2. LOS PRESUPUESTOS DE LAS COMPRENSIONES QUE TRIVIALIZAN EL A PRIORI HISTÓRICO
2.1. Tres interpretaciones del a priori histórico
Initial interest in Cu-Nb system started with using Nb as one of the alloying constituents for elemental Cu, to increase the low intrinsic strength of Cu while retaining its high electrical conductivity for high field magnet applications [42–44]. Cu-Nb was reported to show the best mechanical properties among the remaining potential alloying constituents in the family of bcc elements (e.g. Cr, W, Ta, Nb, Mo, V) [42]. Cu-Nb system has negligibly small solubility in thermodynamic equilibrium. As the melting temperature for Cu and Nb differ from each other significantly (Tm (Cu) = 1085oC and Tm (Nb) = 2469oC) (Fig. 3.1), the synthesis of alloys from
this highly immiscible system mainly relied on rolling, extrusion and mechanical alloying (MA). Due to high-energy impacts during milling, MA was reported to provide alloys with good
23
compositional and microstructural homogeneity and with extended solid solubility of alloy elements [43].
The interest in Cu-Nb system increased when experimental investigations showed that Cu/Nb multilayer composites (with individual layer thickness < 100 nm) manufactured by a bottom-up technique, physical vapor deposition (PVD), resulted in well-controlled lamellar structures that exhibit high thermal stability, strength, ductility, resistance to shock damage and resistance to radiation damage; significantly outperforming bulk Cu and Nb. For example, self-supported sputter deposited Cu-Nb thin films with 75 nm individual layer thickness showed stable layer structure after annealing at 800oC [45]. Moreover, when Cu/Nb multilayers were subjected to high dose implantations of 33 keV and 150 keV He+ at room temperature [46] and at elevated temperatures [47], these interfaces showed superior radiation resistance compared to single-crystal Cu or Nb, while no amorphization, large changes in the interface structure or atomic scale defects such as Frank loops or stacking fault tetrahedra were detected.
These desirable properties displayed by Cu/Nb nanocomposites were tied to the unusually high density of Cu/Nb interfaces that makes up a significant volume fraction of these multilayers (with layer thickness as small as 4 nm) and low density of defects accumulated at the interface when subjected to environmental conditions mentioned above. It is observed that Cu/Nb multilayer composites grown using PVD predominantly show Kurdjumov Sachs (K-S) [48] orientation, in which the Cu and Nb layers are oriented with respect to each other such that <110> direction of Cu is parallel to <111> direction in Nb in the interface plane, while close packed planes of Cu (111) and Nb (110) form an interface. Atomic simulation studies predict that this interface is atomically flat, and has a low energy configuration [49]. Additionally, it was claimed that these nanostructures have a self-healing interface, in the sense that they contain a high density of atomic
24
sites that can effectively promote the recombination and annihilation of point defects [45, 49–51]. The remarkable properties observed in Cu/Nb multilayers suggest that such nanocomposite materials are attractive candidates for a range of structural applications such as next generation nuclear power reactors and military applications. However, with PVD process, only thin films can be produced, and manufacturing large scale composites is prohibitively time-consuming. Several SPD techniques were used in the attempt of successfully synthesizing bulk samples. Cold drawing is employed as one of the SPD applications for producing nanocomposite Cu/Nb [52, 53]. This process stabilized highly textured Cu grains and Nb fibers and semi-coherent Cu/Nb interfaces. A slight tilt (≈ 4o) was observed between <111> Cu and <110> Nb [52, 54].
Fig. 2.7. Cu/Nb samples fabricated with ARB processing. The individual layer thicknesses range from 1 mm to 500 nm (upper left photo), as low as ~ 9 nm (TEM micrographs) [55].
Very recently, Cu-Nb multilayered composites were successfully synthesized from monolithic Cu and Nb metals sheets with the use of ARB [55, 56], with individual layer thicknesses
25
as small as 9 nm, as seen in Fig. 2.7 [55]. It was observed that under rolling strains, these Cu-Nb nanocomposites (below layer thickness ≈ 1 μm) show predominantly K-S orientation relationship between adjacent Cu and Nb grains, where {112}Cu ||{112}Nb ||ND, <111>Cu ||<110>Nb ||RD, and
<110>Cu ||<111>Nb ||TD, with ND, RD, and TD being the direction normal to the rolling plane, the
rolling direction, and the transverse direction, respectively, see Fig. 2.8 [57].
Fig. 2.8. Illustration of {112} K-S interface orientation [57].
The stabilization of this unusual orientation relationship, which is abbreviated as {112} K-S, has been rationalized by the favorable conditions for slip transfer across the Cu/Nb interfaces for these interfaces [55]. It is noteworthy that this orientation is distinct from the traditional K-S orientation relationship found when Cu/Nb multilayers are grown by PVD, in which case the habit plane of the interfaces is {111}Cu ||{110}Nb.
Since the report of {112} K-S orientation, several studies focused on understanding this interface at atomic and microscopic scale. Atomic simulation modeling studies indicate that
26
{112} K-S interface has an ordered but stepped atomic structure with high formation of energy (~ 820 – 886 J/m2) compared to Cu/Nb interfaces synthesized with PVD [57, 58]. Additionally
while K-S interfaces obtained with PVD have low shear strength [51], computer simulations predicted that {112} K-S interfaces possess high shear strength [57].
In order for bulk ARB nanolayered Cu/Nb composites to be used in structural applications mentioned above, it is crucial that these materials exhibit high thermal stability as well as superior mechanical properties. Carpenter et al. [59] investigated bulk nanolamellar Cu/Nb composites with 18 nm individual layer thicknesses, processed by ARB, after annealing at temperatures up to 900oC in vacuum. It was found that Cu/Nb layer structure was maintained at least up to 700oC. At this temperature, the high energy crystal orientations was found to disappear, resulting in sharpening of the existing predominant {112} K-S texture [59].
In this dissertation, the behavior of Cu/Nb nanocomposites initially produced by ARB technique with material properties briefly described above, were investigated after being subjected to severe shear deformation via HPT. The principles of HPT were given in Section 2.1.2. As mentioned previously, this SPD technique introduces a different deformation mode (simple shear) than ARB (pure shear) while enabling quantification of deformation applied on the material. Large strains can be applied via HPT under high hydrostatic pressure, which prevents cracking, therefore, failure of the sample.
Prior studies of HPT mainly rely on investigating elemental metal systems. In the scope of this dissertation, we are interested in behavior of pure Cu and pure Nb under severe shear deformation, as the deformation behavior of these metals is expected to differ from the Cu-Nb nanocomposite layers studied in this research. This dissertation and the paper published on the first topic of this dissertation (Chapter 4) [60], are the first published findings on Cu-Nb system
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subjected to HPT. However, effect of SPD on pure Cu by HPT has been investigated in several studies [61–64] and it is therefore useful to briefly summarize the key findings on pure Cu. Hebesberger et al. [61] ran a systematic parameter study on the resulting microstructure of pure Cu deformed by HPT with ranging hydrostatic pressures of 0.8, 2 and 8 GPa and experiment temperatures of 25oC, 120oC and 200oC. It was found that the microstructure reached to a steady state with substructures as small as 0.4 μm at strain γ = 250 at room temperature. It was reported that the length-scale of steady state microstructure decreases with applied strain but increases with increasing temperature. Later, it was suggested by Edalati et al. [62], that the grain refinement seen in pure Cu samples subjected to HPT at room temperature may be achieved by dislocation coalescence at the early stages of deformation leading to formation of subgrain boundaries. However, as deformation proceeds, competition between recrystallization and increasing dislocation density takes place, which may lead to disappearance of dislocations, until a balance between these two dynamical competition reaches to balance. Lastly, a detailed study on texture evolution of pure Cu subjected to HPT at room temperature [63] up to 10 rotations revealed an uncommon texture at strains ~ 440. It was reported that pure Cu showed strong 𝐵/𝐵̅ texture, which was explained by a deviation from simple shear deformation during HPT. Due to the lack of intensity in both 𝐴/𝐴̅ and cube texture components, one can exclude the possibility that significant dynamic crystallization took place. Mechanisms such as dynamic recovery or/and grain boundary sliding were suggested to be responsible for such uncommon texture evolution.
A number of studies focused on behavior of pure Nb when subjected to HPT [65–70]. Popova et al. [68] investigated the degree of deformation on the structure and thermal stability of pure Nb subjected to HPT. It was reported that Nb grain size (~100 nm) reached to steady state after 5 revolutions. Increased straining had not resulted in further grain refinement. These deformed
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samples were found to be thermally stable up to temperatures of 600oC, and showed intense recrystallization at 700oC. It was reported that further deformation results in earlier thermal
destabilization in deformed Nb, after 10 anvil rotations resulting in microstructures being destabilized at 500oC. Moreover, at steady state, the hardness of nanocrystalline Nb after HPT deformation was found to be similar to high-strength Cu/Nb composite materials. These findings on microstructural evolution of pure Nb subjected to HPT were supported by other studies [67, 69–71]. After application of HPT at cryogenic temperatures (80 K), Nb crystallite sizes as low as 75 nm and hardness of ~4800 MPa were reported [69]. Although the microstructures obtained at cryogenic temperatures were observed to be stable at room temperature, their thermal stability at high temperatures was found to be lower than that for Nb deformed at room temperature via HPT.
2.4. REFERENCES
[1] National School of the Condensed Matter Group, Fundamental properties of nanostructured
materials: [proceedings of the] National School of the Condensed Matter Group, Rimini, Italy, September 20-25, 1993. Singapore ; River Edge, NJ: World Scientific, 1994.
[2] Y. T. Zhu, T. C. Lowe, and T. G. Langdon, “Performance and applications of nanostructured materials produced by severe plastic deformation,” Scr. Mater., vol. 51, no. 8, pp. 825–830, 2004. [3] H. Gleiter, “Nanocrystalline materials,” Prog. Mater. Sci., vol. 33, no. 4, pp. 223–315, 1989. [4] C.C. Koch, “Processing-structure-property relationships in ultrafine grain and nanocrystalline materials,” J. Phys. Conf. Ser., vol. 144, p. 012081, 2009.
[5] C. Suryanarayana, Ed., Non-equilibrium processing of materials, 1st ed. Amsterdam ; New York: Pergamon, 1999.
[6] C. Suryanarayana, “Mechanical alloying and milling,” Prog. Mater. Sci., vol. 46, no. 1–2, pp. 1–184, 2001.
29
tool for grain refinement,” Prog. Mater. Sci., vol. 51, no. 7, pp. 881–981, 2006.
[8] A. M. Russell, L. S. Chumbley, and Y. Tian, “Deformation Processed Metal-Metal Composites,” Adv. Eng. Mater., vol. 2, no. 1–2, pp. 11–22, 2000.
[9] Y. Saito, H. Utsunomiya, N. Tsuji, and T. Sakai, Acta Mat., 47, 2, pp. 579–583, 1997
[10] A. Zhilyaev and T. Langdon, “Using high-pressure torsion for metal processing: Fundamentals and applications,” Prog. Mater. Sci., vol. 53, no. 6, pp. 893–979, 2008.
[11] R. Valiev, “Nanostructuring of metallic materials by SPD processing for advanced properties,” Int. J. Mater. Res., vol. 100, no. 6, pp. 757–761, 2009.
[12] S. Mizunuma, “Large Straining Behavior and Microstructure Refinement of Several Metals by Torsion Extrusion Process,” Mater. Sci. Forum, vol. 503–504, pp. 185–192, 2006.
[13] J. Cave and J. Williams, “Mechanism of cold pressure welding by rolling,” J. Inst. Met., vol. 101, no. JUL-A, pp. 203–207, 1973.
[14] P. W. Bridgman, “On Torsion Combined with Compression,” J. Appl. Phys., vol. 14, no. 6, p. 273, 1943.
[15] A. Zhilyaev, G. Nurislamova, B.-K. Kim, M. Baró, J. Szpunar, and T. Langdon, “Experimental parameters influencing grain refinement and microstructural evolution during high- pressure torsion,” Acta Mater., vol. 51, no. 3, pp. 753–765, 2003.
[16] R. Z. Valiev, R. K. Islamgaliev, and I. V. Alexandrov, “Bulk nanostructured materials from severe plastic deformation,” Prog. Mater. Sci., no. 45, pp. 103–189, 2000.
[17] A. Vorhauer and R. Pippan, “On the homogeneity of deformation by high pressure torsion,”
Scr. Mater., vol. 51, no. 9, pp. 921–925, 2004.
[18] A. Hohenwarter, A. Bachmaier, B. Gludovatz, S. Scheriau, and R. Pippan, “Technical parameters affecting grain refinement by high pressure torsion,” Int. J. Mater. Res., vol. 100, no. 12, pp. 1653–1661, 2009.
[19] G. Sakai, K. Nakamura, Z. Horita, and T. G. Langdon, “Developing high-pressure torsion for use with bulk samples,” Mater. Sci. Eng. A, vol. 406, no. 1–2, pp. 268–273, 2005.
[20] R. B. Figueiredo and T. G. Langdon, “Development of structural heterogeneities in a magnesium alloy processed by high-pressure torsion,” Mater. Sci. Eng. A, vol. 528, no. 13–14, pp. 4500–4506, 2011.
[21] R. B. Figueiredo, M. T. P. Aguilar, P. R. Cetlin, and T. G. Langdon, “Analysis of plastic flow during high-pressure torsion,” J. Mater. Sci., vol. 47, no. 22, pp. 7807–7814, 2012.
30
[22] R. Pippan, S. Scheriau, A. Hohenwarter, and M. Hafok, “Advantages and Limitations of HPT: A Review,” Mater. Sci. Forum, vol. 584–586, pp. 16–21, 2008.
[23] G. Martin, “Phase stability under irradiation: Ballistic effects,” Phys. Rev. B, vol. 30, no. 3, pp. 1424–1436, 1984.
[24] T. Klassen, U. Herr, and R. S. Averback, “Ball Milling of Systems with Positive Heat of Mixing: Effect of Temperature in Ag-Cu”, Acta Mater., vol. 45, no. 7, pp. 2921–2930, 1997. [25] X. Quelennec, A. Menand, J. M. Le Breton, R. Pippan, and X. Sauvage, “Homogeneous Cu– Fe supersaturated solid solutions prepared by severe plastic deformation,” Philos. Mag., vol. 90, no. 9, pp. 1179–1195, 2010.
[26] M. Pouryazdan, D. Schwen, D. Wang, T. Scherer, H. Hahn, R. S. Averback, and P. Bellon, “Forced chemical mixing of immiscible Ag-Cu heterointerfaces using high-pressure torsion,”
Phys. Rev. B, vol. 86, no. 14, 2012.
[27] S. N. Arshad, T. G. Lach, M. Pouryazdan, H. Hahn, P. Bellon, S. J. Dillon, and R. S. Averback, “Dependence of shear-induced mixing on length scale,” Scr. Mater., vol. 68, no. 3–4, pp. 215–218, 2013.
[28] P. Bellon and R. Averback, “Nonequilibrium Roughening of Interfaces in Crystals under Shear: Application to Ball Milling,” Phys. Rev. Lett., vol. 74, no. 10, pp. 1819–1822, 1995. [29] S. Zghal, R. Twesten, F. Wu, and P. Bellon, “Electron microscopy nanoscale characterization of ball milled Cu-Ag powders. Part II: Nanocomposites synthesized by elevated temperature milling or annealing,” Acta Mater., vol. 50, no. 19, pp. 4711–4726, 2002.
[30] S. Zghal, M. J. Hÿtch, J.-P. Chevalier, R. Twesten, F. Wu, and P. Bellon, “Electron microscopy nanoscale characterization of ball-milled Cu–Ag powders. Part I: Solid solution synthesized by cryo-milling,” Acta Mater., vol. 50, no. 19, pp. 4695–4709, 2002.
[31] N. Q. Vo, J. Zhou, Y. Ashkenazy, D. Schwen, R. S. Averback, and P. Bellon, “Atomic Mixing in Metals Under Shear Deformation,” JOM, vol. 65, no. 3, pp. 382–389, 2013.
[32] E. Botcharova, J. Freudenberger, and L. Schultz, “Mechanical alloying of copper with niobium and molybdenum,” J. Mater. Sci., vol. 39, no. 16/17, pp. 5287–5290, 2004.
[33] J. Xu, U. Herr, T. Klassen, and R. S. Averback, “Formation of supersaturated solid solutions in the immiscible Ni–Ag system by mechanical alloying,” J. Appl. Phys., vol. 79, no. 8, p. 3935, 1996.
31
of Nb precipitates in Cu during severe plastic deformation,” Acta Mater., vol. 62, pp. 276–285, 2014.
[35] J. Xu, J. H. He, and E. Ma, “Effect of milling temperature on mechanical alloying in the immiscible Cu-Ta system,” Metall. Mater. Trans. A, vol. 28, no. 7, pp. 1569–1580, 1997.
[36] M. J. Demkowicz, J. Wang, and R. G. Hoagland, “Chapter 83 Interfaces Between Dissimilar Crystalline Solids,” in Dislocations in Solids, vol. 14, Elsevier, 2008, pp. 141–205.
[37] S. Odunuga, Y. Li, P. Krasnochtchekov, P. Bellon, and R. Averback, “Forced Chemical Mixing in Alloys Driven by Plastic Deformation,” Phys. Rev. Lett., vol. 95, no. 4, 2005.
[38] F. Delogu, “Forced chemical mixing in model immiscible systems under plastic deformation,”
J. Appl. Phys., vol. 104, no. 7, p. 073533, 2008.
[39] A. Lund and C. Schuh, “Driven Alloys in the Athermal Limit,” Phys. Rev. Lett., vol. 91, no. 23, 2003.
[40] N. Q. Vo, R. S. Averback, Y. Ashkenazy, P. Bellon, and J. Wang, “Forced chemical mixing at Cu-Nb interfaces under severe plastic deformation,” J. Mater. Res., vol. 27, no. 12, pp. 1621– 1630, 2012.
[41] Y. Ashkenazy, N. Q. Vo, D. Schwen, R. S. Averback, and P. Bellon, “Shear induced chemical mixing in heterogeneous systems,” Acta Mater., vol. 60, no. 3, pp. 984–993, 2012.
[42] D. G. Morris and M. A. Morris, “Mechanical alloying of copper-bcc element mixtures,” Scr.
Metall. Mater., vol. 24, no. 9, pp. 1701–1706, 1990.
[43] E. Botcharova, M. Heilmaier, J. Freudenberger, G. Drew, D. Kudashow, U. Martin, and L. Schultz, “Supersaturated solid solution of niobium in copper by mechanical alloying,” J. Alloys
Compd., vol. 351, no. 1–2, pp. 119–125, 2003.
[44] L. G. Fritzemeier, “High strength, high conductivity composites,” Nanostructured Mater., vol. 1, no. 3, pp. 257–262, 1992.
[45] A. Misra, R. G. Hoagland, and H. Kung, “Thermal stability of self-supported nanolayered Cu/Nb films,” Philos. Mag., vol. 84, no. 10, pp. 1021–1028, 2004.
[46] T. Höchbauer, A. Misra, K. Hattar, and R. G. Hoagland, “Influence of interfaces on the storage of ion-implanted He in multilayered metallic composites,” J. Appl. Phys., vol. 98, no. 12, p. 123516, 2005.
[47] M. J. Demkowicz, Y. Q. Wang, R. G. Hoagland, and O. Anderoglu, “Mechanisms of He escape during implantation in Cu Nb multilayer composites,” Nucl. Instrum. Methods Phys. Res.
32
Sect. B Beam Interact. Mater. At., vol. 261, no. 1–2, pp. 524–528, 2007.
[48] G. V. Kurdjumov and G. Sachs, Z Phys, vol. 164, p. 325, 1930.
[49] M. J. Demkowicz, D. Bhattacharyya, I. Usov, Y. Q. Wang, M. Nastasi, and A. Misra, “The Effect of Excess Atomic Volume on He Bubble Formation at fcc–bcc interfaces,” Appl. Phys. Lett., vol. 97, no. 16, p. 161903, 2010.
[50] X. Zhang, N. Li, O. Anderoglu, H. Wang, J. G. Swadener, T. Höchbauer, A. Misra, and R. G. Hoagland, “Nanostructured Cu/Nb multilayers subjected to helium ion-irradiation,” Nucl. Instrum.
Methods Phys. Res. Sect. B Beam Interact. Mater. At., vol. 261, no. 1–2, pp. 1129–1132, 2007.
[51] M. Demkowicz, R. Hoagland, and J. Hirth, “Interface Structure and Radiation Damage Resistance in Cu-Nb Multilayer Nanocomposites,” Phys. Rev. Lett., vol. 100, no. 13, 2008. [52] E. Snoeck, F. Lecouturier, L. Thilly, M. . Casanove, H. Rakoto, G. Coffe, S. Askénazy, J. . Peyrade, C. Roucau, V. Pantsyrny, A. Shikov, and A. Nikulin, “Microstructural Studies of in Situ Produced Filamentary Cu/Nb Wires,” Scr. Mater., vol. 38, no. 11, pp. 1643–1648, 1998.
[53] L. Thilly, M. Véron, O. Ludwig, and F. Lecouturier, “Deformation mechanism in high strength Cu/Nb nanocomposites,” Mater. Sci. Eng. A, vol. 309–310, pp. 510–513, 2001.
[54] F. Dupouy, E. Snoeck, M. J. Casanove, C. Roucau, J. P. Peyrade, and S. Askenazy, “Microstructural characterization of high strength and high conductivity nanocomposite wires,”
Scr. Mater., vol. 34, no. 7, pp. 1067–1073, 1996.
[55] I. J. Beyerlein, N. A. Mara, J. S. Carpenter, T. Nizolek, W. M. Mook, T. A. Wynn, R. J. McCabe, J. R. Mayeur, K. Kang, S. Zheng, J. Wang, and T. M. Pollock, “Interface-driven microstructure development and ultra-high strength of bulk nanostructured Cu-Nb multilayers fabricated by severe plastic deformation,” J. Mater. Res., vol. 28, no. 13, pp. 1799–1812, 2013. [56] J. S. Carpenter, S. C. Vogel, J. E. LeDonne, D. L. Hammon, I. J. Beyerlein, and N. A. Mara, “Bulk texture evolution of Cu–Nb nanolamellar composites during accumulative roll bonding,”
Acta Mater., vol. 60, no. 4, pp. 1576–1586, 2012.
[57] M. J. Demkowicz and L. Thilly, “Structure, shear resistance and interaction with point defects of interfaces in Cu–Nb nanocomposites synthesized by severe plastic deformation,” Acta Mater., vol. 59, no. 20, pp. 7744–7756, 2011.
[58] K. Kang, J. Wang, and I. J. Beyerlein, “Atomic structure variations of mechanically stable fcc-bcc interfaces,” J. Appl. Phys., vol. 111, no. 5, p. 053531, 2012.
33
“Thermal stability of Cu–Nb nanolamellar composites fabricated via accumulative roll bonding,”
Philos. Mag., vol. 93, no. 7, pp. 718–735, 2013.
[60] E. H. Ekiz, T. G. Lach, R. S. Averback, N. A. Mara, I. J. Beyerlein, M. Pouryazdan, H. Hahn, and P. Bellon, “Microstructural evolution of nanolayered Cu–Nb composites subjected to high- pressure torsion,” Acta Mater., vol. 72, pp. 178–191, 2014.
[61] T. Hebesberger, H. P. Stüwe, A. Vorhauer, F. Wetscher, and R. Pippan, “Structure of Cu deformed by high pressure torsion,” Acta Mater., vol. 53, no. 2, pp. 393–402, 2005.
[62] K. Edalati, T. Fujioka, and Z. Horita, “Microstructure and mechanical properties of pure Cu processed by high-pressure torsion,” Mater. Sci. Eng. A, vol. 497, no. 1–2, pp. 168–173, 2008. [63] N. A. Enikeev, E. Schafler, M. J. Zehetbauer, I. V. Alexandrov, and R. Z. Valiev, “Observations of Texture in Large Scale HPT-Processed Cu,” Mater. Sci. Forum, vol. 584–586, pp. 367–374, 2008.
[64] J. Gubicza, S. V. Dobatkin, E. Khosravi, A. A. Kuznetsov, and J. L. Lábár, “Microstructural stability of Cu processed by different routes of severe plastic deformation,” Mater. Sci. Eng. A, vol. 528, no. 3, pp. 1828–1832, 2011.
[65] J. Alkorta, C. J. Luis Pérez, E. N. Popova, M. Hafok, R. Pippan, and J. Gil Sevillano, “Microstructure and Indentation Size-Effect in Pure Niobium Subjected to SPD via ECAP and HPT,” Mater. Sci. Forum, vol. 584–586, pp. 215–220, 2008.
[66] S. Lee and Z. Horita, “High-Pressure Torsion for Pure Chromium and Niobium,” Mater.
Trans., vol. 53, no. 1, pp. 38–45, 2012.
[67] T. Nishizaki, S. Lee, Z. Horita, T. Sasaki, and N. Kobayashi, “Superconducting properties in bulk nanostructured niobium prepared by high-pressure torsion,” Phys. C Supercond., vol. 493, pp. 132–135, 2013.
[68] E. N. Popova, V. V. Popov, E. P. Romanov, and V. P. Pilyugin, “Thermal stability of nanocrystalline Nb produced by severe plastic deformation,” Phys. Met. Metallogr., vol. 101, no. 1, pp. 52–57, 2006.
[69] V. V. Popov, E. N. Popova, A. V. Stolbovskiy, and V. P. Pilyugin, “Thermal stability of nanocrystalline structure in niobium processed by high pressure torsion at cryogenic temperatures,” Mater. Sci. Eng. A, vol. 528, no. 3, pp. 1491–1496, 2011.
[70] V. V. Popov, E. N. Popova, A. V. Stolbovskiy, V. P. Pilyugin, and N. K. Arkhipova, “Nanostructurization of Nb by high-pressure torsion in liquid nitrogen and the thermal stability of
34
the structure obtained,” Phys. Met. Metall., vol. 113, no. 3, pp. 295–301, 2012.
[71] B. Jóni, E. Schafler, M. Zehetbauer, G. Tichy, and T. Ungár, “Correlation between the microstructure studied by X-ray line profile analysis and the strength of high-pressure-torsion processed Nb and Ta,” Acta Mater., vol. 61, no. 2, pp. 632–642, 2013.
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