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Zhi-qiang Liu, Pei-lei Zhang, Shao-wei Li, Di Wu, and Zhi-shui Yu, Wire and arc additive manufacturing of 4043 Al alloy using a cold metal transfer method, Int. J. Miner. Metall. Mater., 27(2020), No. 6, pp.783-791. https://dx.doi.org/10.1007/s12613-019-1930-6
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Cold metal transfer plus pulse (C + P) arc was applied in the additive manufacturing of 4043 Al alloy parts. Parameters in the manufacturing of the parts were investigated. The properties and microstructure of the parts were also characterized. Experimental results showed that welding at a speed of 8 mm/s and a wire feeding speed of 4.0 m/min was suitable to manufacture thin-walled parts, and the reciprocating scanning method could be adopted to manufacture thick-walled parts. The thin-walled parts of the C + P mode had fewer pores than those of the cold metal transfer (CMT) mode. The thin- and thick-walled parts of the C + P mode showed maximum tensile strengths of 172 and 178 MPa, respectively. Hardness decreased at the interface and in the coarse dendrite and increased in the refined grain area.
Nowadays, metal additive manufacturing has been studied extensively. Compared with the traditional manufacturing process, metal additive manufacturing is more efficient and produces parts with more complex geometry [1]. After years of development, many other types of metal additive manufacturing technologies have emerged, such as selective laser melting, electron beam selective melting, and direct deposition of metal [2]. For the manufacture of large components, arc additive technology is efficient, low cost, and simple. In wire and arc additive manufacturing (WAAM), the wire is melted by a strongly discharged arc and deposited along the planned path to a layer-by-layer fashion [3]. The essence of arc additive manufacturing is a buildup welding process. Given their excellent mechanical properties, large Al alloy components have been widely used in aviation, automobile, railway, and other fields [4]. WAAM is a promising technology to manufacture large aluminum components [5].
During welding, various welding defects such as pores, hot cracks, and oxidization may appear in an Al alloy welding seam [6]. Therefore, an appropriate process must be adopted to ensure that large Al alloy components produced by arc additive manufacturing possess good properties. Cold metal transfer (CMT), a further development of the gas metal arc welding that relies on a controlled dip-transfer-mode mechanism, delivers beads with excellent quality and low heat input nearly without spatter [7]. CMT technology has four arc modes: CMT, CMT advanced (C + A), CMT advanced pulse (C + P + A), and CMT pulse (C + P) [8]. Many studies focused on additive manufacturing using CMT technology and CMT soldering. Cong et al. [6] researched the weld-bead geometry and porosity of AA2219-T851 Al alloy welds with different arc modes and concluded that using the C + P mode can reduce the porosity in the weld bead. Pang et al. [9] investigated the arc characteristics and metal transfer behavior of C + P welding and analyzed the characteristics of welding Al alloy. They found that the C + P mode can provide greater heat input than the CMT mode; thus, the C + P mode has greater penetration and smaller contact angle of the weld bead than the CMT mode. Gu et al. [10] found that Al‒Mg4.5‒Mn alloy parts prepared using CMT WAAM have more refined grains and better plasticity than forged Al‒Mg alloy parts. In sum, the many advantages of CMT confer great prospect in the application of WAAM. Almost no research can be found on the use of the C + P mode to manufacture thin and thick-walled parts of 4043 Al alloy.
In the present study, CMT and C + P arc were applied to WAAM. The C + P arc mode was the research focus. The scan path for manufacturing thick-walled parts was discussed. The metallurgical defects, microstructures, and mechanical properties of the parts being manufactured were analyzed. The anisotropy of mechanical properties was also discussed.
The experiments were conducted on 6061 Al alloy substrates. ER4043 with a 1.2 mm diameter was used as the research material. The wire composition is shown in Table 1 [11].
WAAM of the 4043 Al alloy was carried out using the CMT WAAM system, as shown in Fig. 1. In this system, a CMT power source was used with a VR 7000-CMT 4R/G/W/F++ wire feeder and an ABB IRB 4600 robot to provide processing movement. RobotStudio software from ABB was used to plan the path of additive manufacturing in the experiment.
| Si | Fe | Cu | Mn | Mg | Cr | Zn | Ti | Others | Al |
| 5.60 | 0.80 | 0.30 | 0.05 | 0.05 | 0.05 | 0.10 | 0.02 | 0.08 | Bal. |
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A C + P arc was developed based on the CMT arc. It was a combination of the pulse arc and the conventional CMT arc. As shown in Fig. 2, the C + P arc has several more pulse currents than the CMT arc during a welding cycle. The pulse arc increases the heat input and stirs the molten pool. The CMT and C + P modes were used in WAAM. The substrates were dried and mechanically cleaned with acetone before the experiment. The 4043 Al alloy wire was also dried. The forming path of the thin-walled part is shown in Fig. 3. Thick-walled parts of the 4043 Al alloy were prepared by the welding gun in different paths with the C + P mode.
The cross-section of the cut specimens was ground and polished. Then, the internal formation and defects of the samples were observed. The microstructures of the cross sections of thin-walled and thick-walled parts were analyzed. The samples to be analyzed were sanded, polished, and etched. The etching solution was 0.5vol% hydrofluoric acid solution. Optical microscopes were used to observe the microstructure of the treated specimen.
Tensile specimens were obtained from parts prepared using different modes. The cut mode and sample morphology of the tensile specimen are shown in Fig. 4. The obtained tensile specimen was polished to remove surface scratches. The tensile test was carried out using a Zwick/Reoll Z020 lifter manufactured by Zwick/Roell, Germany, with a stretch rate of 2 mm/min. The S-3400N scanning electron microscope (Hitachi Corp., Japan) was used to observe the fracture morphology. Hardness was determined using a HXD-1000TMC microhardness tester.
The CMT and C + P arc modes were used to manufacture thin-walled Al alloy parts. Test parameters are shown in Table 2. The thin-walled parts were manufactured as shown in Fig. 5. Difference existed in the manufacture of thin-walled parts because of various parameters.
| Specimen No. | Arc mode | Welding speed / (mm·s−1) | Wire feeding speed / (m·min−1) | Average current / A | Average voltage / V | Inter-layer waiting time / s |
| 1# | C + P | 6 | 3.3 | 66 | 14.7 | 60 |
| 2# | C + P | 6 | 4.0 | 82 | 15.1 | 60 |
| 3# | C + P | 6 | 4.2 | 96 | 15.7 | 60 |
| 4# | C + P | 7 | 4.0 | 83 | 14.4 | 60 |
| 5# | C + P | 8 | 4.0 | 83 | 14.5 | 60 |
| 6# | CMT | 8 | 4.0 | 74 | 10.9 | 60 |
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When the welding speed was constant, the wall thickness of parts 1#–3# increased with the wire-feed speed, and the effective width increased from 6.75 to 9.30 mm. As shown in Fig. 5, the weld pool showed a significant downward flow trend at the beginning and end of welding. Thus, the upper surface of the thin-walled parts was not parallel to the substrates. An angle θ (Fig. 5(a)) existed between the upper surface of the thin-walled parts and the substrates. The collapse of parts was analyzed by the following equation:
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C=2θπ×100% |
(1) |
where C is the collapse rate.
Serious collapses occurred on both sides of parts 1#–3#, and the collapse rate ranged from 13% to 20%. At a constant wire feeding speed, the effective width of parts 2#, 4#, and 5# decreased from 8.0, 7.8 to 7.4 mm as the welding speed increased. In addition, the collapse rate on both sides of the parts decreased from 16% to 3.5% as the welding speed increased. Therefore, the shape of part 5# was satisfactory.
The CMT welding machine uses synergic welding, and the current and voltage changed with the wire feeding speed. The line energy increased with the wire feeding speed. Heat was mostly conducted on both sides of the component, and the heat could be quickly transferred to the substrate. When the number of weld beads was large and parts were quite high, heat dissipation became heat conduction on both sides of the parts. Therefore, excessive heat input combined with a single heat-dissipation method resulted in the serious collapse of 1#–3# sides of thin-walled parts. Reducing the wire feeding speed can reduce the heat input. The thickness of thin-walled parts became non-uniform when the wire feeding speed was reduced. Parts 1#−3# had the same number of weld beads. As shown in Fig. 5, the height of the parts decreased with increasing wire feeding speed, which is inconsistent with the theoretical result. An increase in heat input decreased the wetting angle, and the good wetting ability of a weld bead caused molten liquid metal to flow to the side wall of the part and reduce the height of the new weld bead. The welding speed gradually increased. When the welding line energy was constant, part 5# obtained the minimum heat input. The appearance of part 6# was good, with uniform stratified streaks, no collapse on either sides, and the wall thickness was uniform. We analyzed the properties of thin-walled parts 5# and 6#.
The process parameters for preparing thick-walled parts were welding speed of 6 mm/s, wire feeding speed of 5.0 m/min, and C + P mode. Bai et al. [12] proposed an effective method to calculate the distance between adjacent weld beads. The spacing between weld beads is calculated as
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D=VSπD2S4VWH |
(2) |
where VS and DS are the wire feeding speed and wire diameter, respectively; VW is the welding speed; D is the spacing between adjacent beads; and H is the height of a single bead. In this experiment, H was 3.1 mm. Thus, D was 5.0 mm. On the basis of the C + P mode, three scanning paths were used to fabricate the thick-walled Al alloy structure. The path is shown in Figs. 6(a)–6(c). Figs. 6(d)–6(e) show the thick-walled parts made with three scanning paths. As shown in Fig. 6(e), the shape of the thick-walled part manufactured using the reciprocating scanning method was good. The parts manufactured by the method of Fig. 6(a) are shown in Fig. 6(d). A hump appeared in the thick-walled part of the arc ignition position. When welding torches were scanned perpendicular to each other between each layer, collapse occurred in parts of the arc-closing position around the thick-walled part in Fig. 6(f). Thus, the scanning paths in Figs. 6(a) and 6(c) are unsuitable for making thick-walled parts. The surface of the beads was not uniform, and some distinct ripples appeared in the front and tail of the beads [13]. These corrugations can affect the formation of thick-walled parts. The reciprocating scanning method can eliminate the ripple effect.
As shown in Fig. 7(a), the cross-section of the thin-walled part 5# formed well with no obvious defects, such as blowholes. This result has two reasons: (1) When the heat input increased, the surface area of the weld beads increased and the solidification time was prolonged, which facilitated blowhole overflows. (2) The agitation of the liquid metal with the pulse mode promoted the rate of pore overflow. The bonds between beads were perfect, and no cracks appeared between the beads. There were fuzzy fringes between each layer of weld bead. As shown in Fig. 7(b), the microstructure of a single bead of part 5# can be divided into fine- and coarse-grained regions. In addition, the widths of dendrite arms in the fine- and coarse-grained regions were 10.93 and 18.09 µm, respectively. Moreover, a small number of pores appeared in the coarse-grained regions. The relative content of pores was 0.57%. WAAM was a layer-by-layer accumulation process. The welding arc re-melted a part of the metal of the previous weld beads to form a molten pool. The pulse arc stirred the pool, and then the pore in the pool was more likely to overflow. Then, the coarse-grained region was not remelted, and the pores were left. The appearance of the fine-grained region was related to the C + P arc characteristics. In the short-circuit transition phase of the C + P pulsed arc welding process, a sudden decrease in the heat inputs caused supercooling of the liquid metal, resulting in surface nucleation and grain refinement. In addition, Wang et al. [14] reported that the high arc pressure caused by the pulsed arc mode can generate enough oscillations to cause the dendrite arms to split, thereby providing more heterogeneous nucleation sites and refining the grain size. Due to the high heat input of C + P, coarsening of the microstructure occurred in the lower half of a single bead. As shown in Fig. 7(c), the fine-grained region was mainly composed of fine columnar crystals with a rich dendritic structure. A small number of equiaxed crystals appeared between the columnar crystals. The coarse-grained regions were mainly coarse columnar crystals, and the columnar crystals grew vertically. Epitaxial growth of dendrites was observed in the bonding zone, which improved the mechanical properties of thin-walled parts.
Fig. 8 is a cross-section of thin-walled part 6#, showing that the interface of the thin-walled parts had many macroscopic and microscopic pores. The relative content of pores was 1.29%. Fracture of thin-walled parts usually occurred at these air-hole locations. The microstructure of the thin-walled part produced in the CMT mode was similar to that of the fine-grained area of the thin-walled part produced in the C + P mode, and the microstructure coarsened at the stratified fringes.
Fig. 9 shows the formation and microstructure of the cross-section of a thick-walled part. The bonding between layers of the thick-walled structure and each bead was tight. Bright stripes were found between beads. In addition, significant pores, mainly hydrogen pores, appeared at the boundary of the weld bead. The beads of the previous layer were oxidized and contaminated, which is the main reason for the generation of pores. As shown in Fig. 9(e), the microstructure of a single bead has four forms. Two coarse areas of microstructure can be found on both sides of the weld bead. Microscopically coarsened areas also existed at the upper and lower interfaces of the weld. At the lower interface, an abnormal increase in some coarse-grained regions occurred because of the superposition of two weld beads. The coarse area of these microstructures corresponded to the fusion zone (FZ) of the welded joint. The microstructure of a single bead is divided into two parts: upper fine-grained area and lower coarse-grained area. This finding is similar to the results of thin-walled parts prepared in the C + P mode.
The percentage of anisotropy of the tensile strength was calculated from the tensile test data as follows:
|
a=σB−σbσB×100% |
(3) |
where a represents the percentage of anisotropy, σB is the maximum value of the tensile strength of the tensile sample, and σb is the minimum value. Fig. 10 shows the tensile strength and anisotropy of the thin and thick-walled parts.
As shown in Fig. 10, all transverse tensile specimens had lower tensile strength than the longitudinal tensile specimens. For the manufactured thin-walled parts, an FZ appeared between beads. In the FZ region, the transverse mechanical tensile strength of the manufactured thin-walled member decreased because of the coarsening of pores and crystal grains near the FZ. In addition, the tensile strength of part 5# was greater than that of part 6#. As discussed in Section 3.2, many macroscopic pores were found inside thin-walled part 6#. Thin-walled part 5# had no obvious pores. Kobayashi et al. [15] reported that pores have a negative effect on the strength of Al alloys because the porous area diminishes the properties of the metal. The tensile load-bearing capacity of the material is reduced, stress concentration occurs near the void, and premature fracture may occur.
According to the tensile results, the difference between the transverse tensile strength and the longitudinal tensile strength of the thin-walled parts was calculated. Among them, the difference of part 5# was 7 MPa, and that of part 6# was 9 MPa. In addition, the calculated anisotropy percentage of part 5# is 4%, and that of part 6# was 5%. According to Qi et al. [3], the manufactured thin-walled parts show no anisotropy, which is important for mechanical parts.
For the manufactured thick-walled parts, the longitudinal tensile strength was 18 MPa greater than the transverse tensile strength, and the anisotropy was 10%. The mechanical properties of the thick-walled parts of the manufactured Al alloy showed anisotropy. The large number of coarsened grains in the interface region of the transversely stretched weld beads caused anisotropy.
SEM images of the fracture appearance of tensile samples are shown in Fig. 11. Longitudinal fractures had fewer pores than transverse fractures. Obviously, more pores led to lower tensile strength of transverse tensile specimens. Stratified stripes can be clearly observed from the transverse stretch-like fracture, proving that the tensile samples were fractured in the interface FZ. Stratified striations are another major reason for the reduction in the mechanical properties of transverse stretched specimens. As shown in Fig. 11, the fractured pores in thin-walled part 6# was larger in size and greater in quantity compared with those in the others. In Figs. 11(c)–11(d), a small number of dimples appeared on the surface of the tensile fracture surface, but the fracture surface was characterized by cleavage fracture as a whole. The fracture may first occur in the coarse-grained region where pores exist. The tensile fractures of Figs. 11(a)–11(b) and 11(e)–11(f) contained a large number of dimples, which indicate ductile fracture.
As shown in Fig. 12(a), the microhardness of thin-walled part 5# was between HV 45 and HV 57, and the fluctuation of the microhardness was relatively large. The microhardness of thin-walled part 6# was between HV 44 and HV 50. The microhardness test results of the thin-walled parts are related to their microstructure.
The thin-walled part 6# had a relatively uniform microstructure. Thus, their microhardness was relatively uniform and their fluctuation was small. The drop in the microhardness of the thin-walled parts only occurred in grain-coarsening FZ. As shown in Fig. 12(a), the area where the microhardness decreased was only a small percentage. The cross-sectional microstructure of thin-walled part 5# alternated between the coarsened and refined grains. Thus, the microhardness value between two indentations of a thin-walled part can differ greatly. Moreover, Fig. 12(a) shows that the microhardness of some indentations was considerably reduced because of the presence of FZ.
Fig. 12(b) shows the microhardness of the thick-walled parts at different locations. The microhardness value of the longitudinal interface was between HV 43 and HV 55. The two adjacent microhardness values fluctuated greatly. The longitudinal microhardness value was between HV 49 and HV 58, with variation similar to that of the thin-walled parts. Its value was affected by the microstructure. The test results of transverse and longitudinal microhardness were similar. Microhardness was lower at the interface. The microhardness value was greater at the non-interface.
This study investigated the influence of process parameters on the additive manufacturing of thin-walled parts and scanning paths on additive-manufactured thick-walled parts. The microstructure and mechanical properties of the manufactured parts were studied. The following conclusions can be drawn.
(1) The optimum process parameters included a welding speed of 8 mm/s, a wire feeding speed of 4.0 m/min, and a C + P pattern, which can produce well-formed thin-walled parts. With reciprocating scanning, good thick-walled parts can be obtained.
(2) In the thin-walled parts of the C + P process, the microstructure consisted of thick and fine dendrites, and few pores appeared in the cross-section. In the thin-walled parts of the CMT process, the microstructure was dendritic and many pores appeared in the cross-section. The microstructure of the thick-walled parts was similar to that of the C + P thin-walled parts, but the interface had more pores.
(3) The longitudinal tensile strength of the manufactured thin-walled parts was greater than the transverse tensile strength. The difference in pores was the main cause of the difference in tensile strength. There was almost no anisotropy in the tensile strength of the thin-walled parts. However, this anisotropy existed in the thick-walled parts. The microhardness of the thin-walled parts manufactured by the two arc modes varied from HV 44 to HV 57. The microhardness of a part reduced at the interface area of the weld bead. The microhardness in the fine-grained zone was the greatest.
This research was financially supported by the National Natural Science Foundation of China (Nos. 51605276 and 51905333), Shanghai Sailing Program (No. 19YF1418100), Shanghai Science and Technology Committee Innovation Grant (Nos. 17JC1400600 and 17JC1400601), Karamay Science and Technology Major Project (No. 2018ZD002B), and Aid for Xinjiang Science and Technology Project (No. 2019E0235).
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| Si | Fe | Cu | Mn | Mg | Cr | Zn | Ti | Others | Al |
| 5.60 | 0.80 | 0.30 | 0.05 | 0.05 | 0.05 | 0.10 | 0.02 | 0.08 | Bal. |
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| Specimen No. | Arc mode | Welding speed / (mm·s−1) | Wire feeding speed / (m·min−1) | Average current / A | Average voltage / V | Inter-layer waiting time / s |
| 1# | C + P | 6 | 3.3 | 66 | 14.7 | 60 |
| 2# | C + P | 6 | 4.0 | 82 | 15.1 | 60 |
| 3# | C + P | 6 | 4.2 | 96 | 15.7 | 60 |
| 4# | C + P | 7 | 4.0 | 83 | 14.4 | 60 |
| 5# | C + P | 8 | 4.0 | 83 | 14.5 | 60 |
| 6# | CMT | 8 | 4.0 | 74 | 10.9 | 60 |
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