Interfacial Behavior of Copper/Steel Bimetallic Composites Fabricated by CMT-WAMM

Abstract

Copper/steel bimetallic composites were made by using cold metal transfer wire and arc additive manufacturing (CMT-WAAM) with 1.2 mm diameter ER120S-G high-strength steel and 1.2 mm diameter ERCuSi-A silicon bronze welding wires. Based on the optimal tensile strength, the optimal CMT additive parameters of the copper layer were determined by the single-factor method under the conditions of the fixed steel layer process parameters of a 100 A welding current and 550 mm/min welding speed. The interfacial behavior of copper/steel bimetallic composites with the optimum parameters was investigated in particular. The results show that the optimum CMT additive process parameters for depositing a copper layer on a steel layer are a welding current of 100 A and a welding speed of 500 mm/min. The steel side consists mainly of martensite and ferrite, and the copper side consists of α-Cu matrix, Cu₃Si, and Cu₁₅Si₄ reinforcing phases. The composite interfacial region is mainly composed of the FeSi₂ reinforcing phase. At the optimum parameters, the ultimate tensile strength of the composites can reach 404 MPa with a ductile fracture on the copper side. Under the optimum parameters, the microhardness of the composites declines gradually from the steel side to the copper side, and the microhardness at the interface is higher than that at copper side, reaching 190 HV. In addition, the corrosion current density of the copper-side metal is 2.035 × 10⁻⁶ A·cm⁻², and the corrosion current density of the steel-side metal is 7.304 × 10⁻⁶ A·cm⁻². The corrosion resistance of the copper-side metal is higher than that of the steel-side metal. The CMT-WAAM process can produce copper/steel bimetallic composites with excellent comprehensive performance. The advantage of material integration makes it a broad application prospect.

1. Introduction

In recent years, copper/steel bimetallic composites have attracted great attention due to their excellent properties and wide application prospects [1,2]. Copper has good electrical and thermal conductivity, and steel boasts superior mechanical strength and durability [3,4,5]. Combining the two materials, it is possible to create composites that can withstand high loads and have good electrical and thermal conductivity properties, which are particularly important in the automotive, electrical energy, and aerospace industries.

According to the investigation, traditional manufacturing processes for copper/steel bimetallic composites included explosion composites, powder metallurgy, and centrifugal casting [6,7,8]. However, with the rise of rapid prototyping technology, some new techniques such as selective laser melting (SLM) [9,10,11], electron beam additive manufacturing (EBAM) [12,13,14], and wire and arc additive manufacturing (WAAM) [15,16,17] have gradually attracted extensive attention from scholars. Among them, Liu et al. [18] used the SLM technique to fabricate 316L stainless steel and UNS C18400 copper alloy gradient materials. The results showed good metallurgical bonding at the copper/steel interface, and significant diffusion of Fe and Cu elements was observed in the interface. The tensile strength of the copper/steel material could reach 328 MPa, and due to the high porosity of the copper-side metal, the fracture location was in the copper-side metal region. Osipovich et al. [19] manufactured copper/steel bimetallic materials using EBAM technology to achieve a gradient transition from 304 stainless steel to C11000 copper. Compared with these high-energy beam additive manufacturing processes, the preparation of bimetallic materials by arc additive manufacturing technology can overcome a series of technical difficulties, such as the large size of additive structural parts, the high cost of additive materials, and difficult additive processes [20,21,22]. Zhang et al. [23] fabricated copper/steel bimetallic composites using the gas metal arc welding (GMAW) technique. The study found that the interface between low-alloy steel and silicon bronze was prone to microcracking due to the presence of the melting unmixed zone of the α-Fe and ε-Cu two-phase structure, leading to interface failure. The introduction of Cu-Ni alloy as a transition layer improved the stability of the interface and the gradient transition, resulting in an improvement in the mechanical properties of the material with a tensile strength of 345.2 MPa. This paper demonstrated the effectiveness of utilizing the GMAW technique to prepare bimetallic materials of low-alloy steel and silicon bronze. However, in conventional arc additive manufacturing technology, the high-temperature effect of the arc often produces a large amount of spatter, which seriously affects the forming performance of the additive parts [24,25,26]. CMT technology features low heat input, no spattering, and good forming quality [27,28,29]. The introduction of CMT into arc additive manufacturing can effectively overcome the problems of severe welding distortion and low forming accuracy in traditional arc additive manufacturing [30,31]. In addition, over the long service life of the equipment, CMT welding equipment can significantly reduce material, labor, and energy costs compared to conventional arc consumables due to the potential benefits of increased efficiency, reduced material and shielding gas consumption, and lower maintenance costs. Zhang et al. [32] successfully prepared austenitic stainless steel parts with excellent tensile properties using the CMT-WAAM process. The ultimate tensile strength and yield strength of the obtained parts reached 933.3 MPa and 692.7 MPa, respectively. Wang et al. [33] also prepared austenitic stainless steel 316L parts using the CMT-WAAM process and found that the non-equilibrium organization within the layers would affect the mechanical properties of 316L parts. Nikam et al. [34] developed a duplex steel wall piece by the CMT technique using ER2594 as the welding wire. The results showed that no significant defects such as voids and unfused areas were observed in the wall piece. However, a higher ferrite content, which was 7.8% higher than the average ferrite content, was found in the root region of the re-welded wall. Wen et al. [35] produced Cu-Ni-Al-Mn-Fe aluminum bronze alloys using CMT technology for additive manufacturing. The results showed that dense, defect-free, and thin-walled aluminum bronze alloy specimens with good corrosion resistance could be obtained by using appropriate process parameters. When the temperature was increased from 20 °C to 60 °C, the stabilized potential ER of the material decreased from −0.2540 V to −0.2745 V. The self-corrosion current density increased from 2.84 × 10⁻⁶ A/cm² to 5.149 × 10⁻⁶ A/cm², indicating that it has good corrosion resistance in solution. The previous academic research studies only focused on the preparation of steel or copper monolithic material by CMT-WAAM, and there are few studies on the preparation and properties of copper/steel bimetallic materials.

Based on the above background and technological advances, the main objective of this study is to manufacture copper/steel bimetallic composites by CMT-WAAM technology and focus on the interfacial behavior of copper/steel bimetallic composites under the optimum parameters. The study of the interfacial behavior mainly includes the microstructure and properties of the interface. It is expected to provide a theoretical and practical basis for copper/steel bimetallic composites in demanding applications such as the automotive, energy, and aerospace industries.

2. Experimental Materials and Methods

A 4.0 mm thick Q235 steel plate was used as the additive substrate, and an ER120S-G high-strength steel welding wire with Φ = 1.2 mm and ERCuSi-A welding wire with Φ = 1.2 mm were used as the deposition materials. The chemical compositions of the substrate and the steel layer welding wire are shown in

Table 1. Chemical compositions of Q235 steel plate and ER120S-G welding wire (wt.%).

Materials	Main Chemical Components
	C	Si	Mn	P	S	Ni	Mo	Cr	Fe
Q235	0.17	0.30	1.30	0.001	0.001	/	/	/	Bal.
ER120S-G	0.07	1.78	0.74	0.004	0.006	2.3	0.59	0.33	Bal.

, and the chemical composition of the copper layer welding wire is shown in

Table 2. Chemical compositions of ERCuSi-A welding wire (wt.%).

Materials	C	Si	Mn	P	S	Zn	Ni	Fe	Al	Cu
ERCuSi-A	0.02	3.50	1.30	0.050	0.001	0.4	/	0.50	0.02	Bal.

. The main mechanical properties of the Q235 steel plate, ER120S-G welding wire, and ERCuSi-A welding wire are shown in

Table 3. The primary mechanical properties of Q235 steel plate, ER120S-G welding wire, and ERCuSi-A welding wire.

Materials	Tensile Strength/MPa	Yield Strength/MPa	Elongation/%
Q235	375–460	235	26
ER120S-G	1035	912	17
ERCuSi-A	345	236	40

.

Steel and copper were deposited onto a Q235 carbon steel substrate via a Fronius CMT-TPS3200 welding machine. The experiment started with 10 layers of high-strength steel stacked in layers on a steel substrate, followed by 10 layers of copper on top of the steel layers, for a total of 20 layers of metal deposited. A mixture of 80% Ar + 20% CO₂ and pure Ar gas were used to deposit the steel and copper layers, respectively, at a flow rate of 20 L/min to prevent material oxidation. The additive manufacturing process was controlled by the Mach3 CNC programmer, which is responsible for controlling the material build-up by directing the welding torch along a pre-programmed trajectory and for controlling the residence time between layers by automatically managing the starting and quenching of the torch arc. After depositing each layer, the residence time between layers was set to 5 min and the torch was then raised 2 mm to deposit the next layer. The additive manufacturing scheme is shown in Figure 1.

The process parameters were determined by mass pre-tests, with the optimum additive parameters for the ER120S-G high-strength steel layer being a 100 A welding current and 550 mm/min welding speed, and the CMT process parameters for the copper alloy layer ranged from a 70 A to 110 A welding current and a 400 mm/min to 800 mm/min welding speed.

The metallographic specimens with dimensions of 10 mm × 4 mm × 1.5 mm were taken at the copper/steel bond. After the specimens were set, ground, and polished, the steel side was corroded with 7% nitric acid alcohol solution for 8 s, the corroded metallographic specimens were cleaned with an anhydrous ethanol solution, then the copper side was corroded with aqueous ferric chloride hydrochloric acid solution for 8 s, and finally washed and dried. The microstructure of the metallographic specimens was analyzed using an OLYMPUS-CK40M optical microscope (OM, Olympus, Tokyo, Japan), S-4800 scanning electron microscope (SEM, Hitachi, Tokyo, Japan) and an accompanying energy dispersive spectrometer (EDS, Hitachi, Tokyo, Japan). The physical phase composition of the copper/steel bimetallic material was analyzed using D/max-RB X-ray diffraction (XRD, Bruker, Billerica, MA, USA) with the scanning angles ranging from 20° to 90°.

The tensile specimen was centered on the copper/steel bond, and the size of the tensile specimen is shown in Figure 2 [36]. The tensile specimens were subjected to tensile tests at room temperature using a WDW-100B universal testing machine (Rambo era, Jinan, China) with a loading force of 600.0 kN and a loading speed of 2.0 mm/min. The fractured specimens were retained after stretching and the fracture morphology was observed using SEM. The microhardness of specimens was tested using a 402MVD digital Vickers hardness tester (Wolpert Wilson, Norwood, OH, USA) with a loading force of 100 gf and a loading time of 15 s. The hardness tests were mainly concentrated near the copper/steel interface, and hardness measurements were performed at 300 μm intervals starting from the steel side towards the copper layer, with a total of 30 measurements. Finally, the electrochemical corrosion test specimens were tested for corrosion resistance using the CHI604E electrochemical workstation (Chenhua, Shanghai, China). Before the experiment, the specimens were connected to the wires and encapsulated with epoxy resin, and the surface of the specimens was polished to a mirror finish. The corrosion resistance test was carried out in 3.5 wt.% NaCl solution and the experimental temperature was set to 35 °C. The scanning interval was set to ± 250 mV, the scanning rate was 10 mV/min, and the open-circuit time was 1800 s.

3. Results and Discussion

3.1. Optimization of Process Parameters Based on Single-Factor Approach

A welding current of 100 A and a welding speed of 550 mm/min were set as the process parameters for depositing the steel layer. The additive manufacturing parameters for the good formation of copper on steel layers were obtained by pre-experiments in the welding current range of 70 A–110 A and welding speed range of 400 mm/min–800 mm/min. The optimum process parameters for the CMT additive manufacturing of copper/steel bimetallic composites were determined by the single-factor experimental method using tensile strength as a criterion. The tensile strength of copper/steel bimetallic composites under different parameters is shown in

Table 4. Tensile strength of copper/steel bimetallic composites under different process parameters.

	Welding Current	70 A	80 A	90 A	100 A	110 A
Welding Speed
400 mm/min		287 MPa	-	-	-	-
500 mm/min		278 MPa	281 MPa	385 MPa	404 MPa	352 MPa
600 mm/min		249 MPa	401 MPa	400 MPa	398 MPa	358 MPa
700 mm/min		269 MPa	391 MPa	362 MPa	398 MPa	352 MPa
800 mm/min		253 MPa	342 MPa	327 MPa	393 MPa	367 MPa

Note: “-” denotes additive parts with collapsed molten metal.

.

According to the results presented in Table 4, the tensile strength of the copper/steel bimetallic composites is lower when the welding current is set at 70 A. This decrease in performance can be attributed to the lower welding current, resulting in less melting of the cladding metal and ultimately reducing the tensile strength of the additive parts. In addition, when the welding speed is reduced to 400 mm/min, the speed becomes too slow, causing excessive deposition of the copper alloy cladding material on the steel layer. This leads to uneven spreading of the molten cladding material and results in collapse defects. Thus, the copper/steel bimetallic composites are generally well-formed under all the process parameters except for the 70 A welding current or 400 mm/min welding speed. Additionally, the tensile properties of the additive parts are better than those of the silicon-bronze base material, which has a tensile strength of 345 MPa. Figure 3 displays the fracture locations of the tensile specimens under six different process parameters. The fractures occur on the copper side of the metal, indicating that the interface between the silicon bronze alloy and the steel material is well-bonded. As can be seen from Table 4, under the condition that the process parameters of the fixed steel layer are a 100 A welding current and 550 mm/min welding speed, the optimum CMT additive process parameters of the copper layer are a 100 A welding current and 500 mm/min welding speed, because the copper/steel bimetallic composite under these process parameters have the best tensile strength, reaching 404 MPa.

3.2. Microstructure of Copper/Steel Bimetallic Composites in Different Regions

The microstructure of copper/steel bimetallic composites at a welding current of 100 A and a welding speed of 500 mm/min is shown in Figure 4. Figure 5 shows the morphology of copper/steel bimetallic composites under high magnification at the process parameters. Figure 4a and Figure 5a show that the microstructure of the copper side consists mainly of dendritic α-Cu. The formation of such dendritic microstructures is due to the rapid cooling rate and solidification process of the molten pool. The combination of an appropriate temperature gradient and growth rate during the solidification process causes the crystals to grow preferentially in a particular direction. As shown in Figure 4b and Figure 5b, the predominant microstructure on the steel side is martensite, which also contains a small amount of ferrite. During the CMT-WAAM process, due to the rapid cooling rate of the molten pool, the microstructure is transformed from austenite to coarse lath martensite, and a small amount of residual austenite becomes ferrite. The transformation can increase the hardness and strength of the steel due to the large number of dislocations in the martensitic structure which increases the resistance to movement [37,38]. As shown in Figure 4c and Figure 5c, at the interface between copper and steel, there is a clear interfacial separation, which is not only due to the differences in the chemical and physical properties of the two materials but is also related to the solidification process of both during deposition. Under a high-magnification microscope, it can be seen that the microstructure of the copper/steel interface near the steel side is significantly different from that near the copper side. This difference is mainly due to the different diffusion rates of Fe and Cu atoms at the interface, resulting in the formation of different microstructures.

3.3. Phase Analysis

The XRD spectra of copper/steel bimetallic composites at a 100 A welding current and 500 mm/min welding speed are shown in Figure 6. Figure 6b shows a magnified view of Figure 6a for a diffraction angle 2θ from 40° to 60°. It can be seen that there are main peaks in the copper/steel bimetallic composite with α-Fe as a matrix and α-Cu as a matrix. Solid solution strengthened phases containing Cu₃Si, Cu₁₅Si₄, FeSi₂, and Fe₃Si are also observed and play an important role in strengthening the material. During the CMT additive process, the CuSi-A alloy welding wire melts rapidly in the molten pool, which is cooled almost immediately due to the extremely short duration of the high-temperature effect of the arc. The solubility of silicon in copper decreases dramatically under rapid cooling conditions, resulting in the precipitation of Si atoms in the solid phase and their combination with Cu atoms to form specific alloy phases such as Cu₃Si and Cu₁₅Si₄. These phases are formed by eutectic reactions and are diffusely distributed in the α-Cu matrix, thus effectively enhancing the mechanical properties of the copper alloy.

The phase diagram of the Cu-Si binary system is shown in Figure 7. At 852 °C and with a silicon content below 14.5 wt.%, the Cu-Si binary phase diagram shows that the α-Cu solid solution structure begins to precipitate. The α-Cu alloys have good plasticity and can be cold and hot worked. However, when the Si content is between 14.5 wt.% and 36.1 wt.%, various Cu-Si solid solution phases are formed at elevated temperatures through a series of complex peritectic reactions, peritectoid reactions, and eutectic reactions. In the actual CMT-WAAM process, due to the rapid cooling rate, when the α-Cu solid solution precipitates, the Si element may undergo microsegregation during the rapid solidification process, resulting in the precipitation of other Cu-Si solid solution phases in the interdendritic space. Therefore, the second phase in the matrix mainly exists as the η″ phase, i.e., the Cu₃Si phase. The precipitation of the Cu₃Si phase improves the strength of the alloy but also reduces its plasticity.

It is evident from Figure 7 that the copper/steel bimetallic composites exhibit different preferential orientations of the phase structures. Specifically, the α-Fe crystals grow along the (110) plane, while the α-Cu crystals grow along the (111), (200), and (220) planes. Due to the different growth directions of α-Fe and α-Cu crystals, the interfacial phases form dislocation entanglements during remelting and recrystallization, which enhances the pinning effect between α-Fe and α-Cu crystals, thus improving the mechanical properties of the composite interface.

The EDS line scanning results of the copper/steel bimetallic composites at a 100 A welding current and 500 mm/min welding speed are shown in Figure 8. The left side shows the silicon bronze alloy side, while the right side shows the steel side. The results show that the interface is dominated by three elements, Fe, Cu, and Si, with the Si content at the interface being significantly higher than that on either side. Both the silicon bronze and high-strength steel welding wires contain Si elements, so the number of Si atoms at the interface reach the maximum. Si can replace the C atoms in the steel to promote the formation of martensite and can also combine with the Mn atoms to form a solid solution reinforced phase distributed in the steel matrix to strengthen steel, so that the steel has better mechanical properties. Silicon bronze has a relatively small crystallization temperature range. This property enables silicon bronze to form a more homogeneous and denser microstructure during solidification, reducing the generation of internal defects. As a result, silicon bronze exhibits higher mechanical properties than other copper alloys.

The presence of Fe atoms in the silicon bronze alloy welding wire leads to a higher proportion of Fe and Si atoms in copper alloys locally. The Fe and Si atoms form a Fe-Si solid solution under the thermal stress of welding, which influences the crystal arrangement in the silicon bronze alloy and locally strengthens the copper matrix. Additionally, the Fe element near the copper side has a higher content than the Fe atoms near the steel side due to the Kirkendall effect at the interface [40]. Fe and Cu atoms have different diffusion rates during nucleation and recrystallization at the interface between copper alloys and steel. The excess Fe atoms cause an expansion of the spatial dot matrix of Cu atoms on the Cu alloy side, and a contraction of the dot matrix occurs at the steel side due to a decrease in the number of Fe atoms at that position. As a result, the copper/steel interface is offset to the steel side, as evidenced by a greater number of Fe atoms migrating to the copper side than to the steel side. The number of Cu atoms on the steel side is also greater than that on the copper side. Due to the different diffusion rates, the migration number of Fe atoms is greater than that of Cu atoms. A similar phenomenon was found by Zhang et al. in their study of the atomic diffusion behavior of Fe-Cu in copper/steel composites [41]. It was concluded that the diffusion energy of Cu atoms hinders the further diffusion of Cu atoms into the Fe lattice due to the formation of atomic clusters. Therefore, the diffusion of Fe atoms into the Cu lattice is greater than that of Cu atoms.

The EDS surface scans performed at a 100 A welding current and 500 mm/min welding speed are shown in Figure 9. The analysis results show that the FeSi₂ reinforced phase is easily formed at the interface due to the limited solid solubility of Si in the copper alloy and the higher content of Fe atoms than Cu atoms at the interface.

Combining the EDS and XRD results, it can be concluded that the interfacial microstructure is martensite and ferrite near the Cu side and the α-Cu phase near the steel side. At the interface, there are different contents of Fe-Si and Cu-Si strengthened phases, which are manifested as Cu₃Si, Cu₁₅Si₄, FeSi₂, and Fe₃Si phases. The Fe penetration phenomenon is observed at the copper/steel interface, where the number of Fe atoms transitioning to Cu is greater than the number of Cu atoms diffusing into Fe. The presence of the Fe penetration phenomenon is conducive to the improvement in the mechanical properties of the copper/steel bimetallic composites at the interface. Combined with the XRD results, the intensity of the FeSi₂ diffraction peaks is greater than that of the Fe₃Si and Cu-Si solid solution phases, indicating that FeSi₂ is the main strengthening phase at the interface. In addition, due to the low content of Si elements on the steel side, it can be judged that the Fe₃Si phase originates from the steel side of the material, while FeSi₂ is the main form of existence of the Fe-Si solid solution phase at the interface. The mechanical properties of the interface can be improved by solid solution strengthening and dispersion strengthening.

3.4. Tensile Properties Testing and Fracture Morphology Analysis

The stress–strain curve of the copper/steel bimetallic composite at a 100 A welding current and 500 mm/min welding speed are shown in Figure 10. The tensile strength of the copper/steel bimetallic composite at this condition reaches 404 MPa and the material yields at 352 MPa. The tensile strength of this bimetallic material increases by 17% compared to silicon bronze welding wire.

The fracture morphology of the copper/steel bimetallic material under SEM is shown in Figure 11. The decrease in the fracture cross-sectional area is observed and the presence of fracture necking is found. This indicates that during the tensile deformation process, the ability of the material to resist deformation is enhanced due to the presence of the second phase within the material. Figure 11a shows the fracture morphology of the copper/steel bimetallic composites at 100× SEM magnification. There are tearing edges on the surface of the fracture, which is greyish-white and cup-cone. Figure 11b shows the fracture morphology of Figure 11a at a further magnification of 25× and it can be seen that there are a large number of dimples in the fracture. The fracture morphology of Figure 11a,b is characterized by toughness fractures, so it is judged that the fracture mode of the CMT additive of copper/steel bimetallic composites is toughness fracture. Figure 11c shows the results of detecting the content of the second phase elements in point 1 of Figure 11b, and it can be seen that the composition of the second phase is Cu-Si solid solution. Combined with the previous XRD results, the specific phase structure of the Cu-Si solid solution is Cu₃Si and Cu₁₅Si₄. In general, the formation of dimples in copper/steel bimetallic composites is related to the performance of the Cu-Si alloy, and the CMT process parameters that directly determine the metal performance of the copper side are the welding current and welding speed. Under stable heat input conditions, the second phase in α-Cu can grow uniformly, effectively improving the mechanical properties of the alloy. Therefore, the formation of dimples on the Cu alloy side is related to the precipitation of the second phase. The fracture mode of the material is affected by the size, morphology, and distribution of the second phase.

3.5. Microhardness Test

The results of microhardness measurements in different zones of the copper/steel bimetallic composites at a welding current of 100 A and a welding speed of 500 mm/min are shown in Figure 12. As can be seen in Figure 12, the average microhardness of the steel side is 350 HV, while that of the copper side is approximately 150 HV. This difference in hardness is mainly due to the fact that the high-strength steel material is enriched with a large amount of martensite on the steel side, resulting in a high microhardness, and that the steel has a higher density than the copper alloy material. The microhardness of the copper/steel interface region is better than that of the Cu side, up to 190 HV. Compared to the copper side metal, the microhardness of the copper/steel interface region is increased by 26%. This indicates that the combined mechanical properties of the microstructure at the interface are better than those of copper under the current process parameters, and the copper/steel bimetallic composites produced by CMT-WAAM have good mechanical properties at the interface.

The microhardness of the copper/steel bimetallic composites gradually decreases from the steel side to the copper side. α-Cu and Fe do not repulse or form brittle intermetallic compounds, indicating that there is a good metallurgical bond in the copper/steel interface region. The CMT-WAAM technology can produce copper/steel bimetallic components without the need for bonding methods or the introduction of transition layer metals at the interface, enabling the production of integrated structural composites, in line with today’s requirements for green, environmentally friendly, and highly efficient manufacturing processes.

3.6. Corrosion Performance Test

Electrochemical corrosion experiments were carried out on the copper/steel composites at a 100 A welding current and 500 mm/min welding speed and polarization curves were measured as shown in Figure 13. The experimental results show that the corrosion current density (Icorr) of the 120S-G high-strength steel side material is 7.304 × 10⁻⁶ A·cm⁻², while the corrosion current density of the CuSi-A silicon-bronze alloy-side material is 2.035 × 10⁻⁶ A·cm⁻². It can be seen that the corrosion current density of the metal on the high-strength steel side is higher than that on the copper alloy side. The higher the corrosion current density, the lower the charge transfer resistance, indicating that the corrosion rate of the high-strength steel materials is slightly higher than that of the copper alloy materials. In addition, the corrosion potential (Ecorr) of the 120S-G high-strength steel-side material is −0.564 V, while the corrosion potential of the CuSi-A silicon-bronze alloy-side material is −0.325 V. The difference in corrosion potential between the two is large. The more negative the corrosion potential, the higher the metal activity and the faster the electrochemical corrosion. CuSi-A silicon bronze alloy has better corrosion resistance than 120S-G high-strength steel. Due to the passivation of copper and copper alloy in seawater, Cu₂O and CuO are formed. The copper alloy metal covering the outside of the steel insulates the steel from the seawater to protect the mechanical properties of the steel material. Therefore, the CMT additive manufacturing process can be used to produce bimetallic structural materials with copper-covered steel surfaces, and the copper/steel metal interface is well combined and fully applicable to the manufacturing of marine seawater filters and other parts.

The experimental results in Figure 13c show that the corrosion current density (Icorr) in the copper/steel bimetallic interface zone is 7.992 × 10⁻⁶ A·cm⁻², and the corrosion potential (Ecorr) is −0.611 V. The results of the polarization curves in Figure 13a,c are much more similar, indicating that the corrosion of the copper/steel bimetallic interface zone tends to occur on the steel side. The principle of the seawater corrosion experiments at the copper/steel interface is based on a corrosion couple consisting of steel and copper, and the copper and steel are the cathode and anode of the corrosion couple, respectively. During the corrosion process, the corrosion of the steel material is accelerated. In addition, FeSi₂ as the main reinforcing phase at the copper/steel interface does not play a role in improving the corrosion resistance of the interface. Due to the enrichment of FeSi₂ particles’ reinforcing phase at the interface, the composition inhomogeneity at the interface increases, and the corrosion resistance of copper/steel interface decreases, resulting in corrosion occurring at the copper/steel interface. Therefore, copper/steel bimetallic composites used in the shipbuilding industry should be copper-coated steel materials, completely isolating the steel material from seawater to avoid the formation of copper/steel electrolytic cells to accelerate the corrosion of steel materials.

4. Conclusions

(1)

The optimum process parameters of the CMT-WAAM copper/steel bimetallic composites are as follows: the steel layers have a 100 A welding current and 550 mm/min welding speed, and the copper layers have a 100 A welding current and 500 mm/min welding speed. The ultimate tensile strength of the copper/steel bimetallic composite is 404 MPa and the yield strength is 352 MPa under these process parameters. The tensile strength of this bimetallic material increases by 17% compared to silicon bronze welding wire. The copper/steel bimetallic composites are fractured on the copper alloy side under each process parameter, and the fracture mode is ductile fracture.

(2)

Under the optimum process parameters, the copper/steel bimetallic composites can be divided into the steel metal zone, the copper alloy zone, and the copper/steel composite interface zone. The steel metal zone is mainly composed of martensite and ferrite. The copper alloy zone is based on the α-Cu matrix with Cu₃Si and Cu₁₅Si₄ reinforced phases. The copper/steel composite interface zone is mainly composed of the FeSi₂ reinforced phase structure.

(3)

Under the optimum process parameters, the microhardness of copper/steel bimetallic composites decreases gradually from the steel side to the copper side. The average microhardness of the steel side, copper/steel interface zone, and the copper side is 350 HV, 190 HV, and 150 HV, respectively. The corrosion current density of the CuSi-A silicon-bronze alloy-side material is 2.035 × 10⁻⁶ A·cm⁻² and the corrosion potential is −0.325 V. The corrosion resistance of the CuSi-A silicon-bronze alloy-side material is better than that of the 120S-G high-strength steel-side metal. The corrosion current density in the copper/steel bimetallic interface zone is 7.992 × 10⁻⁶ A·cm⁻² with a corrosion potential of −0.611 V. Therefore, copper/steel bimetallic composites used in the shipbuilding industry should be copper-coated steel materials, completely isolating the steel from seawater to avoid accelerating steel corrosion.