Characteristics of Laser-Remelted Al–Ca–Cu–Mn (Zr) Alloys as a New Material for Additive Manufacturing

Abstract

In this study, prospects of designing new Al–Ca–Cu–Mn (Zr) alloys for additive manufacturing (AM) were evaluated for the example of laser remelting of thin-sheet rolled products. The new as-cast alloys have a hypereutectic structure containing Al₂₇Ca₃Cu₇ primary crystals and ultrafine eutectic particles of (Al,Cu)₄Ca and Al₂₇Ca₃Cu₇ phases in equilibrium with the aluminum solid solution. The solid solutions are additionally strengthened by alloying with Mn and micro additions of Zr, which contribute to the formation of coarsening-resistant phases without compromising the manufacturability of the alloys. Laser remelting, which simulates AM-typical solidification conditions, promotes the formation of a pseudoeutectic cellular structure without the occurrence of undesirable primary Al₂₇Ca₃Cu₇. The size of the dendritic cells and eutectic particles is 10 times smaller (for solidification rates of ~200 K/s) than that of the as-cast state. This structure provides for a higher hardness of the laser-remelted alloy (96 HV) as compared to the as-cast alloy (85 HV). Data for the alloy after 350–400 °C long-term annealing for up to 100 h show that the hardness of the Al–Ca–Cu–Mn–Zr alloys declines relatively slowly by ~7.5% as compared to the Zr-free alloy, whose hardness decreases by ~22%. Thus, one can consider these alloys as a promising candidate for AM processes that require high thermal stability.

1. Introduction

A review of published activity in the field of additive manufacturing (AM) of aluminum alloy products shows that the greatest attention is paid to the optimization of printing modes and the strengthening of powder compositions of conventional casting alloys (Al12Si, AlSI10Mg, and AlSi20) [1,2,3,4,5,6,7,8,9]. These alloys are characterized by higher eutectic fraction, narrow solidification range, low hot tearing tendency, and, as a consequence, improved technological properties both for casting and AM processes. AM processes develop conditions similar to those for casting (or welding), melting, and rapid solidification processes, though in micro-volumes of material.

However, the mechanical properties of these alloys after PBF-LB (Laser Beam Powder Bed Fusion) are in the 350–400 MPa range; in a number of studies, the AlSi10Mg alloy [7,8,9] exhibited a tensile strength of about 450 MPa after printing in an optimized mode. However, even short-term heating at above 300 °C leads to a significant loss of strength due to the degradation of the ultrafine eutectic structure formed during the PBF-LB.

As an alternative to silumins (Al–Si), eutectic aluminum alloys based on new alloying systems, including REM (rare-earth metal) and Ni, can be a basis for the development of new alloy compositions for additive manufacturing in an attempt to improve strength or thermal stability [10,11,12,13,14,15,16,17,18,19,20,21,22]. Thus, Al–Ni [10,12], Al–Ce [13,14,15,16], Al–Ce–Fe [17], Al–Cu–Ce [18,19], Al–Cu–Ni [20], and Al–Ce–Ni [21,22,23] systems have been actively studied, including those additionally alloyed with Mn, Zr, and Sc, for reducing the hot tearing tendency index and improving the mechanical properties during annealing due to dispersion hardening.

In [18], the new Al–3Ce–7Cu alloy was proposed for laser powder-bed fusion. After printing, its ultimate strength, yield strength, and relative elongation are 456 MPa, 274 MPa, and 4.4%, respectively. However, annealing at temperatures above 250 °C leads to a significant decrease in the hardness. In [21], the Al–6Ce–3Ni–0.7Fe alloy having a fine eutectic structure with an about 18% volume fraction of second phases was studied. This structure delivers stability of the alloy’s structure and properties at up to 300 °C. The yield strength after annealing at 300 °C is about 170 MPa. The reduction in the hardness due to annealing at 400 °C for 144 h is within 35%. In the course of further research, the Al–7.5Ce–4.5Ni–0.4Mn–0.7Zr alloy with an about 27% volume fraction of second phases was considered [22]. Mn and Zr alloying enhances the hot cracking resistance and mechanical properties of the alloy during annealing due to dispersion hardening.

However, although one can obtain a relatively successful combination of strength, heat resistance, plasticity, and manufacturability for the alloys of the considered systems via PBF-LB, it should be noted that their widespread use in industry is hindered by the high cost of REM.

According to reports data, calcium is a more attractive alloying element. Earlier results [24,25,26] showed that aluminum–calcium eutectic-based alloys exhibit a good set of such difficultly combinable properties as strength, processability (during casting [24], deformation [24], and welding [25]), and corrosion resistance [26]. These properties are achieved due to the formation of an ultrafine (Al+Al₄Ca) eutectic structure during solidification. Moreover, the magnitude of the properties can be controlled by changing the heterogeneity of the structure due to the formation of multiple calcium-containing phases with alloying elements, as reported earlier [27,28,29,30]. Furthermore, there is the possibility of obtaining a supersaturated aluminum solid solution with transition metals, e.g., Mn, Zr, and Sc, which provides hardening during heterogenization annealing due to the release of secondary metastable L1₂–Al₃(Zr, Sc) aluminides, etc. [31].

To date, certain results have been obtained on the laser processing of Al–Ca alloys in terms of modeling the structure and properties of additive manufacturing products, including powder fusion process modes, analysis of the structure and mechanical properties in the cross section of single tracks, and single studies on the structure and properties of samples obtained by laser powder-bed fusion. For example, single tracks obtained by laser powder-bed fusion showed the possibility of defect-free laser fusion of crushed powder particles of the Al6Ca4La2Mn and Al7Ca4Ni1.5Mn hypereutectic alloys [24,32]. It was shown that processing of these alloys produces an ultrafine (quasi-eutectic) composite structure containing more than 25 vol.% intermetallic phase. This structure delivers high microhardness of the alloys, i.e., 170 and 201 HV, respectively. Moreover, the hardness of the alloys remains high, dropping by only about 20% during isothermal exposure at up to 350–400 °C, unlike the standard Al–Si alloy. In [33], 10 × 10 × 10 mm³ cubic samples of the Al–10 wt.% Ca alloy were, for the first time, obtained by laser powder-bed fusion (the powder was synthesized by gas atomization). The samples exhibited high quality and density and had no cracks. The hardness of the obtained samples was in the 178–188 HV range. The new Al3.0Ce0.9Ca1.9Mn1.2Zr alloy [34] leverages the excellent heat resistance of the Al–Ce alloys, with Ca incorporation lowering both the cost and the density, and Mn and Zr addition improving the formability and increasing the precipitation strengthening of the materials. The hardness of the samples was ~124 HV. After annealing at 425 °C for 8 h, the hardness was ~115 HV.

One of the promising alloying systems for further study is Al–Ca–Cu. The new Al–Ca–Cu system is preferable for the design of aluminum matrix composites with very high contents of insoluble intermetallics, due to several quasi-binary eutectic phase transformations described earlier [30]. The alloy compositions close to the eutectic point (5.6 Ca, 4.5 Cu) have elevated contents of the (Al, Cu)₄Ca and Al₂₇Ca₃Cu₇ intermetallic phases in their structure and can be considered promising as a basis for the additive manufacturing of new materials. The above composition ensures the formation of a ~20% second-phase volume fraction. For comparison, the formation of a heterogeneous structure with a similar second-phase volume fraction in the Al–Ce–Cu alloys can be achieved at ~12% Cu and ~7% Ce. Thus, the new alloys offer a significant advantage and are economically profitable and lightweight.

Various methods are used to evaluate the prospects of new alloys in additive technologies. Along with direct laser-powder bed fusion, a promising method is laser remelting of cast or wrought surface aluminum alloy, providing for solidification conditions typical of additive manufacturing (AM) processes [18,35,36,37,38].

The aim of this work was to evaluate the prospects of new Al–Ca–Cu–Mn (–Zr) system alloys for additive technologies (laser powder-bed fusion) by simulating the structure and properties of laser-processed materials (laser remelting of thin-rolled sheets). Mn and Zr contribute to the saturation of (Al) and the formation of coarsening-resistant phases during high-temperature exposure.

2. Materials and Methods

Two model alloys, i.e., Al5Ca3Cu1.5Mn and Al5Ca3Cu1.5Mn0.4Zr, were prepared in a resistance furnace. Technical purity aluminum (99.9%), copper (99.9%), and Al-10% Ca, Al-20% Mn, and Al-15% Zr binary master alloys were used as the charge. After the melting of the basic components, the metal was cast into a graphite mold with a working cavity measuring 10 × 40 × 180 mm. The casting temperature was ~800 °C, the cooling rate being ~10 K/s. The actual chemical compositions (wt.%) of the alloys (

Table 1. Chemical composition of experimental alloys.

No.	Alloy Designation	Concentration, wt.% (target/actual)
		Al	Ca	Cu	Mn	Zr
1	Al5Ca3Cu1.5Mn	balance	5.0/4.8	3/3.2	1.5/1.4	-/-
2	Al5Ca3Cu1.5Mn0.4Zr	balance	5.0/4.7	3/3.1	1.5/1.4	0.4/0.35

) were measured using scanning electron microscopy (SEM, TESCAN VEGA 3, Brno, Czech Republic), with energy-dispersive X-ray spectroscopy (EDS, Oxford Instruments plc, Abingdon, UK) at an 20 kV accelerating voltage.

Qualitative assessment of the hot tearing tendency (HTT) was carried out with a pencil probe [24]. HTT is defined as the maximum diameter of the cylindrical portion of the sample that still does not reveal any cracks. The higher the HTT, the worse the hot cracking performance of the test aluminum alloy.

The resulting ingot was used as an initial blank to fabricate thin-sheet rolled products for further laser remelting. The sheets were obtained by hot rolling of the initial ingot at 400 °C (without annealing) to a thickness of 1 mm (compression ratio of 90%).

The laser remelting process was carried out using a robotic laser welding machine, which included a 6-axis Fanuc M710iC-50 robot, an IPG FLW D50 welding head, and a 5 kW LS-5 ytterbium fiber laser. To identify the optimal welding parameters, the laser power was varied in the range of 1000–1600 W with a spot size of 1.5 mm on the sample surface. The laser scanning speed was 20 mm/s, resulting in energy densities varying from 33 to 53 J/mm². The optimal welding parameters providing a complete remelting (without cracks, poor penetration, and other visual defects) of the thin-sheets workpiece were the laser power P = 1400 W and the scanning speed v = 20 mm/s. The focal length was 193 mm, the shielding gas (argon) supply flow rate to the remelting zone being 15 L/min.

The heat treatment (HT) of laser-remelted alloys was performed in SNOL furnaces with a temperature maintenance accuracy of 3 °C.

Scanning electron microscopy (SEM, TESCAN VEGA 3, Brno, Czech Republic), with energy-dispersive X-ray spectroscopy (EDS, Oxford Instruments plc, Abingdon, UK), and transmission electron microscopy (TEM, JEM−2100, Tokyo, Japan) were employed to examine the structure and elemental compositions of the alloys. The specimens were prepared by mechanical and electrolytic polishing. The thin foils for TEM were prepared by ion thinning with a Struers device (Struers, Copenhagen, Denmark) and studied at 200 kV.

X-ray diffraction (XRD) data were obtained at room temperature using a DRON-4 diffractometer in Co Kα radiation and analyzed with a software package [39]. XRD profiles were taken in a 2θ range of 10 to 130 deg with a 0.1 deg step and a 5 s exposure.

Thermo-Calc software version TCW5, database TTAL5 (Foundation for Computational Thermodynamics, Stockholm, Sweden), was used for thermodynamic calculations.

Microhardness measurements were conducted using DUROLINE MH-6 (Metkon, Bursa, Turkey). The Vickers indentation loads were 1 H, the dwell time being 5 s.

3. Results

3.1. Experimental Analysis of As-Cast and Hot-Rolled Structure

Figure 1 shows the as-cast microstructure of the Al5Ca3Cu1.5Mn(0.4Zr) (wt.%) hypereutectic alloy containing ~20 μm-sized primary Al₂₇Ca₃Cu₇ particles with a faceted equiaxed morphology and ultrafine (Al,Cu)₄Ca and Al₂₇Ca₃Cu₇ eutectic phase particles (Figure 2 and

Table 2. XRD for the experimental alloy.

Alloy, wt.%	Phase	Pearson Symbol	Volume Fraction, %	Lattice Parameters, Å
				a	c
Al5Ca3Cu1.5Mn0.4Zr	Al	cF4/1	80.4 ± 0.1	4.046	-
	(Al,Cu)4Ca	tI10/1	15.5 ± 0.1	4.298	11.273
	Al27Ca3Cu7	cP37/3	4.1 ± 0.0	8.420	-

) in equilibrium with the aluminum solid solution. According to the XRD data (Figure 2 and Table 2), the volume fraction of the (Al,Cu)₄Ca phase is ~15.5%, and that of Al₂₇Ca₃Cu₇ is ~4.1%. TEM microstructure imaging showed that the size of eutectic intermetallic particles is about 1–2 μm (Figure 1). Quantitative EMPA analysis showed that Mn (~1.3–1.4%) and Zr (~0.35%) are contained in the supersaturated solid solution (Al). A small fraction of manganese is dissolved in the Al₂₇Ca₃Cu₇ phase. The structure of the hot-rolled sheet metal obtained at 400 °C and a relative deformation of approximately 90% is shown in Figure 1c. The primary crystals are substantially refined by deformation treatment. The structural components are elongated in the rolling direction.

The hot tearing tendency of the Al5Ca3Cu1.5Mn0.4Zr alloys was evaluated using the pencil test samples. The experimental alloys showed lower hot cracking propensity as compared to the conventional Al-Si system alloys [24]. The hot tearing index (the minimum as-cast sample diameter, mm) is 4. The general appearance of a hot crack formed during the casting of the Al5Ca3Cu1.5Mn0.4Zr alloy is shown in Figure 3. This elevated hot cracking resistance ensures good casting properties and, as a result, presupposes high technological properties for welding and additive technologies.

Thus, in addition to the good castability, the high volume fraction of secondary phases can provide elevated strength of these alloys obtained by additive manufacturing. Laser remelting of thin-sheet products allows studying the structure and properties formed at elevated solidification rates typical of additive manufacturing.

3.2. Influence of Laser Remelting on the Microstructure

Laser remelting of thin-sheet products was carried out to further study the effect of process conditions typical of AM technologies on the structure and properties of the alloys. The appearance of the obtained 1 mm sheet samples with laser remelting field is presented in Figure 4a. SEM microstructure of track transverse view showing the possible zones formed along the molten pool and tracks quality is presented in Figure 4b. The laser treatment mode with P = 1400 W and V = 20 mm/s allowed the obtainment of a high-quality (crack-free and minimum number of pores) laser remelting field.

The microstructure of the laser remelted zone mainly consists of Al-rich gray-colored cells surrounded by a white-colored intermetallic matrix (Figure 5a,b). The average size of the dendritic cells in the melting zone is 3–4 μm, which is 8–10 times smaller than that observed for the structure obtained at conventional cooling rates, i.e., ~8–10 K/s. Using the known relationship between dendritic cell size (d) and cooling rate V = A/d [25], A equaling 100, one can estimate the cooling rate for solidification at laser welding to be approximately 20 times higher. Also, the high cooling rates heavily refine the eutectic microstructures and eliminate the coarse primary Al₂₇Ca₃Cu₇ particles. The eutectic ellipsoid particles are much finer than in the as-cast alloy, with a size of 200 nm, as measured on TEM micrographs (Figure 5c,d). This structure type is commonly considered pseudoeutectic.

Thus, the structures formed during rapid solidification contribute to a higher hardness (~96 HV) as compared to the as-cast alloy, the hardness of which is 85 HV (Figure 6a).

Another notable feature of the laser-remelted zone is a change in the chemical composition of base alloys within that zone. During laser remelting, the interaction of the high-intensity laser beam with the molten metal produces a vapor-gas channel due to the evaporation of the weld material [40,41]. The difference in the partial pressures of the elements in the alloys results in the deposition of elements on the channel walls. Some of the elements are subsequently returned to the weld pool, while other elements evaporate irreversibly.

Table 3. Chemical composition of basic and laser-remelted (LR) alloys.

Field of Investigation	Concentration, wt.%
	Ca	Cu	Mn	Zr	Al
Al5Ca3Cu1.5Mn (basic)	4.9	3.0	1.35	-	balance
Al5Ca3Cu1.5Mn (LR)	4.7	3.0	1.35	-	balance
Al5Ca3Cu1.5Mn0.4Zr (basic)	4.9	3.0	1.4	0.3	balance
Al5Ca3Cu1.5Mn0.4Zr (LR)	4.6	2.9	1.35	0.3	balance

presents the EMPA data on the chemical composition of the as-cast and laser-remelted aluminum–calcium alloy for the optimal laser treatment mode with P = 1400 W and V = 20 mm/s. It can be seen that the calcium evaporation rate in the laser-remelted field is approximately 4.1–6.1%.

3.3. Influence of Annealing on the Hardness and Microstructure

To evaluate the prospect of these alloys for high thermal stability applications, we studied the effect of heat treatment of the laser-remelted zone on hardness. As can be seen from Figure 6b (stepwise annealing), 250 °C exposure for 3 h is insufficient for changing the hardness. A negligible reduction in the hardness to 94 ± 2 may be associated with the elimination of dislocations and thermal stress relief. Further stepwise annealing at 300–400 °C caused softening of the basic laser-remelted Al5Ca3Cu1.5Mn alloy to 86 HV (~10%). However, the hardness of the laser-remelted Al5Ca3Cu1.5Mn0.4Zr alloy decreased to 92 HV (~4%).

Figure 6c,d shows the hardness evolution as a function of annealing time at 350 °C and 400 °C for the laser-remelted Al5Ca3Cu1.5Mn and Al5Ca3Cu1.5Mn0.4Zr (wt.%) alloys. It can be seen that 350 °C annealing leads to a smooth decrease in the hardness of the Al5Ca3Cu1.5Mn alloy from 96 HV to 75 HV (i.e., by ~22%) in 100 h of exposure (i.e., for the longest time used). A different annealing behavior was observed for the Al5Ca3Cu1.5Mn0.4Zr alloy: a substantially higher annealing temperature was required for a similar sharp decrease in the microhardness. The most intense hardness reduction to 91 HV (~5%) occurred during the first 10 h of exposure. Annealing for more than 40 h increased the hardness to the initial level, i.e., ~96 HV. After that, a transition to the steady stage occurred, characterized by an almost unchanged microhardness of 96 ± 2 for up to 100 h of exposure.

Annealing at 400 °C for 100 h led to the same result for the base Al5Ca3Cu1.5Mn alloy (Figure 6d). The alloy exhibited a gradual decrease in the hardness to 74 HV. However, the hardness of the Al5Ca3Cu1.5Mn0.4Zr alloy during exposure for up to 100 h demonstrated high thermal stability toward softening. The minimum hardness of 89 HV was reached after 100 h of exposure.

Analysis of the SEM and TEM micrographs (Figure 7) explains the above-identified behavior and the heat resistance of the laser-remelted Al5Ca3Cu1.5Mn0.4Zr alloy after long-term 350–400 °C exposure. Annealing causes a significant coarsening of the initial eutectic laser-remelted structure. Figure 7a–c show a growth of the intermetallic particle size to ~700–900 nm as compared to the initial laser-remelted state, for which the size of the eutectic particles was ~200 nm (Figure 5).

The TEM micrographs (Figure 7e,f) show dispersoids formed upon long-term annealing at 400 °C of the laser-remelted Al5Ca3Cu1.5Mn and Al5Ca3Cu1.5Mn0.4Zr alloys due to the decomposition of the supersaturated Mn- and Zr-containing aluminum solid solution (Al). The TEM micrographs of the as-annealed Al5Ca3Cu1.5Mn alloy reveal ~100–200 nm sized round-shaped Al₆Mn particles (Figure 7e). In the primary α-Al region of the Al5Ca3Cu1.5Mn0.4Zr alloy, there are 15 nm sized dispersoids identified as the Al₃Zr–L1₂ phase (Figure 7f).

Thus, the laser-remelted Al5Ca3Cu1.5Mn0.4Zr alloy demonstrates excellent thermal stability at 350–400 °C over the whole annealing time range (0–100 h), due to joint Mn and Zr alloying of the aluminum matrix.

It is reasonable to assume that the decrease in the hardness of the Al5Ca3Cu1.5Mn0.4Zr alloy during the first 10 h of annealing is caused by coarsening of the eutectic particles. A further increase in the hardness leads to the decomposition of the aluminum solid solution and the formation of multiple Al₆Mn and Al₃Zr–L1₂ phase dispersoids. The coarsening of eutectic particles and the decomposition of the aluminum solid solution occur simultaneously; however, the effect of the formation of precipitates allows one to offset the decrease in the hardness due to the coarsening of the eutectic particles.

4. Discussion

We finally discuss the hardening mechanism of Al5Ca3Cu1.5Mn(0.4Zr). The initial sample, laser-remelted and annealed at 400 °C, is suitable for comparative analysis for the evolution of the structure and properties. We assume that the structure of the laser-remelted Al5Ca3Cu1.5Mn0.4Zr alloy consists of 15.5 vol.% (Al,Cu)₄Ca eutectic, 4.1 vol.% Al₂₇Ca₃Cu₇ (according to XRD analysis) in the form of 100–200 nm ellipsoid particles (TEM), and Mn (1.4 wt.%) and Zr (0.35 wt.%) saturated (Al). To calculate the resulting strength of the alloy, we can use the basic additivity principle for composite properties. Based on it, the yield strength of the alloy can be calculated using a simple equation:

σ_YS = σ₀ + σ_eut + σ_SS,

(1)

where σ₀ is the yield strength of pure aluminum (~30 MPa) [21,32], σ_eut is the contribution of eutectic particles, and σ_SS is the contribution of Mn and Zr solute atoms. The strengthening effect caused by the eutectic particles can be assessed using the following Orowan equation [31]:

$$∆{\mathsf{\sigma }}_{\mathrm{e}\mathrm{u}\mathrm{t}}={\displaystyle \frac{\mathrm{M}·0.4·\mathrm{G}·\mathrm{b}}{\mathsf{\pi }·\sqrt{(1-\mathsf{\nu })}}}·{\displaystyle \frac{\ln \left(2{\mathrm{P}}_{\mathrm{e}}/\mathsf{\pi }{\mathrm{r}}_0\right)}{\mathsf{\lambda }}},$$

(2)

where M is the Taylor factor (3.06), G is the shear modulus (25.4 GPa), b is Burger’s vector (0.286 nm), ν is Poisson’s ratio (0.345), P_e is the ellipse perimeter, r₀ = 1.5b, and λ is the interparticle spacing. The interparticle spacing λ is proportional to the particle size d and inversely proportional to the particle volume fraction f_v; it can be calculated using the following equation [42]:

$$\mathsf{\lambda }=\mathrm{R}·(\sqrt{{\displaystyle \frac{2\mathsf{\pi }}{3·f_v}}}-{\displaystyle \frac{\mathsf{\pi }}{2}}),$$

(3)

Although Equation (3) applies to spherical precipitates, the application of the effective radius R of the eutectic ellipsoid particles for calculations using this equation is believed to be an acceptable approximation, provided f_v and λ are known, because the resulting approximate R values are used in Equation (2), which has a low logarithmic sensitivity to R. The same approach was previously used for the laser-processed Al–Ce alloy [13,35].

The main numerical parameters used for the calculation strength increments produced by the Orowan looping mechanism for as-laser-remelted (LR) and as-annealed (400 °C/100 h) states are resented in

Table 4. The main numerical parameters used for the calculation strength increments produced by the Orowan looping mechanism for as-laser-remelted (LR) and as-annealed (400 °C/100 h) states.

Alloy	State	M	G, GPa	b, nm	ν	Eutectic Particles			Al6Mn Precipitates			Al3Zr Precipitates
						Pe, nm	fv, %	λ	d, nm	fv, %	λ	d, nm	fv, %	λ
Al5Ca3Cu1.5Mn	LR	3.06	25.4	0.286	0.345	1153	19.6	407	-	-	-	-	-	-
	Annealed	3.06	25.4	0.286	0.345	1897	19.6	152	150	5.2	358	-	-	-
Al5Ca3Cu1.5Mn0.4Zr	LR	3.06	25.4	0.286	0.345	1153	19.6	407	-	-	-	-	-	-
	Annealed	3.06	25.4	0.286	0.345	1897	19.6	152	150	5.2	358	15	0.6	136

.

The solid solution contribution can be calculated as follows:

$${∆\sigma }_{\mathrm{S}\mathrm{S}}=\sum {\mathrm{k}}_{\mathrm{i}}·{\mathrm{C}}_{\mathrm{i}}^{2/3},$$

(4)

where C_i is the concentration of the i-th element in the solid solution and k_i is the corresponding scaling factor (for Mn k = 80 MPa/wt.%^2/3 [43], for Zr k = 9 MPa/wt.%^2/3 [44,45]). In the as-solidified state, Zr acts in α-Al as a solute whose contribution to solid-solution strengthening is expected to be negligible due to its small concentration.

The alloy structure changes significantly during annealing. Initially, the eutectic particles are significantly coarsened. The average size of the intermetallic particles is 800 nm. Furthermore, the new alloys can be strengthened by Mn- and Zr-containing precipitates. A simple thermodynamic calculation (ThermoCalc) shows that the dissolution of 1.4 wt.% Mn in (Al) should lead to the formation of about 5.2 vol.% Al₆Mn phase at 400 °C (average spherical precipitate diameter ~150 nm). Similarly, the dissolution of 0.35 wt.% Zr in (Al) should lead to the formation of about 0.6 vol.% L1₂-type phase (average precipitate diameter ~15 nm). The resulting strength of the alloy after annealing can be calculated as follows:

σ_YS = σ₀ + σ_EUT + σ_Al6Mn+ σ_Al3Zr,

(5)

The modified Orowan model for spherical precipitates [31] can be used for analyzing the nanoparticle-related increase in the yield strength. According to this model,

$$∆{\mathsf{\sigma }}_{\mathrm{A}\mathrm{l}6\mathrm{M}\mathrm{n}\left(\mathrm{A}\mathrm{l}3\mathrm{Z}\mathrm{r}\right)}={\displaystyle \frac{\mathrm{M}·0.4·\mathrm{G}·\mathrm{b}}{\mathsf{\pi }·\sqrt{(1-\mathsf{\nu })}}}·{\displaystyle \frac{\ln \left(2\mathrm{R}/{\mathrm{r}}_0\right)}{\mathsf{\lambda }}},$$

(6)

where R is the average radius of the circular cross section.

According to the calculations (

Table 5. Strength increments (MPa) produced by various strengthening mechanisms for as-laser-remelted (LR) and as-annealed (400 °C/100 h) states.

Alloy	State	σAl (MPa)	σp (MPa)	σss (MPa)	σ0.2calc (MPa)	σ0.2exp (MPa) 3 HV
Al5Ca3Cu1.5Mn	LR	30	162 1	52 4	244	288
	Annealed	30	138 (68 1/70 2)	-	168	225
Al5Ca3Cu1.5Mn0.4Zr	LR	30	162 1	52 4	244	288
	Annealed	30	230 (68 1/70 2/92 3)	-	260	267

¹ In brackets: contribution of (Al,Cu)₄Ca and Al₂₇Ca₃Cu₇ particles. ² In brackets: contribution of Al₆Mn precipitates. ³ In brackets: contribution of Al₃Zr precipitates. ⁴ In brackets: contribution of Mn-saturated (Al).

), the yield strength of the new Al5Ca3Cu1.5Mn(0.4Zr) alloy in laser-remelted state is 244 MPa. This predicted strength increment is lower than the observed hardness, since hardness (~96 HV) can be roughly related to yield strength as HV = YS/3. However, taking into account the measurement error, the correlation coefficient is still close to the known one: ~2.3–2.9. Meanwhile, calculations show the contribution of each structural element to the hardening. Thus, the eutectic particles make the greatest contribution to the strength of the alloy after laser remelting (~56%), the second contribution (~18%) pertaining to atoms (Mn) dissolved in aluminum. In the as-annealed state, the contribution of the eutectic particles to hardening is almost two times smaller (~25–30%), while the contribution of Al₆Mn (~25–30%) and Al₃Zr (~35%) dispersoids to strengthening is the greatest. Moreover, the presence of zirconium in the alloy is crucial for high-temperature processing.

In comparison with another prospect, Al8Ce0.2Sc0.1Zr [35] aluminum alloy considered for PBF-LB, the new Al5Ca3Cu1.5Mn0.4Zr is preferable due to the use of cheaper and lighter calcium instead of cerium. This combination of cheaper and lighter elements makes it possible to increase the fraction of secondary phases (8.8 vol.% vs. 19.6 vol.%). Such a structure provides a comparable level of strength properties and hardening for as-laser-remelted and as-annealed alloy at higher temperatures (325 °C vs. 400 °C/100 h). There is also no expensive scandium additive in the new alloy.

The results show that the Al5Ca3Cu1.5Mn0.4Zr model alloy is quite promising as a basis for the design of additive manufacturing alloys with high thermal stability toward softening at heating up to 400 °C.

5. Conclusions

In this study, thin sheet products of hypereutectic alloy based on the Al–Ca–Cu–Mn (Zr) system were laser-remelted to simulate the solidification conditions typical of SLM processes. The structure and properties of the material (as-cast, wrought, and laser-remelted in the as-solidified state, and annealed at 250–400 °C for 3 to 100 h) were characterized. The results are as follows:

(1)

The as-cast alloy microstructure consists of the primary Al₂₇Ca₃Cu₇ phase and (Al) + (Al,Cu)₄Ca + Al₂₇Ca₃Cu₇ ternary eutectic regions. The total intermetallic volume fraction is about 19.6%. Mn and Zr are contained in the supersaturated solid solution (Al). Laser remelting induces a much higher cooling rate (estimated as ~200 K/s) and suppresses the formation of primary Al₂₇Ca₃Cu₇ phase crystals during solidification. Laser remelting also produces a much finer α-Al dendritic structure and eutectic (Al,Cu)₄Ca and Al₂₇Ca₃Cu₇ particles. Laser-induced microstructural refinement in the Al5Ca3Cu1.5Mn(0.4Zr) alloy produces dendritic cells and eutectic particles, which are approximately 10 times finer than those typical of the as-cast structure.

(2)

During casting, the alloys demonstrate elevated resistance to the formation of hot tears. The hot tearing tendency of the Al5Ca3Cu1.5Mn0.4Zr alloy is close to that of the conventional aluminum alloys (Al–Si).

(3)

Laser processing changes the chemical composition of the aluminum–calcium alloys in the melting zone. Under optimal laser processing conditions, calcium evaporation in the melting zone can reach ~6%.

(4)

Key approaches to the design of the Al–Ca–Cu–Mn (–Zr)-based alloys for achieving enhanced mechanical properties during laser processing and subsequent annealing were demonstrated. Manganese and zirconium addition lead to hardening during exposure at 350–400 °C due to the decomposition of the (Al) solid solution and the formation of Al₆Mn particles and L1₂–Al₃(Zr) phase nanoparticles. The Al5Ca3Cu1.5Mn0.4Zr alloy demonstrates higher thermal stability of hardness during long-term (100 h) isothermal annealing (hardness drops by ~7% during annealing at 350–400 °C). For comparison, the zirconium-free quaternary alloy shows a hardness drop of up to 22%. This suggests that the test Al5Ca3Cu1.5Mn0.4Zr alloy is a promising candidate for high-temperature applications (350–400 °C).

Thus, the prospects of the new Al5Ca3Cu1.5Mn0.4Zr alloy for additive technology applications were demonstrated by simulating typical solidification conditions of laser-induced microstructural refinement. Key alloying approaches to improving the mechanical properties of alloys via laser remelting and subsequent heat exposure were demonstrated. It can be concluded that the Al–Ca–Cu–Mn–Zr alloys have prospects for AM processes which require high thermal stability. Future work based on the data obtained will include laser processing of metal-powder Al5Ca3Cu1.5Mn0.4Zr alloy compositions by laser powder-bed fusion, printing mode testing, and studying the properties of synthesized samples.