Effect of Beam Power on Intermetallic Compound Formation of Electron Beam-Welded Cu and Al6082-T6 Dissimilar Joints

Abstract

In this work, electron beam welds between Cu and Al plates were formed using different power modes, namely 1800 W, 2400 W, and 3000 W. The structure, microhardness, and tensile strength of the raw materials and the weld seams were studied. The low power of the electron beam resulted in the improper penetration and insufficient depth of the weld seam. The low power resulted in high cooling rates, which hindered the nucleation of the copper and aluminum particles. A number of intermetallic compounds (IMCs) were formed, including the metastable Cu₉Al₄ one. An increase in the power of the electron beam reduced the cooling rate and increased the miscibility between the materials. This resulted in the formation of a mostly homogeneous structure comprising an αAl solid solution and dendritic eutectic CuAl₂ intermetallic compounds. A preferred crystallographic orientation of the aluminum phase was detected regarding the sample prepared using a power of 3000 W, forming a specific texture towards the {111} family of crystallographic planes, which is the closest-packed structure. This plane characterizes the highest chemical activity and the highest plasticity. As a result, this sample exhibited the best chemical bonding between the IMCs and the aluminum matrix and the best microhardness and tensile test values.

1. Introduction

The joining of dissimilar materials is very attractive and useful currently, following the ever increasing standards for more efficient manufacturing processes in industrial fields [1,2]. By combining materials with differing thermophysical properties, it is possible to obtain a component possessing the advantages of each of the two materials. Copper and aluminum are highly sought-after materials due to their excellent malleability, electrical and thermal conductivity, high strength-to-weight ratio, and other qualities [3,4]. However, the joining of the two materials is highly limited due to the very low solubility of copper in aluminum, with a maximum value of about 5.65–5.80% [5]. Once oversaturated, the aluminum solid solution releases Cu atoms, which form bonds with the aluminum particles in the form of intermetallic compounds (IMCs) of the Cu_xAl_x type [6]. The most commonly formed IMC is CuAl₂, which is known to be brittle [7]. The presence of IMCs in joints between Cu and Al are a prerequisite for weakening the structure and premature fracture.

The most common method used for the joining Cu and Al is friction stir welding due to the very controlled process of mixing the two materials, which results in a reduced amount of formation of IMCs [8,9]. However, a highly desirable method that can potentially be used for this process is electron beam welding. Its main advantages are high precision, excellent controllability and potential for automation, high purity of the formed weld due to the absence of atmospheric gasses, the high penetration of the electron beam that makes it ideal for welding of butt joints, and so on [10]. A major issue is present with this process, however, which is related to the accelerated formation of IMCs. A well-known strategy used to limit the amount of IMCs formed is the reduction in Cu inclusion in the weld seam by offsetting the electron beam towards the aluminum component, which is also used in laser welding techniques [11]. Applying an offset towards the aluminum’s surface, however, creates another problem related to the rapid over-melting of the aluminum component, which occurs due to the low melting and boiling temperatures of Al. To resolve this issue, a circular beam oscillation can be used, as shown by [12]. Another method for the limitation of formation of IMCs is the application of a filler between the welding surfaces, which limits the incorporation of Cu in the melt pool, and thus the formation of the undesirable IMCs [13].

Regardless of the current advances in electron beam joining of Cu and Al, much more data need to be collected in order to determine the optimal technological conditions required for this operation. Due to this, in the present work, the power of the electron beam was varied in order to observe its influence on the mechanisms of formation of intermetallic compounds during electron beam welding. The structure of the weld seams and their microhardness and tensile strength were also investigated. The results are discussed for the benefit of possible future improvements in the process.

2. Materials and Methods

The metal plates used in the present work have a size of 100 × 50 × 8 mm and are made of Cu (99.8%) and an Al6082-T6 alloy. The chemical composition of the Al6082-T6 plate is as follows (wt%): 98.16% Al, 1.15% Mg, 0.32% Si, and 0.36% Mn. The plates were cut and the working surface areas were milled and ground with abrasive paper in order to reduce the distance between the plates during the welding process as much as possible. The schematic of electron beam welding with an offset and an oscillation of the electron beam is shown in Figure 1b. The electron beam welding was carried out on Evobeam cube 400, Evobeam GmbH, Nieder–Olm, Germany welding machine. Some of the technological conditions are presented in the table shown in Figure 1a. Prior to the welding procedure, both plates were cleaned with acetone and preheated to a temperature of 300 °C in an external furnace. The preheating process was performed in order to reduce the thermal internal stresses caused by the high thermal gradient formed during the EBW process. The temperature used in this work was chosen so that it was specifically lower than the temperature of oxidation of copper in air [14]. It is important to note that the reached maximum temperature of 300 °C was definitely reduced to an extent, influenced by the high thermal conductivity of copper and aluminum and also by the vacuum chamber depressurization time (about 30 s). The welding procedure was carried out using a welding speed of 15 mm/s, a circular oscillation with a radius of 0.2 mm and a frequency of 200 Hz, and an offset of 0.4 mm of the electron beam on the aluminum plate side. The accelerating voltage was constant, with a value of 60 kV. The welding current was varied, with values of 30 mA, 40 mA, and 50 mA, totaling the power of the electron beam to 1800 W, 2400 W, and 3000 W. The specimens were evacuated from the vacuum chamber 20 min after the process, after which they were placed in a furnace (TR 240, Nabertherm GmbH, Lilienthal, Germany) at a temperature of about 100–120 °C, where they were gradually cooled down to room temperature over the course of 60 min. A 2 °C per minute cooling process was determined to be sufficient enough to avoid the formation of cracks in the weld seam caused by rapid cooling. It is also important to note that although different cooling rates could also relieve the stresses during the cooling process successfully, the most favorable possible case was chosen in the current research in order to guarantee the successful evasion of crack propagation along the length of the weld seam.

Metallographic samples of the cross-section of the formed welds and the raw materials were prepared using a sequence of cutting, grounding, and etching. The abrasive paper strategy used was as follows: grit 240, 400, 600, 800, 1200, 2500, and 4000. A combined etching method was used in order to etch both materials using a 10% solution of HF acid for the aluminum side of the sample, and an FeCl₃ + HCl solution for the copper side of the sample. The optical images obtained in this work were taken using a Drawell MIT 300/500 microscope (Drawell Instrument Co., Ltd., Chongqing, China). Additional optical images of the untreated Al6082-T6 alloy and the heat-influenced alloy were taken using the same microscope. The same polishing steps were used as described above; however, in this case, the samples were etched using Tucker’s reagent. It comprised 25 mL dH₂O, 15 mL HF, 25 mL HCl, and 15 mL HNO₃. The samples were swabbed for 10 s each.

To perform a phase analysis of the resulting welded joints, an X-ray diffractometer (XRD), “Bruker D8 Advance” (Bruker, MA, USA), was used. The used method was “Two theta” (Bragg–Brentano). The characteristic X-ray radiation was CuKα, with a wavelength of 1.54 Å. A step of 0.1 degrees and a registration time of 0.5 s per step was applied. All measurements were performed in an area close towards the center of the weld seam. The X-ray beam had a circular shape with a diameter of 1 mm (thus having a scanning area of 0.785 mm²), as shown in Figure 2.

The structures of the welded samples were investigated using a scanning electron microscope (SEM), “LYRA I XMU” (TESCAN, Brno, Czech Republic), in the back-scattered electrons mode, equipped with an energy dispersive X-ray spectroscopy (EDS) “Quantax 200” (Bruker, MA, USA) unit.

The microhardness experiment was performed on a semi-automatic microhardness tester, “ZWICK/Indentec—ZHVμ-S” (ZwickRoell, Ulm, Germany). A load force of 0.5 N was used for all experiments. The microhardness of the weld joints was measured following a linear pattern, starting from the copper side of the samples, moving through the fusion zone, and ending on the aluminum side of the samples. The measurements were performed at a depth of the cross-section of 2 mm below the surface of the samples.

A ZwickRoell Vibrophore 100 unit (ZwickRoell, Ulm, Germany) was used in order to determine the tensile properties of the samples. A static strain method was applied with a strain rate of 30 MPa/s. The dimensions and a visual representation of the tensile test samples are shown in Figure 3. The length of the test zone was 40 mm, and the surface area was 320 mm².

3. Results

3.1. Structure and Properties of the Base Materials

Diffractograms of the aluminum and the copper plates prior to the preheating and welding processes are shown in Figure 4a,b, correspondingly. Studying the results of the aluminum plate, only an aluminum phase was noticed in its typical face-centered cubic structure. This phase was identified using PDF #040787 from the International Center for Diffraction Data’s database (ICDD). The {111}, {100}, {110}, and {311} family of planes were detected. The diffraction maxima corresponding to the {100} family of planes of the aluminum phase has the highest intensity, suggesting that the untreated samples have a texture highly concentrated towards that family of planes. In order to confirm that a specific texture is present in that sample, the pole density P_{hkl} [%] was calculated using the pole density formula presented in [15]. The results for the aluminum phase are presented in Figure 5a, indicating that indeed a preferred crystallographic orientation towards the {100} family of planes is present with a pole density of P_{100} = 72.3%. Typically, the most common family of crystallographic planes observed in polycrystalline aluminum is the {111} one; however, the aluminum bars the plates are prepared from are typically extruded under high temperatures and also heat treated to improve their strength. It is possible that the cause of the re-orientation towards the {100} family of crystallographic planes is the application of the heat treatment process. Previous investigations have observed a similar tendency, where applying heat to aluminum results in a change in the texture of the samples from the {111} family of planes compared to the {100} one. With the further increase in the applied heat a shift from the {100} family of planes towards the {311} family was observed as well [16]. Studying the diffractogram of the copper plate, only a copper phase was detected using PDF #040836 from the ICDD. A polycrystalline structure was observed with no specific texturing, as suggested by the performed pole density calculations presented in Figure 5b.

Figure 6a,b depicts the structure of the copper plate before and after the preheating process, correspondingly. It is evident that this process did not affect the structure of the copper plate in any way. This hypothesis is supported by the measured average particle size, which, in the case of the raw copper plate, was 128 ± 5.1 μm and, after the preheating process, 141 ± 6.6 μm. This is also supported by the measured microhardness, which was 114 ± 2.2 HV_0.05 and 115 ± 1.2 HV_0.05 in both cases, as suggested by the microhardness experiment results presented in Figure 6c and Figure 6d, respectively.

Figure 7a,b shows the structure of the aluminum plate before and after the preheating process, correspondingly. Since no visible particles were noticed, their sizes were not measured; however, the microhardness results, shown in Figure 7c,d, show a decreasing tendency after the preheating process. Initially the microhardness was 99 HV_0.05, after which it was reduced to 66 HV_0.05. The Al6082-T6 plate was heat-treated (T6), which process is performed in order to reduce the amount of defects in the structure to a minimum and induce precipitation hardening. The latter hinders the mobility of dislocations (defects) in the structure, thus increasing the strength of the material, which is in agreement with the Orowan mechanism of strengthening [17]. Applying additional heat with a high enough value to the material after the heat treatment process has been known to reverse its effect [18]. During the preheating stage, relaxation processes occur, which stimulate the relocation of nano- and micro-sized defects in the structure of the aluminum alloy. Figure 8a shows the microstructure of the Al6082-T6 plate before any heat was applied to it. Figure 8b shows the microstructure of the same sample after the preheating process. It was determined that the black regions are comprising a mixture of defects in the form of micro-pores and clusters of precipitates. It is obvious that, in the case of the untreated sample, a much denser structure is observed with a much higher concentration of precipitates. This indicates that the preheating applied to the sample also results in migration of the precipitates, most probably comprising the alloying elements. It is unclear at the moment what those precipitates are composed of, but it is also possible that the migration of defects and precipitates is also propagated by an increase in the size of the microvolumes. This is a known process and a factor that determines the strength of aluminum alloys [19]. As a result, the alloy’s structure changes after the preheating, accompanied by a noticeable decrease in the microhardness.

3.2. Structure and Properties of the Weld Joints

The XRD-patterns of the investigated samples are shown in Figure 9a, Figure 9b, and Figure 9c, corresponding to the three power modes, namely 1800 W, 2400 W, and 3000 W. It is important to note that no amorphous-like halos were observed during the processing of the data. A polycrystalline structure was detected with diffraction maxima corresponding to a number of phases. The last were identified using PDFs #260016, #250012, #040787, and #040836, taken from the International Centre for Diffraction Data’s database. In the case of the first two samples, Al and Cu phases were detected along with the CuAl and CuAl₂ intermetallic phases. The Al and Cu phases have a face-centered cubic crystal structure, while the CuAl phase has a monoclinic structure and the CuAl₂ phase has a body-centered tetragonal crystal structure. This is also confirmed by the authors of [7]. In the case of the sample prepared using a high beam power of 3000 W, only the CuAl₂ phase was detected along with the Al phase. No pure copper or CuAl intermetallic compounds were registered. Also, a noticeable reduction in the intensity of the intermetallic diffraction maxima was observed. This may suggest that the high power of the electron beam resulted in the formation of predominantly the CuAl₂ phase, accompanied by full melting and integration of the Cu particles amongst the aluminum ones and complete bonding between them.

As it is clear that the highest contribution in the melt pool is related to the incorporation of aluminum, the pole density of the aluminum phase was calculated, and the results are shown in Figure 10. Figure 10a shows the results corresponding to a beam power of 1800 W, and Figure 10b shows the results corresponding to a beam power of 2400 W. Evidently, no specific texturing can be noticed when considering those samples. Increasing the power of the electron beam to 3000 W led to the formation of preferred crystallographic orientation of the Al crystal phase towards the {111} family of crystallographic planes with a pole density of P_{111} = 64.8%. A similar tendency was observed in a previous work [12]; however, the texture was not investigated in that particular case. It is important to note that the phase composition and crystallographic properties of the preheated sample were not investigated; however, that should not have a major effect in the final values reported when studying the welded samples.

Figure 11 shows the microstructure of the weld seam formed using a beam power of 1800 W, whereas Figure 11a shows the border between the copper plate and the fusion zone (FZ), Figure 11b shows the FZ, and Figure 11c shows the borders between the FZ and the aluminum alloy plate. Evidently, some defects in the fusion zone are visible close to the border between the FZ and the Cu plate in the form of pitting and porosity. Well-visible borders were noticed while examining this sample. A process of formation of intermetallic compounds has obviously occurred, however, due to the low power of the electron beam and the rapid solidification of the melt pool, small intermetallic agglomerates formed. Figure 12 shows SEM images of the obtained weld seam using a beam power of 2400 W. Figure 12a shows the border between the Cu plate and the fusion zone. A clearly defined border between the two zones was formed, however, no visible cracks were present. Figure 12b shows the center of the fusion zone, where, as a result of increasing the electron beam’s power to 2400 W, the solidification speed decreased. This led to the formation of a denser intermetallic secondary phase amongst the primary matrix, which consists mostly of an αAl solid solution. Figure 12c shows the border between the fusion zone and the Al6082-T6 plate. A well-defined border was visible without any cracks. Better, more defined borders between the plates and the fusion zone was noticed in this case compared to the sample obtained with a beam power of 1800 W. Figure 13 shows SEM images of the obtained microstructure of the sample prepared with a power of the electron beam of 3000 W. Figure 13a shows the border between the copper plate and the fusion zone. In this case no cracks were noticed as well. A microhardness indentation was also noticed. Figure 13b shows the middle of the fusion zone, where, as a result of applying a power of 3000 W, the densest secondary phase was observed, which completely encompasses the grains of the solid solution. The formation of such a dense structure is propagated by the high heat input and slow solidification speed, which favors the nucleation processes. Figure 13c shows the border between the fusion zone and the Al6082-T6 plate, where also no cracks or major defects were noticed.

The cross-section of the weld formed using a power of the electron beam of 1800 W is shown in Figure 14a. Apparently, the applied power was not sufficient to melt through the plates, and only a partial joint formed. The maximum penetration depth in that case was about 4.5 mm. A zoomed image of the structure of the weld seam is shown in Figure 14b. A secondary phase dispersed through the aluminum solid solution is visible. To obtain knowledge on the chemical composition of the observed area, an EDS analysis was performed in points 1 and 2. Point 1 suggests that 66.5 at.% of Cu is present in that area, and 33.5 at.% of Al corresponding to the Cu₉Al₄ phase. The dark area investigated in Point 2 suggests that this area is entirely composed of the aluminum solid solution (95.72 at.% Al).

The cross-section of the sample welded with a power of 2400 W is shown in Figure 14c. Evidently, in this case, the power of the electron beam was also insufficient to fully penetrate the copper and aluminum plates. Despite this, most of the cross-section melted and formed a weld joint comprised predominantly of aluminum and copper agglomerates. The chemical composition of points 3 and 4 suggests that the observed area comprises the CuAl₂ phase and the aluminum solid solution.

The cross-section of the sample prepared with 3000 W of power is shown in Figure 14e. This power setting was enough to completely melt the aluminum and copper plates, and partial leakage of the melt pool was also noticed. Studying the surface of the sample closer, as shown in Figure 14f, reveals that, just like in the case of the other samples, it comprises Al- and Cu-rich intermetallic agglomerates. Points 5 and 6 suggest that, in this case, the composition of the sample is predominantly composed of an aluminum solid solution with 97.74 at.% and the CuAl₂ IMC with 27.21 at.% of Cu. The microhardness of the samples obtained using a beam power of 1800 W, 2400 W, and 3000 W is shown in Figure 15a, Figure 15b, and Figure 15c, respectively. In the case of the sample prepared with a power of 1800 W, an average microhardness of 190 HV_0.05 was measured in the fusion zone. In the case of the sample prepared with a power of 2400 W, a microhardness of 176 HV_0.05 was observed, and in the case of using a beam power of 3000 W, a value of 166 HV_0.05 was obtained, once again in the fusion zone. Evidently, the microhardness of the first sample is the highest. Regarding the copper and aluminum plates, an average microhardness of 50–60 HV_0.05 was measured in all cases. The mechanism of reduction in the mechanical properties of aluminum was already explained. As far as the copper substrate is concerned, it is possible that, due to the high heat applied using the electron beam and due to the very high thermal conductivity of copper, an increase in the size of the copper particles occurs, potentially with the formation of nano- and micro-sized defects within its structure. This could explain the observed reduction in microhardness; however, no current proof of this was observed.

As mentioned above, tensile tests were performed. The copper plate exhibited an yield strength (YS) of 262 MPa and an ultimate tensile strength (UTS) of 278 MPa. The maximum force applied before a fracture occurred was 22.68 kN. In the case of the aluminum plate, a YS of 226 MPa was observed along with a UTS of 345 MPa, with a maximum applied force of 32.33 kN. Subsequently, the tensile properties of the welded samples were investigated. It is important to note that the tensile properties of the raw materials were investigated prior to the preheating stage. In the case of the samples prepared using a power of the electron beam of 1800 W and 2400 W, a reduced cross-section of the working area was observed. Due to this, theoretically, some error regarding the obtained results should be expected. However, this issue was negated by the breaking of those samples immediately after applying a force of just about 200 N. In the case of the third sample, a higher breaking force was necessary to fracture the sample of about 5.5 kN, resulting in an ultimate tensile strength of about 50 MPa.

4. Discussion

During the welding process, a number of physical phenomena are present, such as melting of the solid state materials, formation of defects during the liquid stage, cooling processes, followed by solidification of the melt pool, formation and rearrangement of solid state crystals, migration of defects during the recombination processes, potential plastic deformation, and more. Each of these processes can have a major effect on the output structure of the weld seam. During the solidification stage, the priority of the molten phase is to crystallize with minimal internal strains (achieving the highest possible equilibrium). This process is problematic even when welding similar materials, particularly those that change their phase composition due to the high temperature (such as steel, titanium and others). Welding aluminum, on the other hand, leads to the formation of defects and their migration such that clusters of vacancies form. This causes the separation of some of the aluminum crystals and the formation of hollow cavities (pores). This type of defect is commonly referred to in the literature as a solidification pore [20]. It is obvious from the SEM analysis that a number of solidification pores are present in both the first and second samples formed with a beam power of 1800 W and 2400 W. No such defects are observed in the case of the last sample formed with a power of 3000 W. Instead, a clearly defined secondary phase is observed, mostly consisting of the CuAl₂ phase, as suggested by both the XRD and SEM analysis. It is obvious from the SEM images that a highly inhomogeneous weld seam was formed with clusters of different phases. The area observed by both the XRD and SEM analyses shows a specific preferred crystallographic orientation towards the {111} family of planes. It is possible that such an occurrence in this region was caused by the high chemical activity of this family of planes. Previous research categorizes the chemical activity of the different crystallographic planes with the {111} family of planes being the most chemically active due to its low free surface energy, and the {100} family as the least chemically active due to its highest free surface energy [21]. Another study confirms this and also suggests that the {111} family of planes exhibits the highest wettability potential [22], which improves the chemical bonding, in this case, between the intermetallic compounds and the αAl matrix. Thermodynamically, this family of planes should suffice for the formation of the highest amount of chemically and energetically stable bonds with a low amount of internal stresses. Of course, the latter were not studied in the present research, and this is just a hypothesis.

Electron beam processes, particularly welding processes, are characterized by a high heat input derived from the high density of the accelerated beam of electrons [23]. Due to the rapid and highly focused melting of the materials, high thermal gradients are observed [24]. The theory of solidification processes suggests that thermal gradients (i.e., cooling rates) are of exceptionally high importance to the resultant structure and mechanical properties of the formed non-separable joint [25]. The higher the cooling rate is, the lower the nucleation time is. This hinders the formation of secondary structures within the primary solid solution, such as eutectics and precipitates [26]. In the present case, with the increase in the power of the electron beam, a noticeable increase in the input heat was observed, which resulted in slower cooling processes. This was confirmed by the SEM images, which show that, in the case of the lower power of 1800 W, poorly defined dendritic eutectic structures have formed. The increase in the input heat intensified and propagated the formation of intermetallic compounds of a similar, but yet much more pronounced structure. This is confirmed by the XRD experiments where the CuAl phase was detected in the case of the first two samples, and the EDS experiments which also detected the presence of the metastable Cu₉Al₄ phase. Such phases were absent during the investigations performed regarding the sample prepared using a beam power of 3000 W. Of course, it is crucial to mention that both the XRD and the EDS analyses were performed in highly localized areas, such that the absence or the presence of other intermetallic phases cannot be completely denied. The absence of intermetallic phases other than the CuAl₂ phase in the structure of the sample prepared with the highest power, however, definitively suggests that, even if other IMCs are present, their quantity is low enough to not be detected by either analysis.

The authors of [27] have investigated the possibility of electron beam welding of copper and aluminum bimetallic plates with a similar power of the electron beam (3240 W). They report an accelerated formation of the CuAl₂ phase, which hinders the mechanical properties of the weld. They also suggest that the fracture mechanism is such that due to poor nucleation between the IMCs and the aluminum solid solution the fractures that form in the structure of the samples occur at the border between the IMCs and the crystals of the primary phase. Such results were also observed in the present work, where all fractures occurred between the fusion zone and the heat affected zone of either the aluminum or the copper plates. Li et al. [28] have studied electron beam welded joints using a beam power of 1950 W. Unlike in the present work, they only observed the formation of the CuAl₂ phase and did not detect the presence of other IMCs. Their investigations, however, were performed on Cu and Al plates with a thickness of 3 mm, so full penetration of the electron beam was achieved in all cases. Furthermore, they report similar tensile properties, highly comparable to the ones obtained in this work. Otten et al. [29] have also investigated the possibility of electron beam welding of 3 mm thick copper and aluminum plates with similar technological conditions as the ones investigated in [28]. They, however, study the possibility of reducing the formation of IMCs by applying varying offsets of the electron beam towards the Al plate and achieve great success. However, the parasitic metastable Cu₉Al₄ phase was also present within the structure of the joints. A highly known problem of electron beam welding is the formation of pores either due to low miscibility or due to the breaking apart of Cu_xO_x or Al_xO_x phases, thus causing the release of oxygen. This effect was studied in detail by the authors of [30], who report that different copper wires can have different oxygen absorption affinity, leading to the formation of pores within the weld seam. In the present work, no presence of gas cavities was detected while examining the samples using SEM.

The presence of intermetallic phases within the weld seam theoretically should increase its microhardness proportionally to their concentration. Previous investigations were performed [31], and the hardness of different IMCs was determined. Evidently, the lowest possible microhardness belongs to the CuAl₂ phase, with a value of 324 HV. Considering the obtained microhardness of the weld seams during the present research, it is safe to assume that most of it is comprised specifically of the CuAl₂ phase, since both the average and the momentary microhardness values do not surpass 260 HV_0.05. One of the current hypotheses is that, in the case of the sample prepared with a power of 3000 W, the best miscibility was achieved, resulting in the successful formation of stable compounds in the form of the CuAl₂ phase and a minimal quantity of other IMCs. This was, as mentioned, supported by the XRD and EDS analyses, and it is also supported by the lowest microhardness of that sample, amongst others.

The effect of the formation of IMCs on the resultant microhardness and tensile strength is well visible. As mentioned above, the tensile properties of the first two samples prepared with a power of the electron beam of 1800 W and 2400 W was essentially nonexistent. Gathering accurate data on the exact tensile properties, particularly the ultimate tensile strength, the yield strength, and elongation was impossible due to the only partial penetration of the electron beam and the formation of small, shallow weld seams. Despite the poor miscibility between the aluminum matrix and the copper inclusions, in the case of using the lower values of the power of the electron beam, some intermetallic compounds were still formed. This increased the hardness of the weld seams, but did not result in the formation of cracks, neither along its length nor in the areas between the weld seams and the base materials. Regardless, the formation of IMCs is known to have a negative effect on the mechanical strength of weld seams [32]. The bonding energy between the IMCs and the aluminum matrix is low, in comparison to the high bonding energy of the intermetallic compounds themselves [33]. In addition, the presence of the latter, particularly inhomogeneously, along the cross-section of the weld seam also increases the internal stresses within the weld seam [34]. On top of that, high stresses are normally generated in the border between the weld seam and the base materials due to the rapid change in the crystal structure. All of these factors have effects on the tensile properties of the formed welds, which, in the case of the samples prepared with a power of the electron beam of 1800 W and 2400 W, led to their immediate fracture, due to which, no measurable values were obtained. A much more stable (although still inhomogeneous) structure was observed in the case of the sample prepared with a beam power of 3000 W. Full penetration was observed, along with a preferred crystallographic orientation towards the {111} family of planes. The last is known to have the closest-packed structure with the lowest free surface energy, thus propagating the mobility of defects within the structure of the sample [35]. This is a prerequisite for increased mobility of defects such as dislocations in the structure of the welded joint and, therefore, better plastic properties. As mentioned, in that case the highest chemical activity, the best bonding between the Al matrix and the IMCs is thus theoretically observed. As a result, the ductility of the sample prepared with a power of 3000 W is the highest. That, accompanied with the highest surface area of the weld seam, led to the formation of the strongest weld seam amongst the sample. The ultimate tensile strength of this sample is low, but it still might find some possible applications in the field of electrical engineering as far as the electrical properties (which were not studied in this research) are satisfactory. From a mechanical standpoint, the formed joint does not have satisfactory mechanical properties at all and is considered completely unsuccessful.

The present work studies the influence of the power of the electron beam during electron beam welding of dissimilar metals, in particular Cu and Al6082-T6 alloys. The structure of the weld seams was investigated, along with some mechanical properties. The results, as mentioned, were mostly unsatisfactory; however, a clear correlation between the formation of IMCs and their distribution along the cross-section of the formed weld seams was established. Evidently, the application of a higher-power electron beam is imminent. This has potential promise for the formation of high-strength Cu-Al joints. Since the relationship between the formation of IMCs and the power of the electron beam was established, and previous investigations established the optimal offset and oscillation radius, the only technological condition left for experimentation is the welding speed. Future studies need to be conducted in order to establish the best approach in that regard and its influence on the resultant microstructure and properties. Additionally, as hinted, Cu-Al joints can have great applications in the field of electrical engineering. For this, studies regarding the longevity of such joints, their electrical conductivity, and potential to corrosion resistance and galvanic corrosion should also be conducted. In terms of mechanical strength, it is highly possible that reducing the concentration of IMCs in the weld seam, and also improving their spread across its cross-section, would lead to an improvement in the mechanical properties. One such way this could be achieved is the application of a filler material between the copper and the aluminum plates. Some such studies have already been proposed [13,36].

5. Conclusions

In the present work, the effect of the power of the electron beam on the microstructure, particularly formation of IMCs, microhardness, and tensile strength of Cu-Al weld seams was investigated. Three different power modes were chosen: 1800 W, 2400 W, and 3000 W. The structure and properties of the base materials were also studied. The results can be summarized in a few short postulates as follows:

(1)

Preheating the samples does not affect the structure of the base copper plate, but does affect the Al6082-T6 plate by reversing the heat treatment process.

(2)

Obviously, the application of a higher-power electron beam leads to a higher penetration, and, as a result of which, at the highest beam power of 3000 W, full penetration was achieved when considering 8 mm thick plates.

(3)

At low power, poor bonding between copper and aluminum and the formed Cu_xAl_x IMCs and the aluminum matrix was observed.

(4)

Increasing the power of the electron beam to 3000 W led to a more homogeneous spread of the formed IMCs, increased the bonding between Cu and Al, thus propagating the formation of IMCs. The last also had improved bonding with the aluminum matrix, leading to a change in the preferred crystallographic orientation of the matrix towards the {111} family of crystallographic planes.

(5)

As a result of the better bonding and the change in the preferred crystallographic orientation, improved mechanical properties of the weld seam formed at 3000 W were observed. The highest reached UTS was only 50 MPa. Much more is to be desired from the mechanical properties of the weld seams between Cu and Al6082, which could be the focus of future investigations.

Overall, the present work established the relationship between the formation of IMCs and the power of the electron beam during the electron beam welding of dissimilar metals (Cu and Al6082). This work can hopefully increase the data pool in the scientific literature, which would lead to optimization of this process and the formation of high-strength joints with excellent mechanical and electrical properties.