Electrical Conductivity of Ultrafine-Grained Cu and Al Alloys: Attaining the Best Compromise with Mechanical Properties

Abstract

In recent years, the severe plastic deformation community has developed an interest for metallic materials with high strength and high electrical conductivity, with special focus in Cu- and Al-based alloys, including composite materials. Several processing and metallurgical strategies have been applied to control the influence of microstructure features such as grain refinement, grain boundary condition, defect structures and segregation of secondary phases, over the electrical and mechanical properties. This work summarizes an important body of literature where several strengthening mechanisms and methods to restore the electrical conductivity have been applied to produce ultrafine-grained or nanostructured Cu and Al alloys, mainly by intense imposed strain. A wide variety of alloy systems were studied for their industrial applications in the electrical and electronic market. It can be concluded that the balanced combination of alloying element selection and processing route (mainly attainable under high hydrostatic conditions) could provide high strength with high conductivity and thermal stability materials.

1. Introduction

The modern goal in the engineering of electrically conductive materials consists of attaining a satisfactory combination of high strength, good ductility, formability and thermal stability, with the least possible sacrifice of electrical conductivity.^1–6) Weight is also important for development of structural applications such as high-voltage power transmission lines. The electrical conductivity, σ, is an intrinsic property of a material defined as the inverse of the electrical resistivity ρ – that quantifies the capability of a material to allow the flow of electrons under an applied electric field.^1,7)   

\begin{equation} \sigma = \frac{1}{\rho} \end{equation} (1)

The SI unit of electrical conductivity is the Siemens per meter (S/m), and of electrical resistivity, the Ohm-meter (Ω·m). The International Annealed Copper Standard (IACS%) is a volume-based electrical conductivity measurement used by the metals industry as a percentage of the electrical conductivity of well-annealed Cu (σ = 5.8 × 10⁷ S/m at 293.15 K). The IACS% of the four highest conducting metals: Ag, Cu, Al and Au, is 111%, 101%, 78% and 64%, respectively. However, in cases where the weight is to be considered, the specific (mass-based) electrical conductivity is a more adequate measure to select a material. The specific conductivities of Ag, Cu, Al and Au are 95%, 101%, 211% and 36%, respectively. Thus, Al is a preferred choice, as opposed to Cu, in terms of cost and conductivity per unit of weight.

In principle, the first scattering mechanism in metals results from the collision of conduction electrons with the vibrating lattice. Therefore, in ideal pure metals the electrical resistivity has a strong temperature dependence.^1,7)   

\begin{equation} \rho_{\textit{lattice}} = AT \end{equation} (2)

Where A is a constant independent of temperature. The above equation is called the lattice-scattering-limited conductivity and is valid for a wide range of metals above room temperature, except for magnetic ones such as Fe and Ni. The electrical conductivity of pure metals is also influenced by microstructural features, such as dislocations, vacancies, solute atoms and grain boundaries, because all of these scatters the conduction electrons.¹⁾ Such features create a local distortion of the lattice, increasing the scattering cross section and effectively hindering electron motion. The contribution from solute atoms and crystal defects is called residual resistivity, and it is not significantly dependent of temperature. It can be summed to the lattice contribution:   

\begin{equation} \rho = \rho_{\textit{lattice}}(T) + \rho_{\textit{residual}} \approx AT + B \end{equation} (3)

Equation (3) is a well-known approximation called Matthiessen’s rule.^1,7,8) This approximation is also valid for many alloys. However, the morphology and distribution of second phases and precipitates can affect the electrical conductivity.^1,9) Most alloys follow a basic rule of mixtures, using the volume fraction (χ) and resistivity (ρ) of each phase. For a mixture of two phases α and β, the total resistivity of the alloy is:   

\begin{equation} \rho_{\textit{alloy}} = \chi_{\alpha}\rho_{\alpha} + \chi_{\beta}\rho_{\beta} \end{equation} (4)

There are deviations from this rule if either the matrix or the dispersed phase is much more resistive that the other. For composites, the physical-mathematical model of percolation theory⁷⁾ can describe the electrical conductivity of composites and combined systems of conducting/non-conducting material. The percolation threshold, p_c, defines the critical concentration where the conducting phase of the mixture forms a continuous network. The electrical conductivity is then:   

\begin{equation} \sigma = \sigma_{0}(p - p_{c})^{t} \end{equation} (5)

Where σ₀ and p are the conductivity and concentration of the conducting phase, respectively, and t is a factor between 1.3–3, dependent on the size and shape of the network.

Severe Plastic Deformation (SPD) provides bulk ultrafine-grained (UFG) or nanostructured (NS) materials with enhanced mechanical, chemical and physical properties, due to significant grain refinement and high density of lattice defects.^9–20) High-Pressure Torsion (HPT),^11,12) Equal-Channel Angular Pressing (ECAP)¹³⁾ and Accumulative Roll-Bonding (ARB)¹⁴⁾ are among the most studied techniques of SPD. Other SPD methods such as High-Pressure Sliding (HPS),²¹⁾ ECAP with Parallel Channels (ECAP-PC)²²⁾ Rotating Shear Plane (ECAP-R),²³⁾ Repetitive Corrugation and Straightening (RCS),²⁴⁾ Constrained Studded Pressing (CSP)²⁵⁾ and High-Pressure Torsion Extrusion (HPTE)^26–28) have been developed to expand the sample size for industrial application. There is an increasing demand of the electrical and electronic industries for materials of very good mechanical performance, low rates of failure and a balanced combination of thermal stability and electrical properties. Such requirements can be satisfied by metals and alloys processed by SPD, which provides a huge potential for development of innovative processes.^{15–18,29–33)} Researchers have studied extensively the effect of SPD on several alloys of Cu,^34–36) Al,^37,38) Ni,³⁵⁾ Ti and Zr.³⁹⁾

Processing through SPD produces significant increases in the tensile strength of fine-grained metals, ranging from 25–150%, in comparison with their coarse-grained (CG) counterparts. Although the conductivity is mildly affected by a reduction in grain size,^40,41) there is a marked sensitivity to other microstructural features due to the scattering of electrons by distortions in the crystal structure, thermal vibrations, impurities and lattice defects. Therefore, mechanical strength and electrical conductivity in metallic materials are inversely proportional.⁴²⁾ Pure metals with high conductivity, such as Cu and Al, are relatively soft in annealed condition. Thus, strategies to produce strengthening, while retaining ductility and electrical conductivity, become important for industrial application as conductors. This is a difficult task, since hardening by the addition of solute elements and by deformation decrease the ductility and the electrical conductivity. Therefore, strategies involving the recovery of defects, changing the structure of grain boundaries by twinning or the precipitation of secondary phases can aid to retain a high fraction of ductility and electrical conductivity while keeping high strength. Several research groups have reported high strength while retaining a higher fraction of electrical conductivity in Cu and Al alloys, based on SPD methods combined with heat treatments.^{37,38,43–45)}

This overview summarizes an important body of work where severe deformation techniques are used to produce bulk UFG and/or NS metallic materials, in conditions where high electrical conductivity is attainable after deformation processing, for possible industrial applications in the electrical and electronic markets.

2. Copper Alloys

As it is well known, pure Cu and Cu-based alloys have very attractive mechanical, electrical and thermal properties for industrial purposes. As it was previously stated, strengthening processes based on introduction of crystal structure defects, generally decrease the electrical conductivity, especially in pure Cu. Yang et al.⁶⁾ classified high strength and high conductivity Cu alloys as alloys with tensile strength 1.5 to 4 times that of pure Cu (UTS of 300–800 MPa) and electrical conductivity of 50–95 IACS%. The most desirable Cu conductors should have tensile strength higher than 600 MPa and electrical conductivity higher than 80 IACS%. High performance of both mechanical and electrical properties can be achieved through the addition of alloying elements, especially at elevated temperatures. This metallurgical strategy allows grain refinement and deformation strengthening, while tailoring the strength-conductivity trade-off by precipitation or other fine dispersed phases. In the last decades, Cr and Zr have been used as the most common alloying elements for this purpose. The following sections show newer strategies to increase the strength and retain a electrical conductivity using (1) grain refinement by SPD in high-purity Cu, cast and wrought Cu alloys with or without post-processing annealing, (2) combination of deformation processing with heat treatment in age hardenable Cu alloys and (3) development of Cu with fine grains and nanoscale twins.

2.1 Grain refinement by SPD in high-purity Cu, cast and wrought Cu alloys with or without post-processing annealing

One of the most common strategies to increase the strength of Cu is based on grain refinement down to the ultrafine grained structure. UFG alloys produced by SPD exhibit superior mechanical performance in comparison with their CG counterparts.^9–21) Partial recovery of defects, dynamic recrystallization^46,47) and precipitation phenomena^48,49) can occur during or after SPD, which aids to recover the electrical conductivity.

In the work of Edalati et al.,⁴³⁾ high-purity Cu (99.99%) in disc and ring shape was processed by HPT under 2 GPa. The Vickers microhardness increased from ∼50 HV in the annealed condition (873.15 K for 1 h), up to ∼130 HV at the steady state grain size, which was reached after N = 1 revolution, but was more homogeneous after N = 10 revolutions. Consequently, the tensile strength increased from ∼200 MPa up to ∼440 MPa after N = 1 revolution. The microstructural evolution by TEM showed formation of subgrains, with an initial average size of ∼1.0 µm after N = 1/8 turn in the disks and decreased to ∼0.2 µm in the steady state with increasing density of dislocations within the grains. In the steady-state region of the sample processed for N = 1 revolution, some grains were observed with low density of dislocations, attributed to recrystallization during the HPT processing. It was later shown by Edalati et al.,³⁶⁾ that the microhardness and electrical resistivity of the same high-purity Cu processed by ECAP and HPT increases with equivalent strain at an initial stage and then saturates to a steady state after equivalent strains above ∼20. Figure 1 shows that the electrical resistivity experimented an increase of ∼12% at the steady-state with respect to the annealed condition. In this state, the increase in the hardness is close to 3 times (∼270%) the hardness in the annealed condition. The mechanical properties achieved were comparable with the UFG-high purity copper samples from the study in Ref. 43). The microstructure observation by TEM showed that the average grain size is ∼150 µm after annealing, and then it is refined to 200–400 nm in the steady state for both ECAP and HPT, having all high angles of misorientation, demonstrating that the microstructure does not change at the saturated level with further straining. Further annealing of the microstructure after HPT at 433.15 K and 513.15 K revealed the microhardness decreased to about 1/2 of the hardness in the saturated level (∼160% of the annealed condition), grains sizes were in the range 1–4 µm for 433.15 K annealing and 2–6 µm for 513.15 K annealing, with many visible twins, and the electrical conductivity was recovered to 97%. It was shown in these studies that the proper combination of strength and electrical conductivity can be controlled by the amount of deformation and by annealing after HPT.

Fig. 1

Electrical resistivity and conductivity as function of equivalent strain (ε) from samples processed by ECAP, HPT³⁶⁾ and ARB.⁴⁴⁾

The results in Fig. 1 also include data of severely deformed oxygen-free Cu (OFC; 99.99%) by ARB cycles using 50% of reduction per cycle (1 to 8 cycles for equivalent strain up to 6.4) and produced a grain size of ∼200 nm, which maintained the electrical conductivity above ∼95% and improved the 0.2% proof stress to 350 MPa. Deoxidized low phosphorous Cu (DLP; Cu–0.02 mass%P–0.017 mass%Pb), and Cu containing iron (PMC90; Cu–1 mass%Fe–0.017 mass%Pb) were also processed to achieve grain refinement between 200–400 nm. The 0.2% proof stress increased exponentially with the number of ARB cycles 170 by a factor of 5, without the sacrifice of the electrical conductivity.⁴⁴⁾

A good understanding of metallurgical and electrical performance of the UFG structure can be attained through measuring accurately the electrical resistivity and grain sizes in a wide range of temperatures, the variation of both grain size, grain boundaries and temperature influence the electrical conductivity. To this purpose, Islamgaliev et al.⁴⁶⁾ processed Cu with a purity of 99.98% by HPT to a true strain of 7 and annealed at temperatures from 323.15 K to 1173.15 K. Measurements of electrical conductivity were performed in liquid nitrogen temperature by potentiometric (compensation) methods. Similar studies were conducted on Cu with a purity of 99.99% by Islamgaliev et al.^50,51) Figure 2 shows the behaviour of the relative electrical resistivity (ρ/ρ0) of HPT-processed Cu with respect to post-HPT annealing temperature. The data points were converted to match the relative electrical resistivity with respect to HPT-processed Cu (as opposed to monocrystalline Cu) and the grain size from temperature dependent curves, according to the data in Refs. 50, 51). Thus, the corresponding grain sizes to the annealing temperature can be associated from the top and bottom axes in Fig. 2. The relative electrical resistivity of severely deformed Cu decreases in a nonlinear fashion with the increase in annealing temperature, as shown by the solid line. The data points of ρ/ρ0 with respect to 1/d in Fig. 2 were fit by the model of Mayadas-Shatkez^46,50–53) developed for thin films. The model is based on the consideration of grain boundaries as potential barriers where conduction electrons are partially reflected during their motion within the crystal. According to the model, the relative electrical resistivity can be determined from the solution of the Boltzmann kinetic equation.⁴⁶⁾   

\begin{equation} \rho 0/\rho = 1 - (3/2)\alpha + 3\alpha^{2} - 3\alpha^{3} \ln (1 + \text{L}/\alpha), \end{equation} (6)

Where α = (L/d)R/(1 − R), ρ0 is the electrical resistivity of the single crystal, L is the free path of electrons in the single crystal, d is the average grain size, R is the coefficient of reflection of conduction electrons from grain boundaries. The experimental results plotted for ρ/ρ0 versus 1/d in Fig. 2 fit well with the calculations in this model for LR/(1 − R) = 0.097 µm. Before 450 K the reduction in electrical resistivity is only of ∼7–15%, which is mostly associated with recovery of lattice dislocations and grain boundary relaxation. The decrease is more pronounced from 450 K onwards, where static recrystallization and grain growth become the more prominent mechanisms, since the final electrical resistivity is similar to the typical value of CG Cu. This result was confirmed from TEM observations, XRD measurements of lattice parameters, microstrain and texture indices, as well as DSC studies. The microhardness was also stable up to 450 K, then decreased in a similar fashion as the relative electrical resistivity.⁵⁰⁾ Figure 2 also shows the data of relative electrical resistivity plotted with respect to the inverse average grain size (1/d) from Ref. 46). The dependence of grain size with temperature follows a similar trend of relative electrical resistivity in the range of 0.2 to 13 µm. An equiaxed grain structure of ∼100 nm average size with high angles of misorientation was reported for the as-HPT condition, with a dislocation density of ∼10¹⁴ m⁻².^50,51) The grain size of ∼140 nm was achieved by HPT followed by annealing at 323.15 K for 1 h and this value remained relatively stable up to 473.15 K for 4 h, which is consistent with the sharp decrease of electrical resistivity after this temperature. High-angle grain boundaries remained in this condition but with low dislocation densities within the grains. After annealing at higher temperatures, the average grain size was in the micrometer range. It can be said that the relative electrical resistivity is only slightly affected by grain sizes above ∼1 µm, but then sharply increases in the submicron region. The data is not enough to determine an accurate trend in the nanoscale, but the reports suggest the grain size does not decrease further by HPT. These studies show the UFG microstructure of Cu is thermally stable up to 450 K, which is consistent with the results in Ref. 36). Similar results on the thermal stability of electrical resistivity were obtained for high purity Ni (99.99%) processed by HPT.⁵²⁾

Fig. 2

Ratio of electrical resistivity of annealed Cu samples after HPT (ρ) relative to the electrical resistivity of as-HPT condition (ρ0) plotted against annealing temperature and inverse average grain size (1/d). Data extracted from Refs. 9, 46, 50, 51).

In Ref. 53) application of the Mayadas-Shatkez model coupled with the Grüneisen equation to the experimental electrical resistivity measurements, allowed to determine the grain boundary width of the UFG-Cu structure prepared by HPT (mean grain size 109 nm). Three contributions to the electrical resistivity of UFG-Cu were clearly established as main factors: (1) electrical resistivity of dislocations ρ_d, (2) electrical resistivity of grain boundaries ρ_b(T) and (3) electrical resistivity of the crystal lattice ρ_m(T). The sum of these terms as ρ(T) = ρ_d + ρ_b(T) + ρ_m(T), which is in the form consistent with Matthiessen’s rule, determines an approximation of the electrical resistivity of UFG-Cu and its temperature dependence. In Ref. 51), the electrical resistivity of Cu subjected to HPT at room temperature for N = 5 turns was measured using a DC four-probe method in the cryogenic state, which was previously evacuated and filled with helium. The electrical resistivity of UFG-Cu and CG-Cu at lower temperatures in Fig. 3, shows that with decreasing temperature the electrical resistivity of the samples decreases. The value of electrical resistivity for UFG Cu at 77.15 K is significantly different from the value of the CG state. The difference between the values of resistivity at 273.15 K and 293.15 K of the CG and UFG states is reduced with respect to 77.15 K. This result is consistent with the data in Fig. 2. The values of electrical resistivity of Cu at room temperature are comparable with values of electrical resistivity from GB.⁵⁴⁾ The increased electrical resistivity for UFG Cu was explained using a model of tunnel conductivity in solids.⁵⁵⁾ Based on this model, the width of the potential barrier where conduction electrons are scattered was estimated as ∼2.1 nm for Cu. This width is larger than the conventional thickness of grain boundary, which is below 1 nm (based on electron microscopy observation). Therefore, it is argued that the width of this zone includes an area near the boundary area in which conduction electrons are partially scattered during their motion in the material. This is attributed to the presence of an elastically distorted layer at near GB, and it leads to an increased scattering of conduction electrons.

Fig. 3

Temperature dependence of electrical resistivity (ρ): UFG Cu, CG Cu, NT Cu and NC Cu sample from 2 to 296.15 K. Data extracted from Refs. 51, 76, 77).

The electrical resistivity measurements in these studies can be used to determine the best strategies to attain high strength while maintaining the high electrical conductivity, even at elevated temperatures, by tailoring the processing route and heat treatment to produce fine grains and a favourable density of electro scattering sites. The latter is particularly important for cast Cu alloys without heat treatment, where a good combination of electrical conductivity and high strength is also sought. ECAP processing, using routes C and BC, of Cu–9Fe–1.2Cr cast alloy⁵⁶⁾ revealed that fragmentation of Fe–Cr dendrites occurred as opposed to elongation by using route A. Although drawing or extrusion may be more appropriate to generate high strength in this alloy because of the formation of a lamellar nanograined structure, ECAP for 8 passes produced a significant increase of the hardness from a more equiaxed structure with average grain size ∼0.2 µm, which is more efficient to retain higher electrical conductivity. The same strategy has been proven valid for age hardenable Cu alloys, since fine dispersion of either precipitates or disperse particles increases the mean free path of electrons in the matrix and can increase the thermal stability of the fine-grained microstructure.

2.2 Age-hardenable copper alloys

Interest in age-hardenable copper alloys has increased in the last few decades, particularly in designing the chemical composition to optimize the strength and electrical conductivity for industrial applications.⁶⁾ Due to the low solubility of Cr in Cu at room temperature, such properties of alloys of the Cu–Cr system can be enhanced by aging treatments.^9–11) On the other hand, the introduction of strain through SPD in these alloys gives rise to grain refinement and consequently higher strength when compared to coarse-grained Cu–Cr. However, plastic deformation produces the loss of some electrical conductivity. Fortunately, electrical conductivity can be restored significantly with an adequate post-processing heat treatment or processing at elevated temperatures. Therefore, the combination of SPD with selected heat treatments is an effective processing route to strengthen and recover electrical properties of UFG Cu–Cr,^57–61) Cu–Cr–Zr,^{40,49,62–71)} Cu–Cr–Hf^72–74) and Cu–Cr–Mg.⁷⁵⁾

Islamgaliev et al.⁵⁹⁾ produced a Cu–Cr alloy with an UFG structure and nanosized second phase precipitates. The initial discs were solution treated at 1323.15 K for 1 h and quenched in water. HPT was conducted under 6 GPa of pressure for N = 10 revolutions at ω = 1 rpm. The samples were either aged dynamically during HPT, or statically for 30 min after HPT in a range of temperatures from 293.15 K to 773.15 K. Figure 4 shows the electrical conductivity with respect to the heat treatment temperature for Cu–0.5%Cr samples: (1) processed by HPT at room temperature and subsequently aged for 30 min, (2) processed by HPT at 573.15 K and subsequently aged and (3) processed by HPT at the respective temperatures. After HPT processing and aging, the electrical conductivity increased with temperature, more so for the samples processed by HPT at elevated temperatures, followed by the sample processed by HPT at room temperature. The sample processed by HPT at 573.15 K and subsequently annealed achieved the highest value of electrical conductivity, and there was no effect of the post-HPT aging on the microhardness or electrical conductivity up to 773.15 K, which shows the microstructure was stable up to this temperature. Dynamic aging effect was observed in the sample processed by HPT at elevated temperatures after 573.15 K, suggesting the contribution of both grain refinement and precipitate particles, achieving a similar level of electrical conductivity than that of the samples processed at 573.15 K and aged. Meanwhile, the samples processed below 573.15 K had similar behaviour as the sample processed at room temperature and aged below that same temperature, as evidenced in the distributions shown in Fig. 4. The authors showed that the electrical conductivity was restored after 573.15 K due to the lowering of the alloying element content in the grains, with simultaneous precipitation and migration to grain boundaries. The TEM studies revealed a homogeneous UFG microstructure with grain size ∼200 nm without precipitates in the sample processed at RT. Meanwhile, the precipitation of the chromium-rich particles with size ∼10 nm appeared in the sample processed by HPT at 573.15 K.

Fig. 4

Temperature dependence of electrical conductivity for: Cu–0.5%Cr, Cu–0.75%Cr, Cu–0.18%Zr, Cu–0.9%Hf, Cu–0.5%Cr–0.08%Zr and Cu–0.7%Zr–0.9%Hf. Data extracted from Refs. 59, 60, 73).

Dobatkin et al.⁶⁰⁾ also explored the behaviour of Cu–Cr alloys by SPD and aging treatments based on the increasing Cr concentration: Cu–0.75%Cr, Cu–9.85%Cr and Cu–27%Cr. Two heat treatments were performed prior to HPT processing: annealing at 1273.15 K for 2 h and cooled in air (annealed samples) or in water (quenched samples). HPT processing was conducted at room temperature under 4 GPa of pressure for 5 revolutions at ω = 1 rpm. Combination of increasing Cr content and HPT processing decreased the electrical conductivity and provided higher levels of microhardness due to the reduction in average grain size, from ∼209 nm to ∼40 nm. To restore the electrical conductivity an aging treatment was performed in a range of temperature from 323.15 K to 873.15 K, which provided similar results than the previous work from Islamgaliev et al.⁵⁹⁾ Figure 4 shows that the electrical conductivity increased considerably above 523.15 K for the sample with 0.75%Cr. Meanwhile, higher content of Cr exhibited a similar trend but lower steady-state level of electrical conductivity. In all cases the increase of electrical conductivity after aging treatment is attributed to the grain boundary relaxation, decomposition of the Cu–Cr solid solution and Cr-rich particle precipitation.

Recently, Chu et al.⁶¹⁾ produced UFG Cu–0.5%Cr through the effective combination of 4 passes of ECAP, Deep Cryogenic Treatment for 12 h and Artificial Aging (ECAP+DCT+AA). The microstructure examination revealed nanosized banded grains (200–300 nm) with Cr-rich particles uniformly precipitated at grain boundaries and within grains. The tensile strength and electrical conductivity reached 587 MPa and 69% IACS after aging at 698.15 K for 1 h. Aksenov et al.⁵⁷⁾ also studied the Cu–Cr alloy with 0.2 and 1.1 mass% Cr by ECAP and then aged at 723.15 K. Cu–1.1%Cr experimented higher thermal stability than Cu–0.2%Cr, mainly attributed to banded grains with slightly finer nano-sized widths and higher degree of fragmentation of second phase particles. Cu–1.1%Cr microstructure also exhibited more non-equilibrium boundaries between bands, mostly dislocation cell walls and a higher fraction of second phase particles, which provided a UTS of 485 MPa and maximum electrical conductivity of 76% IACS.

SPD technologies change the kinetics of phase transformations and precipitation in both binary and ternary systems, such as Cu–Cr, Cu–Zr and Cu–Cr–Zr alloys, inducing nucleation of precipitates at lower temperatures than expected from their CG counterparts.⁵⁷⁾ In addition, Zr additions improve thermal stability, reduce the diffusion activity of Cr in the alloy and contribute to the formation of smaller particles.⁵⁷⁾ Vinogradov et al.^40,65) showed by detailed DSC studies in a Cu–0.44Cr–0.21Zr alloy, that the precipitation of Cr at 713.15 K and Cu₃Zr at 793.15 K became affected by ECAP processing, reversing the sequence of precipitation. The Cu₃Zr nucleated at 643.15 K and Cr at 698.15 K. In both cases, the temperatures were lower than in the non-deformed state, suggesting that the presence of high density of dislocations could promote the atomic mobility for nucleation of precipitates. Vinogradov et al.⁶⁵⁾ also processed Cu–0.36Cr and Cu–0.8Cr–0.05Zr alloys through ECAP and applied an isothermal aging. These group of alloys exhibited similar precipitation conditions after SPD processing than in the previous work. The electrical conductivity increased after the combination of ECAP and aging from an initial level of 30–40 IACS% to a final level of 80–93 IACS%.

Shangina et al.⁷²⁾ studied the thermal stability of UFG structures in Cu–0.7%Cr and Cu–0.7%Cr–0.9%Hf alloys through SPD processing and heat treatments. The initial microstructure was produced by two heat treatments prior to HPT processing: water quenching from 1173.15 K and air cooling from 873.15 K. HPT processing was carried out at room temperature under 4 GPa for 5 revolutions at ω = 1 rpm. Followed by annealing from 323 K to up to 823 K, which promoted the additional strengthening, recovery and partially restoration of electrical conductivity. The additional strengthening was caused by particles of Cu₅Hf in the structure. Dobatkin et al.⁷³⁾ performed further studies of similar alloys with the addition of low fraction of alloying elements: Cu–0.7%Cr, Cu–0.18%Zr, Cu–0.9%Hf, Cu–0.5%Cr–0.08%Zr and Cu–0.7%Cr–0.9%Hf. HPT was performed under 4 GPa, for 5 revolutions at ω = 1 rpm, as well as ECAP using route BC for 10 passes. Both HPT and ECAP were carried out at room temperature, followed by aging treatment to restore the electrical conductivity and promote the precipitation of particles from the solid solution to achieve higher levels of hardening. Figure 4 shows the increase in electrical conductivity with aging temperature in these alloys. As expected, Cu–0.7%Cr alloy remained thermally stable below 573.15 K. After aging at 573.15 K, the particles precipitate in the UFG structure, Meanwhile, Cu–0.18%Zr and Cu–0.9%Hf alloys exhibit precipitation of Cu₅Zr and Cu₅Hf particles, allowing higher degree of strengthening by suppressing grain growth of the UFG structures during aging, consistent with mechanisms described in Refs. 40, 59–74).

Li et al.⁷⁵⁾ explored electrical conductivity of Cu–0.28%Cr–0.19%Mg through Powder Metallurgy (PM), Cold Rolling (CR) and aging treatment. The sintered ingot was solution treated (ST) at 1223.15 K for 4 h and water quenched. The ST material was processed through CR with 50%, 70% and 80% reductions and then aged to restore electrical properties. The microstructure of the sample with 50% reduction and aged at 583.15 K was characterized by the presence of uniformly distributed nano-scale precipitates in a matrix with fine grains. A decrease in the diffusion activation energy after aging, promoted higher fraction of precipitated particles and achieved an improved UTS of 647 MPa and electrical conductivity of 85 IACS%.

2.3 Cu with fine grains and nanoscale twins

Cu strengthening mechanisms most of the times decrease in high proportion the electrical conductivity. However, a third approach for Cu strengthening without losing electrical conductivity is based on the combination of grain refinement and nanoscale twinning in the microstructure. Twin boundaries (TBs) block the dislocation motions and minimize the scattering of conduction electrons.^76–81) Lei Lu et al.⁷⁶⁾ produced high-purity Cu with nanoscale twins through pulsed electrodeposition (PED). This process provided grains with high density of fine twins at grain boundaries or triple junctions (TJs), as shown in Fig. 5(a) and 5(b). The grain size range was stablished between ∼100 nm to ∼1 µm with an average of ∼400 nm. Meanwhile, the strength was ∼1068 MPa UTS and ∼900 MPa YS, which is considerably higher than that of CG Cu.

Fig. 5

TEM micrographs with electron diffraction pattern insets showing (a) twin boundary (b) common twins pattern ATATA sequence in as-deposited Cu⁷⁶⁾ and (c) a twinned region in as-deposited NT-Cu-15.⁷⁷⁾

Figure 3 shows the electrical resistivity (ρ) measurements of Cu with nanoscale twins (NT Cu) and coarse-grained Cu (CG Cu). The resistivity showed a thermal dependency with a linear relation for values of temperature higher than 70 K. The NT Cu exhibited a behaviour close to the CG microstructure, whereas the measurements at room temperature showed an electrical resistivity of ρ_293K = 1.75 ± 0.02 × 10⁻⁸ Ωm, in comparison with 1.69 ± 0.02 × 10⁻⁸ Ωm for the CG one. The increase in resistivity of NT Cu is attributed to the contributions of grain boundaries and twin boundaries, as stated by Matthiessen’s rule. Thus, the electrical conductivity at room temperature was 96.9 ± 1.1 IACS%. The temperature dependence of electrical conductivity for sputtered nanocrystalline Cu (NC Cu), with a grain size of ∼15 nm, was also plotted in Fig. 3.

Chen et al.⁷⁷⁾ also prepared high-purity ultrafine-grained copper with nanoscale growth twins via PED, and different concentrations of twins were set by the adjustment of the process parameters. Besides, SPD was performed via CR processing at room temperature by a twin-roller apparatus with 50 mm of diameter and deform the Cu foils for a maximum strain of ∼40%. The average grain size attained was 500 nm and growth twins with average of twin lamellar spacing: 15 nm (NT-Cu-15), 35 nm (NT-Cu-35) and 90 nm (NT-Cu-90). Figure 5(c) shows a TEM micrograph of a twinned region with an inhomogeneous SAED pattern. On the other hand, the electrical resistivity trend plotted in Fig. 3 exhibited similar behaviour than the nanotwins produced by Lu et al.⁷⁶⁾ However, the increase of twin spacing from 15 nm to 90 nm produces a remarkable increase in the electrical resistivity from 1.75 × 10⁻⁸ to 2.12 × 10⁻⁸ Ωm, as the temperature increases up to RT. Therefore, the conductivity of NT-Cu-15 was ∼97 IACS%, similar to the CG. After CR the electrical conductivity at 293.15 K decreased in comparison with the initial as-deposited condition: from 100 to 98 IACS% in CG Cu, from 97 to 88 IACS% in NT-Cu-15, from 88 to 82 IACS% in NT-Cu-35 and from 80 to 77 IACS% for NT-Cu-90, as plotted in Fig. 3.

Due to the exceptional combination of electrical and mechanical properties achieved by the Cu NTs, Zhang et al.⁷⁸⁾ produced them in bulk Cu foils through magnetron sputtering technique. To accomplish a nanoscale growth twin with an average twin spacing of ∼5 nm, a thick cooper layer (20 µm) was deposited on Si (100). The average of columnar grain sizes increased up to 80 nm from an initial value of 43 nm using deposition rates in the range 1.8 to 3.0 nm/s. For this particular case, Cu foils achieved a tensile strength of 1.2 GPa, reported as the maximum ever achieved in NC Cu.

The strengthening provided by the twin boundaries established the basis for developing new combination of processes. Zhang et al.⁷⁹⁾ synthesized bulk nanograined Cu embedded with nanoscale twins through DPD at liquid nitrogen temperature (LNT). The DPD Cu samples were characterized by the presence of nanograins with an average size of ∼66 nm and high density of nanoscale twins, with an average lamellar thickness of ∼44 nm. The DPD Cu samples achieved an electrical conductivity of 95 IACS% and yield strength of 610 MPa at room temperature. To determine the effect of nanotwins over the mechanical and electrical properties, the DPD Cu samples were subjected to CR at ambient temperature with a rolling strain of 50%. After CR, the volume fraction of nanotwins was reduced, which decreased the yield strength below 545 MPa and increased the electrical conductivity up to 97 IACS%. Similar work was performed by Sun et al.,⁸⁰⁾ who produced bulk nanostructured Cu–Cr–Zr disks with a thickness of 2.3 mm by LNT-DPD with a strain ε = 2. The microstructure was composed by the mixture of high density of nanotwins with an average thickness of 25 ± 14 nm and nanograins with an average grain size of 47 ± 19 nm. A tensile strength of 700 MPa and electrical conductivity of 78.5 IACS%, could be obtained by processing without additional heat treatment.

2.4 Cu composites

Nanostructuring Cu composites through SPD also allowed synergistic increase of mechanical and electrical properties. Islamgaliev et al.⁸²⁾ produced a metal matrix nanocomposite of Cu–0.5 mass% Al₂O₃ through HPT processing under 2.8 GPa and 6 GPa. The microstructure from samples processed at 6 GPa was composed by small grains of Cu with a size of ∼80 nm and highly dispersed particles of Al₂O₃ with an average size of ∼20 nm. The tensile strength reached 680 MPa. For samples processed at 2.8 GPa grain size was of ∼200 nm and electrical conductivity was ∼86 IACS% at room temperature. After annealing at 673.15 K for 1 h, the electrical conductivity increased to 93 IACS%.

Xiong et al.⁸³⁾ developed a strengthened super-aligned carbon nanotube reinforced Cu matrix (SACNT/Cu) through rolling at room temperature. The combination of 2.5% vol% SACNT in pure Cu and 40% of rolling strain provide an enhanced electrical conductivity of ∼98 IACS% and a tensile strength of 470 MPa, but the efficiency of the SACNT addition was discussed since the electrical conductivity slightly decreased as the SACNT fraction increased. After rolling the electrical conductivity was partially restored due to the formation of ultra-low dislocation density regions strengthened by SACNTs which suppressed grain rotation, within the high-dislocation density regions in the Cu matrix: pure Cu had 20% higher dislocation density than Cu–2.5%SACNT.

Figure 6 shows the summary plot of tensile strength (YS or UTS) and electrical conductivity of several Cu alloys processed discussed in this work. Fine-grained Cu alloys and composites exhibit an equilibrated combination of mechanical strength at least higher than 300 MPa and electrical conductivity above 50 IACS%, which is the current baseline for conventional thermomechanical treatments, as indicated by the lower shaded area in Fig. 6.⁶⁾ On the other hand, conventionally processed alloys show strength higher than 400 MPa, provided by the alloying elements Cr, Ni, Fe and Si, but lower electrical conductivity below 70 IACS%. The Cu–Al₂O₃ and Cu-SACNT nanocomposites achieved attractive combinations of yield strength and electrical conductivity. Cu–Cr–X based alloys show similar strength with enhanced electrical conductivity after SPD complemented with an aging treatment, as detailed in Fig. 4. The introduction of nanotwins in UFG Cu by DPD or laser deposition seem to attain the best combination of strength and conductivity. These latter two processes can provide potential high-performance strength above 600 MPa and electrical conductivity above 80 IACS%, as indicated by the upper shaded portion of Fig. 6.

Fig. 6

Electrical conductivity vs. yield strength of severely deformed Cu and Cu alloys. Data extracted from Refs. 44, 76, 78, 80, 82, 83).

3. Aluminium Alloys

The increasing demand of lightweight conductor materials with high mechanical and electrical performance projects Al-based alloys as an economical option to replace Cu alloys in industrial applications, such as structural components in the automobile industry,^84–86) parts with good corrosion resistance in structures and electronic systems,⁸⁶⁾ and more recently commercial anodes for Al-air batteries,⁸⁷⁾ hydrogen and electricity production.⁸⁸⁾ The versatility of Al alloys through the addition of alloying elements is a good option to enhance the mechanical properties. However solid solution and dislocation strengthening sacrifice the electrical conductivity in the Al matrix in ambient conditions, especially in the case of miscible elements, when saturation of such defects is reached. In the case of the immiscible elements, the morphology of secondary phases can also produce the loss of electrical conductivity.^89–93) The expected service properties of Al-alloys for high-performance conductive purposes are at least a UTS of 300 MPa and electrical conductivity of 52 IACS%, which are already difficult to achieve by conventional thermomechanical methods. For example, pure Al (99.6%) after hot extrusion and multipass cold drawing achieved an electrical conductivity of 62.59 IACS% and UTS 204 MPa.⁴⁾ On the other hand, alloys of Al–Mg–Si system processed by solution treatment, cold drawing and artificial aging can exhibit a UTS in the range of 255–330 MPa and electrical conductivity of 57.5 and 52.5% IACS%, respectively.⁹⁴⁾ For instance, the fabrication of Al 6201RE wire 3.35 mm in diameter produced through hot forging, in combination with artificial aging at 448.15 K for 4 h, followed by 10 passes cold drawing produced UTS of 352.3 MPa and electrical conductivity of 56.0 IACS%. The addition of RE elements and this novel processing route resulted in properties closer to a high-performance conductor.⁴⁾ The introduction of other immiscible elements such as increasing the Sc fraction in Al–X%Sc–0.2Zr system, allowed to attain UTS up to ∼180 MPa with a corresponding decrease of the electrical conductivity to 59.4 IACS%. The precipitation of Al₃Sc particles by artificial aging at 553.15 K also provided higher thermal stability of the microstructure.⁴⁾

Several techniques of SPD combined with post-processing heat treatment become an effective strategy to produce UFG Al alloys with high strength and high fraction of the initial electrical conductivity with shorter aging times.^90–105) Recent studies provide attractive results of Al alloys with addition of Mg, Si, Cu, Fe, Zr, Zn and rare-earth elements (RE) such as La and Ce. Figure 7 shows the electrical conductivity of Al alloys with different combinations of processing and heat treatment, with respect to their chemical composition. A general decreasing trend in the electrical conductivity can be observed with increasing the total fraction of alloying elements, especially higher than 2 mass%. After SPD the initial states exhibit conductivities as low as ∼40 IACS%, but that can be increased by aging to ∼55 IACS% or more, while keeping the strength. The following sections show important results in Al alloys developed to produce balanced strengthening with competitive electrical properties divided in two groups: (a) alloys with miscible components; mainly Al–Mg–Si and others^31,90–100) and (b) alloys with immiscible components, such as Al–Fe,^{38,101–107)} Al-RE^108–110) and Al–Zr.^111–116)

Fig. 7

Electrical conductivity of Al alloys processed by SPD as function of total fraction of alloying elements.

3.1 Aluminium alloys with miscible components

Bobruk et al.⁹³⁾ subjected Al 6060 (Al–0.55Mg–0.50Si–0.10Cu–0.10Mn–0.05Cr–0.30Fe) and Al 6063 (0.60Mg–0.45Si–0.24Fe) to solution treating at 803.15 K during 2 h and quenched in water. The Al 6060 alloy was subjected to HPT processing at room temperature and at 453.15 K for grain refinement down to an average size of ∼180 nm and ∼350 nm, respectively. During HPT processing at 453.15 K higher fraction of a globular metastable β-phase (Mg₂Si) with grain size ∼20–40 nm precipitated as result of decomposition of the solid solution and dynamic strain aging (DSA). DSA at 453.15 K provided a balanced combination of properties in comparison with the solution treated state, increasing the electrical conductivity up to 58.1 IACS% and the UTS to ∼347 MPa. The Al 6063 was processed through ECAP with parallel channels (ECAP-PC) at 373.15 K complemented by an aging treatment at 403.15 K, the resultant UFG-structure was of ∼500 nm average grain size. Metastable β-phase precipitates ∼10 nm in size were also observed in areas near boundaries and inside the grains. Aging treatment allowed the enhancement of both UTS and electrical conductivity to ∼308 MPa and ∼58.6 IACS%.

Murashkin et al.³⁷⁾ solution treated Al–0.60Mg–0.45Si alloy at 808.15 K, then processed by ECAP-PC at 373.15 K with an angle 110° with accumulative strain of ∼1.6 for 1 pass, 3.2 for 2 passes and 9.6 for 6 passes and then quenched in water. After the last pass the samples were subjected to artificially aging for 24 h at 403.15 K. Decomposition of solid solution occurred after 2 passes of ECAP-PC, leading to precipitation of a spherical β′ phase with higher Mg:Si fraction at the grain boundaries, as shown in Fig. 8(a). The artificial aging allowed the precipitation of needle-type β′′ phase in all the treated samples. Both precipitates enhanced the electrical conductivity due to the existent coherency between β′, β′′ and the matrix. Figure 7 shows the effect of the aging treatment on the enhancement of the electrical conductivity, which increased to 52.93, 53.28 and 56.17 IACS% for samples aged after 1, 2 and 6 passes, respectively. The corresponding UTS values reached 241, 310 and 289 MPa, respectively.

Fig. 8

Spherical second-phase precipitates formed in Al–Mg–Si alloys: (a) AA6063 after ECAP-PC at 100°C for 6 passes and aged,³⁷⁾ (b) AA6201 after HPT at 180°C for 20 revolutions⁹⁴⁾ and (c) AA6101 after HE to Φ3 mm and aged at 180°C with (d) STEM image of precipitate and EDS spectra compared with Al matrix.⁹¹⁾

Valiev et al.⁹⁴⁾ carried out HPT in Al 6201 for N = 1 revolution at room temperature followed by N = 20 revolutions at elevated temperatures of 403.15 K, 453.15 K and 503.15 K. Dynamic aging occurred during HPT processing, producing metastable nano-sized precipitates, as can be observed in Fig. 8(b). This experimental strategy provided an increase in electrical conductivity with respect to the temperature of processing: 55.6 IACS% for 403.15 K, 58.4 IACS% for 453.15 K and 59.0 IACS% for 503.15 K.

The formation of nanostructures through SPD has become an outstanding metallurgical path to increase electrical and mechanical properties in Al alloys for replacing Cu alloys¹¹⁷⁾ and thereby the reduction of weight and costs in fabrication of electrical wires. Majchrowicz et al.⁹¹⁾ explored the fabrication of long wires made of UFG Al 6101. A solution heat treatment at 783.15 K followed by HE to reductions of ϕ10 mm (HE10), ϕ3 mm (HE3) and ϕ4 mm (HE4). An aging treatment at 453.15 K was also applied to all samples for 1–24 h. Figure 8(c) shows spherical particles β′ precipitated at the grain boundaries after HE3 and aging at 453.15 K for 2 h. The HE4 after 7 h of aging attained strength of 329 MPa and conductivity of 58.04 IACS%. Pakiela et al.⁹⁵⁾ developed similar studies in rods of Al 6101 and AA6201 with an initial diameter of 20 mm. The samples were annealed for 1 h at 783.15 K and quenched in water. Followed by HE at RT for two passes (total true strain 1.39) with a final diameter of the extruded wires of 5 mm (HE5). However, the electrical conductivity enhancement was poor for AA6101 and not effective for AA6201 due to the absence of heat treatment after HE. The results of these studies are also plotted in Fig. 7.

Meagher et al.⁹⁶⁾ processed Al 6101 via continuous ECAP. The experimental work was developed in two regimens of temperature: the first pass was conducted with an initial temperature lower than 323.15 K and the second one in a range from 393.15 K to 423.15 K. For the final state of the microstructure, an aging at 463.15 K during 8 h was applied to the processed billets. After ECAP the microstructure exhibited high density of dislocations and grain refinement from ∼2.2 µm of the as-received wire to ∼0.89 µm. After aging treatment, the strength increased to 238 MPa and electrical conductivity to 57 IACS%.

Murashkin et al.⁹⁸⁾ fabricated Al 6101 alloy wires using ECAP-Conform (ECAP-C) to produce long-length UFG billets and then drawing to ϕ3.2 mm. The initial state was solution treated at 823.15 K to produce a grain size of ∼10 µm, followed by the ECAP-C processing for 6 cycles at 403.15 K and artificial aging (AA) in two different regimes: 443.15 K for 1–6 h or 463.15 K for 1–12. After ECAP-C at 403.15 K, the UFG structure with grain size 400–600 nm had spherical particles 3–20 nm in size dispersed within grains, corresponding to β and metastable β′ phases as result of dynamic aging and decomposition of supersaturated solid solution, as mentioned in similar works described in Refs. 91–100). The electrical conductivity in the initial state was 50.4 IACS%, after ECAP-C + 443.15 K aging was 57.1 IACS%, after ECAP-C + 463.15 K aging was 57.4 IACS% and in the final wire 56.4 IACS%. The UTS increased from 195 MPa to 288 MPa, 304 MPa and 364 MPa, respectively.

3.2 Aluminium alloys with immiscible components

The Al–Fe system has been widely studied due to the capability to produce high strength and conductive alloys by tailoring the size and distribution of several nano-sized Al–Fe intermetallic particles that can form as function of Fe fraction.^{38,90,102–107,118–120)}

Cubero-Sesin et al.¹⁰²⁾ produced micro-sized wires of Al–2%Fe composed by nanosized discontinuous lamellae of Al₃Fe, very fine fibers of Al₆Fe and UFG α-Al matrix with a balanced mechanical and electrical properties. As-cast Al–2%Fe in the shape of 20 mm outer diameter and 3 mm width rings was processed by HPT for N = 1 at R.T under 3 GPa, followed by wire drawing and an aging treatment at 473.15 K in air and ice water quenching. The minimum diameter attained for wire drawing was 0.008 mm. Such treatment provided an effective increase in hardness of ∼75 HV and electrical conductivity up to 54 IACS% due to grain refinement and fragmentation of intermetallic after aging. Further studies accomplished by Cubero-Sesin et al.^38,90) on Al–2%Fe and Al–4%Fe described more detailed results from HPT at room temperature on discs and rings, combined with aging treatments at 473.15 K. Al–2%Fe alloys showed extraordinary balance of strengthening with UTS in the range 450 to 650 MPa for samples processed between 10 and 75 revolutions, with electrical conductivity around 55 to 52 IACS%, respectively. The thermal stability of this material was very high, with no significant changes for at least 12 h at 473.15 K.

More recent studies in Al–Fe system were carried out by Medvedev et al.^104–107) In Ref. 104) Al–0.5Fe and Al–2.5Fe (mass%) were prepared by electromagnetic casting (EMC), the microstructure was mainly composed of cell-like Al rich phase and Al₂Fe and Al₁₃Fe₄ intermetallics in the cell walls, with Al–2.5Fe showing a higher fraction of intermetallic particles. Both alloys were processed by ECAP (4 passes) and ECAP+CR (85% reduction). The samples were also artificially aged at 503.15 K and 553.15 K. Al–0.5Fe ECAP+CR+AA at 503.15 K attained UTS of 175 MPa and EC of 59.6 IACS%, whereas Al–2.5Fe a UTS of 275 MPa and EC of 49.3%. The thermal stability was verified up to 503.15 K, since at 553.15 K the strength decreased ∼20% with a slight increase in conductivity. In Ref. 105) as-cast Al–2Fe and Al–4Fe (mass%) were processed by HPT for 20 revolutions. The Al–2Fe had UTS of 649 MPa and EC of 40.4 IACS%, which was more favourable in terms of ductility, due to the more refined particle morphology after HPT.

In Ref. 106) the HPT-processed Al–2Fe (N = 20 rev.) was further studied by (1) annealing at 473.15 K for 8 h and (2) HPT processing at 473.15 K (N = 5 rev.) The electrical conductivities attained for each case were 49.3 IACS% and 52.3 IACS%, with UTS of 335 MPa and 327 MPa, respectively. The latter was a more effective method to produce a good balance of properties, due to an accelerated decomposition of supersaturated solid solution by dynamic aging, without much loss in the strength. In Ref. 107) an optimized composition of Al–1.7 mass%Fe alloy was prepared again by EMC+ECAP+CR and additional aging as in Ref. 104). The as-processed sample had UTS of 298 MPa and EC of 51.3 IACS%, whereas after aging at 503.15 K and 553.15 K, the samples had a UTS of 306 MPa and 235 MPa, with conductivities of 53.2 IACS% and 55.3 IACS%, respectively. In addition, the thermal stability was very similar to case in Ref. 104). Figure 7 shows that alloys of the Al–Fe system performed well in comparison with other alloys.

Murashkin et al.¹⁰⁸⁾ and Medvedev et al.¹⁰⁹⁾ processed Al–X mass% RE alloys, by HPT at room temperature for 20 revolutions under 6 GPa, in combination with annealing treatments in several regimes: 553.15 K and 673.15 K for Al-8.5RE¹⁰⁸⁾ and 503.15 K, 553.15 K and 673.15 K for Al-2.5RE, Al-4.5RE and Al-8.5RE.¹⁰⁹⁾ The formation of spherical Al₁₁RE₃ nanoparticles embedded in the UFG Al-matrix was observed by TEM, as shown by red arrows in Fig. 9. After the annealing, the intermetallic particles did not grow significantly, and the density of dislocations was reduced. After HPT processing the samples exhibited a considerable increase in the mechanical properties in comparison with C.P Al. The combination of high density of defects and the increase of RE concentration decreased the electrical conductivity significantly, as observed in Fig. 10. The best combination of electrical conductivity and mechanical strength after HPT occurred with additions of RE between 3.5–4.5 mass% and subsequent annealing between 523.15–553.15 K for 1 h. The annealing was successful to partially restore the electrical conductivity from 56.6 to 60.2 IACS% in Al-2.5RE,¹⁰⁸⁾ 45.5 to 57.4 IACS% in Al-4.5RE¹⁰⁹⁾ and 39.7 to 52.4 IACS% in Al-8.5RE,^108,109) as shown also in Fig. 7.

Fig. 9

Microstructure after HPT processing and subsequent annealing at 200°C and 400°C for 1 h: (a), (b) Al-4.5RE and (c), (d) Al-8.5RE.¹⁰⁹⁾

Fig. 10

Electrical conductivity of severely deformed Al alloys with respect to the UTS.

Mavlyutov et al.¹¹¹⁾ and Orlova et al.¹¹²⁾ processed Al–0.4Zr by HPT at room temperature under 6 GPa for 10 revolutions. After HPT aging treatment at 363.15 K, 423.15 K, 473.15 K, 503.15 K,¹¹²⁾ and 673.15 K for 1 h¹¹¹⁾ and at 503.15 K for 3 h¹¹²⁾ to restore the electrical conductivity due to the formation of a uniform UFG structure with Al₃Zr nanoparticle dispersion promotes higher strength and lower conductivity. The maximum strength was achieved after aging for 3 h. The aging provided relaxation of HAGBs with lower dislocation density, which allowed a balanced combination of microhardness and electrical conductivity in the ranges 508–530 HV and 54.9–58.8 IACS%. The rest of the results reported for this alloy are also plotted in Fig. 7. Figure 10 summarizes the magnitudes of the UTS and electrical conductivity attained by several of the Al alloys compiled in this review, as well as the range of properties desired for industrial application, and the comparison with alloys processed by conventional thermo-mechanical treatment. Special attention can be drawn to be cases closer to the upper-right hand corner of the plot, particularly the Al–Mg–Si series for miscible systems, and the Al–Fe for immiscible systems. HPT can provide higher levels of mechanical strengthening, but with more significant reduction of electrical conductivity. Thus, dynamic aging or post-processing aging must be used to reduce the density of lattice scattering sites, especially solute atoms and dislocations. Fine precipitates can strengthen the conducting Al matrix and provide thermal stability. The continuous processing can be an attractive technique to produce wire material with balanced properties for industrial scale up.¹¹⁷⁾

4. Innovation Potential of UFG Cu and Al Alloys for Electrical Conductors

SPD methods provide an effective path for fine-structuring a wide range of materials including the capability for powder consolidation and alloy design with multiple applications in several research fields and industry.^30,118,119) Electrical and mechanical properties of Cu and Al alloys can be balanced by factors such as thermal-mechanical processing, strengthening mechanisms,¹²⁰⁾ alloy thermodynamics (precipitation of second phases accelerate grain refinement),^14,122) grain boundary phenomena^{30,103,120–133)} and microstructural evolution.^14,30,128) All these factors can be partially controlled by SPD processing complemented with aging treatments, or even, dynamic aging to produce high conductive and strengthened bulk-nanostructured alloy.^30,134) However, the potential application for commercialization of UFG materials is limited by the fabrication capabilities and the sample size. Therefore, the SPD community has been constantly in the development of processing strategies, new methods and technologies for scaling up sample size.¹³⁰⁾ Recent studies^{123–127,131–133)} describe SPD arrangements for scaling up sample dimensions, experimental processing and thermodynamic simulations to optimize continuous deformation processes and increase the quantity/volume of nanostructured material. ECAP-Conform (ECAP-C) is one of the most promising technologies to produce high strengthened rods and wires in continuous extrusion.³³⁾ Murashkin et al.⁹⁸⁾ produced wires of Al–Mg–Si with 3.2 mm diameter and UTS of 364 MPa and EC of 56.4 IACS%. Cubero-Sesin et al.¹⁰²⁾ produced micro-wires 0.008 mm in diameter, by cold drawing ring-HPT processed samples of Al–2%Fe with an expected strength of 225 MPa and conductivity of 54 IACS%. On the other hand, Masuda et al.¹⁰³⁾ prepared nanostructured wires of Al–2%Fe with 50 mm length and 2.8–3 mm diameter by Continuous High-Pressure Torsion (C-HPT), which was firstly developed by Edalati et al.¹²⁵⁾ C-HPT is designed with symmetrical grooves in the upper anvil after passing through the inlet and outlet guiding holes. This continuous processing produced an UFG Al–2Fe wire with 400 nm grain size and UTS of 280 MPa. In addition, C-HPT capabilities can be set from one to four wires simultaneously. Fujioka et al.²¹⁾ developed High-Pressure Sliding (HPS) to process simultaneously two sheets of 5 mm × 100 mm and 0.8 mm thickness. Other strategies that have been proposed include scaling up the machine capacity (pressure that can enable increases in sample size for HPT, C-HPT and HPS), the use of larger ring samples or incremental feeding (IF) in HPT and HPS for sheet products.¹³⁵⁾ In addition, Planar High-Pressure Torsion (P-HPT)¹³⁶⁾ was developed for processing metal sheets of 220 mm × 110 mm. These methods are based on the common HPT or HPS setup with several anvil modifications and continuous sample feeding mechanisms.

5. Concluding Summary

In this work, several publications were presented to explore processing routes and deformation mechanisms employed to produce UFG and NS Cu and Al alloys, with a balanced trade-off relationship between mechanical and electrical properties.

The first approach demonstrates that SPD processing can produce strengthening via grain refinement to ultrafine or nano-levels, which in most cases decreases the electrical conductivity. Thermodynamic interactions at grain boundaries, grain size and the thermal dependence of the electrical resistivity were stablished as the main factors for variation of both mechanical and electrical properties in Cu. Thus, annealing or aging treatments after SPD recover high fraction of electrical conductivity. For example, high-purity Cu (99.99%) processed by HPT and annealed at 433.15 K achieved a microhardness increase of ∼160% and electrical conductivity ∼97 IACS% with a microstructure free of dislocations with a grain size 1–4 µm.

The second approach, as described for Cu–Cr and Cu–Cr–X alloy system, two processing strategies could be followed: (1) SPD at room temperature and post processing aging or annealing and (2) SPD processing at high temperature. Both can achieve acceptable strength - electrical conductivity relationship, associated with precipitation of metastable phases and simultaneous control of grain size and lattice defects. The same mechanism was proven effective in Al alloys with miscible and immiscible alloying elements. Static or dynamic aging increases electrical conductivity by grain boundary relaxation, decomposition of solid solution and precipitation of metastable particles, usually with enhanced kinetics due to the prior severe deformation of the microstructure. On the other hand, the nanostructured Al-alloys are guided mainly by the influence of alloying elements, their chemical nature (miscible/immiscible) and the characteristics of secondary metastable phases (coherent/incoherent). A decreasing trend of electrical conductivity was also observed as the total fraction of alloying components increase to contents higher than 2 mass%. Therefore, the immiscible components provided higher magnitudes of strengthening, but can be controlled aided by the high density of nucleation sites provided by SPD. The attractive combination of properties exhibited by the bulk nanostructured Al alloys, with a wide range of strength from 300 to 680 MPa and electrical conductivity as high as ∼55 IACS% increases the number of applications for Al alloys in the electrical conductor industries, where weight and cost are of concern.

The third approach described the combination of grain refinement and nanoscale twinning in the microstructure. The DPD Cu samples achieved an enhancement of tensile strength close to ∼610 MPa and conductivity of 95 IACS%. The formation of nanotwins in the material block the dislocation motions and minimize the scattering of conduction electrons. However, the industrial scalability of this approach, along with the thermal stability of this material, suggest the previous two approaches have higher probability of implementation.

Acknowledgements

The authors would like to thank Vicerrectoría de Investigacion y Extension, Instituto Tecnológico de Costa Rica for financial support from grant VIE-CF-1490033. JGH would like to acknowledge Direccion de Posgrado, Instituto Tecnológico de Costa Rica for a doctoral program scholarship.

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