Cu合金在電弧熔絲增材制造過程中的組織與力學(xué)性能演化

1 Introduction

Compared with the laser additive manufacture (LAM) using alloy powders, the wire and arc additive manufacture (WAAM) shows the advantages of low cost, wider application range, good flexibility and very high efficiency in manufacture of the larger sized components [1-3]. Besides, due to the high cooling rate, some mechanical properties of the components manufactured by WAAM are higher than that prepared through casting and machining [4-5]. Because of these advantages, WAAM is quite promising in manufacture of some complex aerospace components, such as large frames, wallboards and pressure vessels [2, 6-7]. Whereas, the thickness of each deposition layer is usually much higher in WAAM than that in the LAM, thus a solidification segregation in a relative large range will probably occur, and the heat input of the later deposition passes can affect the former deposition passes, which will obviously change the microstructure and may results in nonuniform in the mechanical property of the components. Therefore, it is necessary to reveal the evolution in microstructure and mechanical properties of the WAAM parts during the manufacture process.

In recent years, there have been a lot of reports on WAAM using different alloys, mainly focusing on the Ni alloys, Ti alloys, Al alloys and so on [8-12]. It was found that residual stress and segregation can easily occur in the WAAM components, and the microstructure and mechanical properties of the components are affected by the thermal effect and usually not uniform [9, 11, 13]. As a kind of important structural materials, the Cu alloys show high corrosion resistance, and the strength and microhardness of the Cu alloys can be improved and conveniently adjusted through multi elements alloying [14]. For the components service in a corrosive environment for a long time, like the propeller, ship condenser and pipeline, the Cu alloys are preferred materials [15-16]. Also, the Cu alloys are favored in manufacture of the spacecraft components due to their high thermal conductivity. As the Cu has a high reflectivity to laser, it is difficult to be melted by laser and additive manufactured. In contrast, the arc weldability of the Cu alloy is quite superior and the oxidation resistance of Cu is very high [17], which makes the Cu alloys suitable to be manufactured by WAAM. Therefore, WAAM of the Cu-based alloys has been conducted, and the microstructure and properties of some Cu alloy parts have been preliminarily investigated [4, 18-20], while investigation on evolution in microstructure and properties of the Cu alloy parts during the WAAM process is relatively lacked. In a WAAM test carried out by the authors, it was found that the microstructure at the last deposited part of the Cu alloy components is not uniform. As the last deposited part usually locates at the surface, its properties are important for the component, while the evolutions in microstructure and mechanical properties of the CuAl alloy during the WAAM process have not been deeply revealed.

For the reasons above, in this study, a thin-walled CuAl alloy plate specimen was prepared using the Cu-8Al wire by WAAM, and the microstructure, composition, phase, microhardness of the specimens from the top to the inner of the plate were characterized firstly, and then the tensile properties, deformation and fracture behaviors at different locations of the plate were in-situ observed. Based on that, the evolutions mechanisms in microstructure and mechanical properties of the CuAl plate were proposed. It is hoped that this study can reveal the evolutions in microstructure and property of the WAAM Cu alloy plate, and provide some guidance for optimization of the manufacture parameters of WAAM, the post heat treatment method promotes a wider application of WAAM in manufacture of the Cu alloy parts.

2 Experimental

2.1 Specimen preparation

The material used in this study is a Cu-8Al (wt%) wire with a diameter of 1.0 mm produced by Ningbo Boway Alloy Materials Co., Ltd., which was manufactured from a horizontal continuous casted ingot of 8.0 mm in diameter through multi mode wire drawing. A WAAM equipment (Arc 405, GEFERTEC GmbH, Germany) with a Fronius cold metal transfer (CMT) melt inert-gas welding machine was employed to prepare the thin-walled CuAl alloy plate. The size of the plate is 150 mm×120 mm (additive direction)×5 mm, and it was prepared on a steel block with a size of 240 mm×180 mm×20 mm. A CAM (computer-aided manufacturing) software was used to convert the data into sliced up individual digital layer. The thickness of each layer is about 1.5 mm, and there is a pause of 10 min after each 6 layers. The WAAM parameters were as follows: current ~100 A, voltage 12.5 V, travel speed 350 mm/min and wire feeding rate 5 m/min. The temperature of the melting pool is estimated to be about 1300 ℃ and the total temperature gradient from the melting pool to the inner of the plate is about 1000 ℃. An argon side shielding gas was used to suppress the surface oxidation. The appearance of the as-fabricated plate specimen is presented in Figure 1(a), in which the interfaces between different layers can be seen and there is no cracking. No further heat treatment was conducted after the WAAM.

Figure 1  Appearance of the Cu-8Al alloy plate manufactured by WAAM (a) and sampling location and size of the tensile specimens (b)

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2.2 Microstructure and phase characterization

To characterize the microstructure and phase of the WAAM plate, a small block was sliced from the top of the plate, as presented in Figure 1(b), and then mechanically ground and carefully polished. The phase constitution of the polished sample was identified by X-ray diffraction (XRD, D8 Advance, Germany). After that, the specimen was etched with a 10% HNO3 (volume ratio) for tens of seconds at room temperature to show the alloy grains. Microstructure of the samples was observed by an optical microscope (OM, NMM-800RF) and a field emission scanning electron microscopes (FE-SEM, FEI Sirion 200 and Quanta 250), and the compositions at some locations of the specimens were analyzed using energy dispersive spectrometer (EDS). The microhardness from the top to the inner of the plate specimen was tested using a Wilson VH3300 microhardness tester at a load of 0.5 kg for 15 s, and repeated for 3 times to get the average value.

2.3 In-situ tensile test

As the thin-walled CuAl plate is only about 6 mm thick, and the microstructure transformation only exists in a narrow range at the top of the plate, the tensile specimen should be small enough to show the difference in tensile behaviors. Therefore, the tensile specimens were sliced parallel to the upper top surface, and an in-situ tensile stage was used to conduct the tensile test. The slice location and size of the in-situ tensile specimen is also shown in Figure 1(b), with the thickness of the specimen about 0.2 mm. An electric spark wire cutting machine was used for the slicing. The surfaces of the sliced tensile specimens were ground by SiC abrasive paper, and one side surface was polished for observation of the deformation morphologies. The in-situ tensile tests were conducted using a Deben Microtest 200 N tensile stage equipped on a Quanta 250 FE-SEM, and the cross-beam speed was set to be 0.5 mm/min. For the location with a certain distance to the top, three specimens were tested. Deformation morphologies of the specimens at certain strains were observed by SEM, and the fracture surfaces were also observed to comprehensively reveal the fracture mechanisms.

3 Results and discussion

3.1 Microstructure and composition

The grain structure at different locations of the longitudinal section of the CuAl plate is shown in Figure 2, with the distance from the top of the plate to the center of the images marked at the top right corner of each image. It can be found that the as-deposited CuAl alloy at the top of the plate is composed of coarse equiaxed dendrites, as presented in Figure 2(a), and the dendrites located a little longer from the top surface are columnar. Besides, there are dendrite segregations in the as-deposited region. The red lines and values show the primary dendrite arm spacing, from which it can be found that the primary dendrite arm spacing of the dendrites at the inner is much lower. With increasing distance from the top, the dendritic microstructures gradually disappear and the grains transform into columnar grains, with the axes of the columnar grains parallel to the deposition direction of the plate, as presented in Figures 2(b) and (c), the columnar grains are much finer than the dendrites. With the increase in the distance from the top, the width of the columnar grains increases a little bit, but becomes stable when the distance is over 5 mm, as shown in Figures 2(d)-(f). The blue values show the length of the blue lines and the number of grains passed through by these lines. The width of the columnar grains increases obviously with increasing distance from the surface, while the width of the columnar grains is much lower than that of the dendrites. In addition, the clear boundaries (fusion lines) between the deposition layers and the non- uniformity within each deposition layer are not so obvious as that of the parts prepared by laser additive manufacturing (LAM) or some other high melting point alloy parts prepared by WAAM [21-23], indicating that the CuAl alloy part manufactured by WAAM is more uniform in the microstructure and composition within each cladding pass.

Figure 2  Grain structures at different locations of the longitudinal section of the CuAl plate: (a) Dendrite; (b) Dendrite-columnar transition; (c, d) Fine columnar grains; (e, f) Columnar grains, with the red lines and values show the primary dendrite arm spacing, the blue values show the length of the blue lines and the number of grains crossed by these lines, and the distance from the top of the plate to the center of the image are marked at the right corner

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The longitudinal section and cross-section images at the locations with different distance from the plate top are exhibited in Figure 3, also the distances from the top are marked at the top right corner. At the location close to the plate top, the grains show a dendrite appearance, and there are high density micro pores in the dendrites, as shown in Figure 3(a). Since there is serious dendritic segregation in the rapid cooled CuAl alloy [20, 24], making the content of Al at the gaps of the dendrite arms much higher, the micro pores are preliminarily inferred to be induced by dendrite segregation of Al. In Figure 3(b), a transition state from dendrite to fine columnar grain can be seen, and the density of the micro pores decrease obviously. After the dendrites transform into columnar grains, few micro pores exist (see Figure 3(c)). The mechanisms for the disappearance of the micro pores should be the dissolution of the Al segregation under heat effect. Figures 3(d)-(f) show the cross-sectional images of the grain structures at different locations, which were observed on the surface of the tensile specimens, with higher magnification. As the columnar grains were transversally “cut”, all the grains are “equiaxed” in appearance, while the dendrite structure can still be observed at the location close to the top (see Figure 3(d)), and obvious increase in the grain size can be found through comparing Figures 3(e) and (f). Since the Cu-Al intermetallic compounds (IMCs) have high corrosion resistance [25-26], if there are Cu-Al IMCs at the Al-rich zone, there should be particles left after corrosion. In contrast, the corrosion resistance of Al is low, and the Al segregation zone is easier to be corroded if there is no Cu-Al IMC. As there are high density micro pores but few protruding precipitates in Figure 3, it is preliminary judged that there is few precipitates with high corrosion resistance in the dendrites.

Figure 3  Longitudinal section images (a)-(c) and cross-sectional images (d)-(f) at different locations of the CuAl plate, with the distance from the top of the plate to the center of the image marked at the right corner

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The EDS point analysis results of the Al element on the polished and corroded CuAl alloy surfaces of the dendrites are shown in Figure 4, with the analysis locations marked by “+”. For the polished surface, there is no micro pores, indicating that the micro pores are formed by the corrosion rather than during the solidification process. Besides, the contents of Al at the randomly selected locations are a little lower than 8% (see Figure 4(a)). For the corroded surface, micro pores can be observed, and the Al contents at the micro pores are commonly higher than 8%, while the Al contents in the matrix is lower than 8%, as shown in Figure 4(b). It is proved that there is concentration of the Al at the gap between the dendrite arms, but the content of Al at the gaps are only 2%-3% higher than that in the matrix, so the segregation is not very serious. Besides, no precipitate particles can be observed in Figure 4.

Figure 4  Contents of Al elements at some locations of the CuAl alloy surfaces obtained by EDS point analysis: (a) Polished surface; (b) Corroded surface

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3.2 Phase transformation

Figure 5 shows the XRD patterns of three locations from the surface to the inner of the CuAl plates, which were characterized from the side surface of the tensile specimens. It can be found that the patterns at the three locations are similar. For the location with a little longer distance from the surface, there is no Al segregation, and the Cu-Al alloy plate can be predicated to be a single solid solution phase, while that of the locations close to the surface and with Al segregation is not so simple, because the possibility of the existence of the        Cu-Al IMCs should be further clarified. Early investigations on the CuAl alloys reveal that there are CuAl2 and Cu9Al4 IMCs in the CuAl and CuAlSi alloys [18, 24, 27], while the content of Al is lower in this study. Although there is dendritic segregation of Al, the content of Al at the Al segregation zone is much lower than that in CuAl2 and Cu9Al4, no Cu-Al IMCs can be found at the Al segregation zone, and the peaks of the CuAl2 and Cu9Al4 phases were not detected by the XRD. Therefore, there should be little Cu-Al IMCs, and the peaks are marked with FCC Cu. The angles of the peaks at the three locations are similar, which demonstrates that the alloy is composed of the grains with only a few dominate orientations and there is no recrystallization under the heat effect. Specifically, the peaks of the (111) planes are stronger, and the peaks of the other orientations are much weaker. As the XRD patterns were characterized from the upper surface of the tensile specimens, it can be predicated that the <111> crystal orientation is parallel to the major axes of the columnar grains, i.e., vertical to the heat flow direction. Besides, grain coalesce occurs during the coarsen process of the alloy, and some grains are absorbed by their neighboring grains with dominant orientations, resulting in a little change in grain orientation distribution.

Figure 5  XRD patterns of the side surfaces of the tensile specimens from different locations of the CuAl plate

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3.3 Microhardness evolution

The microhardness from the surface to the inner of the sample is presented in Figure 6. As shown in the figure, the microhardness decreases sharply from about HV0.5160 at the surface to about HV0.590 when the distance from surface is over 1 mm. After that, the microhardness shows a slight decrease trend with increasing depth, which fits with the evolution in microstructure. In fact, with the WAAM process goes on, the earlier solidified region experienced an heating process similar to annealing when the latter deposition was conducted, making the grain coarsen and the Al segregation dissolve, thus the microhardness decreases obviously. When a region is relatively far from the WAAM melting pool (a few millimeters or so), the heat effect on this region will be quite slight and its microstructure and microhardness will become constant. For the arc welding seams of the Cu alloy, the post welding heat treatment temperature is usually around 350 ℃ to 970 ℃ [28-29], and the grain size is a little different after annealed at different temperatures. The quantitative influences of the heat treatment temperature on the grain size of the WAAM CuAl parts should be further revealed. The indentation morphologies are also shown in Figure 6, from which obvious increase in indentation size can be found from the surface to the inner.

Figure 6  Microhardness from the top to the inner of the CuAl plate

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3.4 In-situ tensile deformation and fracture behaviors

The in-situ tensile stress-strain curves of three specimens with different distances from the top are shown in Figure 7. As there are pauses during the in-situ tensile tests, slight drop in the tensile stress occurs. It is obvious that there are some differences in the yield strength, ultimate tensile strength (UTS) and elongation of the three curves. For the specimen close to the top (specimen 1), the yield strength is a little higher, while the elongation and the UTS are lower than the others, because the region close to the top solidified in a higher rate. Also, there is Al segregation in the as-deposited CuAl alloy, which can act as strength phase but decrease the ductility. With increasing distance from the top, the yield strength decreases but the elongation and the UTS increase. Whereas, the increase and decrease proportions are not very high. Although the strengthening effect disappeared after the Al segregation was eliminated, whereas, more Al atoms dissolve into the Cu matrix and the solid solution strengthening increase, making the decrease in yield strength and increase in UTS not so serious. Besides, the very top zone of the plate has much higher microhardness, but it is not included in the tensile specimen, thus the strength of specimen 1 is not so high. Since the yield strengths are only 160-170 MPa, there is still a large room for further improvement of the yield strength through adding a little precipitation strengthening elements.

Figure 7  Tensile stress-strain curves of specimens from different locations of the CuAl plate

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To further reveal the evolution in deformation and fracture behaviors, deformation morphologies of the specimens were in-situ observed. Figure 8 shows the surface deformation morphologies of the grains in specimens 1 (Figures 8(a)-(c)) and 3 (Figures 8(d)-(f)), with the strains marked at the top right corner of each figure, and the tensile direction is indicated by arrows. Since the side surfaces of the in-situ tensile specimens show the cross-section morphologies of the grains, the shapes of the grains are similar, only the grain sizes are different. In general, it can be found that the slip bands appear on the surface of the grains after yielding of the specimens, which stop at the grain boundaries, and become wider and deeper with increasing strain. Moreover, cross-slip and obvious grain rotation were also observed at a relative higher strain. In macro scope, the grain rotation makes the side surface of the tensile specimen quite uneven. For specimen 1, the slip bands are not so serious as that of specimen 3, while the grain rotation is more obvious (see the grain marked by G1), because the grain in specimen 1 is finer but the dislocations in it are more difficult to slip. Due to the dislocation pile-up, microcracks will appear at the grain boundaries, especially the triple grain boundaries, and result in the fracture [30-31], thus the elongation of specimen 1 is a little lower.

Figure 8  Surface deformation morphologies of specimen 1 (a)-(c) and specimen 3 (d)-(f) in Figure 7 at different strains, with the tensile strain marked at the right corner and the arrows show the tensile direction

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Figure 9 shows the fracture surfaces and side surface of specimen 1. In Figure 9(a), the fracture shrinkage is not very serious, while the plastic deformation around the fracture location is quite obvious. Whereas, both the deep, shallow dimples and a mixture of deep and shallow dimples can be observed at different locations of the the fracture surface, as shown in Figures 9(b)-(d), indicating that the plasticity at different locations is a bit different. In fact, although the thickness of the specimen is only about 0.2 mm, there is still nonuniformity in the microstructure for the specimen close to the plate top. For the region with higher microhardness, the ductility is lower and the dimples are more shallow, because the shallow dimples usually formed in alloys with lower ductility. For the other regions, the ductility is better and the dimples are deeper.

Figure 9  Morphologies of the side surface (a), macroscopic (b) and microscopic fracture surfaces (c, d) of the tensile specimen 1

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The fracture surfaces and side surface of specimen 3 are shown Figure 10. From the macro side surface deformation morphologies shown in Figure 10(a), it is obvious that the plastic deformation and shrinkage around the fracture location is more serious than that in specimen 1, i.e., the ductility is better. Meanwhile, few shallow dimples can be found (Figure 10(b)), indicating that the mechanical property within the specimen is relatively uniform, which fits with the microstructure and the microhardness. At higher magnification, the size of the dimples was found a little larger than that in Figure 9, as presented in Figures 10(c) and (d). As the grains are larger in specimen 3, the specimen shows better plasticity and more ductile fracture pattern.

Figure 10  Morphologies of the side surface (a), macroscopic (b) and microscopic fracture surfaces (c, d) of the tensile specimen 3

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3.5 Discussion

Based on the observation results and analysis above, a summary on the microstructure transformation of the CuAl alloy during the WAAM process is presented in Figure 11. For the as-deposited Cu-8Al alloy, dendrite structure consisting of Cu solution and Al dendrite segregation is formed due to the rapid solidification. In fact, a dendrite is usually composed of fine arms connected at their bottom, and the arms are separated by the last solidified zone with higher low melting point element or even phase boundaries. During the further deposition process, the heat effect will make the previously formed Al segregation dissolve into the Cu matrix and transform the dendrites into columnar grains, resulting a decrease in yield strength, microhardness and a little increase in ductility.

Figure 11  Schematic diagram on evolution in microstructure and mechanical properties of the WAAM prepared CuAl alloy plate

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The transformation from dendrite to columnar grain should not be recrystallization but grain boundary migration between the arms of the original dendrites, because the change in grain orientation is not serious, and further grain boundary migration will coarsen the columnar grains. As the grain structure transformation is induced by the heat effect, the major axes of the columnar grains are parallel to the heat flow direction. Besides, the melting point of the CuAl alloy is relatively low, thus the heat effect of the later deposition is enough to homogenize the microstructure, and the fusion lines and nonuniformity in microstructure within each deposition pass can be eliminated. Similar grain structure evolution has also been found in WAAM manufactured Hastelloy, duplex stainless steel [32-34], while the nonuniformity in microstructure can hardly be thoroughly eliminated in these alloys. For the other alloys, the heat input may be adjusted to make the heat effect sufficient to homogenize the microstructure and composition.

The solidification segregation is usually negative for mechanical properties and corrosion resistance of the alloys. The results of this study reveal that for the CuAl alloys, it is possible to make the microstructure at the top (or surface) similar to the microstructure at the inner through some proper heat treatment, and the shape and size of the columnar crystals can be adjusted by changing the heat dissipation conditions or further optimized through specific heat treatment [35-37]. Whereas, for the applications require high surface microhardness or a higher hardness at some locations, the surface dendrite structure can be retained.

Precipitation strengthening elements such as Cr, Zr, Ni and Si are suggested to be added into the CuAl alloy to improve the strength after aging. Through in-situ metallurgical reaction, the alloy elements can be added and their contents can be adjusted to get proper mechanical property [38]. As the heating effect during WAAM will result in a solid solution, a low temperature aging treatment after WAAM is suggested to improve the strength of the WAAM parts through promoting precipitation.

4 Conclusions

The microstructures and mechanical properties of the CuAl part prepared by WAAM were investigated in this study. Based on the experimental results and discussions, the following conclusions can be drawn:

1) The as-deposited CuAl alloy is composed of coarse dendrites, and the dendrites gradually transform into relatively fine columnar grains by heat effect of the later deposition passes, then the size of the columnar grains increases and becomes stable at a distance of a few millimeters from the top.

2) The WAAM CuAl is mainly a single FCC Cu solution phase, with some dendrite segregation of Al but no Cu-Al IMCs in the as-deposited zone, which can be eliminated under the heat-effect and the composition and microstructure can be homogenized.

3) With the transformation in microstructure from the top to the inner of the CuAl plate, the microhardness decreases from HV160 to HV80, while the plate shows high tensile strength

(~400 MPa) and high elongation (~50%), and fail a dimple mode. Addition of some precipitation strengthening elements is suggested to further improve the mechanical property of the CuAl alloy.