Influence of the Chromium Content in Low-Alloyed Cu–Cr Alloys on the Structural Changes, Phase Transformations and Properties in Equal-Channel Angular Pressing

Abstract

The quantitative concentration of alloying elements in low-alloyed copper alloys is an important factor in forming electrical and mechanical characteristics. It is known that severe plastic deformation is accompanied by both a substantial refinement of the structure and changes in the kinetics of phase transformations during the deformation and the post-deformation thermal treatment. This paper presents the results of a comparative analysis of the Cu–0.2Cr and Cu–1.1Cr alloys subjected to equal-channel angular pressing at room temperature. The analysis was performed for the grain structure, solid solution, and second-phase particles using transmission electron microscopy, scanning electron microscopy, X-ray crystal analysis, and the small-angle diffraction method. It was found that the level of structure refinement and mechanical characteristics after equal-channel angular pressing was almost the same for both studied alloys. Post-deformation aging of the Cu–0.2Cr alloy leads to the development of polygonization and re-crystallization within it. The aging of the Cu–1.1Cr alloy shows a better thermal stability than that of the Cu–0.2Cr alloy. In the Cu–1.1Cr alloy, after aging, in comparison with Cu–0.2Cr, a denser-packed ensemble of fine particles with an average size of 54 ± 2 nm is formed. In this case, the average size of fragments is 270 ± 15 nm and the ultimate tensile strength reaches 485 MPa.

1. Introduction

Copper has a high electrical conductivity but a low strength, which limits its area of application. One of the approaches to increase the material strength is its alloying. However, this leads to a reduction of electrical conductivity. Low-alloyed precipitation-hardening bronzes of Cu–Cr system demonstrate a good combination of strength and conductivity [1,2,3,4,5]. It has been found that the Zr content on the order of several hundredths has a beneficial effect on the strength of a Cu–Cr alloy, but the electrical conductivity of an alloy is somewhat reduced. It has also been found that zirconium reduces the diffusion activity of chromium in an alloy, which contributes to the formation of smaller dispersed particles. The strength of a ternary system alloy can exceed 500 MPa, depending on the method of deformation processing and the accumulated strain. In this case, the electrical conductivity is about 80% IACS (International Annealed Copper Standard) [2,5,6,7,8,9,10].

Another efficient method of hardening is deformation treatment that results in a substantially increased density of dislocations and a deep refinement of the initial structure. The methods of severe plastic deformation (SPD) developed during the recent decades have shown that due to the refinement of the structure to the ultra-fine-grained (UFG) and nano scale, strength of metals and alloys could substantially exceed the strength achieved by conventional methods of cold deformation [11,12,13]. In the papers [14,15], the methods of sideways and forward extrusion are presented, which make it possible to produce billets of the required profile with a fine-grained structure where the proportion of high-angle boundaries is over 40%. For producing long-length products, the methods of equal-channel angular pressing (ECAP) combined with Conform and Multi-ECAP-Conform [16,17] have been developed, which make it possible to produce rods with an ultra-fine-grained structure of about 400–500 nm.

A combination of the precipitation and deformation methods of hardening allows for a more than two-fold increase in the strength of precipitation-hardening copper alloys [1,6,18,19,20,21]. Initially, hardening occurs due to an increased number of defects and micro-structure refinement during the deformation treatment and then by means of precipitation hardening during metal aging. When using SPD technologies, the kinetics of phase transformations changes, which leads to a high level of functional properties in low-alloyed copper alloys. The authors of [21,22,23,24,25,26,27] found that under SPD conditions, the second-phase particles could undergo the process of deformation-induced dissolution. Thereby, under SPD conditions, there can be additional supersaturation of the solid solution (SS), which increases the lifetime of precipitation hardening in the case of subsequent post-deformation heat treatment.

Despite the fact that the addition of Zr improves thermal stability and somewhat increases the strength of a chromium bronze, a Cu–Cr alloy has a couple of advantages over the ternary system. For example, the production of alloying compositions and the production of a ternary-system bronze is more complex and expensive than the production of a binary-system alloy. In addition, there is a large problem of homogenization of these bronzes in the manufacture of long-length products (>> 100 m), for example, for the production of overhead wires for railway lines. Since Zr reduces the diffusion activity of chromium, this aspect can play an important role in the formation of structure and properties.

Therefore, scientists [1,6,13,19,21,22,28,29,30] paid great attention to investigating the structure and the kinetics of phase transformations in copper alloys with chromium concentration of 0.64–0.74 wt.% at the threshold of maximum solubility and at 960–1070 °C. These papers study the processes of the mutual influence of the dislocation structure and second-phase particles during deformation. For example, it is found that dislocations are the predominant location where particles are generated, but the dispersed nanoparticles released are obstacles to dislocation migration, which decelerates relaxation of the deformed structure [6,7,13,21]. The dimensions and nature of the distribution of second-phase particles have a substantial effect on the mechanical properties of the material.

It should be noted that these papers show no complete dissolution of second-phase particles during high-temperature treatment.

In connection with the change in the kinetics of phase transformations under SPD conditions, studies of copper alloys with a chromium content exceeding both the dissolution limit and the eutectic point become relevant. Some papers present the study results for the effects of the weight fraction of alloying elements on the strength of copper alloys. For example, Martynenko et al. in [31] studied the specifics of structural and phase transformations in a Cu–Cr alloy with a chromium concentration over 7 wt.%. It was shown that after ECAP processing, an ultra-fine-grained structure was formed having a grain size of ~200 nm, but the hardening annealing (aging) showed a high degree of non-equilibrium of such structure since grain growth was observed along with the electrical conductivity increase to 70% IACS. In this case, the processes of relaxation of the structure occur faster than the process of decomposition of the supersaturated solid solution, thus the efficiency of decomposition of the solid solution and, accordingly, the level of hardening decreases.

At the same time, insufficient attention was paid to the effect of SPD processing on chromium bronzes with a chromium content significantly lower than the solubility limit, when a deep state of solid solution is formed at a quenching temperature (~ 1000 °C). In addition, reducing the concentration of the alloying element in the alloy can have a positive economic effect. Moreover, the behavior under SPD conditions of alloys with a chromium content close to the eutectic point (~ 1.1–1.2 wt.%) has not been sufficiently studied. Despite the fact that a complete dissolution will not occur, the fracture of the initial particles and a change in the state of the phase boundaries under SPD conditions will effectively affect the structure formation and phase transformations. At the same time, a moderate chromium content (<1.2 wt%) will not lead to a significant decrease in the electrical conductivity of the material.

The above information shows that the weight content of alloying elements is the most important value in creating the most reasonable electrical and mechanical properties of low-alloyed copper alloys after mechanical heat treatment.

The task of studying the diffusion processes and the evolution of the dislocation structure under SPD conditions and a subsequent heat treatment for copper alloys with different chromium contents remains relevant and of scientific interest.

In this paper, the study is devoted to the processes of structure formation and changes in the physical and mechanical properties in alloys with chromium contents of 0.2% and 1.1% by weight (Cu–0.2Cr and Cu–1.1Cr) subjected to SPD by equal-channel angular pressing and a subsequent heat treatment.

2. Materials and Methods

In this study, we used low-alloyed electrically conductive Cu–Cr alloys with chromium concentrations of 0.2 and 1.1 wt.%. The initial diameter of the rods was 14 mm. The initial condition of these alloys was state of the supersaturated SS produced by holding at 1000 °C for 2 h in a Nabertherm L15/11/180 (Nabertherm GmbH, Lilienthal, Germany) furnace and further water quenching. To remove the oxidized layer after the high-temperature treatment, the tempered specimens were turned to the diameter of 10 mm. The length of the workpieces for ECAP processing was 55 mm.

Deformation by ECAP with the channel alignment angle of 90° was carried out at room temperature with a deformation rate of 10 mm/s. One deformation cycle was completed for each alloy. There are prototypes of industrial Conform units based on ECAP, so it becomes relevant to study the first cycle of the deformational treatment.

Post-deformational treatment (aging) was performed in a shaft furnace by immersing specimens into a salt bath of ½ KNO + ½ NaNO. The temperature was 450 °C.

All structural studies of the specimens after ECAP processing were performed in the cross-section. Studies of the structure and second-phase particles at the meso-scale were conducted using a Jeol JSM-6490LV (JEOL Ltd., Tokyo, Japan) scanning electron microscope (SEM). Micro-structural studies were performed using a Jeol JEM-2100 (JEOL Ltd., Tokyo, Japan) transmission electron microscope (TEM). Specimens were prepared as thin foils by electrolytic polishing in a Tenupol-5 unit. The electrolyte composition was methanol 70%, nitric acid 30%. The temperature was −20 °C.

X-ray patterns were obtained using a Rigaku Ultima IV diffractometer (RIGAKU CORPORATION, Tokyo, Japan) in the Bragg-Brentano geometry with the help of CuKα radiation generated at 40 kV and a current of 40 mA. Vertically (2/3 deg.) and horizontally (10 mm) limiting slits and a Soller slit 5° were used on the primary beam, and a vertically limiting slit (2/3 deg.), horizontal slit (0.6 mm), and a Soller slit 5° were used on the reflected beam. A graphite monochromator was placed in front of the detector. X-ray patterns were obtained within the range from 35° to 145° in the continuous scanning regime with a rate of 0.25°/sec.

The calculation of the lattice parameter, the distribution of coherent scattering regions by size, the density of marginal and helical dislocation, and the effective radius of dislocation were performed by analyzing X-ray patterns under the Whole Powder Pattern Modelling (WPPM) approach implemented in the PM2K (ver. 2.12) software (University of Trento, Trento, Italy) [32].

The quantitative and phase analysis of second-phase particles was completed using the low-angle diffraction method. Dispersion curves were obtained using a Rigaku Ultima IV diffractometer with a low-angle add-on. CuKα radiation obtained at 40 kV and 40 mA in the X-ray tube was used. The recording was performed in the parallel beam mode implemented in the Cross Beam Optics (RIGAKU CORPORATION, Tokyo, Japan) add-on using a parabolic multi-layer mirror, Soller 5.0° and 0.5° slits on the primary and secondary beams, respectively. Dispersion on air fluctuations was minimized by vacuuming the optical path between the specimen and the detector. Measurements of dispersion curves were taken within the vector variation range of 0.05 nm⁻¹ to 0.1 nm⁻¹. The radiated surface diameter was ~4 mm.

Specific electrical conductivity was measured using a VE-27NC (Sigma LLC, Ekaterinburg, Russia) eddy current sensor for non-ferrous metals and alloys. The obtained values of electrical conductivity are given in the units of the International Annealed Copper Standard (IACS). The electrical conductivity of annealed commercially pure copper of 58 MSm/m was taken as 100% IACS. Measurements were taken at the room temperature of 22 °C.

Micro-hardness was measured using a Micromet 5101 micro-hardness tester (Buehler Ltd., Lake Bluff, Illinois, IL, USA) with 10 N loading. Mechanical tests were performed at room temperature using an Instron 8862 (Instron, Norwood, MA, USA) pull test machine at the deformation speed of 10⁻¹ s⁻¹. For elongation testing, proportional cylindrical specimens with a gage part diameter of 3 mm and an initial rate length of 15 mm were used.

3. Research Results

The initial condition was the state of the supersaturated (SS) produced during high-temperature treatment by holding at 1000 °C for 2 h and a subsequent water quenching. After quenching, a coarse structure was formed in specimens with a residual content of non-dissolved inclusions in the form of particles (Figure 1).

The results of structural analysis are given in

Table 1. Average grain size and lattice parameter of Cu–Cr alloys in the initial state.

Alloy	Average Grain Size, µm	Lattice Parameter, Ǻ
Cu–0.2Cr	65 ± 4	3.6158 ± 0.0003
Cu–1.1Cr	130 ± 10	3.6170 ± 0.0005

. An increased Cr percentage led to an increased SS concentration during high-temperature treatment, which is proved by an increased lattice parameter (Table 1).

Particle analysis in the initial state by scanning electron microscopy (SEM) (at the meso-scale) (Figure 2) showed that Cu–0.2Cr predominantly had fine particles below 0.4 µm located mainly along grain boundaries. Large particles of elongated shape with a transverse size of about 2 µm and a shape elongation coefficient of about k = 1 : 3 were observed. Cu–1.1Cr had a more uniform distribution of particles below 1.6 µm in size.

Deformation using the ECAP method led to the formation of a band-type structure (Figure 3). The average transverse size of bands was 35 ± 3 µm for the Cu–0.2Cr alloy and 30 ± 3 µm for the Cu–1.1Cr alloy. Analysis of second-phase particles at the meso-scale (SEM) showed that during ECAP processing, the Cu–0.2Cr alloy preserved a high share of particles with a size below 0.4 µm. In the Cu–1.1Cr alloy, the share of particles about 0.8 µm in size was increased, and the share of particles above 1.2 µm in size decreased, which can be related to the process of particles breakdown. A more complete analysis of the processes related to particles requires studies at various structural levels.

Transmission electron microscopy showed that fragments of band type were formed (Figure 4). The average transverse size of fragments for the Cu–0.2Cr and Cu–1.1Cr alloys was 210 ± 12 and 170 ± 8 nm, respectively. Differences in the nature of the boundaries of the observed fragments must be noted. The specimen of the Cu–0.2Cr alloy had fine and formed boundaries, while the specimen of the Cu–1.1Cr alloy had more non-equilibrium boundaries, primarily including dislocation walls. The body of the bands in the Cu–0.2Cr alloy showed a cellular structure, indirectly indicating that there was a process of polygonization. The observed particles in both alloys were located along dislocations and the boundaries of fragments.

The lattice parameter analysis by X-ray diffraction (XRD) showed that during ECAP processing, the lattice parameter was restored insignificantly (

Table 2. XRD (X-ray diffraction) results.

State	Lattice Parameter, Ǻ	Coherent Scattering Region (CSR), nm	Lattice Distortions, %	Dislocation Density, 1014 m−2
Cu–0.2Cr 1000 °C	3.6158 ± 0.0003	162	0.06	0.18
Cu–0.2Cr ECAP 1p.	3.6154 ± 0.0004	49	0.18	4.61
Cu–0.2Cr ECAP 1p. + 450 °C 1 h	3.6151 ± 0.0002	52	0.09	0.87
Cu–1.1Cr 1000 °C	3.6170 ± 0.0004	157	0.11	0.67
Cu–1.1Cr ECAP 1p.	3.6164 ± 0.0004	46	0.18	5.13
Cu–1.1Cr ECAP 1p. + 450 °C 4 h	3.6149 ±0.0003	53	0.12	2.46

). The results are in the domain of statistical error, but based on the previous studies of alloys of this class [1,19,21,22], it can be suggested that there is a process of deformation-stimulated decay of the supersaturated SS. The Cu–0.2Cr and Cu–1.1Cr alloys showed almost the same level of lattice restoration during ECAP processing (~0.0006 Å). The value of areas of coherent scattering regions after ECAP processing in the Cu–0.2Cr and Cu–1.1Cr alloys was almost the same, which shows a comparable level of micro-structure refinement in both alloys.

Both in the initial state and after ECAP processing, the highest dislocation density is observed in the Cu–1.1Cr alloy. The micro-distortions of the lattice after ECAP processing were the same for both alloys and were almost twice as high as in the initial state, which is related to an increased number of defects in the alloys.

Additional analysis of second-phase particles by XRD using low-angle diffraction provided data for the distribution by sizes and the distance between fine (below 100 nm) second-phase particles. ECAP led to an increase in average size of second-phase particles in both alloys, while the dimensional area of the observed particles grew, which is indicated by an increased peak width (Figure 5). It can be noted that the average particle size in the Cu–0.2Cr alloy was smaller than that in the Cu–1.1Cr alloy (

Table 3. Average particle size and the distance between particles in the initial state and after ECAP.

Alloy	Average Particle Size, nm			Average Distance between Particles, nm
	1000 °C	ECAP	ECAP+ aging	1000 °C	ECAP	ECAP+ aging
Cu–0.2Cr	7 ± 1	22 ± 2	28 ± 2	440 ± 20	380 ± 20	330 ± 20
Cu–1.1Cr	14 ± 1	43 ± 2	54 ± 2	320 ± 20	280 ± 15	260 ± 15

). The distance between particles was smaller in the Cu–1.1Cr alloy. This size distribution and changes in the distance between particles can be related to the breakdown of large particles that were observed by SEM and the difference in the chromium concentration in each of the studied alloys.

The phase analysis performed using low-angle diffraction indicated that the crystal lattice of fine-dispersed particles belongs to the Im-3m (229) space group, i.e., they have a bcc lattice. The particles thus identified by X-ray analysis are chromium.

3.1. Aging Kinetics of the Cu–0.2Cr and Cu–1.1Cr Alloys

The specimens were subjected to aging at a temperature of 450 °C [10,19,21,22], which is more characteristic for this class of alloys. Aging up to one hour was completed with an increment of 30 min, and further aging up to 5 h was performed with an increment of 1 h. Micro-hardness and electrical conductivity were measured at each increment. Post-deformation aging (Figure 6) showed that the thermal stability of the Cu–0.2Cr and Cu–1.1Cr alloys greatly differed. Aging in the Cu–0.2Cr alloy did not result in hardening, and after 1 h, micro-hardness decreased, which shows that over-aging takes place. The electrical conductivity of this alloy was below 67 ± 2% IACS even after 5 h of aging. In the case of aging for 4 h, hardness rose from 1120 ± 50 to 1660 ± 50 MPa in the Cu–1.1Cr alloy, and then this value stabilized. Electrical conductivity was substantially restored in 30 min and reached 76% IACS after 4 h and then stabilized.

3.2. Aging of the Cu–0.2Cr and Cu–1.1Cr Alloys

In this study, the Cu–0.2Cr and Cu–1.1Cr alloys were subjected to aging after quenching and ECAP in the conditions selected as per the results of the thermal stability experiments and ensuring the best physical and mechanical characteristics to be obtained. In this manner, the Cu–0.2Cr alloy was heat-treated at 450 °C for 1 h, and Cu–1.1Cr was heat-treated for 4 h.

The structural analysis by TEM (Figure 7) showed a substantial difference in changes in the initial structure of these alloys after ECAP and aging. The Cu–0.2Cr alloy contained large re-crystallized grains up to ~5 µm in size (Figure 7b). The average transverse size of non-crystallized fragments was 215 ± 12 nm. In the Cu–1.1Cr alloy, the average transverse size of fragments rose to 270 ± 15 nm. Boundaries became more perfect. There was a developed dislocation structure in the form of cells in the body of fragments. Second-phase particles were attached to dislocations, which is explainable because dislocations are predominant areas where second-phase particles are generated.

The X-ray structural analysis showed that the parameters of the crystalline lattice were restored in both alloys, which shows the SS decay process. However, in the Cu–0.2Cr alloy, the density of dislocations fell down almost to the initial value, while in the Cu–1.1Cr alloy, the density of dislocations was reduced two-fold but still remained of the same order (Table 2). Analysis of second-phase particles by XRD showed the increased number of increased average particle size to 28 ± 1 nm in the Cu–0.2Cr alloy and 54 ± 2 nm in the Cu–1.1Cr alloy (Table 3). Aging led to increased scattering of particles by size and the further rising average distance between them (Figure 5). Their location density in the alloy was increased.

3.3. Physical and Mechanical Properties

Substantial refinement of the structure after the first cycle of ECAP led to increased micro-hardness and strength of the studied alloys (

Table 4. Physical and mechanical properties of the Cu–0.2Cr and Cu–1.1Cr alloys.

State	Microhardness (HV), MPa		Ultimate Tensile Strength (UTS), MPa		Yield Strength (YS), MPa		Electrical Conductivity (G), %IACS
	Cu–0.2Cr	Cu–1.1Cr	Cu–0.2Cr	Cu–1.1Cr	Cu–0.2Cr	Cu–1.1Cr	Cu–0.2Cr	Cu–1.1Cr
Initial	660 ± 30	760 ± 30	240 ± 15	225 ± 15	120 ± 15	120 ± 15	48 ± 2	35 ± 2
ECAP	1180 ± 50	1120 ± 50	370 ± 15	380 ± 15	345 ± 15	360 ± 15	49 ± 2	36 ± 2
ECAP + aging	1200 ± 50	1660 ± 60	360 ± 15	485 ± 20	200 ± 15	440 ± 15	54 ± 2	76 ± 2

). Both alloys showed the same level of hardening after ECAP 360–380 MPa. Electrical conductivity during ECAP had almost no changes.

Due to the development of re-crystallization processes in the Cu–0.2Cr alloy after aging, the SS decay process with the formation of particles allows for preserving the strength level achieved during ECAP, but the yield stress is significantly reduced. The electrical conductivity of the alloy showed an unsatisfactory level for this class of alloys—54% IACS. This effect can be explained by a substantial reduction of the density of dislocations in the material that substantially accelerated diffusive processes [33]. Due to active SS decay in the Cu–1.1Cr alloy, precipitation hardening took place, which allowed for a significant increase in strength to 485 ± 20 MPa and electrical conductivity restoration to 76% IACS. The yield stress rose to 440 ± 15 MPa.

Figure 8 shows the tensile curves for the studied Cu–Cr alloys. The curves are similar, with the only difference being that the relative elongation of the Cu–1.1Cr alloy for all states is slightly higher than that of the Cu–0.2Cr alloy. The coarse-grained state with a supersaturated solid solution is characterized by low strength characteristics and high relative elongation. The state after ECAP is characterized by rapid hardening and a significantly lower relative elongation, but it does not decrease below 9%. Aging leads, on the one hand, to relaxation of stresses, which leads to an increase in the ductility of the material; on the other hand, precipitation hardening makes it possible to increase strength. As noted earlier, re-crystallization in the Cu–0.2Cr alloy during aging also significantly reduces the strength characteristics of the material; therefore, there is no increase in strength in comparison to the state after ECAP.

4. Discussion

As noted above, an important aspect of gaining strength for precipitation-hardening Cu–Cr alloys is the SS concentration in the initial condition that forms a resource for precipitation hardening in further deformational and heat treatment. Using the Matissen rule, the SS concentration can be evaluated in the initial states of the Cu–0.2Cr and Cu–1.1Cr alloys:

ρ = ρ₀ + Δρ_SS + Δρ_IMP

(1)

where ρ0 is the electrical resistance of pure copper, Δρ_SS is the contribution of SS to electrical conductivity, and Δρ_IMP is the contribution from impurities. For the electrical conductivity of 92% IACS of the experimental specimen of commercially pure copper, the contribution of the copper matrix to electrical resistance will be 0.15 × 10⁻⁸ Ohm·m.

Contribution from SS can be represented as follows [30]:

Δρ_SS = mC_SS

(2)

where m is the coefficient (tabular constant) (

Table 5. The values of electrical resistance and the co-efficient m for chromium.

ρ0.2Cr (1000 °C)	2.08 × 10−8 Ohm·m	ΔρIMP	0.15 × 10−8 Ohm·m
ρ 1.1Cr (1000 °C)	2.94 × 10−8 Ohm·m	m	3.9 × 10−8 Ohm·m [34]
ρ0	1.72 × 10−8 Ohm·m

) taking into account the rise of specific resistance per 1 at.% of the alloying element in the SS, and C_SS is the SS element concentration in at.%. Respectively, it follows from Equations (1) and (2) that:

C_SS = (ρ−ρ₀−Δρ_IMP)/m

(3)

The calculations show that the chromium concentration in the Cu–0.2Cr alloy after the high-temperature treatment and quenching is 0.044 wt.%, while the chromium concentration in Cu–1.1Cr is 0.223 wt.% for the same treatment. Taking into account that the chromium solubility limit at 1000 °C is about 0.7 wt.% in the alloy with a substantially lower quantity of chromium, no complete dissolution occurs. A much higher chromium concentration does not lead to dissolution by a value corresponding to the solubility limit.

Separate attention in alloys of this class must be paid to the characteristics of the assembly of second-phase particles. The size and distribution nature of particles affect the strength characteristics of the alloy. Apart from large particles observed using SEM, a large number of nano-sized particles (<20 nm) are in the initial state, whose presence in the conditions of pre-melting temperatures can be related to a fluctuation process of generation-dissolution of particles whose size is close to critical ones.

For electrically conductive low-alloyed copper alloys, electrical conductivity and the lattice parameter are structure-sensitive characteristics. Changes in these parameters show that the SS condition in the solid matrix of the material has changed. As the research results (Table 2 and Table 4) show, these parameters almost do not change during ECAP, which proves that there are almost no SS changes. However, quantitative characteristics of the assembly of particles (Table 3) change substantially. The average size of second-phase particles is increased almost three-fold as compared to the initial tempered state. The average distance between particles is also decreased. These changes can be related to the quasi-brittle breakdown of second-phase large particles [18]. In this manner, ECAP is accompanied by refinement of large particles of sub-micron size, which results in small nano-sized particles being formed. The average size of second-phase particles after ECAP in the Cu–1.1Cr alloy is two times higher than the average size of particles in the Cu–0.2Cr alloy, which can be attributed to the quantitative content of second-phase particles and their sizes in the initial state.

The aging time of alloys when the maximum hardening is achieved directly depends on the SS supersaturation degree both in the initial state and after ECAP. As the changes in micro-hardness and electrical conductivity in Cu–0.2Cr show, the process of structure relaxation primarily related to the decreased dislocation density and grain growth will be more intensively developed than the decay of the supersaturated SS and the formation of new second-phase particles. An assembly of nano-sized particles after ECAP, which is more densely packed and larger than that in the initial state, ensures the long-term thermal stability of the Cu–1.1Cr alloy (Figure 6). Due to a high number of both nano-sized and micron-sized particles during aging, the dislocation density shows almost no decrease. Aging in reasonable conditions does not lead to a substantial growth of the average size of particles, but the width of peaks (Figure 5 a,b) is increased since the size range of particles is increased.

Finding the optimal balance between strength and electrical conductivity for this class of alloys is an important scientific and industrial task. It is known that zirconium contributes to an increase in the strength properties and the level of thermal resistance of the material, but it has a stronger effect on the electrical conductivity of copper. It is also important that the complexity and cost of producing a ternary alloy is higher than that of a binary one. From this point of view, a simple increase in the proportion of the chromium content in the alloy looks reasonable. As it can be seen from comparative

Table 6. Comparative table of characteristics of Cu–Cr and Cu–Cr–Zr alloys subjected to ECAP.

Condition	HV, MPa	UTS, MPa	Elongation, %	G, %IACS	Source
Cu–1.1Cr 1000 °C-ECAP N = 1 + aging	1660 ± 60	485 ± 20	16 ± 2	76 ± 2	This work
Cu–0.75Cr ECAP (120°) N = 6 + aging	1630 ± 70	-	-	57	[35]
Cu–0.87Cr–0.06Zr ECAP N = 1 T = 473 K	-	490		73 ± 2	[36]
Cu–0.1Cr–0.06Zr ECAP N = 1 T = 673K	1300	320	30	80	[37]
Cu–0.36Cr ECAP N = 8 + aging	1620 ± 80	445 ± 5	23 ± 3	77	[38]

, the level of a set of properties achieved for the Cu–1.1Cr alloy in this work is not inferior to the level of properties of the Cu–0.87Cr–0.06Zr alloy and naturally exceeds the level of properties of alloys with a lower chromium content.

5. Conclusions

The studies show that high-temperature pre-treatment of the Cu–Cr alloys does not allow for the dissolution of the initial chromium phase within its solubility limits. Only 0.044 wt.% chromium dissolved in the Cu–0.2Cr alloy during the high-temperature treatment, and only 0.223 wt.% chromium dissolved in the Cu–1.1Cr alloy, while the possible value is ~0.7 wt.% at 1000 °C.

The completed physical experiment showed that the first cycle of ECAP led to a substantially refined structure up to 170 nm with an assembly of larger nano-sized particles to be formed (22 and 43 nm for Cu–0.2%Cr and Cu–1.1Cr, respectively) than in the initial tempered state (1 and 21 nm for Cu–0.2%Cr and Cu–1.1Cr, respectively). The increased average size of particles during ECAP is more likely related to the breakdown of large initial particles.

During aging in the Cu–0.2Cr alloy, the SS decay takes place less intensively than in the Cu–1.1Cr alloy. This explains the development of relaxation processes of polygonization and re-crystallization in the Cu–0.2Cr alloy.

Precipitation hardening during aging in Cu–1.1Cr is ensured by the assembly of particles ~54 ± 2 in size. The primary process taking place during aging is the growth of second-phase particles formed at the stage of ECAP. However, for the Cu–0.2Cr alloy, precipitation hardening during aging is much less intensive.

The Cu–1.1Cr alloy showed a better thermal stability after deformational treatment by ECAP as compared to the Cu–0.2Cr alloy treated in the same manner.

The best properties are achieved by the post-deformational aging of the Cu–1.1Cr alloy, which are ultimate tensile strength 485 ± 20 MPa and electrical conductivity 76% IACS.