Home Improvement on interfacial properties of CuW and CuCr bimetallic materials with high-entropy alloy interlayers via infiltration method
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Improvement on interfacial properties of CuW and CuCr bimetallic materials with high-entropy alloy interlayers via infiltration method

  • Xiaohong Yang EMAIL logo , Xiaoyong Sun , Zhe Xiao , Baocha Zhang , Peng Xiao and Shuhua Liang
Published/Copyright: July 30, 2024
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Science and Engineering of Composite Materials
From the journal Science and Engineering of Composite Materials Volume 31 Issue 1

Abstract

In order to achieve metallurgical bonding in the form of solid solution at the Cu/W interface and avoid the formation of intermetallic compounds, the novel high entropy alloys (HEAs) were designed on the basis of the mature high entropy alloy criteria. The CuCrCoFeNi x Ti high entropy alloys interlayers were applied to weld the CuW and CuCr bimetals by sintering–infiltration technology. Scanning electron microscope, energy dispersive spectrometry, and X-ray diffraction were used to explore the interfacial microstructure evolutions and strengthening mechanism of CuW/CuCr joints with applied HEA interlayers. The interfacial characterization results show that HEAs were diffused and dissolved into bimetallic materials, and a diffusion solution layer of 2–3 μm thickness was formed at the Cu/W phase interface, and there is no new phase generated at the CuW/CuCr interface. When CuCrCoFeNi1.5Ti interlayer was infiltrated into the CuW/CuCr interface, the electrical conductivity of CuCr side is 71.6%IACS, and the interfacial tensile strength reaches 484.5 MPa. Compared with the CuW/CuCr integral material without interlayer, the interfacial bonding strength is increased by 43.1%. And the SEM fracture morphology presents a larger amount of cleavage fractures of W particles. It indicates the appropriate solid solution layer on edge of W skeletons formed at the Cu/W phases interface. The Cu/W phase interface is strengthened, and it can effectively transfer and disperse the external load. Tungsten phase with higher elastic modulus endures a large amount of load, resulting in enhancing the CuW/CuCr interfacial bonding strength. When CuCrCoFeNi2Ti high entropy alloy interlayer was applied, the W skeleton near the CuW/CuCr interface was eroded, the imperfect W skeleton cannot withstand the tensile load effectively, resulting in decrease in the CuW/CuCr interfacial bonding strength. In the interfacial fracture, appears some fragmentations of W particles, and fewer W particles occur at the cleavage fracture.

Keywords: bimetallic interface; interlayer; HEAs; interfacial bonding strength; microstructures

1 Introduction

CuW and CuCr bimetallic materials are widely used in high-voltage electrical appliances. These bimetallic materials not only possess excellent resistance to arc erosion and fusion welding of W–Cu pseudo alloy but also present favorable conductivity of copper alloy [1,2,3,4,5]. Cu–W composites exhibit excellent ablative resistance owing to the high melting point of tungsten (3,410°C). CuCr alloy (Cr content is usually less than 1.0 wt%) is a typical precipitation strengthening alloy, which has high strength and high conductivity properties [6,7]. Switches, which combined the superior properties of Cu–W composites and CuCr alloy, are well suited to serve as electrical contacts in high-voltage witches.

Cu–W (70 wt% W) composites are usually prepared by powder metallurgy sintering–infiltration technique [8,9,10,11]. In our previous work, we have successfully connected Cu–W composite with low melting point bimetal by sintering–infiltration process [1,12]. However, the brittle Fe7W6 phase is inclined to W skeletons due to appending the excessive Fe element addition, and results in the lower interfacial strength of bimetallic materials [13]. High entropy alloys (HEAs) can form simple solid solutions and inhibit the production of intermetallic compounds, which can greatly improve the comprehensive performance of materials [14,15,16,17]. Especially, a number of research are focused on using HEA interlayers to connect different materials [18,19,20,21]. Li et al. [18] developed AlCoCrCuNi2 high-entropy alloy interlayer and used for vacuum diffusion bonding of TC4 titanium alloy to 316L stainless steel at 1,010°C. The results show that the desirable metallurgical bonding was obtained at the interface and the mechanical properties of the joint are enhanced. Tian et al. [19] used Al0.3FeCoNiCr HEA to diffusion bond DD5 single crystal and FGH98 superalloy at 1,130°C for 1 h. It is found that the elements at the interface are fully diffused and the reliable interfacial bonding is achieved. Azhari-Saray et al. [20] designed Al0.5FeCoCrNi HEA interlayer for dissimilar resistance spot welding between the sheets of 6061-T6 aluminum alloy and St-12 low carbon steel. Compared to interlayer-free weldment, the mechanical properties of the joint with Al0.5FeCoCrNi HEA interlayer are significantly improved. Ding et al. [21] reported that copper and titanium were bonded by utilizing vacuum solid-state diffusion method with an interlayer of CoCrFeMnNi HEA, and revealed the interfacial diffusion reaction mechanism. The results show that only FCC phase exists at the Cu/CoCrFeMnNi interface, and the Cu/CoCrFeMnNi/Ti joint was bonded firmly in the diffusion temperature range of 800–900°C.

On the other hand, the heat treatment process is often conducted to improve the microstructures and performance alloys. It not only can enhance the machinability, mechanical properties, wear resistance, and corrosion resistance of alloys but also can improve the internal quality of the material, stability, and surface quality to meet different engineering requirements. Modassir et al. [22,23,24] studied the effect of heat treatment on the heat affected zone (ICHAZ) of boron modified P91 steel (P91B). The results show that the least deformation was observed for re-normalized and re-tempered ICHAZs restoring grain boundary hardening, and type IV cracks can be avoided by re-normalizing and re-tempering P91B steel. It was also found that the addition of boron element made the parent metal have large austenite grains and fine nano-precipitations, resulting in precipitation strengthening.

Therefore, in this study, high-entropy mixed powder interlayers were introduced into the interface between the W–Cu composites and CuCr alloy, and the integral bimetallic materials were fabricated by sintering–infiltration technology. The metallurgical diffusion mechanism of high-entropy alloy interlayer into the bimetallic materials of CuW and CuCr was explored. The interfacial bonding strength was not only related to the diffusion of elements, but also to the CuCr side, so we studied the precipitation strengthening of CuCr side after heat treatment. By deeply analyzing the microstructures in bimetallic interfaces, the effects of high-entropy alloy interlayer on the formation of element diffusion solution layer in the CuW/CuCr interfaces were evaluated and discussed in detail.

2 Materials and methods

Elemental powders of Cu, Cr, Co, Fe, Ni, and Ti (purity 99.9%) were used as experimental raw materials, according to the designed mole ratio n(Cu):n(Cr):n(Co):n(Fe):n(Ni):n(Ti) = 1:1:1:1:x:1, where x is 0.5, 1, 1.5, and 2, respectively. In the V-type mixer, Cu, Cr, Co, Fe, Ti, and Ni powder with different contents were separately mixed for 8 h to obtain the mixed homogeneous powders. Then, the HEAs mixed powder was filled into in a steel mold and pressed as a circular green compact with a thickness of 0.5 mm. The bimetallic material studied in this work consists of Cu-70 wt% W composites and Cu–Cr alloy (Cr content is less than 1 wt%). These bimetallic raw materials were processed into round rods and then cleaned with acetone. After the raw materials dried, Cu–W composites, HEA interlayer, and CuCr alloy rods were stacked in the corundum crucible in the prescribed order. The crucibles loaded different integral samples were then placed in a sintering furnace with a nitrogen gas protective atmosphere. Initially, the furnace was purged of air by hydrogen for duration of 40 min. Subsequently, when the temperature was heated to 950°C at a rate of 20°C/min, the hydrogen was changed into nitrogen. After sintering for 40 min at 950°C, the furnace was heated to 1,380°C at a rate of 10°C/min, infiltrated for 60 min and cooled to room temperature in the furnace. The CuW/CuCr bimetallic samples were subjected to heat treatment in an electric resistance furnace. The solution treatment of these samples was performed at 950°C for 1 h, followed by water quenching. All the aging treatments of bimetallic samples were carried out at 450°C for 4 h, and the experimental flow chart is shown in Figure 1.

Figure 1 The flow chart of the experiment.
Figure 1

The flow chart of the experiment.

The electrical conductivity of the CuCr alloy near the bimetallic interface was tested using an FQR-7501 vortex conduct-meter after the infiltration process and subsequent heat treatment, respectively. The final tested value represents the average of five measured values obtained in different regions for each sample. In order to examine the CuW/CuCr interfacial microstructures, the samples were cut by the electrical discharge machining, and the cross sections perpendicular to the interface were polished for metallographic examination. And then they were etched in a solution of 5 ml FeCl3 + 25 ml HCl + 100 ml H2O for 3–8 s. The microstructures of the bimetallic interface were characterized by a JSM-6700F scanning electron microscope (SEM). The different microscopic areas at the interface were examined by an Oxford INCA type energy dispersive spectrometry (EDS). X-ray diffraction (XRD) was used to confirm the phase crystal structure near the bimetallic interfacial regions. The scanning speed was maintained at 2°/min in the range from 20° to 80°, and then Jade 5.0 software was utilized to analyze the XRD patterns and determine the lattice distortion. As the strength of CuW composite is much higher than that of CuCr alloy, the fracture phenomenon often occurs on the CuCr alloy when the tensile bars adopt the equal diameter standard rods. Therefore, all tensile tested bars were machined into the variable diameter samples, as shown in Figure 2. It can ensure that the fracture occurs at the CuW/CuCr bimetallic bonding interface rather than the CuCr alloy side, so the interfacial bonding strength of the bimetallic materials with different HEA interlayers can be obtained. Tensile tests were performed using the HT-2402-100 KN computer control material testing machine, which has a speed control accuracy of ±0.2%. SEM was used to examine the interfacial fracture morphologies of various bimetallic materials after the tensile testing.

Figure 2 The samples of CuW/CuCr bimetallic material used for tensile tests. (a) Dimensions of tensile sample and (b) the photograph of tensile samples.
Figure 2

The samples of CuW/CuCr bimetallic material used for tensile tests. (a) Dimensions of tensile sample and (b) the photograph of tensile samples.

3 Results and discussion

3.1 Selection of alloying elements and components design

According to the characteristics of CuW and CuCr bimetals, a novel CuCrCoFeNi x Ti HEA interlayer was designed. Table 1 gives the characteristics of selected elements. It can be seen that these elements are located in adjacent positions of the fourth period in the periodic table. They have similar atomic radii and similar properties, so it is easy to form simple solid solution structures with each other. On the other hand, during the preparation of CuW composites, alloying elements such as Cr, Co, Fe, Ni, and Ti are usually utilized as activation elements in the activation sintering of W skeleton. These elements have large solid solubility in W, it is conducive to forming a reliable metallurgical bond at the Cu/W interface. In addition, the solid solubility of these elements into Cu matrix is lower except for Ni, which has an excellent precipitation hardening effect. After the solution and aging treatment, the mechanical properties of the CuCr side can be enhanced, and the effect of these elements on the conductivity of the CuCr side can be declined.

Table 1

Characteristics of selected elements [25]

Elements Atomic number Atomic weight (g/mol) Density ρ (g/cm3) Melting point (K) Crystal structure Atomic radius (10−10 m) Electronegativity
Cu 29 63.55 8.96 1,357 fcc 1.28 1.90
Cr 24 52.00 7.19 2,133 bcc 1.28 1.66
Co 27 58.93 8.90 1,768 hcp 1.25 1.90
Fe 26 55.85 7.86 1,809 bcc + fcc 1.27 1.83
Ti 22 47.90 4.5 1,688 bcc + hcp 1.46 1.54
Ni 28 58.69 8.9 1,726 fcc 1.25 1.91

Due to the infinite solution of Cu and Ni [26], element Ni is hard to precipitate from the copper matrix after solution aging treatment, which can decrease the electrical conductivity of the bimetallic material interface and CuCr alloy side. Therefore, CuCrFeCoNi x Ti (x = 0.5, 1.0, 1.5, and 2) HEAs with different Ni contents were designed in this work, and the effects of Ni contents on the interfacial microstructural evolution and properties of CuW/CuCr bimetallic materials were studied.

In CuCrCoFeNi x Ti HEAs, random solid solutions are formed instead of Intermetallic compounds due to the high-entropy effect. The first HEA was first put forward by Yeh et al. [27,28], which overcomes the restriction of the single principal element in traditional alloy. The HEA is composed of five or more elements with a content ranging from 5 to 35%, and it does not have a dominant principal element. The configurational entropy of this alloy is at least 1.5 times the gas constant (R). The configurational entropy per mole can be determined using R ln ( n ) which is derived from Boltzmann’s hypothesis.

(1) ∆ S conf = − k ln ω = − R 1 n ln 1 n + 1 n ln 1 n + ⋯ + 1 n ln 1 n = − R ln 1 n = R ln n ,

where R is the gas constant, 8.314 J/K mol, and n is the number of elements. The configurational entropy of CuCrCoFeNi x Ti HEA is ≥1.5 R, which meets a criterion proposed by Yeh et al. [27,28] In addition, the other formation criteria of HEAs have been proposed by Zhang et al. [29,30,31] and Guo et al. [32,33]. They found that the formation of phases in multi-principal systems is intensely depended on several factors, such as the valance electron concentration VEC, atomic radial difference δ, mixing enthalpy ∆H mix, and entropy of mixing ∆S mix. And δ, ∆H mix, VEC, and ∆S mix can be calculated by the following formulas, respectively:

(2) δ = ∑ i = 1 n C i 1 − r i r ̅ 2 , r ̅ = ∑ i = 1 n C i r i ,

(3) ∆ H mix = 4 ∑ i = 1 , i ≠ j n ∆ AB mix C i C j ,

(4) VEC = ∑ i = 1 n C i ( VEC ) i ,

(5) ∆ S mix = − R ∑ i = 1 n C i ln C i ,

where r ̅ is the average atomic radius, C i , r i , and ( VEC ) i are the atomic percentage, atomic radius, and valance electron concentration of ith element, respectively, n is the number of principal elements, and ∆ AB mix is the mixing enthalpy of binary A-B system. The molar mixing enthalpy of binary is given in Table 2.

Table 2

Molar mixing enthalpy of binary alloys [34]

Cu Cr Co Fe Ti Ni
Cu 0 12 6 13 −9 4
Cr 0 −4 −1 −7 −7
Co 0 −1 −28 0
Fe 0 −17 −2
Ti 0 −35
Ni 0

Based on the formation rules proposed by Guo et al. [32,33], solid solution can be formed if δ , ∆ H mix , and ∆ S mix satisfy 0 ≤ δ ≤ 8.5, −22 ≤ ∆ H mix ≤ 7 kJ/mol, and 11 ≤ ∆ S mix ≤ 19.5 J/(K mol). In addition, FCC solid solution can be stably formed when VEC ≥ 8, and BCC solid solution will form when VEC < 6.87. The principal parameters of δ , VEC, ∆ H mix , and ∆ S mix of CuCrCoFeNi x Ti HEAs are listed in Table 3. All the parameters are consistent with the formation criterions of HEAs proposed by Guo et al. [32,33], even in accordance with the stricter rules reported by Jiang et al. [35] As a result, the phase composition in CuCrCoFeNi x Ti alloy is inferred to be a single FCC solid solution.

Table 3

Principal parameters of δ, VEC, ∆H mix, and ∆S mix of CuCrCoFeNi x Ti HEAs

Ni content (x) 0.5 1.0 1.5 2.0
ΔS mix J/Kmol 14.70 14.90 14.78 14.53
ΔH mix J/Kmol −7.40 −8.44 −9.09 −9.47
δ (%) 5.76% 5.65% 5.53% 5.42%
VEC 14.254 14.119 13.936 13.729

3.2 Characterization of CuW/CuCr interfacial microstructures

When the designed CuCrCoFeNi x Ti HEAs above were introduced into the interface of CuW and CuCr bimetals, the integral CuW/CuCr dissimilar materials were papered by infiltration method. After the solid solution and aging treatment, the microstructures of CuW/CuCr interface are shown in Figure 3. Figure 3(a) presents SEM image of the specimen without interlayer, it can be seen that there is no evident diffusion phenomenon at Cu/W interface. For the CuW/CuCr integral material, different HEA interlayers were applied as shown in Figure 3(b)–(e), it is obvious that the dark gray metallurgical diffusion solution layer with thickness of 2–3 μm was anchored in the boundaries of light gray tungsten skeleton. It indicates that the diffusion and dissolution phenomenon occurred on the CuW/CuCr interface in the infiltration process, resulting in metallurgical bonding between Cu and W phases. This metallurgical bonding at Cu/W phase interface is beneficial for the enhancement of the CuW/CuCr bonding strength. In addition, it can be clearly seen that when the interlayer is CuCrCoFeNi0.5Ti HEAs, the diffusion depth from the CuW/CuCr interface to CuW composite is only 3–4 μm. However, with the increase in Ni content in HEAs interlayer, the depth of metallurgical diffusion layer in CuW side is gradually extended. Figure 3(e) shows the interface microstructure of CuCrCoFeNi2Ti interlayer. It is observed that the depth of diffusion layer in CuW composite was more than 20 μm. At the same time, it is also found that the W skeleton near the CuW/CuCr interface was eroded by the molten HEA, and several tiny tungsten particles were fallen from tungsten skeletons. As for the reason for the detachment, we can use EDS for further analysis and explanation later. Based on the characterization results of the interfacial microstructure above, the following two useful pieces of information are worth paying attention to. On one hand, when the HEA interlayers infiltrated into the CuW/CuCr interface, the metallurgical diffusion layer on the edge of the W skeleton was formed, while its longitudinal depth continuously extends toward the CuW side. It is extremely beneficial to strengthen the bonding strength of the Cu/W phase interface and to enhance CuW/CuCr interfacial strength further. On the other hand, the disintegration of the W skeleton can inhibit the further improvement of the Cu/W phase interfacial strength.

Figure 3 The interfacial microstructures of the CuW/CuCr dissimilar materials with different HEAs interlayers. (a) Without interlayer, (b) CuCrCoFeNi0.5Ti, (c) CuCrCoFeNiTi, (d) CuCrCoFeNi1.5Ti, and (e) CuCrCoFeNi2Ti.
Figure 3

The interfacial microstructures of the CuW/CuCr dissimilar materials with different HEAs interlayers. (a) Without interlayer, (b) CuCrCoFeNi0.5Ti, (c) CuCrCoFeNiTi, (d) CuCrCoFeNi1.5Ti, and (e) CuCrCoFeNi2Ti.

In order to explore the elemental distributions of HEAs interlayer at CuW/CuCr interface after infiltration and heat treatment, some areas of EDS analyses were carried out at the interface as shown in Figure 4. And Table 4 lists the analysis results of different micro areas near interface, such as CuCr alloy side (Area A), copper matrix of CuW composite side (Area B), the metallurgical diffusion layer of W skeleton (Area C), and the center of W skeleton (Area D). It can be clearly seen from the Table 4 that the alloying elements in the HEAs interlayer diffused into the interfacial bimetallic materials. In addition, the amount of alloying elements diffused into the CuW composites was greater than that into the CuCr alloy side. Especially, the content of alloying elements in the region C at the interface is much higher than that of other areas, and the content of Cr, Fe, and Ti in the boundaries of tungsten skeleton is more prominent. It indicates that Cr, Fe, Ti, and other alloying elements are easier to dissolve into W phase in the sintering–infiltration process. The reason is that the solid solubility of these alloying elements in W phase is greater than that in Cu phase. The appropriate diffusion depth on the edge of the W skeleton is beneficial to improve the bonding strength of the CuW/CuCr bimetallic material. However, as shown in Figure 4, when the molar ratio of element Ni in HEAs interlayer is 2, mixing entropy ∆S mix of CuCrCoFeNi2Ti alloys is decreased, the segregation tendency of elements was increased as listed in Table 3, and the sluggish diffusion effect of HEAs was inhibited during infiltration [16,19]. Therefore, the W skeleton was dissolved and eroded by molten HEAs, the enrichment of alloying elements and segregation phenomenon at the boundaries of tungsten skeleton were intensified, and the sintered necks among W particles were destroyed. It causes the W skeletons to be eroded at interface, several tiny tungsten particles detached from tungsten skeletons, and Cu/W phase interfacial bonding characteristic was damaged.

Figure 4 The spot scanning analyses of the interface of CuW/CuCr dissimilar material with CuCrCoFeNi2Ti interlayer.
Figure 4

The spot scanning analyses of the interface of CuW/CuCr dissimilar material with CuCrCoFeNi2Ti interlayer.

Table 4

EDS analysis results of different micro regions in Figure 4

Area Cu (wt%) Cr (wt%) Co (wt%) Fe (wt%) Ni (wt%) Ti (wt%) W (wt%)
A 96.41 0.29 0.03 0.12 0.42 0.07 2.66
B 90.22 1.44 0.15 0.33 0.62 0.43 6.84
C 6.87 3.50 1.06 2.99 0.86 2.60 82.12
D 2.81 0.69 0.36 0.48 0.13 0.60 94.93

To further explore the phase compositions of CuW/CuCr bimetallic materials applied with HEAs interlayer, the interfaces were characterized by XRD analysis. Figure 5 presents the XRD pattern of the CuW/CuCr interface infiltrated by the CuCrCoFeNi2Ti interlayer at 1,380°C for 1 h, and the thickness of alloying interlayer was 0.5 mm. From Figure 5, it is found that only Cu and W phases exist at the CuW/CuCr interface, no intermetallic compound phase was generated, and the presence of other components in the HEA interlayer was also not detected. It indicates that the alloying elements in the HEAs interlayer were not reacting with W and Cu phases in the infiltration process, but were presented in the solid solution of W and Cu phases at the interface in the form of solute atoms.

Figure 5 XRD pattern of the CuW/CuCr interface infiltrated with CuCrCoFeNi2Ti interlayer.
Figure 5

XRD pattern of the CuW/CuCr interface infiltrated with CuCrCoFeNi2Ti interlayer.

3.3 Variations in the electrical conductivity

When the CuW/CuCr bimetallic materials are utilized as electrical contact materials, the electrical conductivity becomes an important performance. To evaluate the effect of HEAs addition on the electrical properties of these materials, the conductivity of different CuW/CuCr bimetallic materials applied with different interlayers were measured after the infiltration, solid solution, and aging treatment, respectively. Figure 6 presents the variations in electrical conductivity of the CuCr alloy with Ni contents in the HEAs interlayers under different states. It can be seen that the electric conductivity decreased gradually with the increase in the Ni content in HEAs. In addition, compared to the solution and initial infiltrated state, the electric conductivity of CuCr alloy is increased significantly after aging treatment.

Figure 6 Variations in electrical conductivity of bimetallic materials with different interlayers in different states.
Figure 6

Variations in electrical conductivity of bimetallic materials with different interlayers in different states.

Figure 7 presents the XRD patterns of CuCr alloy side near CuW/CuCr interface applied with CuCrCoFeNiTi HEAs interlayer in different stages. From the XRD analysis results, it is observed that there are only the diffraction peaks of Cu phase, and no new phase is detected after experiencing the different heat treatment. By comparing the diffraction peaks of Cu phase in different states, it is found that the diffraction peaks of Cu(111) and Cu(200) are significantly shifted to the right after aging treatment, while the intensity of Cu(220) diffraction peaks are significantly increased.

Figure 7 XRD patterns of the CuCr alloy side near CuW/CuCr interface with CuCrCoFeNiTi interlayer in different states.
Figure 7

XRD patterns of the CuCr alloy side near CuW/CuCr interface with CuCrCoFeNiTi interlayer in different states.

From the XRD results, it can be inferred that the main reason for the decrease in electrical conductivity is that the alloying elements dissolved into the Cu phase during the solution treatment, resulting in the larger lattice distortion of Cu matrix. It makes the electronic scattering phenomenon increase, and hinders the directional movement of free electrons in the Cu lattice. As shown in Figure 6, the electrical conductivity of CuCr alloy near CuW/CuCr interface with the CuCrCoFeNi0.5Ti interlayer is higher. However, with the increase in Ni content in HEAs interlayer, the alloying elements can sufficiently dissolve and diffuse into the CuW and CuCr materials from the interface, and then Ni elements dissolved in the CuCr side was also increased. Due to the infinite solid solution of Cu and Ni [26], element Ni is easy to dissolve into the crystal lattice of Cu, resulting in the increase in Cu lattice distortion. Therefore, the conductivity of the CuCr alloy decrease gradually with the increase in Ni content in the HEAs interlayer.

However, after the aging treatment, the Cu diffraction peaks shift toward the right as shown in Figure 7. It indicates that the lattice distortion of Cu is greatly decreased. This may be due to the precipitation of initial alloying elements after dissolving into the Cu matrix. Therefore, the integrity of the copper lattice has been recovered, and the scattering effect of solid solution atoms and copper lattice distortion on free electrons become weak. As a result, the electrical conductivity of the CuCr alloy was enhanced after the aging treatment. In addition, the mechanical property of CuCr alloy side can be heightened owing to precipitation strengthening.

3.4 Interfacial bonding strength and fracture morphology

To further explore the interfacial strengthening mechanism of CuW and CuCr bimetallic materials applied with HEAs interlayer, the interfacial bonding strengths of CuW/CuCr materials were tested through tensile experiment. Figure 8 presents the tensile stress–strain curves of CuW/CuCr bimetallic materials applied with different HEA interlayers. Compared with CuW/CuCr material without interlayer, it can be clearly observed that the bimetallic materials with HEAs interlayer exhibit higher bonding strength. In addition, with the increase in Ni content in the HEA interlayer, the interfacial strength first increases and then decreases. And the interfacial strength of bimetallic material reaches the maximum value of 484.5 MPa when CuCrCoFeNi1.5Ti interlayer was applied to the interface. Compared with the bimetallic material without interlayer, the interfacial strength is increased by 43.1%.

Figure 8 Tensile stress–strain curves of CuW/CuCr dissimilar materials with different HEA interlayers.
Figure 8

Tensile stress–strain curves of CuW/CuCr dissimilar materials with different HEA interlayers.

In order to clarify the failure mechanisms in the CuW/CuCr interface applied with HEA interlayers, the fracture morphologies of various tensile samples were examined by SEM. Figure 9 shows the interfacial fracture images of CuW/CuCr material and bimetallic materials with different interlayers. It is found that the fracture location of these bimetallic materials occurs at the CuW/CuCr interface. From Figure 9(a), it can be seen that the fracture surface is relatively smooth, and some ductile tearing ridges of copper phase appear on the fracture image. It can be deduced that the debonding phenomenon primarily occurred at the Cu/W interphase when the tensile load was applied at CuW/CuCr interface. As the CuW/CuCr bonding interface is mainly Cu/W interface phase, and the content of W in the CuW composite is 70 wt%, the interfacial strength depends on the bonding strength of the Cu/W interphase. Due to the mutual insolubility and non-reaction between W and Cu [12], the bond of Cu/W interphase is weak. The Cu/W interphase is prone to debonding, and then more flat areas remained on the fracture surface. Therefore, the interfacial strength of the CuW/CuCr bimetallic material without interlayer is relatively low, as shown in Figure 8.

Figure 9 Fracture morphology of CuCr side of CuW/CuCr integral material containing different HEA interlayers (a) without interlayers, (b) CuCrCoFeNi0.5Ti, (c) CuCrCoFeNiTi, (d) CuCrCoFeNi1.5Ti, and (e) CuCrCoFeNi2Ti.
Figure 9

Fracture morphology of CuCr side of CuW/CuCr integral material containing different HEA interlayers (a) without interlayers, (b) CuCrCoFeNi0.5Ti, (c) CuCrCoFeNiTi, (d) CuCrCoFeNi1.5Ti, and (e) CuCrCoFeNi2Ti.

As for CuW/CuCr bimetallic materials applied with CuCrCoFeNi x Ti (x = 0.5, 1.0, 1.5, 2.0) HEAs interlayer, the interlayer was sufficiently dissolved and diffused into CuW composite and CuCr alloy in the infiltration process. Owing to the compliance with HEA design criteria, the high-entropy effect can be generated at the interface. When the HEAs were infiltrated into the CuW/CuCr interface, the metallurgical diffusion layer was formed on the edge of W skeletons, and the longitudinal depth extended toward the CuW side up to 20μm, as shown in Figure 3. So, the Cu/W phase interface achieves the metallurgical bonding, Cu/W bonding strength is enhanced, and it can play a role in transmitting and dispersing tensile loads. Therefore, with the increase in Ni content in the HEA interlayer, the fracture position of the integral material is obviously inclined to the CuW side, and the characteristic morphologies of the fracture are gradually changed. From Figure 9(a), it can be seen that the fracture image is composed of dimples and some W particles, which were pulled out by copper phase tearing. Because the diffusion depth from the CuW/CuCr interface to CuW composite is only 3–4 μm as shown in Figure 3(b), Cu/W phase interface cannot effectively transfer load, resulting in the Cu/W interphase debonding. Whereas, CuCrCoFeNiTi and CuCrCoFeNi1.5Ti HEAs have larger mixing entropy, as listed in Table 3. The high entropy effect improves the interfacial diffusion and infiltration behavior, and then the appropriate solid solution layer on edge of W skeletons was formed at the Cu/W phases interface. In addition, the depth of metallurgical diffusion layer at the CuW/CuCr interface is also deeper, and the W skeletons are intact as shown in Figure 3(c) and (d). Because the Cu/W phase interface is strengthened, it can effectively transfer and disperse the external load. When the bimetallic material is subjected to tensile load, the external load is first shared by the continuous Cu phase in the interface, and subsequently the load can be distributed and transferred smoothly to the W phase by the Cu/W phase interface. Owing to the higher elastic modulus of tungsten phase, the bimetallic material interface can resist greater tensile stress. Therefore, the CuW/CuCr integral material applied with CuCrCoFeNi1.5Ti interlayer exhibits the highest interfacial bonding strength. In addition, the fracture morphology presents a larger amount of cleavage fractures of W particles as shown in Figure 9(d). This shows that W phase with higher elastic modulus endures a large amount of load and occurs cleavage fracture.

For the CuW/CuCr integral material, CuCrCoFeNi2Ti interlayer was introduced as shown in Figure 9(e), the main characteristics of fracture morphology is similar to the two mentioned above. It can be seen that the number of W particles with the transgranular cleavage decreased significantly, but more importantly, the fragmentation of W particles appeared, as shown by the arrow in Figure 9(e). The main reason is that the W skeleton near the CuW/CuCr interface dissolved and was eroded by the molten CuCrCoFeNi2Ti interlayer during infiltration, and tiny tungsten particles fell from the W skeletons as shown in Figure 3(e). Therefore, when the CuW/CuCr interface was subjected to tensile load, the imperfect W skeleton can no longer withstand the tensile load effectively, resulting in decrease in the CuW/CuCr interfacial bonding strength.

4 Conclusion

  1. On the basis of the mature HEA formation rules, the novel HEAs, CuCrCoFeNi x Ti HEAs, were designed to connect CuW composites and CuCr alloys. With the increase in the Ni content of the CuCrCoFeNi x Ti (x = 0.5, 1.0, 1.5, 2.0) HEAs, the mixing enthalpy of the alloys system first increases and then decreases, the mixing enthalpy of the system decreases gradually, the atomic radial difference and valance electron concentration also decreases gradually. All the parameters are consistent with the formation criterions of HEAs.

  2. CuW/CuCr bimetallic materials applied with different CuCrCoFeNi x Ti (x = 0.5, 1.0, 1.5, 2.0) HEA interlayers were prepared by sintering infiltration technology. The HEA interlayers can be completely dissolved and diffused into both side materials, and the diffusion and dissolution quantity of HEA interlayers into the CuW composites is greater than that into the CuCr alloy side. It is obvious that the dark gray metallurgical diffusion solution layer with thickness of 2–3 μm was anchored in the boundaries of tungsten skeletons in the infiltration process, resulting in metallurgical bonding between Cu and W phases. XRD results indicate there is no new phase generated at the CuW/CuCr interface.

  3. With the increase in the Ni content of the CuCrCoFeNi x Ti (x = 0.5, 1.0, 1.5, 2.0) HEAs interlayer, the electrical conductivity of the CuCr alloy decreases gradually. Compared with the CuW/CuCr integral material without interlayer, the electrical conductivity of CuCr alloy side is decreased by 19%. After the aging treatment, the conductivity of CuCr alloy is increased compared with that of the initial infiltration state and solution state.

  4. Compared with CuW/CuCr material without interlayer, the bimetallic materials applied with CuCrCoFeNi x Ti (x = 0.5, 1.0, 1.5, 2.0) HEAs interlayer exhibit higher interfacial bonding strength. When CuCrCoFeNi1.5Ti was applied to CuW/CuCr interface, the interfacial bond strength reaches a maximum value of 484.5 MPa, which is increased by 43.1%. And the fracture morphology presents a larger amount of cleavage fractures of W particles. It indicates that appropriate solid solution layer on edge of W skeletons was formed at the Cu/W phases interface. In addition, the depth of metallurgical diffusion layer at the CuW/CuCr interface is deeper, the Cu/W phase interface is strengthened, and it can effectively transfer and disperse the external load. So, tungsten phase with higher elastic modulus endures a large amount of load, resulting in enhancing the CuW/CuCr interfacial bonding strength. Whereas, when CuCrCoFeNi2Ti HEA interlayer was applied, the interfacial fracture surfaces appear with some fragmentations of W particles, and the cleavage fracture occurred in fewer W particles, resulting in the decrease in the CuW/CuCr interfacial bonding strength.

  1. Funding information: This work was supported by the National Natural Science Foundation of China (No. 52274366).

  2. Author contributions: All authors have accepted responsibility for the entire content of this manuscript and consented to its submission to the journal, reviewed all the results and approved the final version of the manuscript. XY designed the experiments and ZX carried them out. BZ, XS, ZX, PX and SL performed data collation and review. XY and XS conducted writing, revision and editing of manuscript.

  3. Conflict of interest: Authors state no conflict of interest.

  4. Data availability statement: All data generated or analyzed during this study are included in this published article.

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Received: 2023-10-11
Revised: 2024-05-31
Accepted: 2024-06-24
Published Online: 2024-07-30

© 2024 the author(s), published by De Gruyter

This work is licensed under the Creative Commons Attribution 4.0 International License.

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https://doi.org/10.1515/secm-2024-0029
Keywords for this article
bimetallic interface; interlayer; HEAs; interfacial bonding strength; microstructures
Creative Commons
BY 4.0

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