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. 2021 Oct 8;8(5):331–339. doi: 10.1089/3dp.2020.0321

Influence of Silicon and Magnesium on the Mechanical Properties of Additive Manufactured Cu-Al Alloy

Yanhu Wang 1,2, Sergey Konovalov 1,2, Xizhang Chen 1,2,✉, Arvind Singh Ramachandra 1, Jayalakshmi Subramanian 1
PMCID: PMC9828624  PMID: 36654935

Abstract

By using cold metal transfer technique, Cu-Al alloy with addition of silicon (Si) and magnesium (Mg), viz (1) Cu-6.5% Al and (2) Cu-6.5% Al-1.2% Si-0.5% Mg were manufactured additively. Four samples were deposited: Cu-6.5% Al alloy as samples 1 and 2, and Cu-6.5% Al-1.2% Si-0.5% Mg alloy as samples 3 and 4. The alloys were homogenized by heat treatments: (1) 800°C (2 h) for sample 2 and (2) sample 4, to improve their mechanical properties. Detailed microstructural investigation conducted using scanning and transmission electron microscopies showed formation of various intermetallic phases. Results revealed that (1) the addition of Si and Mg increases the strength properties and ductility and (2) heat treatments improved strength properties but reduce the ductility of the alloys. The article discusses the correlation of the identified microstructure and the evaluated mechanical properties of the additively manufactured alloys.

Keywords: additive manufacturing, Cu-Al alloy, CMT, intermetallic compounds, mechanical properties

Introduction

Cu-Al alloy has high strength, good ductility, and resistance to wear, and therefore, in recent years, much attention has been given to this alloy in terms of research and industrial applications.1–3 To date, various approaches by processing have been proposed to improve the strength and ductility of the alloy, such as cold rolling, equal-channel angular pressing, high-pressure torsion, composition design, and addition of microalloying elements.4,5 These processes can give rise to strengthening mechanisms such as solid solution strengthening, grain boundary strengthening, and precipitation strengthening that restrict dislocation movement and thereby increase the strength properties of the alloy.6 Severe plastic deformation brings forth changes such as dislocation sliding and twinning to achieve grain refinement and form nanocrystalline particles to improve the hardness and strength properties.7–9

In Cu-Al alloy, Al is a solid solution strengthener, which enhances hardness and yield strength (YS).10,11 Furthermore, the addition of microalloying elements such as Mn, silicon (Si), Ni, Zr, Cr, and Ag can produce second-phase particles by precipitation strengthening.12–15 Both the solid solution strengthening and precipitation strengthening mechanisms improve the mechanical properties by deformation twin formation and increase in dislocation density.16–19 Considering the microalloying elements, Si and magnesium (Mg) have sufficient solid solubility in copper and aluminum and significantly change the precipitation behavior and kinetics in Cu-Al alloys.20 The content of magnesium at 0.3–1% in the Cu-Al alloy can improve oxidation resistance, thermal conductivity, and electrical conductivity. Silicon addition increases the fluidity of the molten metal and reduces the formation of intermetallic phases or intermetallic compounds (IMCs).21,22

Wire arc additive manufacturing (WAAM) is a type of rapid prototyping method that provides high deposition rate and material utilization rate and saves lead time and costs, which is advantageous in manufacturing large-scale components. WAAM process has been successfully applied to fabricate Ti-Al,23 Fe-Al,24 and Cu-Al1 series alloys. Since spatter (arc splash) poses a serious problem when manufacturing Cu-Al alloy by WAAM process, an improvement in method of manufacturing is essential to overcome this problem. The cold metal transfer (CMT) inherently works with low heat input, avoids the droplet penetration, and realizes the spatter-free metal transfer. Thus, CMT-WAAM is very appropriate for the fabrication of Cu-Al alloy.

In this work, the effect of (1) addition of Si and Mg and (2) heat treatments on the microstructure and mechanical properties of Cu-Al alloy made by WAAM based on CMT was investigated. The article discusses the correlation of the microstructure and the mechanical properties of the alloys.

Experimental Procedure

The setup of the CMT-based WAAM process used to deposit models is shown in Figure 1. During the manufacturing process, an auxiliary wire feeder was used. The two wires were fed into a single molten pool under independent control of feed rate. For model 1, 99.9% Cu wire was fed by an Advanced 4000R NC CMT Welder and 99.9% Al wire was fed by a WPC-600 feeder machine. For model 2, the SAFRA-CuSi3 wire was fed by an Advanced 4000R NC CMT Welder and AlSi5-ER4043 wire was fed by a WPC-600 feeder machine. The elemental composition of these two standard welding wires is given in Table 1, as provided by the manufacturer. The welding process parameters employed for deposition are shown in Table 2. The two models were deposited as shown in Figure 2a and b. Each model is divided in two parts. The as-fabricated part of Cu-6.5% Al alloy as sample 1 and the part of Cu-6.5% Al alloy heat treated at 800°C for 2 h as sample 2. The as-fabricated part of Cu-6.5% Al-1.2% Si-0.5% Mg alloy is denoted as sample 3 and Cu-6.5% Al-1.2% Si-0.5% Mg alloy heat treated at 800°C for 2 h as sample 4. Heat treatments, that is, solution treatments, were conducted to the deposited alloys using an SX2-2.5-10 box resistance furnace.

FIG. 1.

FIG. 1.

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(a) Image of the WAAM setup. (b) Schematic of WAAM process. WAAM, wire arc additive manufacturing.

Table 1.

Elemental Composition of the Welding Wires (wt %)

  Sn Mn≤ Si Zn Cu  
SAFRA CuSi3 0.1 1.0 3.0 0.1 bal.  
  Si Mg Fe Zn Cr Mn Cu Al
AlMg5-ER5356 0.25 5 0.4 0.1 0.05 0.05 0.10 bal.
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Mg, magnesium; Si, silicon.

Table 2.

Wire Arc Additive Manufacturing Process Parameters

Parameter Model 1
 Model 2
Value Value
Deposition current, A 98 98
Deposition voltage, V 10.4 10.4
SAFRA CuSi3 speed, m/min — 8.5
AlSi5-ER4043 speed, m/min — 1.4
Cu wire, m/min 8.5 —
Al wire, m/min 1.4 —
Travel speed, m/min 0.35 0.35
Pure argon, L/min 24 24
Dwell time between deposition layers, s 40 40
Angle between the torch and filler wire 45° 45°
Distance between the torch and workpiece, mm 22 22
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FIG. 2.

FIG. 2.

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(a) The manufactured model 1. (b) The manufactured model 2. (c) Schematic showing the locations in the deposited/heat-treated alloy.

The locations in the deposited/heat-treated alloys from where samples were taken for metallographic analysis, microhardness test, and tensile tests are shown in Figure 2. Samples for hardness test, chemical composition, and metallographic observation were taken from the vertical region on the left side of the deposited sample, as shown in Figure 2c. Vickers microhardness test was conducted at locations with 2.5 mm interval between two locations, along the height of the sample starting from the bottom. A HV-1000 microhardness tester was used to measure the microhardness values at the test load of 100 g. The metallographic samples were prepared as per standard metallographic procedure, and the polished samples were etched using hydrochloric acid and ferric chloride solution (120 mL H2O + 30 mL HCl +10 g FeCl3). The microstructure of the deposited sample was observed using an optical microscope and a scanning electron microscope (SEM). The chemical composition analysis of the deposited sample was performed using energy-dispersive spectroscopy in a Phenom XL Desktop SEM. The dislocations and IMCs were analyzed using a JEOL-2100F-Cryo-transmission electron microscope (TEM). Samples for TEM analysis were taken from the rectangular location on the left side of the alloys, as shown in Figure 2c. Schematic of a tensile specimen with its dimension is shown in Figure 2c.

Results and Discussion

Phase characterization

The X-ray diffraction (XRD) spectra of the samples are shown in Figure 3. In sample 1, there are three phases, which are Cu in (111) and (200) planes, CuAl2 in (402) and (521) planes, and Cu9Al4 in (820) planes. In sample 2, there are two phases, which are Cu in (111) and (200) planes and CuAl2 in (402) and (521) planes. In sample 3, there are five phases, including SiO2 in the (111) plane, Mg2Si in the (220) plane, Cu in (111) and (200) planes, CuAl2 in (400), (402), and (521) planes, and Cu9Al4 in the (600) and (820) planes. There are four phases in sample 4, which are SiO2 in (111) plane, Mg2Si in (220) plane, Cu in (111) and (200) planes, and CuAl2 in (400), (402), and (521) planes. There are CuAl2 and Cu9Al4 phases in samples 1 and 3, whereas Cu9Al4 phase is absent in samples 2 and 4. The CuAl2 and Cu9Al4 phases are generated during the Cu-Al alloy deposition. Heat treatment can affect the intermetallic phases. During heat treatment, CuAl2 and Cu9Al4 phases enter into Cu-Al solid solution25 Upon homogenization annealing processes, the intermetallic Cu9Al4 phase is absent, whereas the CuAl2 phase decreases. When compared with Cu-Al alloy, after adding silicon and magnesium, the number of intermetallic phases increases, that is, phases of silicon and magnesium such as SiO2 and Mg2Si have appeared. The presence of these intermetallic phases can increase the hardness of the alloy.

FIG. 3.

FIG. 3.

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The XRD patterns of the samples: (a) sample 4, (b) sample 3, (c) sample 2, and (d) sample 1. XRD, X-ray diffraction.

Microstructure evolution

The copper-rich side of the Cu-Al binary equilibrium diagram is shown in Figure 4. When the Al content is <12%, the phase in the alloy is mainly α-phase. It is a solid solution based on copper and is the basic phase of Cu-Al alloy. From the Cu-rich part of the binary Al-Cu phase diagram, there are most α-phases. As the samples in the current work are Cu-rich, α-phase exists in samples 1–4.

FIG. 4.

FIG. 4.

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The binary Al-Cu phase diagram.

The microstructures at the bottom, middle, and top sections of the samples are shown in Figure 5. Columnar grains are observed in the samples. The columnar grains grow along the deposition direction, that is, from bottom to top. The directional growth of the grains is due to the multi-pass WAAM process. As the Cu-Al alloy is deposited layer by layer, the grains form and grow in the bottom layer. The temperature gradient during each layer of deposition provides the driving force for the columnar grain growth. As the build-up process continues, the melted metal in the subsequent deposition layers continues to nucleate on the previously formed columnar grains and further contributes to the grain growth. The columnar grains are larger in samples 1 and 2 than in samples 3 and 4. When compared with Cu-Al alloy, there is an increase in the intermetallic phase after addition of silicon magnesium (as shown in Fig. 3). In addition, silicon addition promotes the fluidity of the Cu-Al alloy, thus leading to the change in the structure.

FIG. 5.

FIG. 5.

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Optical micrographs for the samples, sample 1: (a) top section, (b) middle section, and (c) bottom section; sample 2: (d) top section, (e) middle section, and (f) bottom section; sample 3: (g) top section, (h) middle section, and (i) bottom section; sample 4: (j) top section, (k) middle section, and (l) bottom section.

In samples 3 and 4, black, slender, and block-shaped phase and irregular white block phase are observed that may be intermediate compounds of copper and aluminum. In sample 3, the black blocks and strips are located at the junction in between layers. The strips could be silicon compound. When compared with sample 3, the black blocks and strips are significantly less in sample 4. This may be due to the reason that after solution treatment, the black and block-shaped phases get refined into small globules and get dispersed in the copper matrix, as shown in Figure 5l.

To further illustrate the grain refinement after heat treatment. The microstructures of samples 1, 2, 3, and 4 from middle sections are shown in Figure 6. By comparing sample 1 with sample 2, and sample 3 with sample 4, there is a clear grain refinement tendency in the microstructures after heat treatment. From Figure 6, the specific grain sizes are measured according to the ASTM E23 standard. Roughly, the average grain size of sample 1 is 300 μm, sample 2 is 170 μm, sample 3 is 200 μm, and sample 4 is 130 μm.

FIG. 6.

FIG. 6.

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An enlarged metallographic diagram of a region in the middle of the sample: (a) sample 1, (b) sample 2, (c) sample 3, and (d) sample 4.

TEM was used to further analyze the organizational structure in samples 3 and 4. Second-phase particles can be seen in the TEM images (Fig. 7). In addition to the ellipsoidal and rod-like precipitates, there are elongated features in sample 3. However, only ellipsoidal precipitates and fine particles are observed in sample 4, and no elongated features are observed in this sample (as observed in sample 3) (Fig. 7a).

FIG. 7.

FIG. 7.

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(a) TEM images of the second-phase particles in sample 3. (b) TEM images of the second-phase particles in samples 3 and 4. TEM, transmission electron microscope.

X-ray microspectral analysis was used to identify the composition of the ellipsoidal particle and rod-like precipitate (i.e., second-phase particles shown in Fig. 7a) formed during the deposition of sample 3. Figure 8 shows the results of X-ray microanalysis of the second-phase particles. The ellipsoidal particle (as shown by the arrow in Fig. 8) mainly contains silicon and oxygen elements, with depleted aluminum, magnesium, and copper elements. It is most likely a compound of silicon and oxygen. The rod-like precipitate mainly contains aluminum and copper elements. It is most likely a compound of aluminum and copper, such as CuAl2 and Cu9Al4 compounds.

FIG. 8.

FIG. 8.

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Electron microscopic image of the ellipsoidal particle seen in Fig. 7a, obtained in a BF. BF, bright field.

X-ray microspectral analysis was used to identify the composition of ellipsoidal particle (i.e., second-phase particles shown in Fig. 7b) formed after the annealing of sample 3, that is, sample 4. Figure 9 shows the results of X-ray microanalysis of the ellipsoidal particle. The ellipsoidal particle is rich in silicon and magnesium and is almost free of copper and aluminum. In the XRD results, there is also Mg2Si phase, and hence, it is most likely a compound of magnesium and silicon. The Mg2Si is a strengthening phase that is reported to improve the hardness and wear resistance of Cu-Al alloy.26

FIG. 9.

FIG. 9.

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Electron microscopic image of the ellipsoidal particle seen in Fig. 7b, obtained in a BF. BF, bright field.

Hardness and composition

The measured hardness values of the samples are mentioned in Table 3 and are plotted in Figure 10. Comparing the microhardness values of samples 1 and 2, it could be observed that the hardness of the Cu-6.5% Al alloy can be increased after heat treatment. The increase in hardness is due to the intermetallic CuAl2 and Cu9Al4 formation, grain refinement, and solid solution strengthening effect.25 However, upon comparing the microhardness values of samples 3 and 4, it is observed that the hardness of sample 4 has reduced. The reason for the decrease in hardness in the Al-1.2% Si-0.5% Mg alloy after solution treatment may be that the Cu9Al4 and Mg2Si phases disappear or decrease during the solution treatment. Comparing with the Cu-6.5% Al alloy, the hardness of the Cu-6.5% Al-1.2% Si-0.5% Mg alloy with and without heat treatment is higher. Magnesium is relatively active, and in the Cu-6.5% Al-1.2% Si-0.5% Mg alloy, it easily reacts with silicon and aluminum to form compounds such as Mg2Si.27 These compounds have higher hardness and could increase the hardness of the alloy.28

Table 3.

Hardness of the Samples

No. Hardness range (Hv) Average hardness (Hv)
Sample 1 75–106 95
Sample 2 90–115 110
Sample 3 245–280 270
Sample 4 225–254 240
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FIG. 10.

FIG. 10.

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Measured microhardness values with increasing height from the bottom of the deposited (samples 1 and 3) and heat-treated alloys (samples 2 and 4).

Elemental analysis was conducted at along the height of the samples starting from the bottom. The distance between each spot for the elemental spot analyses is ∼3.5 mm. The results are shown in Figure 11. The aluminum content in the samples differs by a small amount. The mean values of the Al content for samples 1, 2, 3, and 4 are 6.0%, 6.2%, 5.9%, and 6.4%, respectively. The Si and Mg content between samples 3 and 4 vary slightly. Although the same deposition parameters were used to deposit these two samples, the slight variation may be due to the arc force, heat input, preheating effect, and cooling rate, which will influence the deposited alloy composition.

FIG. 11.

FIG. 11.

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The Al, Si, and Mg contents of the deposited (samples 1 and 3) and heat-treated alloys (samples 2 and 4) with increasing height from the bottom. Mg, magnesium; Si, silicon.

Tensile test

The room temperature ultimate tensile strength (UTS), 0.2% YS, and elongation (EL) of the deposited (samples 1 and 3) and heat-treated alloys (samples 2 and 4) are shown in Figure 12. Their UTS, 0.2% YS, and EL (%) average values are listed in Table 4.

FIG. 12.

FIG. 12.

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Tensile properties of the deposited (samples 1 and 3) and heat-treated alloys (samples 2 and 4).

Table 4.

The Tensile Properties of the Samples

No. UTS (MPa) 0.2% YS (MPa) EL (%)
Sample 1 244 60.2 33.2
Sample 2 264 65.8 34.8
Sample 3 456 70.7 7.1
Sample 4 507 79.5 8.7
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EL, elongation; UTS, ultimate tensile strength; YS, yield strength.

It can be observed from the results that the UTS, 0.2% YS, and EL increased slightly after the heat treatments. The improvement in the tensile properties is mainly due to the grain refinement. Finer the grains, more the grain boundaries, which provide resistance to the movement of dislocations. When the dislocation movement is hindered, the material strength increases. Furthermore, due to the heat treatments, solid solution strengthening occurs, the CuAl2 and Cu9Al4 transformed into Cu-Al solid solution alloy, which improves the tensile properties. Upon the addition of silicon and magnesium to the Cu-Al alloy, the UTS increases significantly, the 0.2% YS increases slightly, but the EL decreases significantly. The change in the tensile properties upon the addition of silicon and magnesium is due to the formation of hard and brittle SiO2 and Mg2Si phase, which leads to an increase in alloy strength and a decrease in ductility. Also, as heat treatment relieves internal stresses, the heat-treated samples show a slightly higher EL values when compared with the nonheat-treated samples.

SEM images of the fractured surfaces of the samples are shown in Figure 13. Samples 1 and 2 show ductile fracture features (dimples) (Fig. 13a, b). The fractured surfaces of samples 3 and 4 are relatively flat and brighter, indicating brittle fracture (Fig. 13c, d). Small flakes are distributed on the fracture surface of sample 3 (Fig. 13c) but is not observed in sample 4 (Fig. 13d). The flakes in the sample without solution treatment are likely to be silicon compound, which would initiate cracks in the alloy matrix due to their irregular morphology. When subjected to tensile stress, the increased stress concentration at the corners gives rise to microcracks. These microcracks get interconnected and propagate to cause material fracture.

FIG. 13.

FIG. 13.

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Scanning electron microscope (SEM) micrographs of fractured surfaces: (a) sample 1, (b) sample 2, (c) sample 3 (Small flakes as shown by the arrow), and (d) sample 4.

In the tensile test analysis, the performance of sample 2 was better than sample 1, and the performance of sample 4 was better than sample 3, which proves that heat treatments improve the performance of the Cu-6.5% Al and Cu-6.6% Al-1.2% Si-0.5% Mg alloys. Therefore, homogenization annealing process is beneficial in improving the strength properties of the alloys fabricated by the WAAM process.

Conclusions

In this work, Cu-6.5% Al and Cu-6.6%Al-1.2%Si-0.5% Mg alloys were additively manufactured by WAAM process. The effect of (1) Si and Mg addition and (2) solution treatments on microstructural evolution and mechanical properties of the alloys was investigated. The conclusions are as follows:

  • (1)

    The addition of Si and Mg to the Cu-6.5% Al alloy increased the hardness and tensile strength due to precipitation strengthening by the formation of intermetallic phases.

  • (2)

    The presence of intermetallic phases reduced the ductility.

  • (3)

    Homogenization annealing process of the alloys was beneficial in improving the strength properties of the alloys fabricated by the WAAM process.

Author Disclosure Statement

No conflict of interest exists in the submission of this manuscript, and manuscript is approved by all authors for publication. I would like to declare on behalf of my co-authors that the work described is original research, has not been published before, and not under consideration for publication elsewhere, in whole or part.

Funding Information

This work was sponsored by the National Natural Science Foundation of China under the Grant number 51975419.

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