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body-centered-cubic interface mixing during severe plastic deformation
Journal of Alloys and Compounds 586 (2014) 16–21
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Dislocation assisted face-centered-cubic/body-centered-cubic interface mixing during severe plastic deformation Jiabin Liu a,b, Lihui Zhang a, Lunan Song b, Liang Meng b, Yuewu Zeng b, Miao Wang c, Youtong Fang a, Jien Ma a,⇑ a b c
College of Electrical Engineering, Zhejiang University, Hangzhou 310027, China Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, China Department of Physics, Zhejiang University, Hangzhou 310027, China
a r t i c l e
i n f o
Article history: Received 25 May 2013 Received in revised form 9 September 2013 Accepted 11 September 2013 Available online 20 September 2013 Keywords: Transmission electron microscopy Nanostructure materials Diffusion Interface Electrical transport
a b s t r a c t Abundant Cu/Cr interface was produced to study the interfacial diffusion at face-centered-cubic/bodycentered-cubic interface during heavy cold drawing. Notable interface diffusion with a width of 2– 4 nm at Cu/Cr interface was detected. The appearance of the interdiffusion is well explained by the mechanical mixing mechanism with dislocations gliding through the Cu/Cr interface. The greater inelastic scattering of electrons at the interdiffusion zone was responsible to the more rapid increase of electrical resistivity with the increase of drawing strains. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Phase interfaces are key structural features in a wide range of engineering materials [1] and play a key role in controlling the mechanical, electrical, magnetic and optical properties of multiphase materials. Two kinds of typical structures are the face-centered-cubic/face-centered-cubic (FCC/FCC) and face-centered-cubic/body-centered-cubic (FCC/BCC) phase interfaces. Compared to the FCC/FCC phase interface, the FCC/BCC phase interface exhibits more interesting properties mainly due to the different crystalline characteristics between FCC and BCC lattices. The preferred slip system in FCC lattice is {1 1 1}/h1 1 0i while they are {1 1 0}/h1 1 1i and {1 1 2}/h1 1 1i in BCC lattice. The FCC phase mainly develops h1 1 1i and h1 0 0i texture in the drawing direction while the BCC phase prefers a h1 1 0i texture after heavy drawing [2,3]. A direct consequence is that the FCC–BCC alloys are commonly more difficult to undergo cold deformation than FCC–FCC alloys The morphological evolution of the embedded BCC phase in FCC matrix was found to adopt particularly convoluted cross-sectional shapes upon deformation to large strains by wire drawing [4–6].More pronounced Hall–Petch hardening has been observed in the FCC–BCC alloys, as compared to the FCC–FCC alloys [7,8]. Since deformation compatibility across the interface should be ⇑ Corresponding author. Tel./fax: +86 57187953134. E-mail address: [email protected] (J. Ma). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.09.061
accommodated by different dislocation slip systems, i.e. {1 1 1}/ h1 1 0i for FCC phase and {1 1 0}/h1 1 1i for BCC phase, the two phases between the interface should coordinate each other to realize the co-deformation. High density of dislocation may be generated at the interface as the geometrically necessary dislocations due to the inherently strain incompatibility between FCC and BCC phases [9]. It is noted that FCC–BCC alloys can undergo strain-induced chemical mixing (mechanical alloying) or even amorphization during heavily plastic deformation [10,11]. Amorphization has widely found for the binary metallic systems such as Ni–Ti and Ni–Nb, under ball milling. It has been proposed that the amorphization might be due to the local melting and rapid solidification among powder particles, as the result of intensive plastic deformation and fast heat conduction [12]. However, the peak temperature during milling was calculated to be well below the melting temperature of the starting materials. As a compromise, it was suggested that the large amount of defects caused by the severe plastic deformation should be responsible for the amorphization [13]. Although it is not necessary to form an amorphous phase for every binary metallic system after ball milling, the measured solubility of each element is notably higher than the equilibrium solubility. For example, the solubility limit of Ti in FCC Ni was about 28% for the Ni–Ti alloys after ball milling, which was significantly larger than that in annealed Ni [14]. The study of the Cu–15 wt.% Cr prepared by elemental powders also showed that there was a
J. Liu et al. / Journal of Alloys and Compounds 586 (2014) 16–21
significant diffusion zone with a range of 40 nm between Cu and Cr phases [15]. Compared to ball milling, the drawing process seems much more gentle and only produces little amorphous zone at high strains (above 9 for Cu–Nb–Ag alloys [10]). Each phase is gradually elongated during drawing and the density of FCC/BCC interface increases with the drawing strain. It is also noted that the mutual bulk solubility shows considerable deviation from the respective equilibrium values in heavily drawn specimens. Sauvage et al. [11] have documented large deviations from the negligible equilibrium solubility of Cu and Nb within each other in the vicinity of Cu/ Nb interface by using the 3D atom probe technique. Similar results were observed by Ohsaki et al. [16] in Cu–Ag systems. The enhanced interface diffusion was thought to be responsible for the capillary driven process of the Cr filaments during annealing [17]. The FCC/BCC interface should also play important role in the strain coordinating and electrons transmitting from one phase to the other phase. In this work abundant Cu/Cr interface was produced to study the interfacial diffusion at FCC/BCC interface during heavy cold drawing and the corresponding electrical properties were analyzed based on the microstructural investigation.
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shielded atmosphere at a pressure of 2.7 104 Pa. The as-cast ingots were machined to about 18 mm diameter in order to remove the surface defects. Then the ingots were cold drawn at room temperature with the help of lubricant (KLUBER ISOFLEX NBU 15). Drawing reduction was presented in term of the logarithmic strain and referred as drawing strain.
g ¼ lnðA0 =AÞ
ð1Þ
where A0 and A were the cross-section areas of the original and drawn specimens. The drawing reduction per step was less than 0.2. The microstructure of the specimens at various drawing strains was investigated by high resolution transmission electron microscope (HRTEM, FEI F20) operated at 200 kV. Energy-dispersive spectrometer (EDS, Oxford INCA) equipped in F20 was used to perform composition analysis. The nominal maximum error for the EDS is 0.5% at an operating voltage of 200 kV for HRTEM and a dwelling time of 60 s. EDS test was performed under scanning transmission electron microscopy (STEM) mode with a step size of 1 nm or 1.5 nm and a dwelling time of 3 s. Before EDS test, TEM samples were tilted to realize edge-on of Cu/Cr phase interface to avoid overlapping. TEM samples were sliced from longitudinal section of tested wires and mechanically thinned to about 60 lm followed by ion thinning at 3 kV with an incidence angle of 6°. The electrical resistivity of the specimens at various drawing strains was tested by using the standard four-probe method at room temperature. A digital resistance tester (SB2230) working at direct current mode was used to test the resistance of Cu–Cr wires. A bridge fixture (DQ-1) was used to fix the Cu–Cr wires with a standard length of 1 m.
2. Experiment details
3. Results and discussions
Cu–3 wt.% Cr was prepared by using vacuum induction furnace. Electrolytic Cu with 99.99% purity and metallic Cr with 99.9% purity were used as starting materials. Cylindrical ingots of 20 mm diameter were cast into copper mould under Ar
Fig. 1a shows the as-cast Cu–3 wt.% Cr alloy with Cr particles uniformly distributed in Cu grains. The average Cr particle size is
Fig. 1. Microstructure of tested alloy at different drawing strains (a) g = 0 and (b) g = 6.5. Insets are corresponding EDS results. Some dislocations penetrating the Cu/Cr interface are pointed by the coarse white arrows, (c) HRTEM image of the Cu/Cr interface at g = 6.5, the corresponding planes are indexed, (d) electron diffraction patters of Cu/Cr interface at g = 6.5. Inset is the schematic illustration of the unit cells of Cu and Cr phases, and (e) HRTEM image of the Cu/Cr interface, the dotted line indicates the location of Cu/Cr interface and the arrows point out the interface dislocations.
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J. Liu et al. / Journal of Alloys and Compounds 586 (2014) 16–21
2 lm. During the drawing process, both Cu and Cr phases are elongated along the drawing direction. Plenty of dislocations are generated at the Cu/Cr interface and some dislocations are observed to penetrate the Cu/Cr interface (Fig. 1b). After heavy drawing, both Cu and Cr phases evolve into filaments with diameters of a few nanometers. The diameters of Cu and Cr filaments decrease and the density of Cu/Cr interface increases with the drawing strains. HRTEM observation shows that there is no amorphous region at the Cu/Cr interface (Fig. 1c). From the electron diffraction patterns of Cu/Cr interface (Fig. 1d), Cu and Cr phases are indexed to have a modified Kurdjumov–Sachs orientation 1Þ ==ð1 1 0Þ with a misorienrelationship: h1 1 0iCu//h1 1 1iCr, ð1 1 Cu Cr ==ð1 1 0Þ with a misorientation tation about 6.3° and ð1 1 1Þ Cu Cr 1Þ about 5.1°. The nearly paralleled relationship between ð1 1 Cu and (1 1 0)Cr ensures the Cu and Cr phase keeping almost coherent at the interface. As shown in Fig. 1e, most of the Cu planes match well with the Cr planes except at the site of few dislocations. The equilibrium solubility of Cr in Cu at room temperature is less than 0.03 at.% while that of Cu in Cr is nearly zero [18]. EDS tests performed at Cu/Cr interface in the deformed Cu–3 wt.% Cr show that the interdiffusion region exists and widens as the increase of drawing strains in this study. For the Cr sphere in the undeformed Cu–3 wt.% Cr, the EDS results do not show obvious
interdiffusion region at the Cu/Cr interface, or the content of Cr element in Cu phase is too low to be detected by EDS (Fig. 2a and b). For the Cr phase deformed into ellipse, the EDS results show that there is also not obvious interdiffusion region at the Cu/Cr interface (Fig. 1a and b in supply materials). This result suggests that small drawing strain could not produce obvious interdiffusion of Cu and Cr elements. As the Cr was significantly refined into filament with an average diameter of about 90 nm, notable interdiffusion was observed on the Cu/Cr interface (Fig. 2c and d). To evaluate the interdiffusion degree, the interdiffusion region is defined as the region containing 2–98 at.% Cr or 2–98 at.% Cu in this study. Therefore, the width of interdiffusion region could be measured from the composition distribution curve. A width of 2 nm was observed in Fig. 2c and d. As Cr fibers were refined to 50 nm in diameter, it is widened to about 4 nm (Fig. 1c and d in supply materials). Further refining the Cr fibers to only 10 nm no longer increases the width of the interdiffusion region (Fig. 2e and f). It shows a platform of a maximum width of 4 nm for the interdiffusion region when the Cr fibers are refined below 50 nm. A significant diffusion zone was also detected in a range of 40 nm in Cu–15 wt.% Cr prepared by ball milling and hot extrusion [15]. Molecular dynamics simulation found that the interdiffusion behavior was different between the FCC–FCC system and
Fig. 2. (a) STEM image of Cu matrix and Cr sphere in the tested alloy before deformation, (b) distribution of Cu and Cr relative content along the line ‘‘AB’’ in (a), (c) STEM image of Cu matrix and Cr ribbon in the tested alloy at g=4.5, (d) distribution of Cu and Cr relative content along the line ‘‘AB’’ in (c), (e) STEM image of Cu matrix and Cr fiber in the tested alloy at g = 6.5 and(f) distribution of Cu and Cr relative content along the line ‘‘AB’’ in (e).
J. Liu et al. / Journal of Alloys and Compounds 586 (2014) 16–21
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Fig. 4. Variation of electrical resistivity of the tested alloy dependent on drawing strain.
Fig. 5. Schematic comparison of electrical resistivity of Cu–Cr alloy (a) without and (b) with intermixing. Fig. 3. Schematic illustration of interface diffusion based on the dislocation movement (a) the original configuration of Cu/Cr interface. Cu and Cr phases have the orientation relationship of [1 1 0]Cu//[1 1 1]Cr and ð1 1 1ÞCu // ð1 0 1ÞCr with a misfit of 6°. The Cu/Cr interface is sharp and there are some dislocations at the lower left corner moving towards the Cu/Cr interface. These dislocations are supposed to glide on the ð1 1 1ÞCu along [1 0 1]Cu in Cu phase and on the ð1 0 1ÞCr along [1 1 1]Cr in Cr phase. (b) The configuration of Cu/Cr interface after those dislocations gliding through. The Cu/Cr interface becomes bent and there are some dislocations at the right moving towards the Cu/Cr interface. These dislocations are 1ÞCu along supposed to glide on the (1 1 0)Cr along [1 1 1]Cr in Cr phase and on ð1 1 the [1 0 1]Cu in Cu phase, and (c) final configuration of Cu/Cr interface after those dislocations gliding through. Some Cr atoms are embedded in Cu lattice and the interdiffusion realizes.
FCC–BCC system with a model of particle embedded in Cu matrix [19]. For FCC–FCC system, such as Cu–Ag and Cu–Ni, the FCC particles were observed to undergo superdiffusive mixing with dislocation crossing the particle–matrix interface. For FCC–BCC system, such as Cu–Nb and Cu–V, dislocations did not cross into
the particles and develop amorphous shells at the particle–matrix interface. The Cu–Cr system in this study behaves similarly to FCC– BCC system initially but more like FCC–FCC system at high drawing strains. At beginning, Cr particles deformed little and lots of dislocations pile up at the Cu/Cr interface. At high drawing strains, Cr particles have undergone considerable shear and are gradually elongated into filaments, and interdiffusion region appears and widens. Both of elongation of the Cr filaments and the appearance of the interdiffusion region are the result of plastic deformation in this study. The elongation of the Cr filaments increases the density of Cu/Cr interface. The interdiffusion region mainly occurs at the Cu/Cr interface since there is severe plastic deformation at the interface and there are plenty of dislocations penetrating the interface. Several physical mechanisms have been proposed to explain the deformation induced interdiffusion. The thermodynamic destabilization mechanism suggests that the high interface energy and high
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J. Liu et al. / Journal of Alloys and Compounds 586 (2014) 16–21
density of crystalline defect are the driving forces for the interdiffusion. This mechanism has successfully explained the interdifusion in ball milled Cu–Fe powers [20,21]. The local temperature following impacts of steel balls might be high enough to promote diffusion during ball milling. However, the temperature does not significantly increase during cold drawing. Since the atomic mobility of Cu in Cr and Cr in Cu is not significant at room temperature, thermal diffusion cannot account for the enhanced atomic mobility and even form interdiffusion zone. It is noted that plenty of dislocations forming and penetrating the Cu/Cr interface (Fig. 1b), which suggests mechanical mixing mechanism to be dominated for the Cu/Cr interfaces. Shear of atomic planes by dislocations will lead to atomic shuffle across the Cu/Cr interface. Consequently significant Cu–Cr intermixing may occur near Cu/Cr interface. A schematic illustration is proposed to describe the atomic interdiffusion process at Cu/Cr interface (Fig. 3). As shown in Fig. 1d, Cu and Cr phases have the 1Þ ==ð1 1 0Þ with a misorienrelationship: h1 1 0iCu//h1 1 1iCr, ð1 1 Cu Cr ==ð1 1 0Þ with a misorientation tation about 6.3° and ð1 1 1Þ Cu Cr about 5.1°. The Cu/Cr interface is sharp at first and some dislocations are moving towards the Cu/Cr interface in the Cu phase (Fig. 3a). After dislocations gliding through the Cu/Cr interface, the Cu/Cr interface becomes bent. Some Cu atoms seem across the original Cu/Cr interface (Fig. 3b). Then some dislocations move towards the Cu/Cr interface in Cr phase. After those dislocations glide through the Cu/Cr interface, the Cu/Cr interface seems to be recovered. At the same time, there are some Cr atoms fully embedded in the Cu phase and therefore the interdiffusion realizes (Fig. 3c). Repeating this process leads to more Cr atoms embedded in Cu phases and vice versa. It seems that each transition of a lattice burgers vector through the Cu/Cr interface transports Cu or Cr atoms into the counter phase, which thus mechanically enhances the fraction of solved atoms and forms an intermixing zone at the Cu/Cr interface. The width of intermixing zone increases with the drawing strains at first and then tends to keep constant when the Cr filaments reduce below 10 nm. This mechanism has been discussed as an important contribution to explain the plastic compatibility of the heavily drawn Cu–FCC or Cu–BCC alloys [17]. Although there are a notable diffusion region as wide as 4 nm at the Cu/Cr interface, there is no amorphous zone formed, which is different from the Cu–Nb and Cu–V [10,19]. The reason may be that the Cr atom has similar single-bond covalent radii with Cu atom (0.122 nm for Cr and 0.112 nm for Cu [22]). Similar situation was also observed in Cu–Fe by either molecular dynamics simulation or high pressure torsion due to the same reason [19,23]. The interface evolution also affects the electrical resistivity, as shown in Fig. 4. At g < 4.5, the electrical resistivity increases with the drawing strain at a moderate rate. When g > 4.5, the electrical resistivity increases at a much faster rate with the drawing strain. For Cu–Cr alloys at different drawing strains, the photon scattering and solute atom scattering almost remain constant when there is no significant interface diffusion. As the increase of drawing strain, the densities of dislocation and Cu/Cr interface increase, resulting in the increase of electrical resistivity of Cu–Cr alloys. For g > 4.5, the density of dislocation is saturated at large drawing strains [24,25], which should not account for increasing rate of electrical resistivity with drawing strains. The appearance of the interdiffusion zone leads to more pronounced electrons scattering at the Cu/Cr interface. Fig. 5 shows the comparison of the electrical resistivity in the Cu–Cr alloys with and without intermixing phase.
4. Conclusions The original Cr particles evolved into filaments with a diameter below 10 nm after heavy cold drawing. There was not notable
interdiffusion at the Cu/Cr interface for the Cu–Cr alloy at g < 4.5. Obvious interdiffusion was observed for the Cu–Cr alloy at g > 4.5 and the width of the interdiffusion zone increased with the drawing strain at first and then tended to keep constant. The interdiffusion between Cu and Cr phases was caused by the mechanical mixing mechanism with plenty of dislocations gliding through the Cu/Cr interface. The electrical resistivity of the Cu–Cr alloy increased with the drawing strain and the existence of interdiffusion is responsible for the more rapid increase of the electrical resistivity at large drawing strains. Acknowledgements The project is financially supported by National Natural Science Foundation of China (No. 11202183) and Zhejiang Provincial Natural Science Foundation of China (Grant nos. Y4100193 and Z1110057). Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.jallcom. 2013.09.061. References [1] J.P. Hirth, R.C. Pond, R.G. Hoagland, X.Y. Liu, J. Wang, Interface defects, reference spaces and the Frank–Bilby equation, Prog. Mater. Sci. 58 (2013) 749–823. [2] B.Z. Cui, Y. Xin, K. Han, Structure and transport properties of nanolaminate Cu– Nb composite foils by a simple fabrication route, Scr. Mater. 56 (2007) 879– 882. [3] U.F. Kocks, C.N. Tome, H.R. Wenk, Texture and Anisotropy: Preferred Orientations in Polycrystals and their Effect on Materials Properties, Cambridge University Press, UK, 2000. [4] C.W. Sinclair, J.D. Embury, G.C. Weatherly, Basic aspects of the co-deformation of bcc/fcc materials, Mater. Sci. Eng. A272 (1999) 90–98. [5] K. Adachi, S. Tsubokawa, T. Takeuchi, H.G. Suzuki, Strengthening mechanism of cold-drawn wire of in situ Cu–Cr composite, J. Japan Inst. Metals 61 (1997) 397–403. [6] J.C.M. Kampe, T.H. Courtney, Transverse kinking of BCC fibers in drawn FCC(matrix)–BCC(dispersoid) in situ composites, Scr. Metall. 23 (1989) 141– 145. [7] W.A. Spitzig, A.R. Pelton, F.C. Laabs, Characterization of the strength and microstructure of heavily cold worked Cu–Nb composites, Acta Metall. 35 (1987) 2427–2442. [8] G. Frommeyer, G. Wassermann, Microstructure and anomalous mechanical properties of in situ-produced silver–copper composite wires, Acta Metall. 23 (1975) 1353–1360. [9] P.D. Funkenbusch, T.H. Courtney, On the strength of heavily cold worked in situ composites, Acta Metall. 33 (1985) 913–922. [10] D. Raabe, S. Ohsaki, K. Hono, Mechanical alloying and amorphization in Cu– Nb–Ag composite wires studied by transmission electron microscopy and atom probe tomography, Acta Mater. 57 (2009) 5254–5263. [11] X. Sauvage, L. Renaud, B. Deconihout, D. Blavette, D.H. Ping, K. Hono, Solid state amorphization in cold drawn Cu/Nb wires, Acta Mater. 49 (2001) 389–394. [12] A.Y. Yermakov, Y.Y. Ymrchikov, V.A. Barinov, Magnetic properties of amorphous powders of Y–Co alloys produced by gringding, Phys. Met. Metall. 52 (1981) 50. [13] R.B. Schwarz, C.C. Koch, Formation of amorphous alloys by the mechanical alloying of crystalline powders of pure metals and powders of intermetallics, Appl. Phys. Lett. 49 (1986) 146–148. [14] R.B. Schwarz, R.R. Petrich, C.K. Saw, The synthesis of amorphous Ni–Ti alloy powders by mechanical alloying, J. Non-Cryst. Solids 76 (1985) 281–302. [15] J.L. Liu, E.D. Wang, Z.Y. Liu, L.X. Hu, W.B. Fang, Phases interface in deformation processed Cu–15 wt.% Cr composite prepared by elemental powders, Mater. Sci. Eng. A382 (2004) 301–304. [16] S. Ohsaki, K. Yamazaki, K. Hono, Alloying of immiscible phases in wire-drawn Cu–Ag filamentary composites, Scr. Mater. 48 (2003) 1569–1574. [17] D. Raabe, J. Ge, Experimental study on the thermal stablility of Cr filaments in a Cu–Cr–Ag in situ composite, Scr. Mater. 51 (2004) 915–920. [18] T.B. Massalski, Binary Alloys Phase Diagrams, second ed., ASM International, Metals Park, OH, USA, 1990. [19] Y. Ashkenazy, N.Q. Vo, D. Schwen, R.S. Averback, P. Bellon, Shear induced chemical mixing in heterogeneous systems, Acta Mater. 60 (2012) 984–993.
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[23] X. Quelennec, A. Menand, J.M.L. Breton, R. Pippan, X. Sauvage, Homogeneous Cu–Fe supersaturated solid solutions prepared by severe plastic deformation, Philos. Mag. 90 (2010) 1179–11195. [24] S.I. Hong, Yield strength of a heavily drawn Cu–20% Nb filamentary microcomposite, Scr. Mater. 39 (1998) 1685–1691. [25] S.I. Hong, M.A. Hill, Microstructural stability and mechanical response to Cu– Ag microcomposite wires, Acta Mater. 46 (1998) 4111–4122.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Dislocation assisted face-centered-cubic/body-centered-cubic interface mixing during severe plastic deformation Jiabin Liu a,b, Lihui Zhang a, Lunan Song b, Liang Meng b, Yuewu Zeng b, Miao Wang c, Youtong Fang a, Jien Ma a,⇑ a b c
College of Electrical Engineering, Zhejiang University, Hangzhou 310027, China Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, China Department of Physics, Zhejiang University, Hangzhou 310027, China
a r t i c l e
i n f o
Article history: Received 25 May 2013 Received in revised form 9 September 2013 Accepted 11 September 2013 Available online 20 September 2013 Keywords: Transmission electron microscopy Nanostructure materials Diffusion Interface Electrical transport
a b s t r a c t Abundant Cu/Cr interface was produced to study the interfacial diffusion at face-centered-cubic/bodycentered-cubic interface during heavy cold drawing. Notable interface diffusion with a width of 2– 4 nm at Cu/Cr interface was detected. The appearance of the interdiffusion is well explained by the mechanical mixing mechanism with dislocations gliding through the Cu/Cr interface. The greater inelastic scattering of electrons at the interdiffusion zone was responsible to the more rapid increase of electrical resistivity with the increase of drawing strains. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Phase interfaces are key structural features in a wide range of engineering materials [1] and play a key role in controlling the mechanical, electrical, magnetic and optical properties of multiphase materials. Two kinds of typical structures are the face-centered-cubic/face-centered-cubic (FCC/FCC) and face-centered-cubic/body-centered-cubic (FCC/BCC) phase interfaces. Compared to the FCC/FCC phase interface, the FCC/BCC phase interface exhibits more interesting properties mainly due to the different crystalline characteristics between FCC and BCC lattices. The preferred slip system in FCC lattice is {1 1 1}/h1 1 0i while they are {1 1 0}/h1 1 1i and {1 1 2}/h1 1 1i in BCC lattice. The FCC phase mainly develops h1 1 1i and h1 0 0i texture in the drawing direction while the BCC phase prefers a h1 1 0i texture after heavy drawing [2,3]. A direct consequence is that the FCC–BCC alloys are commonly more difficult to undergo cold deformation than FCC–FCC alloys The morphological evolution of the embedded BCC phase in FCC matrix was found to adopt particularly convoluted cross-sectional shapes upon deformation to large strains by wire drawing [4–6].More pronounced Hall–Petch hardening has been observed in the FCC–BCC alloys, as compared to the FCC–FCC alloys [7,8]. Since deformation compatibility across the interface should be ⇑ Corresponding author. Tel./fax: +86 57187953134. E-mail address: [email protected] (J. Ma). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.09.061
accommodated by different dislocation slip systems, i.e. {1 1 1}/ h1 1 0i for FCC phase and {1 1 0}/h1 1 1i for BCC phase, the two phases between the interface should coordinate each other to realize the co-deformation. High density of dislocation may be generated at the interface as the geometrically necessary dislocations due to the inherently strain incompatibility between FCC and BCC phases [9]. It is noted that FCC–BCC alloys can undergo strain-induced chemical mixing (mechanical alloying) or even amorphization during heavily plastic deformation [10,11]. Amorphization has widely found for the binary metallic systems such as Ni–Ti and Ni–Nb, under ball milling. It has been proposed that the amorphization might be due to the local melting and rapid solidification among powder particles, as the result of intensive plastic deformation and fast heat conduction [12]. However, the peak temperature during milling was calculated to be well below the melting temperature of the starting materials. As a compromise, it was suggested that the large amount of defects caused by the severe plastic deformation should be responsible for the amorphization [13]. Although it is not necessary to form an amorphous phase for every binary metallic system after ball milling, the measured solubility of each element is notably higher than the equilibrium solubility. For example, the solubility limit of Ti in FCC Ni was about 28% for the Ni–Ti alloys after ball milling, which was significantly larger than that in annealed Ni [14]. The study of the Cu–15 wt.% Cr prepared by elemental powders also showed that there was a
J. Liu et al. / Journal of Alloys and Compounds 586 (2014) 16–21
significant diffusion zone with a range of 40 nm between Cu and Cr phases [15]. Compared to ball milling, the drawing process seems much more gentle and only produces little amorphous zone at high strains (above 9 for Cu–Nb–Ag alloys [10]). Each phase is gradually elongated during drawing and the density of FCC/BCC interface increases with the drawing strain. It is also noted that the mutual bulk solubility shows considerable deviation from the respective equilibrium values in heavily drawn specimens. Sauvage et al. [11] have documented large deviations from the negligible equilibrium solubility of Cu and Nb within each other in the vicinity of Cu/ Nb interface by using the 3D atom probe technique. Similar results were observed by Ohsaki et al. [16] in Cu–Ag systems. The enhanced interface diffusion was thought to be responsible for the capillary driven process of the Cr filaments during annealing [17]. The FCC/BCC interface should also play important role in the strain coordinating and electrons transmitting from one phase to the other phase. In this work abundant Cu/Cr interface was produced to study the interfacial diffusion at FCC/BCC interface during heavy cold drawing and the corresponding electrical properties were analyzed based on the microstructural investigation.
17
shielded atmosphere at a pressure of 2.7 104 Pa. The as-cast ingots were machined to about 18 mm diameter in order to remove the surface defects. Then the ingots were cold drawn at room temperature with the help of lubricant (KLUBER ISOFLEX NBU 15). Drawing reduction was presented in term of the logarithmic strain and referred as drawing strain.
g ¼ lnðA0 =AÞ
ð1Þ
where A0 and A were the cross-section areas of the original and drawn specimens. The drawing reduction per step was less than 0.2. The microstructure of the specimens at various drawing strains was investigated by high resolution transmission electron microscope (HRTEM, FEI F20) operated at 200 kV. Energy-dispersive spectrometer (EDS, Oxford INCA) equipped in F20 was used to perform composition analysis. The nominal maximum error for the EDS is 0.5% at an operating voltage of 200 kV for HRTEM and a dwelling time of 60 s. EDS test was performed under scanning transmission electron microscopy (STEM) mode with a step size of 1 nm or 1.5 nm and a dwelling time of 3 s. Before EDS test, TEM samples were tilted to realize edge-on of Cu/Cr phase interface to avoid overlapping. TEM samples were sliced from longitudinal section of tested wires and mechanically thinned to about 60 lm followed by ion thinning at 3 kV with an incidence angle of 6°. The electrical resistivity of the specimens at various drawing strains was tested by using the standard four-probe method at room temperature. A digital resistance tester (SB2230) working at direct current mode was used to test the resistance of Cu–Cr wires. A bridge fixture (DQ-1) was used to fix the Cu–Cr wires with a standard length of 1 m.
2. Experiment details
3. Results and discussions
Cu–3 wt.% Cr was prepared by using vacuum induction furnace. Electrolytic Cu with 99.99% purity and metallic Cr with 99.9% purity were used as starting materials. Cylindrical ingots of 20 mm diameter were cast into copper mould under Ar
Fig. 1a shows the as-cast Cu–3 wt.% Cr alloy with Cr particles uniformly distributed in Cu grains. The average Cr particle size is
Fig. 1. Microstructure of tested alloy at different drawing strains (a) g = 0 and (b) g = 6.5. Insets are corresponding EDS results. Some dislocations penetrating the Cu/Cr interface are pointed by the coarse white arrows, (c) HRTEM image of the Cu/Cr interface at g = 6.5, the corresponding planes are indexed, (d) electron diffraction patters of Cu/Cr interface at g = 6.5. Inset is the schematic illustration of the unit cells of Cu and Cr phases, and (e) HRTEM image of the Cu/Cr interface, the dotted line indicates the location of Cu/Cr interface and the arrows point out the interface dislocations.
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2 lm. During the drawing process, both Cu and Cr phases are elongated along the drawing direction. Plenty of dislocations are generated at the Cu/Cr interface and some dislocations are observed to penetrate the Cu/Cr interface (Fig. 1b). After heavy drawing, both Cu and Cr phases evolve into filaments with diameters of a few nanometers. The diameters of Cu and Cr filaments decrease and the density of Cu/Cr interface increases with the drawing strains. HRTEM observation shows that there is no amorphous region at the Cu/Cr interface (Fig. 1c). From the electron diffraction patterns of Cu/Cr interface (Fig. 1d), Cu and Cr phases are indexed to have a modified Kurdjumov–Sachs orientation 1Þ ==ð1 1 0Þ with a misorienrelationship: h1 1 0iCu//h1 1 1iCr, ð1 1 Cu Cr ==ð1 1 0Þ with a misorientation tation about 6.3° and ð1 1 1Þ Cu Cr 1Þ about 5.1°. The nearly paralleled relationship between ð1 1 Cu and (1 1 0)Cr ensures the Cu and Cr phase keeping almost coherent at the interface. As shown in Fig. 1e, most of the Cu planes match well with the Cr planes except at the site of few dislocations. The equilibrium solubility of Cr in Cu at room temperature is less than 0.03 at.% while that of Cu in Cr is nearly zero [18]. EDS tests performed at Cu/Cr interface in the deformed Cu–3 wt.% Cr show that the interdiffusion region exists and widens as the increase of drawing strains in this study. For the Cr sphere in the undeformed Cu–3 wt.% Cr, the EDS results do not show obvious
interdiffusion region at the Cu/Cr interface, or the content of Cr element in Cu phase is too low to be detected by EDS (Fig. 2a and b). For the Cr phase deformed into ellipse, the EDS results show that there is also not obvious interdiffusion region at the Cu/Cr interface (Fig. 1a and b in supply materials). This result suggests that small drawing strain could not produce obvious interdiffusion of Cu and Cr elements. As the Cr was significantly refined into filament with an average diameter of about 90 nm, notable interdiffusion was observed on the Cu/Cr interface (Fig. 2c and d). To evaluate the interdiffusion degree, the interdiffusion region is defined as the region containing 2–98 at.% Cr or 2–98 at.% Cu in this study. Therefore, the width of interdiffusion region could be measured from the composition distribution curve. A width of 2 nm was observed in Fig. 2c and d. As Cr fibers were refined to 50 nm in diameter, it is widened to about 4 nm (Fig. 1c and d in supply materials). Further refining the Cr fibers to only 10 nm no longer increases the width of the interdiffusion region (Fig. 2e and f). It shows a platform of a maximum width of 4 nm for the interdiffusion region when the Cr fibers are refined below 50 nm. A significant diffusion zone was also detected in a range of 40 nm in Cu–15 wt.% Cr prepared by ball milling and hot extrusion [15]. Molecular dynamics simulation found that the interdiffusion behavior was different between the FCC–FCC system and
Fig. 2. (a) STEM image of Cu matrix and Cr sphere in the tested alloy before deformation, (b) distribution of Cu and Cr relative content along the line ‘‘AB’’ in (a), (c) STEM image of Cu matrix and Cr ribbon in the tested alloy at g=4.5, (d) distribution of Cu and Cr relative content along the line ‘‘AB’’ in (c), (e) STEM image of Cu matrix and Cr fiber in the tested alloy at g = 6.5 and(f) distribution of Cu and Cr relative content along the line ‘‘AB’’ in (e).
J. Liu et al. / Journal of Alloys and Compounds 586 (2014) 16–21
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Fig. 4. Variation of electrical resistivity of the tested alloy dependent on drawing strain.
Fig. 5. Schematic comparison of electrical resistivity of Cu–Cr alloy (a) without and (b) with intermixing. Fig. 3. Schematic illustration of interface diffusion based on the dislocation movement (a) the original configuration of Cu/Cr interface. Cu and Cr phases have the orientation relationship of [1 1 0]Cu//[1 1 1]Cr and ð1 1 1ÞCu // ð1 0 1ÞCr with a misfit of 6°. The Cu/Cr interface is sharp and there are some dislocations at the lower left corner moving towards the Cu/Cr interface. These dislocations are supposed to glide on the ð1 1 1ÞCu along [1 0 1]Cu in Cu phase and on the ð1 0 1ÞCr along [1 1 1]Cr in Cr phase. (b) The configuration of Cu/Cr interface after those dislocations gliding through. The Cu/Cr interface becomes bent and there are some dislocations at the right moving towards the Cu/Cr interface. These dislocations are 1ÞCu along supposed to glide on the (1 1 0)Cr along [1 1 1]Cr in Cr phase and on ð1 1 the [1 0 1]Cu in Cu phase, and (c) final configuration of Cu/Cr interface after those dislocations gliding through. Some Cr atoms are embedded in Cu lattice and the interdiffusion realizes.
FCC–BCC system with a model of particle embedded in Cu matrix [19]. For FCC–FCC system, such as Cu–Ag and Cu–Ni, the FCC particles were observed to undergo superdiffusive mixing with dislocation crossing the particle–matrix interface. For FCC–BCC system, such as Cu–Nb and Cu–V, dislocations did not cross into
the particles and develop amorphous shells at the particle–matrix interface. The Cu–Cr system in this study behaves similarly to FCC– BCC system initially but more like FCC–FCC system at high drawing strains. At beginning, Cr particles deformed little and lots of dislocations pile up at the Cu/Cr interface. At high drawing strains, Cr particles have undergone considerable shear and are gradually elongated into filaments, and interdiffusion region appears and widens. Both of elongation of the Cr filaments and the appearance of the interdiffusion region are the result of plastic deformation in this study. The elongation of the Cr filaments increases the density of Cu/Cr interface. The interdiffusion region mainly occurs at the Cu/Cr interface since there is severe plastic deformation at the interface and there are plenty of dislocations penetrating the interface. Several physical mechanisms have been proposed to explain the deformation induced interdiffusion. The thermodynamic destabilization mechanism suggests that the high interface energy and high
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density of crystalline defect are the driving forces for the interdiffusion. This mechanism has successfully explained the interdifusion in ball milled Cu–Fe powers [20,21]. The local temperature following impacts of steel balls might be high enough to promote diffusion during ball milling. However, the temperature does not significantly increase during cold drawing. Since the atomic mobility of Cu in Cr and Cr in Cu is not significant at room temperature, thermal diffusion cannot account for the enhanced atomic mobility and even form interdiffusion zone. It is noted that plenty of dislocations forming and penetrating the Cu/Cr interface (Fig. 1b), which suggests mechanical mixing mechanism to be dominated for the Cu/Cr interfaces. Shear of atomic planes by dislocations will lead to atomic shuffle across the Cu/Cr interface. Consequently significant Cu–Cr intermixing may occur near Cu/Cr interface. A schematic illustration is proposed to describe the atomic interdiffusion process at Cu/Cr interface (Fig. 3). As shown in Fig. 1d, Cu and Cr phases have the 1Þ ==ð1 1 0Þ with a misorienrelationship: h1 1 0iCu//h1 1 1iCr, ð1 1 Cu Cr ==ð1 1 0Þ with a misorientation tation about 6.3° and ð1 1 1Þ Cu Cr about 5.1°. The Cu/Cr interface is sharp at first and some dislocations are moving towards the Cu/Cr interface in the Cu phase (Fig. 3a). After dislocations gliding through the Cu/Cr interface, the Cu/Cr interface becomes bent. Some Cu atoms seem across the original Cu/Cr interface (Fig. 3b). Then some dislocations move towards the Cu/Cr interface in Cr phase. After those dislocations glide through the Cu/Cr interface, the Cu/Cr interface seems to be recovered. At the same time, there are some Cr atoms fully embedded in the Cu phase and therefore the interdiffusion realizes (Fig. 3c). Repeating this process leads to more Cr atoms embedded in Cu phases and vice versa. It seems that each transition of a lattice burgers vector through the Cu/Cr interface transports Cu or Cr atoms into the counter phase, which thus mechanically enhances the fraction of solved atoms and forms an intermixing zone at the Cu/Cr interface. The width of intermixing zone increases with the drawing strains at first and then tends to keep constant when the Cr filaments reduce below 10 nm. This mechanism has been discussed as an important contribution to explain the plastic compatibility of the heavily drawn Cu–FCC or Cu–BCC alloys [17]. Although there are a notable diffusion region as wide as 4 nm at the Cu/Cr interface, there is no amorphous zone formed, which is different from the Cu–Nb and Cu–V [10,19]. The reason may be that the Cr atom has similar single-bond covalent radii with Cu atom (0.122 nm for Cr and 0.112 nm for Cu [22]). Similar situation was also observed in Cu–Fe by either molecular dynamics simulation or high pressure torsion due to the same reason [19,23]. The interface evolution also affects the electrical resistivity, as shown in Fig. 4. At g < 4.5, the electrical resistivity increases with the drawing strain at a moderate rate. When g > 4.5, the electrical resistivity increases at a much faster rate with the drawing strain. For Cu–Cr alloys at different drawing strains, the photon scattering and solute atom scattering almost remain constant when there is no significant interface diffusion. As the increase of drawing strain, the densities of dislocation and Cu/Cr interface increase, resulting in the increase of electrical resistivity of Cu–Cr alloys. For g > 4.5, the density of dislocation is saturated at large drawing strains [24,25], which should not account for increasing rate of electrical resistivity with drawing strains. The appearance of the interdiffusion zone leads to more pronounced electrons scattering at the Cu/Cr interface. Fig. 5 shows the comparison of the electrical resistivity in the Cu–Cr alloys with and without intermixing phase.
4. Conclusions The original Cr particles evolved into filaments with a diameter below 10 nm after heavy cold drawing. There was not notable
interdiffusion at the Cu/Cr interface for the Cu–Cr alloy at g < 4.5. Obvious interdiffusion was observed for the Cu–Cr alloy at g > 4.5 and the width of the interdiffusion zone increased with the drawing strain at first and then tended to keep constant. The interdiffusion between Cu and Cr phases was caused by the mechanical mixing mechanism with plenty of dislocations gliding through the Cu/Cr interface. The electrical resistivity of the Cu–Cr alloy increased with the drawing strain and the existence of interdiffusion is responsible for the more rapid increase of the electrical resistivity at large drawing strains. Acknowledgements The project is financially supported by National Natural Science Foundation of China (No. 11202183) and Zhejiang Provincial Natural Science Foundation of China (Grant nos. Y4100193 and Z1110057). Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.jallcom. 2013.09.061. References [1] J.P. Hirth, R.C. Pond, R.G. Hoagland, X.Y. Liu, J. Wang, Interface defects, reference spaces and the Frank–Bilby equation, Prog. Mater. Sci. 58 (2013) 749–823. [2] B.Z. Cui, Y. Xin, K. Han, Structure and transport properties of nanolaminate Cu– Nb composite foils by a simple fabrication route, Scr. Mater. 56 (2007) 879– 882. [3] U.F. Kocks, C.N. Tome, H.R. Wenk, Texture and Anisotropy: Preferred Orientations in Polycrystals and their Effect on Materials Properties, Cambridge University Press, UK, 2000. [4] C.W. Sinclair, J.D. Embury, G.C. Weatherly, Basic aspects of the co-deformation of bcc/fcc materials, Mater. Sci. Eng. A272 (1999) 90–98. [5] K. Adachi, S. Tsubokawa, T. Takeuchi, H.G. Suzuki, Strengthening mechanism of cold-drawn wire of in situ Cu–Cr composite, J. Japan Inst. Metals 61 (1997) 397–403. [6] J.C.M. Kampe, T.H. Courtney, Transverse kinking of BCC fibers in drawn FCC(matrix)–BCC(dispersoid) in situ composites, Scr. Metall. 23 (1989) 141– 145. [7] W.A. Spitzig, A.R. Pelton, F.C. Laabs, Characterization of the strength and microstructure of heavily cold worked Cu–Nb composites, Acta Metall. 35 (1987) 2427–2442. [8] G. Frommeyer, G. Wassermann, Microstructure and anomalous mechanical properties of in situ-produced silver–copper composite wires, Acta Metall. 23 (1975) 1353–1360. [9] P.D. Funkenbusch, T.H. Courtney, On the strength of heavily cold worked in situ composites, Acta Metall. 33 (1985) 913–922. [10] D. Raabe, S. Ohsaki, K. Hono, Mechanical alloying and amorphization in Cu– Nb–Ag composite wires studied by transmission electron microscopy and atom probe tomography, Acta Mater. 57 (2009) 5254–5263. [11] X. Sauvage, L. Renaud, B. Deconihout, D. Blavette, D.H. Ping, K. Hono, Solid state amorphization in cold drawn Cu/Nb wires, Acta Mater. 49 (2001) 389–394. [12] A.Y. Yermakov, Y.Y. Ymrchikov, V.A. Barinov, Magnetic properties of amorphous powders of Y–Co alloys produced by gringding, Phys. Met. Metall. 52 (1981) 50. [13] R.B. Schwarz, C.C. Koch, Formation of amorphous alloys by the mechanical alloying of crystalline powders of pure metals and powders of intermetallics, Appl. Phys. Lett. 49 (1986) 146–148. [14] R.B. Schwarz, R.R. Petrich, C.K. Saw, The synthesis of amorphous Ni–Ti alloy powders by mechanical alloying, J. Non-Cryst. Solids 76 (1985) 281–302. [15] J.L. Liu, E.D. Wang, Z.Y. Liu, L.X. Hu, W.B. Fang, Phases interface in deformation processed Cu–15 wt.% Cr composite prepared by elemental powders, Mater. Sci. Eng. A382 (2004) 301–304. [16] S. Ohsaki, K. Yamazaki, K. Hono, Alloying of immiscible phases in wire-drawn Cu–Ag filamentary composites, Scr. Mater. 48 (2003) 1569–1574. [17] D. Raabe, J. Ge, Experimental study on the thermal stablility of Cr filaments in a Cu–Cr–Ag in situ composite, Scr. Mater. 51 (2004) 915–920. [18] T.B. Massalski, Binary Alloys Phase Diagrams, second ed., ASM International, Metals Park, OH, USA, 1990. [19] Y. Ashkenazy, N.Q. Vo, D. Schwen, R.S. Averback, P. Bellon, Shear induced chemical mixing in heterogeneous systems, Acta Mater. 60 (2012) 984–993.
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