriev and N.Timofeev, Platinum Metals Rev., 1991, 35, (4), 196

34 P. W. Bridgman “Studies in Large Plastic Flow andFracture: With Special Emphasis on the Effects ofHydrostatic Pressure”, McGraw-Hill, New York, 1952,362 pp

35 D. L. Rohr, L. E. Murr and S. S. Hecker, Metall. Trans.A, 1979, 10, (4), 399

36 H. Hieber, B. L. Mordike and P. Haasen, PlatinumMetals Rev., 1964, 8, (3), 102

37 Kh. Ya. Mulyukov, S. B. Khahizov and R. Z. Valiev,Phys. Status Solidi A, 1992, 133, (2), 447

38 P. B. Hirsch, R. W. Horne and M. J. Whelan, Philos.Mag., 1956, 1, (7), 677

39 J. E. Bailey and P. B. Hirsch, Philos. Mag., 1960, 5, (53),485

40 J. E. Bailey, Philos. Mag., 1963, 8, (86), 22341 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)

Platinum Metals Rev., 2009, 53, (3) 146

The Authors

Peter Panfilov is a Professor ofMaterials Science at the Ural StateUniversity in Ekaterinburg, Russia. Hisscientific interests are closelyconnected with the problems of plasticdeformation, fracture and processingof the platinum group metals.

Alexander Yermakov is the TechnicalDirector of the platinum group metalsmanufacturer INTECH at Ekaterinburg,Russia. He has been a visitingresearcher at the Ural State Universitysince 1987. His research interest is thestudy of the properties and creation ofindustrial technology for the productionof noble metals and their alloys.

Olga V. Antonova is a Senior Scientistat the Institute of Metal Physics of UralDivision of the Russian Academy ofSciences in Ekaterinburg. Her area ofresearch activity is TEM study of defectstructure and phase transformations inhigh temperature alloys andintermetallics.

Vitalii P. Pilyugin is the Head of theHigh Pressure Laboratory at theInstitute of Metal Physics of UralDivision of the Russian Academy ofSciences in Ekaterinburg. His researchcovers deformation behaviour ofmaterials under high pressureincluding high temperature alloys,structural intermetallics, ceramics andbiomaterials. He is also a Professor ofMaterials Science at the Ural StateUniversity in Ekaterinburg, Russia.