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Nb nanocomposite wires
International Journal of Fatigue 128 (2019) 105188
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International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue
Low-cycle bending fatigue and electrical conductivity of high-strength Cu/ Nb nanocomposite wires
T
A.B. Rozhnova, V.I. Pantsyrnyb, A.V. Krayneva, S.O. Rogacheva, , S.A. Nikulina, N.E. Khlebovab, M.V. Polikarpovac, M.Yu. Zadorozhnyya ⁎
a
National University of Science and Technology “MISIS”, 4 Leninsky pr., 119049 Moscow, Russia Research and Production Company “NANOÉLECTRO”, 5a Rogova St., 123060 Moscow, Russia c A.A. Bochvar High-Technology Scientific Research Institute for Inorganic Materials, 5a Rogova St., 123060 Moscow, Russia b
ARTICLE INFO
ABSTRACT
Keywords: Metal-matrix composites (MMCs) Low-cycle fatigue Dynamic mechanical analyzer Fractography Electrical conductivity
The comparative low-cycle fatigue tests of the Cu/Nb nanocomposite wires obtained by the “melt-and-deform” method and pure copper samples were carried out using a dynamic mechanical analyzer (DMA). The fatigue tests were carried out using transverse bending scheme. It has been established that the fatigue fracture resistance characteristics of the Cu/Nb composites significantly higher as compared to that of pure copper samples. It has been found that the fatigue crack propagation in the Cu/Nb composites occurs at two scale levels. The effect of fatigue damage accumulation on the change in the electrical resistance of the Cu/Nb composites was studied.
1. Introduction Binary in situ nanocomposite wires based on a metallic copper matrix with the addition of bcc metals, namely, Cu/Nb, Cu/Fe, Cu/Cr, Cu/ V, are the subject of intensive research in recent decades [1–9]. The strenghtening phases formed during the deformation process have high thermodynamic stability and good adhesion to the matrix. The limited solubility of bcc metals in copper allows the copper matrix to maintain the high electrical conductivity, while bcc metal filaments provide high strength properties [10–12]. The unique combination of high strength and high electrical conductivity, due to the transition of the structure to the nanoscale state, distinguishes these composites into a separate class of materials and allows one to use them as promising materials for the manufacture of mechanically loaded parts of electrical devices. Since the mid-1980s, the Cu/Nb nanocomposite conductors have traditionally been used in the manufacture of windings of high-field pulsed magnets [13]. The manufacturing process of the Cu/Nb composite conductors with large cross sections and high mechanical properties is characterized by some technological features that have much in common with the manufacturing process of low-temperature superconductors by the RRP method (a method of a multi-rod set stacking, its subsequent evacuating and extrusion). As a result of the manufacturing process, there obtained a complex composite with a multi-scale structure. Despite the fact that common mechanisms of the structure
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formation of in-situ composites, as well as their mechanical properties under static loading, are well understood, literature data on mechanical and functional properties of the composites under cyclic loading are scarce [14–16]. Unconventional fatigue test techniques, in particular, performed by a dynamic mechanical analyzer (DMA) can be applied to evaluate the fatigue resistance of small-scale samples with a diameter of less than 1 mm, such as in situ composite wires. In case of fatigue test of small cross-section samples, the DMA installation shows the advantage in the accuracy of the applied load compared to the standard testing machines. Previously, DMA was successfully used for comparative bending fatigue tests of orthodontic wires (NiTi and Cu-NiTi) [17], thin samples of zirconium alloys [18,19] and Cu-Fe and Cu-V nanocomposite wires [15]. In this work, the comparative low-cycle bending fatigue tests of the Cu/Nb nanocomposite wires obtained by the “melt-and-deform” method, as well as of the pure copper samples were carried out using DMA. The change in electrical conductivity of the Cu/Nb nanocomposite wires due the accumulation of fatigue damage was also studied. 2. Material and research methods The samples of a nanocomposite wire with a diameter of 0.46 mm, obtained by the “melt-and-deform” method (the method has been described in more detail in [20]), were investigated. The basis of the
Corresponding author. E-mail address: [email protected] (S.O. Rogachev).
https://doi.org/10.1016/j.ijfatigue.2019.105188 Received 1 April 2019; Received in revised form 1 July 2019; Accepted 4 July 2019 Available online 05 July 2019 0142-1123/ © 2019 Elsevier Ltd. All rights reserved.
International Journal of Fatigue 128 (2019) 105188
A.B. Rozhnov, et al.
Fig. 1. The structure of the Cu/Nb composite in cross section (SEM): (a) general view; (b) fragment of the structure. Table 1 The characteristics of the studied materials. Sample
Diameter, mm
Electrical conductivity, % IACS
Tensile strength (σUTS), MPa
Bending strength, MPa
Cu/Nb Cu
0.463 ± 0.005 0.50 ± 0.01
68.8 ± 0.1 95.4 ± 0.1
1125 ± 10 415 ± 10
1660 ± 20 790 ± 20
composite was a Cu-18% wt.Nb alloy, produced by a consumableelectrode vacuum arc melting. The core of the Cu-18% wt.Nb alloy was placed in a copper shell, then deformed by extrusion and drawing to form a hexagonal rod. The rods were cut to specified lengths, than stacked and assembled with copper shell. Thus, the final composite was fabricated using a double assembly and consisted of hexagonal Cu/Nb sub-elements surrounded by copper layers (Fig. 1). The total degree of cold drawn deformation of the final assembly was ln(A0/A) ∼ 8. For comparison, the wire samples of pure copper with a diameter of 0.50 mm, obtained by cold drawn deformation, were used. The characteristics of the studied materials, their strength and electrically conductive characteristics are presented in Table 1. Low-cycle fatigue tests of the Cu/Nb composite samples and pure copper with length of 30 mm were carried out using DMA Q800 (TA Instrument) according to the previously described method [18]. The tests were carried out using transverse bending scheme in one plane at a constant stress level with an alternating symmetric loading cycle (stress ratio R = −1) using a single cantilever grip of DMA Q800, until the failure of the samples (Fig. 2). During the grip movement in the course of tests, a sample was symmetrically deviated from the initial position to a certain amplitude. A soft loading scheme (at a constant cycle stress amplitude) with a frequency of loading cycles of 10 Hz at a temperature of 30 °C was used. At the end of tests, the number of load cycles to failure (N) was registered (based on 2.5 × 104 cycles). The tests were
carried out at a cycle stress of 500–1330 MPa (for Cu/Nb composite) and 230–710 MPa (for copper), which accounted for 30–80% of the fracture stress of the sample in tests with a stress amplitude increment (bending strength, see Table 1). At least three samples were used for each cycle stress amplitude (σa). The microstructure analysis of the samples in cross section and the analysis of the fatigue fractures was performed using a JSM-6610LV scanning electron microscope (JEOL). Electron microscopic studies of the structure of the Cu/Nb composite samples before and after cyclic tests were performed using a JEM2100 transmission electron microscope (TEM). The samples for electron microscopic studies were cut from hexagonal Cu/Nb sub-elements of the composite in transverse (before fatigue testing) and longitudinal (before and after fatigue testing) sections. The sample preparation was performed by the focused ion beam method using a Strata 201 SIMSmapIII×P scanning ion microscope with a gallium liquid metal ion gun, as described in [21]. 10 images at the same magnification were taken for the microstructure quantitative analysis, that was carried out using the Image Expert Pro software. The measurements of an electrical conductivity (electrical resistivity) were carried out using the Cu/Nb composite samples, which were cyclically loaded at 25; 50 and 75% of life time (where life time is a total number of load cycles to failure) at each cycle stress amplitude. The number of load cycles to failure (N) was determined from the
Fig. 2. The sample in a single cantilever grip of DMA Q800: (a) image of the sample and grip, (b) the sample deviation scheme at fatigue testing. 2
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Fig. 3. TEM photographs of the Cu/Nb composite microstructures before (a, c) and after (b, d) fatigue tests at a maximum cycle stress: (a–c) cross section, (d) longitudinal section.
fatigue curve obtained from fatigue tests. The electrical resistance was determined by the four-probe method using a test bench for measuring the RRR and Tc of superconductors (in accordance with IEC 61788-4: 2016 and IEC 61788-10: 2006). The measurements were carried out at temperatures of 293 K, 77.4 K and 10 K.
Table 2 The characteristics of the Cu/Nb composite microstructure (in the cross section of the wire). The thickness of the Nb-filaments, nm The width of the Nb-filaments, nm The density of the Nb-filaments, pcs/µm2 Area in the cross section, occupied with the Nb-filaments, % Cross-sectional size of copper grains (subgrains), nm
3. Results and discussion 3.1. The analysis of the microstructure
13.6 ± 3.9 81 ± 12 140.0 ± 11.0 23.4 ± 3.1 100 ± 15
obtained from the longitudinal sections of the composite, no notable changes associated with the action of the alternating deformation during cyclic tests were found (Fig. 3c, d). It should be noted that the Cu/Nb samples were obtained by drawing with a very large degree of strain without subsequent annealing. Thus, before the fatigue tests, the Cu/Nb samples are already contain a very high density of dislocation and other lattice defects. Therefore, the dislocation density during fatigue tests could not increase very significantly. The TEM images clearly show that the structure of the Cu/Nb samples contains a very large number of defects both before and after the fatigue tests. However, it is difficult to quantify the change in the density of defects.
Fig. 3 shows TEM images of the microstructures (with electron diffraction patterns) of a single hexagonal sub-element of the Cu/Nb composite in transverse and longitudinal section in original and tested states (at a maximum cycle stress −1330 MPa). A grain-subgrain nanostructure of copper matrix is observed. The “twisted” Nb-filaments are surrounded by copper grains (subgrains) with an average crosssection of 100 nm. The twisting of the filaments in such composites is usually associated with the appearance of a rotational mode in the process of drawing [22]. In our opinion, the twisting process can also be associated with the inhomogeneity of the plastic flow in the material during large plastic deformations due drawing, which leads to the appearance of a turbulent plastic flow and the appearance of the “vortex” defects [23,24]. In the longitudinal section, the Nb-filaments have an elongated shape. The results of the quantitative metallographic analysis of the Cu/Nb composite microstructure, carried out using the crosssectional images, are presented in Table 2. When comparing the microstructure of the Cu/Nb composite in original and tested states, as well as the electron diffraction patterns
3.2. Fatigue strength and fractographic research Fatigue curves (N − σa) of the Cu/Nb composite samples, as well as of the pure copper samples are shown in Fig. 4. The fatigue limit (σRN) of the Cu/Nb composite samples at N = 2.5 × 104 cycles was 490 MPa. The fatigue limit of pure copper samples is much lower, σRN = 390 MPa. The maximum difference in the fatigue strength 3
International Journal of Fatigue 128 (2019) 105188
A.B. Rozhnov, et al.
Fig. 4. The fatigue curves of the studied samples.
striations (Fig. 6c). Elongated dimples were observed in the area of static fracture. The static fracture area is shifted from the center of the fracture as in the case of the Cu/Nb composite. At large cycle stresses, the macro- and microstructure of the fatigue fracture of pure copper samples do not qualitatively change, but more intense secondary cracking is observed along the boundaries of the striations, and the distance between the micro-striations is larger. There are macro-striations at the fracture, the distance between which is 2–13 μm. The static fracture area is shifted to the center of the fracture. Thus, the fatigue crack propagation in the Cu/Nb composite samples, unlike pure copper samples, occurs at two scale levels: within individual Cu/Nb sub-elements with the formation of micro-striations and within the entire volume of the composite with the formation of macro-striations. This fact, obviously, contributes to the increased resistance to fatigue failure of the Cu/Nb composite samples, especially at high cycle stresses. In addition, the Nb-filaments also appear to be additional barriers to the fatigue crack propagation, which also explains the higher resistance to fatigue failure of the Cu/Nb composite compared to pure copper. A possible mechanism of the influence of the filaments in the in-situ composites was previously considered using the Cu/Cr composite [14]. The cracking along the boundaries of sub-elements in the Cu/Nb composite sample under fatigue loading may be due to the fact that the copper interlayers along the boundaries of sub-elements are in a highly deformed state and lose their ductility. This is also confirmed by the rather brittle mechanism of fatigue fracture of pure copper, the degree of strain of which is close to the degree of strain of the copper interlayers in the Cu/Nb composite sample. In [16], when studying the fatigue failure of the Cu-18% (vol.) Nb composite with a Nb-filament thickness of 5–10 nm, it was found that the copper interlayers between the sub-elements are the place of the fatigue crack nucleation, as well as the place of its inhibition. However, it should be noted that the authors [16] removed the copper shell of the composite before fatigue tests, and also carried out high-cycle fatigue tests (i.e., fatigue defects were nucleated mainly in the region of elastic stresses) and used another loading scheme (cyclic tensile), which does not allow to compare the results.
between the Cu/Nb composite samples and pure copper samples was observed at a high stress level of load cycles. The nucleation of fatigue cracks and further fracture of all samples during tests on the DMA were initiated on the sample surface at a distance from the ends of the grips of 1–3 mm. The fractographic analysis of the Cu/Nb composite samples as well as of the pure copper samples revealed similar macrostructure of fatigue fractures, but different mechanisms of fatigue crack propagation (Figs. 5a, 6a). As usual for bending fatigue fracture, two areas of fatigue crack nucleation are seen on the sample surface near the opposite edges of the sample (sometimes only one area of fatigue crack nucleation was observed), then the crack propagated to the center of the sample, where the static fracture occurred. At low cycle stresses, typical fatigue micro-striations were observed on the surface of hexagonal sub-elements of the Cu/Nb composite samples in the area of stable growth of fatigue crack. The orientation of the micro-striations changed from one filament to another (Fig. 5e), which is associated with the branching of fatigue crack under the influence of local deformation conditions [25]. The distance between these striations was about 0.3 μm. Also, in the area of stable growth of fatigue crack, intense cracking was observed along the boundaries of the group of the Cu/Nb sub-elements to the normal direction of fatigue crack growth. In the area of accelerated crack growth, the macrostriations were observed perpendicular to the direction of crack growth, the distance between which was ∼1.5 μm (Fig. 5g). Such striations did not change their direction. The static fracture area is characterized by a fully ductility small-dimple structure (Fig. 5g). At the same time, the static fracture area is shifted from the center of the fracture; this is typical of bending tests of the samples with a low cycle stress and small stress concentrators [25]. At high cycle stresses, the macro- and microstructure of the fatigue fracture of the Cu/Nb composite samples do not qualitatively change, however, more intense cracking was observed along the boundaries of the group of the Cu/Nb sub-elements with an increased distance between micro-striations (Fig. 5d, f). The static fracture area located in the center of the fracture is substantially larger, which is associated with a shorter period of fatigue fracture itself (Fig. 5h). In fatigue fractures of pure copper samples at low cycle stresses, the fatigue striations are observed in the area of the fatigue crack propagation with a distance between them of ∼1.6 μm (Fig. 6e). In local areas, the secondary cracking was observed along the boundaries of the
3.3. The measurements of electrical resistivity Fig. 7 shows the graphs of the change in the electrical resistivity of 4
International Journal of Fatigue 128 (2019) 105188
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Fig. 5. The surface of fatigue fractures of the Cu/Nb composite samples: at low (a, c, e, g) and high (b, d, f, h) cycle stress; 1 – fatigue crack nucleation area; 2 – stable growth of fatigue crack area; 3 – accelerated crack growth area; 4 – static fracture area (arrows in e indicate the direction of the fatigue crack growth).
the Cu/Nb composite samples under cyclic loading depending on the cycle stress amplitude (1330; 1000; 830; 665 and 500 MPa), the number of load cycles (25; 50 and 75% of life time) and temperature of the measurement (293; 77.4 and 10 K). To more clearly reveal the regularities of electrical resistivity, the data presented in Fig. 7 is rearranged in relative units ρi/ρ0, where ρ0 is the electrical resistance before cyclic loading (see Fig. 8). It is seen that, the electrical resistivity of the Cu/Nb composite samples decreases sharply with decreasing temperature, which is
associated with an increase in the mean free path of electron (λ). At the number of cycles equal to 25% of life time, an increase in the electrical resistivity is observed by 7% and 14% at a temperature of 293 K and 77.4 K, respectively. With further increase in the number of cycles a significant increase in electrical resistivity does not occur. This character is traced for all values of the cycle amplitude. At a temperature of 10 K, no apparent dependence of the electrical resistivity is observed either on the number of load cycles or on the stress cycle amplitude. This is due to a change in the nature of the scattering of conduction
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International Journal of Fatigue 128 (2019) 105188
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Fig. 6. The surface of fatigue fractures of pure copper samples: at low (a, c, e) and high (b, d, f) cycle stress; 1 – fatigue crack nucleation area; 2 – stable growth of fatigue crack area; 3 – accelerated crack growth area; 4 – static fracture area.
electrons at the Cu/Nb interface. The same Cu/Nb interface, depending on the value of λ, can be diffusely reflective and mirror reflective, that varies scattering coefficient p from 0 to 1, and, accordingly, affects the increase in electrical resistivity [26–28]. The scattering is considered to be diffuse if the irregularities of the interface have a value comparable with λ. At a temperature of 10 K, when λ can reach values of the order of 1 μm, which far exceeds the parameters of the roughness of the interface, the reflection of conduction electrons at the Cu/Nb interface is close to mirror, and the changes in electrical resistivity with increasing deformation are weakly expressed [26]. In the case of cyclic loading, the change in electrical conductivity (at 25% of life time) is obviously associated with the formation of fatigue defects: the continuity of the material is disturbed, shears are formed in individual grains, which then develop into microcracks. As can be seen, the fatigue failure is initiated on the sample surface, i.e. in the area of severe deformed copper shell. Thus, in the first place, the continuity of
the copper shell is broken. Considering that the electric current mainly flows on the surface of the conductor, the discontinuity of the copper shell leads to a decrease in electrical conductivity. Further accumulation of fatigue damage occurs already in the internal volume of the Cu/ Nb composite wire, which has less impact on the change in electrical conductivity. It should be noted that not only the size of defects, but also the heterogeneity of their location and morphology can influence the change in electrical conductivity. Thus, the absence of an explicit nature of the change in electrical resistivity, measured at 10 K, may be due to the fact that at a given temperature the contribution from electron scattering on fatigue defects to electrical resistivity is significantly less than the contribution from scattering at the Cu/Nb interfaces. You should also take into account the variation of the results characteristic of fatigue tests.
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International Journal of Fatigue 128 (2019) 105188
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Fig. 7. The electrical resistivity of the Cu/Nb composite samples vs the cycle stress amplitude and number of loading cycles: (a) at 293 K, (b) at 77.4 K, (c) at 10 K.
Fig. 8. The normalized electrical resistivity (ρi/ρ0) of the Cu/Nb composite samples vs the cycle stress amplitude and number of loading cycles: (a) at 293 K, (b) at 77.4 K, (c) at 10 K.
4. Conclusion
(3) Fatigue crack propagation in the Cu/Nb composite samples occurs at two large-scale levels: within individual Cu/Nb sub-elements with the formation of micro-striations and within the entire volume of the composite with the formation of macro-striations. (4) The accumulation of fatigue damage leads to a slight decrease in the electrical conductivity of the Cu/Nb composite samples.
(1) Nanocomposite Cu/Nb wires obtains by the “melt-and-deform” method have a complex hierarchically non-uniform structure: the hexagon Cu/Nb sub-elements surrounded by copper interlayers. The Cu/Nb sub-element is characterized by grain-subgrain nanostructure with second phase precipitates (niobium) having a “twisted” shape in cross section and elongated shape in longitudinal section, located mainly along the boundaries of copper grains (subgrains) with an average cross section of 100 nm. (2) It has been established that the fatigue fracture resistance characteristics of the Cu/Nb nanocomposite wires are significantly higher as compared to that of pure copper. The fatigue limit (σRN) at N = 2.5 × 104 cycles for the samples of the Cu/Nb composite and pure copper is 490 and 390 MPa, respectively.
Acknowledgment The work was carried out with financial support from the Ministry of Science and Higher Education of the Russian Federation in the framework of Increase Competitiveness Program of NUST «MISiS» (№ К22019-008), implemented by a governmental decree dated 16th of March 2013, N 211. 7
International Journal of Fatigue 128 (2019) 105188
A.B. Rozhnov, et al.
References
2015;70:92–8. https://doi.org/10.1016/j.compositesb.2014.10.046. [16] Nayeb-Hashemi H, Pourrahimi S. Synthesis, strengthening, fatigue and fracture behavior of high-strength, high-conductivity P/M processed Cu-Nb microcomposite. In: Mitchell MR, Buck O, editors. Cyclic deformation, fracture, and nondestructive evaluation of advanced materials: second volume ASTM International; 1994. p. 87–108. https://doi.org/10.1520/STP18121S. [17] Prymak O, Klocke A, Kahl-Nieke B, Epple M. Fatigue of orthodontic nickel–titanium (NiTi) wires in different fluids under constant mechanical stress. Mater Sci Eng A 2004;378:110–4. https://doi.org/10.1016/j.msea.2003.10.332. [18] Nikulin SA, Markelov VA, Gusev AYu, Nechaykina TA, Rozhnov AB, Rogachev SO, Zadorozhnyy MYu. Low-cycle fatigue tests of zirconium alloys using a dynamic mechanical analyzer. Int J Fatigue 2013;48:187–91. https://doi.org/10.1016/j. ijfatigue.2012.10.019. [19] Nikulin SA, Rozhnov AB, Gusev AYu, Nechaykina TA, Rogachev SO, Zadorozhnyy MYu. Fracture resistance of Zr-Nb alloys under low-cycle fatigue tests. J Nucl Mater 2014;446:10–4. https://doi.org/10.1016/j.jnucmat.2013.11.039. [20] Dubois JB, Thilly L, Lecouturier F, Olier P, Renault PO. Cu/Nb nanocomposite wires processed by severe plastic deformation for applications in high pulsed magnets: effects of the multi-scale microstructure on the mechanical properties. IEEE Trans Appl Supercond 2012;22(3):6900104. https://doi.org/10.1109/TASC.2011. 2174574. [21] Nikulin SA, Pantsyrnyi VI, Rozhnov AB, Rogachev SO, Khlebova NE, Nechaikina TA, et al. Possibilities of scanning and transmission electron microscopy for comparative analysis of the microstructure of in-situ nanocomposite high-strength conductors based on a copper matrix and bcc-metals. Metal Sci Heat Treat 2016;58(3–4):209–13. https://doi.org/10.1007/s11041-016-9990-y. [22] Hosford WF. Microstructural changes during deformation of (011) fiber-textured metals. Trans Metall Soc AIME 1964;230(1):12–4. [23] Kulagin R, Beygelzimer Y, Ivanisenko Y, Mazilkin A, Hahn H. High pressure torsion: from laminar flow to turbulence. IOP Conf Series: Mater Sci Eng 2017;194:012045. https://doi.org/10.1088/1757-899X/194/1/012045. [24] Rogachev SO, Sundeev RV, Khatkevich VM. Evolution of the structure and strength of steel/vanadium alloys/steel hybrid material during severe plastic deformation. Mater Lett 2016;173:123–6. https://doi.org/10.1016/j.matlet.2016.03.044. [25] Terent’ev VF. Fatigue strength of metals and alloys. Moscow: Intermet Inzhiniring; 2002. [in Russia]. [26] Guryev VV, Polikarpova MV, Lukyanov PA, Khlebova NE, Pantsyrny VI. Size effects influence on conducting properties of Cu-Nb alloy microcomposites at cryogenic temperature. Cryogenics 2018;90:56–9. https://doi.org/10.1016/j.cryogenics. 2018.01.010. [27] Verhoeven JD, Downing NL, Chumbley LS, Gibson ED. The resistivity and microstructure of heavily drawn Cu-Nb alloys. J Appl Phys 1989;65(3):1293–301. https://doi.org/10.1063/1.343024. [28] Deng L, Han K, Hartwig K, Siegrist T, Dong L, Sun Z, et al. Hardness, electrical resistivity, and modeling of in situ Cu–Nb microcomposites. J Alloys Comp 2014;602:331–8. https://doi.org/10.1016/j.jallcom.2014.03.021.
[1] Liu Z, Chen X, Li W. Mechanical properties of a composite Cu-Cr alloy obtained insitu by directed solidification. Met Sci Heat Treat 2006;48(9):463–5. https://doi. org/10.1007/s11041-006-0118-7. [2] Biselli C, Morris DG. Microstructure and strength of Cu–Fe in situ composites obtained from prealloyed Cu–Fe powders. Acta Metall Mater 1994;42:163–76. https:// doi.org/10.1016/0956-7151(94)90059-0. [3] Snoeck E, Lecouturier F, Thilly L, Casanove MJ, Rakoto H, Coffe G, Askénazy S, Peyrade JP, Roucau C, Pantsyrny V, Shikov A, Nikulin A. Microstructural studies of in situ produced filamentary Cu/Nb wires. Scr Mater 1998;38:1643–8. https://doi. org/10.1016/S1359-6462(98)00080-3. [4] Raabe D, Hangen U. Simulation of the yield strength of wire drawn Cu-based in-situ composites. Comput Mater Sci 1996;5:195–202. https://doi.org/10.1016/09270256(95)00072-0. [5] Liu K, Jiang Z, Zhao J, Zou J, Lu L, Lu D. Thermal stability and properties of deformation-processed Cu-Fe in situ composites. Metall Mater Trans A 2015;46:2255–61. https://doi.org/10.1007/s11661-015-2791-x. [6] Sauvage X, Genevois C, Costaa GDa, Pantsyrny V. Atomic scale characterization of deformation-induced interfacial mixing in a Cu/V nanocomposite wire. Scr Mater 2009;61:660–3. https://doi.org/10.1016/j.scriptamat.2009.06.007. [7] Deng L, Yang X, Han K, Lu Y, Liang M, Liu Q. Microstructure and texture evolution of Cu–Nb composite wires. Mater Char 2013;81:124–33. https://doi.org/10.1016/j. matchar.2013.04.013. [8] Lu X, Yao D, Chen Y, Wang L, Dong A, Meng L, et al. Microstructure and hardness of Cu-12% Fe composite at different drawing strains. J Zhejiang Univ-Sci A 2014;15(2):149–56. https://doi.org/10.1631/jzus.A1300164. [9] Brokmeier HG, Bolmaro RE, Signorelli JA, et al. Texture development of wire drawn Cu–Fe composites. Phys B 2000;276–278:888–9. https://doi.org/10.1016/S09214526(99)01540-9. [10] Raabe D, Heringhaus F, Hangen U, et al. Investigation of a Cu-20mass%Nb in situ composite. Part I: Fabrication, microstructure and mechanical properties. Zeitschrift für Metallkunde 1995;86:405–22. [11] Raabe D, et al. Metallic composites processed via extreme deformation: toward the limits of strength in bulk materials. MRS Bull 2010;35:982–91. https://doi.org/10. 1557/mrs2010.703. [12] Pantsyrnyi V, Shikov A, Vorobieva A, et al. High strength, high conductivity macroand microcomposite winding wires for pulsed magnets. Phys B 2001;294–295:669–73. https://doi.org/10.1016/S0921-4526(00)00741-9. [13] Foner S. 68.4-T-long pulse magnet: test of high strength microcomposite Cu/Nb conductor. Appl Phys Lett 1986;49:982–3. https://doi.org/10.1063/1.97503. [14] Masuda C, Tanaka Y. Fatigue properties of Cu–Cr in situ composite. Int J Fatigue 2006;28:1426–34. https://doi.org/10.1016/j.ijfatigue.2006.02.051. [15] Nikulin SA, Rogachev SO, Rozhnov AB, Pantsyrnyi VI, Khlebova NE, Nechaykina TA, Khatkevich VM, Zadorozhnyy MYu. Microstructure and fatigue strength of high-strength Cu-Fe and Cu-V in-situ nanocomposite wires. Compos B
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International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue
Low-cycle bending fatigue and electrical conductivity of high-strength Cu/ Nb nanocomposite wires
T
A.B. Rozhnova, V.I. Pantsyrnyb, A.V. Krayneva, S.O. Rogacheva, , S.A. Nikulina, N.E. Khlebovab, M.V. Polikarpovac, M.Yu. Zadorozhnyya ⁎
a
National University of Science and Technology “MISIS”, 4 Leninsky pr., 119049 Moscow, Russia Research and Production Company “NANOÉLECTRO”, 5a Rogova St., 123060 Moscow, Russia c A.A. Bochvar High-Technology Scientific Research Institute for Inorganic Materials, 5a Rogova St., 123060 Moscow, Russia b
ARTICLE INFO
ABSTRACT
Keywords: Metal-matrix composites (MMCs) Low-cycle fatigue Dynamic mechanical analyzer Fractography Electrical conductivity
The comparative low-cycle fatigue tests of the Cu/Nb nanocomposite wires obtained by the “melt-and-deform” method and pure copper samples were carried out using a dynamic mechanical analyzer (DMA). The fatigue tests were carried out using transverse bending scheme. It has been established that the fatigue fracture resistance characteristics of the Cu/Nb composites significantly higher as compared to that of pure copper samples. It has been found that the fatigue crack propagation in the Cu/Nb composites occurs at two scale levels. The effect of fatigue damage accumulation on the change in the electrical resistance of the Cu/Nb composites was studied.
1. Introduction Binary in situ nanocomposite wires based on a metallic copper matrix with the addition of bcc metals, namely, Cu/Nb, Cu/Fe, Cu/Cr, Cu/ V, are the subject of intensive research in recent decades [1–9]. The strenghtening phases formed during the deformation process have high thermodynamic stability and good adhesion to the matrix. The limited solubility of bcc metals in copper allows the copper matrix to maintain the high electrical conductivity, while bcc metal filaments provide high strength properties [10–12]. The unique combination of high strength and high electrical conductivity, due to the transition of the structure to the nanoscale state, distinguishes these composites into a separate class of materials and allows one to use them as promising materials for the manufacture of mechanically loaded parts of electrical devices. Since the mid-1980s, the Cu/Nb nanocomposite conductors have traditionally been used in the manufacture of windings of high-field pulsed magnets [13]. The manufacturing process of the Cu/Nb composite conductors with large cross sections and high mechanical properties is characterized by some technological features that have much in common with the manufacturing process of low-temperature superconductors by the RRP method (a method of a multi-rod set stacking, its subsequent evacuating and extrusion). As a result of the manufacturing process, there obtained a complex composite with a multi-scale structure. Despite the fact that common mechanisms of the structure
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formation of in-situ composites, as well as their mechanical properties under static loading, are well understood, literature data on mechanical and functional properties of the composites under cyclic loading are scarce [14–16]. Unconventional fatigue test techniques, in particular, performed by a dynamic mechanical analyzer (DMA) can be applied to evaluate the fatigue resistance of small-scale samples with a diameter of less than 1 mm, such as in situ composite wires. In case of fatigue test of small cross-section samples, the DMA installation shows the advantage in the accuracy of the applied load compared to the standard testing machines. Previously, DMA was successfully used for comparative bending fatigue tests of orthodontic wires (NiTi and Cu-NiTi) [17], thin samples of zirconium alloys [18,19] and Cu-Fe and Cu-V nanocomposite wires [15]. In this work, the comparative low-cycle bending fatigue tests of the Cu/Nb nanocomposite wires obtained by the “melt-and-deform” method, as well as of the pure copper samples were carried out using DMA. The change in electrical conductivity of the Cu/Nb nanocomposite wires due the accumulation of fatigue damage was also studied. 2. Material and research methods The samples of a nanocomposite wire with a diameter of 0.46 mm, obtained by the “melt-and-deform” method (the method has been described in more detail in [20]), were investigated. The basis of the
Corresponding author. E-mail address: [email protected] (S.O. Rogachev).
https://doi.org/10.1016/j.ijfatigue.2019.105188 Received 1 April 2019; Received in revised form 1 July 2019; Accepted 4 July 2019 Available online 05 July 2019 0142-1123/ © 2019 Elsevier Ltd. All rights reserved.
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Fig. 1. The structure of the Cu/Nb composite in cross section (SEM): (a) general view; (b) fragment of the structure. Table 1 The characteristics of the studied materials. Sample
Diameter, mm
Electrical conductivity, % IACS
Tensile strength (σUTS), MPa
Bending strength, MPa
Cu/Nb Cu
0.463 ± 0.005 0.50 ± 0.01
68.8 ± 0.1 95.4 ± 0.1
1125 ± 10 415 ± 10
1660 ± 20 790 ± 20
composite was a Cu-18% wt.Nb alloy, produced by a consumableelectrode vacuum arc melting. The core of the Cu-18% wt.Nb alloy was placed in a copper shell, then deformed by extrusion and drawing to form a hexagonal rod. The rods were cut to specified lengths, than stacked and assembled with copper shell. Thus, the final composite was fabricated using a double assembly and consisted of hexagonal Cu/Nb sub-elements surrounded by copper layers (Fig. 1). The total degree of cold drawn deformation of the final assembly was ln(A0/A) ∼ 8. For comparison, the wire samples of pure copper with a diameter of 0.50 mm, obtained by cold drawn deformation, were used. The characteristics of the studied materials, their strength and electrically conductive characteristics are presented in Table 1. Low-cycle fatigue tests of the Cu/Nb composite samples and pure copper with length of 30 mm were carried out using DMA Q800 (TA Instrument) according to the previously described method [18]. The tests were carried out using transverse bending scheme in one plane at a constant stress level with an alternating symmetric loading cycle (stress ratio R = −1) using a single cantilever grip of DMA Q800, until the failure of the samples (Fig. 2). During the grip movement in the course of tests, a sample was symmetrically deviated from the initial position to a certain amplitude. A soft loading scheme (at a constant cycle stress amplitude) with a frequency of loading cycles of 10 Hz at a temperature of 30 °C was used. At the end of tests, the number of load cycles to failure (N) was registered (based on 2.5 × 104 cycles). The tests were
carried out at a cycle stress of 500–1330 MPa (for Cu/Nb composite) and 230–710 MPa (for copper), which accounted for 30–80% of the fracture stress of the sample in tests with a stress amplitude increment (bending strength, see Table 1). At least three samples were used for each cycle stress amplitude (σa). The microstructure analysis of the samples in cross section and the analysis of the fatigue fractures was performed using a JSM-6610LV scanning electron microscope (JEOL). Electron microscopic studies of the structure of the Cu/Nb composite samples before and after cyclic tests were performed using a JEM2100 transmission electron microscope (TEM). The samples for electron microscopic studies were cut from hexagonal Cu/Nb sub-elements of the composite in transverse (before fatigue testing) and longitudinal (before and after fatigue testing) sections. The sample preparation was performed by the focused ion beam method using a Strata 201 SIMSmapIII×P scanning ion microscope with a gallium liquid metal ion gun, as described in [21]. 10 images at the same magnification were taken for the microstructure quantitative analysis, that was carried out using the Image Expert Pro software. The measurements of an electrical conductivity (electrical resistivity) were carried out using the Cu/Nb composite samples, which were cyclically loaded at 25; 50 and 75% of life time (where life time is a total number of load cycles to failure) at each cycle stress amplitude. The number of load cycles to failure (N) was determined from the
Fig. 2. The sample in a single cantilever grip of DMA Q800: (a) image of the sample and grip, (b) the sample deviation scheme at fatigue testing. 2
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Fig. 3. TEM photographs of the Cu/Nb composite microstructures before (a, c) and after (b, d) fatigue tests at a maximum cycle stress: (a–c) cross section, (d) longitudinal section.
fatigue curve obtained from fatigue tests. The electrical resistance was determined by the four-probe method using a test bench for measuring the RRR and Tc of superconductors (in accordance with IEC 61788-4: 2016 and IEC 61788-10: 2006). The measurements were carried out at temperatures of 293 K, 77.4 K and 10 K.
Table 2 The characteristics of the Cu/Nb composite microstructure (in the cross section of the wire). The thickness of the Nb-filaments, nm The width of the Nb-filaments, nm The density of the Nb-filaments, pcs/µm2 Area in the cross section, occupied with the Nb-filaments, % Cross-sectional size of copper grains (subgrains), nm
3. Results and discussion 3.1. The analysis of the microstructure
13.6 ± 3.9 81 ± 12 140.0 ± 11.0 23.4 ± 3.1 100 ± 15
obtained from the longitudinal sections of the composite, no notable changes associated with the action of the alternating deformation during cyclic tests were found (Fig. 3c, d). It should be noted that the Cu/Nb samples were obtained by drawing with a very large degree of strain without subsequent annealing. Thus, before the fatigue tests, the Cu/Nb samples are already contain a very high density of dislocation and other lattice defects. Therefore, the dislocation density during fatigue tests could not increase very significantly. The TEM images clearly show that the structure of the Cu/Nb samples contains a very large number of defects both before and after the fatigue tests. However, it is difficult to quantify the change in the density of defects.
Fig. 3 shows TEM images of the microstructures (with electron diffraction patterns) of a single hexagonal sub-element of the Cu/Nb composite in transverse and longitudinal section in original and tested states (at a maximum cycle stress −1330 MPa). A grain-subgrain nanostructure of copper matrix is observed. The “twisted” Nb-filaments are surrounded by copper grains (subgrains) with an average crosssection of 100 nm. The twisting of the filaments in such composites is usually associated with the appearance of a rotational mode in the process of drawing [22]. In our opinion, the twisting process can also be associated with the inhomogeneity of the plastic flow in the material during large plastic deformations due drawing, which leads to the appearance of a turbulent plastic flow and the appearance of the “vortex” defects [23,24]. In the longitudinal section, the Nb-filaments have an elongated shape. The results of the quantitative metallographic analysis of the Cu/Nb composite microstructure, carried out using the crosssectional images, are presented in Table 2. When comparing the microstructure of the Cu/Nb composite in original and tested states, as well as the electron diffraction patterns
3.2. Fatigue strength and fractographic research Fatigue curves (N − σa) of the Cu/Nb composite samples, as well as of the pure copper samples are shown in Fig. 4. The fatigue limit (σRN) of the Cu/Nb composite samples at N = 2.5 × 104 cycles was 490 MPa. The fatigue limit of pure copper samples is much lower, σRN = 390 MPa. The maximum difference in the fatigue strength 3
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Fig. 4. The fatigue curves of the studied samples.
striations (Fig. 6c). Elongated dimples were observed in the area of static fracture. The static fracture area is shifted from the center of the fracture as in the case of the Cu/Nb composite. At large cycle stresses, the macro- and microstructure of the fatigue fracture of pure copper samples do not qualitatively change, but more intense secondary cracking is observed along the boundaries of the striations, and the distance between the micro-striations is larger. There are macro-striations at the fracture, the distance between which is 2–13 μm. The static fracture area is shifted to the center of the fracture. Thus, the fatigue crack propagation in the Cu/Nb composite samples, unlike pure copper samples, occurs at two scale levels: within individual Cu/Nb sub-elements with the formation of micro-striations and within the entire volume of the composite with the formation of macro-striations. This fact, obviously, contributes to the increased resistance to fatigue failure of the Cu/Nb composite samples, especially at high cycle stresses. In addition, the Nb-filaments also appear to be additional barriers to the fatigue crack propagation, which also explains the higher resistance to fatigue failure of the Cu/Nb composite compared to pure copper. A possible mechanism of the influence of the filaments in the in-situ composites was previously considered using the Cu/Cr composite [14]. The cracking along the boundaries of sub-elements in the Cu/Nb composite sample under fatigue loading may be due to the fact that the copper interlayers along the boundaries of sub-elements are in a highly deformed state and lose their ductility. This is also confirmed by the rather brittle mechanism of fatigue fracture of pure copper, the degree of strain of which is close to the degree of strain of the copper interlayers in the Cu/Nb composite sample. In [16], when studying the fatigue failure of the Cu-18% (vol.) Nb composite with a Nb-filament thickness of 5–10 nm, it was found that the copper interlayers between the sub-elements are the place of the fatigue crack nucleation, as well as the place of its inhibition. However, it should be noted that the authors [16] removed the copper shell of the composite before fatigue tests, and also carried out high-cycle fatigue tests (i.e., fatigue defects were nucleated mainly in the region of elastic stresses) and used another loading scheme (cyclic tensile), which does not allow to compare the results.
between the Cu/Nb composite samples and pure copper samples was observed at a high stress level of load cycles. The nucleation of fatigue cracks and further fracture of all samples during tests on the DMA were initiated on the sample surface at a distance from the ends of the grips of 1–3 mm. The fractographic analysis of the Cu/Nb composite samples as well as of the pure copper samples revealed similar macrostructure of fatigue fractures, but different mechanisms of fatigue crack propagation (Figs. 5a, 6a). As usual for bending fatigue fracture, two areas of fatigue crack nucleation are seen on the sample surface near the opposite edges of the sample (sometimes only one area of fatigue crack nucleation was observed), then the crack propagated to the center of the sample, where the static fracture occurred. At low cycle stresses, typical fatigue micro-striations were observed on the surface of hexagonal sub-elements of the Cu/Nb composite samples in the area of stable growth of fatigue crack. The orientation of the micro-striations changed from one filament to another (Fig. 5e), which is associated with the branching of fatigue crack under the influence of local deformation conditions [25]. The distance between these striations was about 0.3 μm. Also, in the area of stable growth of fatigue crack, intense cracking was observed along the boundaries of the group of the Cu/Nb sub-elements to the normal direction of fatigue crack growth. In the area of accelerated crack growth, the macrostriations were observed perpendicular to the direction of crack growth, the distance between which was ∼1.5 μm (Fig. 5g). Such striations did not change their direction. The static fracture area is characterized by a fully ductility small-dimple structure (Fig. 5g). At the same time, the static fracture area is shifted from the center of the fracture; this is typical of bending tests of the samples with a low cycle stress and small stress concentrators [25]. At high cycle stresses, the macro- and microstructure of the fatigue fracture of the Cu/Nb composite samples do not qualitatively change, however, more intense cracking was observed along the boundaries of the group of the Cu/Nb sub-elements with an increased distance between micro-striations (Fig. 5d, f). The static fracture area located in the center of the fracture is substantially larger, which is associated with a shorter period of fatigue fracture itself (Fig. 5h). In fatigue fractures of pure copper samples at low cycle stresses, the fatigue striations are observed in the area of the fatigue crack propagation with a distance between them of ∼1.6 μm (Fig. 6e). In local areas, the secondary cracking was observed along the boundaries of the
3.3. The measurements of electrical resistivity Fig. 7 shows the graphs of the change in the electrical resistivity of 4
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Fig. 5. The surface of fatigue fractures of the Cu/Nb composite samples: at low (a, c, e, g) and high (b, d, f, h) cycle stress; 1 – fatigue crack nucleation area; 2 – stable growth of fatigue crack area; 3 – accelerated crack growth area; 4 – static fracture area (arrows in e indicate the direction of the fatigue crack growth).
the Cu/Nb composite samples under cyclic loading depending on the cycle stress amplitude (1330; 1000; 830; 665 and 500 MPa), the number of load cycles (25; 50 and 75% of life time) and temperature of the measurement (293; 77.4 and 10 K). To more clearly reveal the regularities of electrical resistivity, the data presented in Fig. 7 is rearranged in relative units ρi/ρ0, where ρ0 is the electrical resistance before cyclic loading (see Fig. 8). It is seen that, the electrical resistivity of the Cu/Nb composite samples decreases sharply with decreasing temperature, which is
associated with an increase in the mean free path of electron (λ). At the number of cycles equal to 25% of life time, an increase in the electrical resistivity is observed by 7% and 14% at a temperature of 293 K and 77.4 K, respectively. With further increase in the number of cycles a significant increase in electrical resistivity does not occur. This character is traced for all values of the cycle amplitude. At a temperature of 10 K, no apparent dependence of the electrical resistivity is observed either on the number of load cycles or on the stress cycle amplitude. This is due to a change in the nature of the scattering of conduction
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Fig. 6. The surface of fatigue fractures of pure copper samples: at low (a, c, e) and high (b, d, f) cycle stress; 1 – fatigue crack nucleation area; 2 – stable growth of fatigue crack area; 3 – accelerated crack growth area; 4 – static fracture area.
electrons at the Cu/Nb interface. The same Cu/Nb interface, depending on the value of λ, can be diffusely reflective and mirror reflective, that varies scattering coefficient p from 0 to 1, and, accordingly, affects the increase in electrical resistivity [26–28]. The scattering is considered to be diffuse if the irregularities of the interface have a value comparable with λ. At a temperature of 10 K, when λ can reach values of the order of 1 μm, which far exceeds the parameters of the roughness of the interface, the reflection of conduction electrons at the Cu/Nb interface is close to mirror, and the changes in electrical resistivity with increasing deformation are weakly expressed [26]. In the case of cyclic loading, the change in electrical conductivity (at 25% of life time) is obviously associated with the formation of fatigue defects: the continuity of the material is disturbed, shears are formed in individual grains, which then develop into microcracks. As can be seen, the fatigue failure is initiated on the sample surface, i.e. in the area of severe deformed copper shell. Thus, in the first place, the continuity of
the copper shell is broken. Considering that the electric current mainly flows on the surface of the conductor, the discontinuity of the copper shell leads to a decrease in electrical conductivity. Further accumulation of fatigue damage occurs already in the internal volume of the Cu/ Nb composite wire, which has less impact on the change in electrical conductivity. It should be noted that not only the size of defects, but also the heterogeneity of their location and morphology can influence the change in electrical conductivity. Thus, the absence of an explicit nature of the change in electrical resistivity, measured at 10 K, may be due to the fact that at a given temperature the contribution from electron scattering on fatigue defects to electrical resistivity is significantly less than the contribution from scattering at the Cu/Nb interfaces. You should also take into account the variation of the results characteristic of fatigue tests.
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Fig. 7. The electrical resistivity of the Cu/Nb composite samples vs the cycle stress amplitude and number of loading cycles: (a) at 293 K, (b) at 77.4 K, (c) at 10 K.
Fig. 8. The normalized electrical resistivity (ρi/ρ0) of the Cu/Nb composite samples vs the cycle stress amplitude and number of loading cycles: (a) at 293 K, (b) at 77.4 K, (c) at 10 K.
4. Conclusion
(3) Fatigue crack propagation in the Cu/Nb composite samples occurs at two large-scale levels: within individual Cu/Nb sub-elements with the formation of micro-striations and within the entire volume of the composite with the formation of macro-striations. (4) The accumulation of fatigue damage leads to a slight decrease in the electrical conductivity of the Cu/Nb composite samples.
(1) Nanocomposite Cu/Nb wires obtains by the “melt-and-deform” method have a complex hierarchically non-uniform structure: the hexagon Cu/Nb sub-elements surrounded by copper interlayers. The Cu/Nb sub-element is characterized by grain-subgrain nanostructure with second phase precipitates (niobium) having a “twisted” shape in cross section and elongated shape in longitudinal section, located mainly along the boundaries of copper grains (subgrains) with an average cross section of 100 nm. (2) It has been established that the fatigue fracture resistance characteristics of the Cu/Nb nanocomposite wires are significantly higher as compared to that of pure copper. The fatigue limit (σRN) at N = 2.5 × 104 cycles for the samples of the Cu/Nb composite and pure copper is 490 and 390 MPa, respectively.
Acknowledgment The work was carried out with financial support from the Ministry of Science and Higher Education of the Russian Federation in the framework of Increase Competitiveness Program of NUST «MISiS» (№ К22019-008), implemented by a governmental decree dated 16th of March 2013, N 211. 7
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References
2015;70:92–8. https://doi.org/10.1016/j.compositesb.2014.10.046. [16] Nayeb-Hashemi H, Pourrahimi S. Synthesis, strengthening, fatigue and fracture behavior of high-strength, high-conductivity P/M processed Cu-Nb microcomposite. In: Mitchell MR, Buck O, editors. Cyclic deformation, fracture, and nondestructive evaluation of advanced materials: second volume ASTM International; 1994. p. 87–108. https://doi.org/10.1520/STP18121S. [17] Prymak O, Klocke A, Kahl-Nieke B, Epple M. Fatigue of orthodontic nickel–titanium (NiTi) wires in different fluids under constant mechanical stress. Mater Sci Eng A 2004;378:110–4. https://doi.org/10.1016/j.msea.2003.10.332. [18] Nikulin SA, Markelov VA, Gusev AYu, Nechaykina TA, Rozhnov AB, Rogachev SO, Zadorozhnyy MYu. Low-cycle fatigue tests of zirconium alloys using a dynamic mechanical analyzer. Int J Fatigue 2013;48:187–91. https://doi.org/10.1016/j. ijfatigue.2012.10.019. [19] Nikulin SA, Rozhnov AB, Gusev AYu, Nechaykina TA, Rogachev SO, Zadorozhnyy MYu. Fracture resistance of Zr-Nb alloys under low-cycle fatigue tests. J Nucl Mater 2014;446:10–4. https://doi.org/10.1016/j.jnucmat.2013.11.039. [20] Dubois JB, Thilly L, Lecouturier F, Olier P, Renault PO. Cu/Nb nanocomposite wires processed by severe plastic deformation for applications in high pulsed magnets: effects of the multi-scale microstructure on the mechanical properties. IEEE Trans Appl Supercond 2012;22(3):6900104. https://doi.org/10.1109/TASC.2011. 2174574. [21] Nikulin SA, Pantsyrnyi VI, Rozhnov AB, Rogachev SO, Khlebova NE, Nechaikina TA, et al. Possibilities of scanning and transmission electron microscopy for comparative analysis of the microstructure of in-situ nanocomposite high-strength conductors based on a copper matrix and bcc-metals. Metal Sci Heat Treat 2016;58(3–4):209–13. https://doi.org/10.1007/s11041-016-9990-y. [22] Hosford WF. Microstructural changes during deformation of (011) fiber-textured metals. Trans Metall Soc AIME 1964;230(1):12–4. [23] Kulagin R, Beygelzimer Y, Ivanisenko Y, Mazilkin A, Hahn H. High pressure torsion: from laminar flow to turbulence. IOP Conf Series: Mater Sci Eng 2017;194:012045. https://doi.org/10.1088/1757-899X/194/1/012045. [24] Rogachev SO, Sundeev RV, Khatkevich VM. Evolution of the structure and strength of steel/vanadium alloys/steel hybrid material during severe plastic deformation. Mater Lett 2016;173:123–6. https://doi.org/10.1016/j.matlet.2016.03.044. [25] Terent’ev VF. Fatigue strength of metals and alloys. Moscow: Intermet Inzhiniring; 2002. [in Russia]. [26] Guryev VV, Polikarpova MV, Lukyanov PA, Khlebova NE, Pantsyrny VI. Size effects influence on conducting properties of Cu-Nb alloy microcomposites at cryogenic temperature. Cryogenics 2018;90:56–9. https://doi.org/10.1016/j.cryogenics. 2018.01.010. [27] Verhoeven JD, Downing NL, Chumbley LS, Gibson ED. The resistivity and microstructure of heavily drawn Cu-Nb alloys. J Appl Phys 1989;65(3):1293–301. https://doi.org/10.1063/1.343024. [28] Deng L, Han K, Hartwig K, Siegrist T, Dong L, Sun Z, et al. Hardness, electrical resistivity, and modeling of in situ Cu–Nb microcomposites. J Alloys Comp 2014;602:331–8. https://doi.org/10.1016/j.jallcom.2014.03.021.
[1] Liu Z, Chen X, Li W. Mechanical properties of a composite Cu-Cr alloy obtained insitu by directed solidification. Met Sci Heat Treat 2006;48(9):463–5. https://doi. org/10.1007/s11041-006-0118-7. [2] Biselli C, Morris DG. Microstructure and strength of Cu–Fe in situ composites obtained from prealloyed Cu–Fe powders. Acta Metall Mater 1994;42:163–76. https:// doi.org/10.1016/0956-7151(94)90059-0. [3] Snoeck E, Lecouturier F, Thilly L, Casanove MJ, Rakoto H, Coffe G, Askénazy S, Peyrade JP, Roucau C, Pantsyrny V, Shikov A, Nikulin A. Microstructural studies of in situ produced filamentary Cu/Nb wires. Scr Mater 1998;38:1643–8. https://doi. org/10.1016/S1359-6462(98)00080-3. [4] Raabe D, Hangen U. Simulation of the yield strength of wire drawn Cu-based in-situ composites. Comput Mater Sci 1996;5:195–202. https://doi.org/10.1016/09270256(95)00072-0. [5] Liu K, Jiang Z, Zhao J, Zou J, Lu L, Lu D. Thermal stability and properties of deformation-processed Cu-Fe in situ composites. Metall Mater Trans A 2015;46:2255–61. https://doi.org/10.1007/s11661-015-2791-x. [6] Sauvage X, Genevois C, Costaa GDa, Pantsyrny V. Atomic scale characterization of deformation-induced interfacial mixing in a Cu/V nanocomposite wire. Scr Mater 2009;61:660–3. https://doi.org/10.1016/j.scriptamat.2009.06.007. [7] Deng L, Yang X, Han K, Lu Y, Liang M, Liu Q. Microstructure and texture evolution of Cu–Nb composite wires. Mater Char 2013;81:124–33. https://doi.org/10.1016/j. matchar.2013.04.013. [8] Lu X, Yao D, Chen Y, Wang L, Dong A, Meng L, et al. Microstructure and hardness of Cu-12% Fe composite at different drawing strains. J Zhejiang Univ-Sci A 2014;15(2):149–56. https://doi.org/10.1631/jzus.A1300164. [9] Brokmeier HG, Bolmaro RE, Signorelli JA, et al. Texture development of wire drawn Cu–Fe composites. Phys B 2000;276–278:888–9. https://doi.org/10.1016/S09214526(99)01540-9. [10] Raabe D, Heringhaus F, Hangen U, et al. Investigation of a Cu-20mass%Nb in situ composite. Part I: Fabrication, microstructure and mechanical properties. Zeitschrift für Metallkunde 1995;86:405–22. [11] Raabe D, et al. Metallic composites processed via extreme deformation: toward the limits of strength in bulk materials. MRS Bull 2010;35:982–91. https://doi.org/10. 1557/mrs2010.703. [12] Pantsyrnyi V, Shikov A, Vorobieva A, et al. High strength, high conductivity macroand microcomposite winding wires for pulsed magnets. Phys B 2001;294–295:669–73. https://doi.org/10.1016/S0921-4526(00)00741-9. [13] Foner S. 68.4-T-long pulse magnet: test of high strength microcomposite Cu/Nb conductor. Appl Phys Lett 1986;49:982–3. https://doi.org/10.1063/1.97503. [14] Masuda C, Tanaka Y. Fatigue properties of Cu–Cr in situ composite. Int J Fatigue 2006;28:1426–34. https://doi.org/10.1016/j.ijfatigue.2006.02.051. [15] Nikulin SA, Rogachev SO, Rozhnov AB, Pantsyrnyi VI, Khlebova NE, Nechaykina TA, Khatkevich VM, Zadorozhnyy MYu. Microstructure and fatigue strength of high-strength Cu-Fe and Cu-V in-situ nanocomposite wires. Compos B
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