Study on Microstructure and High-Temperature Mechanical Properties of Al-Mg-Sc-Zr Alloy Processed by LPBF

Abstract

Al-Mg-Sc-Zr alloy processed via laser powder bed fusion (LPBF) is poised for significant application in aerospace, where its high-temperature capabilities are paramount for the safety and longevity of engineered structures. This study offers a systematic examination of the alloy’s high-temperature tensile properties in relation to its microstructure and precipitate phases, utilizing experimental approaches. The LPBF-processed Al-Mg-Sc-Zr alloy features a bimodal microstructure, with columnar grains in the melt pool’s interior and equiaxed grains along its boundary, conferring exceptional properties. The application of well-calibrated processing parameters has yielded an alloy with an impressive relative density of 99.8%, nearly fully dense. Following a thermal treatment of 350 °C for 4 h, the specimens were subjected to tensile tests at both room and elevated temperatures. The data reveal that the specimens exhibit a tensile strength of 560.6 MPa and an elongation of 11.1% at room temperature. A predictable decline in tensile strength with rising temperature is observed: at 100 °C, 150 °C, 200 °C, and 250 °C; the respective strengths and elongations are 435.1 MPa and 25.8%, 269.4 MPa and 20.1%, 102.8 MPa and 47.9%, 54.0 MPa and 72.2%. These findings underpin the technical rationale for employing LPBF-processed Al-Mg-Sc-Zr alloy in aerospace applications.

1. Introduction

Laser powder bed fusion (LPBF) is an innovative additive manufacturing technique that integrates 3D modeling, material design, and automated control from various disciplines. This technology enables the rapid formation of parts from metal powders, eliminating the need for molds, machining, and assembly processes traditionally required in manufacturing. LPBF is a highly material-efficient process. By melting powder layer by layer, it minimizes material waste, which is particularly crucial for expensive and rare alloy materials. Moreover, LPBF technology significantly shortens production cycles and reduces tooling costs, making prototype design and small-batch production more economically feasible. In the aerospace sector, this facilitates rapid iteration and design optimization, accelerating the implementation of new technologies and materials. The characteristic of forming parts in a single step allows for weight reduction through integrated structural design, as demonstrated in recent studies [1,2]. LPBF-processed aluminum alloys, known for their high specific strength, excellent machinability, and superior corrosion resistance, are gaining significant traction in industries such as aerospace and automotive manufacturing due to their broad application prospects [3,4,5,6].

In recent years, an advanced Al-Mg-Sc-Zr alloy has emerged as an ideal candidate for LPBF. This alloy stands out as a multifunctional material, demonstrating exceptional repeatable printability and superior mechanical properties. Derived from an Al-Mg-based alloy modified with the addition of Sc and Zr elements, it has successfully been commercialized in aerospace and automotive industries, exemplified by Scalmalloy^®—a material with the composition Al-4.5Mg-0.75Sc-0.35Zr-0.45Mn [7,8,9,10]. Current research indicates that LPBF-produced Al-Mg-Sc-Zr alloys can achieve a yield strength (YS) of up to 500 MPa at room temperature (RT), while maintaining an elongation at fracture (ε_f) of over 10%. The high strength of the Al-Mg-Sc-Zr alloy is attributed to its microstructure, which features a bimodal grain structure, known as a dual-phase structure. This consists of fine grains (FG) distributed along the melt pool boundaries and coarse grains (CG) within the melt pool interior. The fine grain regions are formed by the precipitation of Al₃(Sc, Zr) phases during solidification at the melt pool boundaries. These precipitates have lattice parameters similar to those of α-Al and a low lattice mismatch, making them ideal nucleation sites for Al grains, thus promoting heterogeneous nucleation [11]. This approach significantly mitigates the issue of cracking, providing effective nucleation sites for the crystallization of α-Al grains, thereby reducing the formation of columnar crystals [12,13]. The heterogeneous nucleation of columnar grains within the melt pool is attributed to the rapid cooling from temperatures above 800 degrees Celsius, which generates a substantial temperature gradient [14,15,16]. Existing research indicates that optimizing LPBF process parameters—such as laser power, scan speed, layer thickness, hatch spacing, powder layer thickness, and scan path—can significantly enhance the alloy’s density and microstructure, thereby improving its mechanical properties [17,18,19,20,21,22,23]. The research demonstrates that Al-Mg-Sc-Zr alloys processed by the LPBF process possess exceptional mechanical performance and robust corrosion resistance under room-temperature conditions. Despite the extensive studies on LPBF-processed Al-Mg-Sc-Zr alloys, further research is needed to understand their tensile properties in high-temperature environments, particularly the microstructural evolution and changes in mechanical properties under varying temperature conditions.

Researching the high-temperature tensile properties of Al-Mg-Sc-Zr alloy processed by LPBF holds significant practical importance and potential industry impact. High-temperature environments are common in the aerospace sector, making it crucial to study the tensile properties of this alloy at elevated temperatures for designing and manufacturing highly reliable aerospace components. Furthermore, understanding the microstructural evolution and mechanical property changes of the alloy under different temperature conditions aids in optimizing LPBF process parameters, further enhancing material performance and manufacturing quality. This study systematically investigates the tensile property anisotropy of LPBF Al-Mg-Sc-Zr alloy in the as-built (AB) state and after annealing at 350 °C for 4 h (as-aged, AA). The 4 h annealing process is also considered as a heat treatment to alleviate residual stresses, which are a known issue for LPBF components. In light of this, the research characterizes the microstructure of the LPBF-formed Al-Mg-Sc-Zr alloy, performs tensile experiments at room and elevated temperatures. It examines the tensile properties and fracture morphology, establishing a relationship between microstructure, anisotropy, and high-temperature mechanical performance, which is significant for the application and advancement of LPBF-processed Al-Mg-Sc-Zr alloys.

2. Materials and Methods

2.1. Powder and LPBF Parameters

The Al-Mg-Sc-Zr powder utilized in this study was graciously supplied by Baohang Advanced Material Co., Ltd. (Nanchang, China) This powder was meticulously produced through gas atomization technology, a method specifically optimized for use with LPBF machinery in 3D printing applications. The scanning strategy employed adheres to the machine’s default checkerboard pattern. A detailed composition of the material is presented in Table 1, which reveals that the measured content of each constituent within the Al-Mg-Sc-Zr powder falls well within the prescribed standard range. For the processing of the Al-Mg-Sc-Zr powder, the HBD-450 3D printing equipment, provided by Hanbang Laser Technology Co., Ltd. (Zhongshan, China), was chosen. The specific printing parameters are delineated in Table 2. Additionally, the volume energy density (VED) is determined through the application of the formula presented below [24]:

$$VED={\displaystyle \frac{P}{V\times h\times t}}$$

P is the laser power, V is the laser scanning speed, h is the hatching distance, and t is the layer thickness.

Table 1. Chemical composition of Al-Mg-Sc-Zr powder.

Elements	Mg	Sc	Zr	Mn	Si	Al
wt%	7.80	0.59	0.31	0.47	0.66	Bal

Table 1. Chemical composition of Al-Mg-Sc-Zr powder.

Elements	Mg	Sc	Zr	Mn	Si	Al
wt%	7.80	0.59	0.31	0.47	0.66	Bal

Table 2. Printing parameters of Al-Mg-Sc-Zr powder.

Laser Power	Scanning Speed	Layer Thickness	Hatching Distance	VED
400 W	1250 mm/s	50 μm	120 μm	53.33 J/mm3

Table 2. Printing parameters of Al-Mg-Sc-Zr powder.

Laser Power	Scanning Speed	Layer Thickness	Hatching Distance	VED
400 W	1250 mm/s	50 μm	120 μm	53.33 J/mm3

The particle size distribution of the powder ranges from 12 to 72 μm, with the specific diameters at 10%, 50%, and 90% cumulative distribution (d10, d50, d90) being 19.53 μm, 35.18 μm, and 57.68 μm, respectively. The morphology of the powder is depicted in Figure 1, revealing a high sphericity and smooth surface overall. However, minor instances of particle agglomeration and a small number of irregularly shaped particles are observed. Such irregularities can impair the powder’s flowability, leading to uneven spreading and consequently reducing the density of the formed parts. Samples were annealed in a muffle furnace at 350 °C for 4 h, as illustrated in Figure 2. This heat treatment results in the supersaturation of Mn within the matrix and the formation of nanoscale Al₃Sc precipitates, which contribute to solution strengthening, grain refinement, and precipitation strengthening.

2.2. Microstructure Characterization

The density of each sample was measured using the Archimedes drainage method. The formed samples were wire-cut and subsequently polished using progressively finer sandpapers to obtain standard metallographic specimens. These specimens were etched for 15 s with Keller’s reagent (1 mL HF, 1.5 mL HCl, 2.5 mL HNO₃, and 95 mL H₂O) for surface morphology and microstructure observation using optical microscopy (OM, Axio Observer, Carl Zeiss AG, Oberkochen, Germany) and scanning electron microscopy (SEM, GEMINI 300, Carl Zeiss AG, Oberkochen, Germany). For electron backscatter diffraction (EBSD) analysis, the samples were electrolytically polished at −50 °C using an electrolyte solution composed of 5% perchloric acid and 95% ethanol. The grain size and crystallographic orientation data were extracted using AZtecCrystal software (v2.0, Oxford Instruments NanoAnalysis, High Wycombe, UK). In this study, high resolution transmission electron microscopy (HRTEM, TF20, FEI Company, Hillsboro, OR, USA) was used to examine the second-phase particles in the alloy’s microstructure. Samples were prepared by wire-cutting 10 mm × 10 mm × 0.3 mm sections, which were then ground to a thickness of 150 μm using sandpaper. Subsequently, a 3 mm diameter disc was punched from the thinned section. The final stage of sample preparation involved electrolytic twin-jet polishing, using an electrolyte consisting of 30% nitric acid (HNO₃) and 70% methanol (CH₃OH), cooled to −30 °C. This process was carried out at an applied voltage of 20 V and a current of 120 mA. The samples were immediately rinsed with anhydrous ethanol after polishing. The prepared specimens were then subjected to HRTEM analysis.

2.3. Experimental Method of Mechanics

To test the tensile anisotropy of the Al-Mg-Sc-Zr alloy processed by LPBF, tensile specimens were categorized into H direction and Z direction based on the building direction, as shown in Figure 3a. The tensile tests were conducted according to the GB/T 228.2-2015 standard, with specimen dimensions illustrated in Figure 3b,c, applying an engineering strain rate of 7 × 10⁻⁵ s⁻¹. During high-temperature tensile tests, the specimens were preconditioned at the specified temperature for 30 min before loading. To ensure data reproducibility, each tensile test was performed three times, and the average value was taken.

3. Results and Discussion

3.1. Microstructure Observations

Figure 4 displays the optical micrographs (OM) of samples from the Al-Mg-Sc-Zr alloy. The relative density of the samples, determined by the Archimedes method, reached 99.8%. The microstructure showed no visible cracks, confirming the successful fabrication of high-density samples using LPBF technology. Figure 5 provides OM images of the etched Al-Mg-Sc-Zr alloy samples. In the XOY and XOZ planes, a distinct dark stripe was observed in the brighter regions, indicating the boundary of the melt pool, while the brighter regions represent the interior of the melt pool. The shape of the melt pool was unevenly distributed in the XOY and XOZ planes. In the XOY plane, the melt pool appeared as irregular ellipses. In the YOZ plane, the melt pool exhibited a fan shape, characterized by a regular, fish-scale-like arrangement. Measurements of the observed surface indicated that the size of the melt pools ranged from 20 μm to 120 μm.

Figure 6a displays the X-ray diffraction (XRD, D8 Advance, Bruker Corporation, Karlsruhe, Germany) patterns of the pristine powder and the sample produced by the laser powder bed fusion (LPBF) technique. The XRD analysis confirms that the α-Al phase constitutes the principal crystalline phase in both the powder and the LPBF-processed samples. Notably, no Mg-containing phases or other phases resulting from processing were identified in the XRD patterns, possibly due to their low abundance below the detection capabilities of XRD or the rapid scanning speed of the XRD equipment, which could hinder the detection of these phases. A significant observation is the increased intensity of the diffraction peak on the (111) crystal plane in the LPBF-processed sample when compared to the original powder. This suggests a pronounced preferred orientation on the (111) plane, likely a consequence of the thermal processes and crystallographic growth during LPBF, which can significantly affect the resulting microstructure and material properties. Figure 6b offers a detailed perspective of the XRD pattern between 37.5° and 39°. It is evident that the (111) diffraction peak of the α-Al phase in the sample prepared by LPBF has undergone a leftward shift relative to the standard α-Al peak at 2θ = 38.446°. This shift results in the (111) diffraction peaks aligning with lower 2θ values [25]. As dictated by Bragg’s law, such a displacement to the left suggests an increase in the lattice parameters. The rapid solidification process characteristic of LPBF does not allow for timely diffusion of the majority of Mg atoms. These atoms are instead quickly assimilated into the α-Al matrix as solutes, causing substantial lattice distortion. This distortion is responsible for the observed leftward shift of the (111) peak.

To further investigate the microstructural properties of the samples processed by LPBF, such as grain size, shape, and texture, EBSD analysis was conducted. Figure 7 shows the EBSD inverse pole figures and pole figures of the XOY and XOZ planes of the alloy. As shown in Figure 7a, the ultrafine equiaxed grains of the Al-Mg-Sc-Zr alloy are distributed around the periphery of the melt pool on the XOZ plane, while elongated grains are mainly found within the melt pool. Figure 7b reveals that the columnar crystals in the melt pool exhibit a preferred orientation along <100>. Additionally, the melt pool depth is approximately 150 μm, significantly exceeding the alloy’s 50 μm layer thickness. This is attributed to the excessively high melt pool temperature during the LPBF process, which causes remelting of previously solidified material during subsequent layer scans [26,27].

The SEM images in Figure 8 offer insights into the microstructure of the Al-Mg-Sc-Zr alloy crafted by LPBF. Figure 8a displays large columnar crystals in the heart of the molten pool, with several black corrosion pits evident at the pool’s edges. A high-magnification SEM image, Figure 8b, captures nanoscale white precipitated phase particles within and encircling the corrosion pits at the pool’s edge. These particles are postulated to be nano-Al₃(Sc, Zr), as per the analysis in Reference [28]. The temperature gradients within the molten pool, under suitable process parameters, promote the precipitation of these thermally stable Al₃(Sc, Zr) particles [29]. As such, these particles precipitate stably and accumulate at the periphery of the molten pool during the repeated cycles of melting and solidification.

The HRTEM-BF-S images in Figure 9, complemented by EDS spectra, provide a detailed examination of the fine grain zone in the 3D printed samples, highlighting the concentration of Sc and Zr elements around this area. The precipitated secondary phase particles are identified as Al₃(Sc, Zr) based on HRTEM and SAED analyses. The LPBF process involves a laser scanning a powder bed, with each layer being elevated and melted in sequence to achieve a layered scanning and melting effect. The formation of a molten pool upon laser interaction with Al powder results in the distinctive fish-scale pattern observable in the alloy’s microstructure (as shown in Figure 5). The cooling rate of these molten pools is exceptionally high, between 10³ and 10⁶ K/s [30]. As the molten pool cools, solidification initiates at the edges due to uneven heat dissipation. The hypereutectic elements Sc and Zr form a fine, dispersed L1₂ phase with Al, known as the Al₃(Sc, Zr) precipitate. This phase is coherent with the Al matrix [31] and is prevalent at the fine grain zone’s grain boundaries, where it serves to pin the boundaries, restrict dislocation movement, and also acts as an effective nucleation site for grain refinement [32]. The solidification process at the molten pool boundary leads to the consumption of the Sc element, thereby eliminating potential nucleation points within the molten pool. Additionally, the nth layer’s molten pool is subject to the influence of neighboring molten pools (as shown in Figure 7, where there is a preference for columnar crystal orientation) and the subsequent n + 1th layer, which collectively restricts solidification within the molten pool. With temperature gradients diminishing from a high to a lower range, solidification within the molten pool initiates at its boundary. The depletion of Sc at the boundary, coupled with a substantial temperature gradient, results in solidification occurring counter to the direction of heat transfer, thus giving rise to columnar crystals. This explains the bimodal microstructure characteristic of the LPBF-processed Al-Mg-Sc-Zr alloy. As shown in Figure 10, the dispersion of nano-sized Al₃(Sc, Zr) precipitates in the alloy’s fine-grained region enhances its strength. The Hall–Petch strengthening model [33] dictates that the strength increment from grain boundaries is related to the inverse square root of the grain size, formulated as σGB ∝ d^(−1/2). Based on reports in the existing body of literature [34], the alloy’s strengthening mechanisms are considered to involve dispersion strengthening and the refinement of grain size.

3.2. Mechanical Properties

Tensile testing was performed on specimens AB in their as-built state and AA in their as-aged state, under both ambient and elevated temperature conditions. The resulting stress–strain curves are displayed in Figure 11, with ‘H’ indicating the specimens oriented perpendicular to the build direction and ‘Z’ indicating those oriented parallel. The mechanical properties, including yield strength (YS), ultimate tensile strength (UTS), and elongation, are summarized in

Table 3. Al-Mg-Sc-Zr tensile test data.

	YS (MPa)	UTS (MPa)	ELO (%)	E (GPa)
AB-Z-RT	321	404	21.3	69.9
AB-H-RT	352	412	23.6	72.2
AA-Z-RT	527	551	11.2	70.7
AA-H-RT	535	560	11.1	65.5
AA-H-100 °C	417	435	25.8	63.6
AA-H-150 °C	214	269	20.1	62.1
AA-H-200 °C	77	102	47.9	38.5
AA-H-250 °C	52	54	72.2	36.2

. The Al-Mg-Sc-Zr alloy fabricated using the LPBF technique demonstrates a higher tensile strength compared to the commercial aluminum alloy sheets (5B70), with a measured value of 421 MPa [35]. This enhanced mechanical performance can be attributed to the crack-free nature and the pronounced grain refinement observed in the microstructure of the LPBF-processed material. In Table 3 and Figure 11, we observe that the elongation of AA-H at 150 °C is lower than that at 100 °C. This may be due to the fact that at 150 °C, internal microvoids and cracks in the material are more likely to propagate, leading to a significant decrease in elongation. At 200 °C, despite the higher temperature, the material likely undergoes more plastic deformation, and the propagation mechanism of microvoids and cracks differs from that at 150 °C.

Fracture usually occurs in the weakest or most stressed part of the material, and the fracture morphology records the relevant information of the fracture process. By analyzing the deformation degree, position, color, precipitates, particle size, and pattern of the fracture, the nature of the fracture can be determined, the cause of the failure can be analyzed, and the mechanism of the fracture can be further studied [36]. In order to better understand the fracture mechanism, the fracture surface of Al-Mg-Sc-Zr samples prepared by LPBF after tensile was observed by SEM. From the macroscopic view, it can be found that the dimples of different sizes can be found in the fracture of the samples, so it can be determined that the tensile fracture of the alloy is a typical ductile fracture. With the increase in tensile temperature, the fracture of the sample is more uneven, which is related to the increase in elongation. As depicted in Figure 12a, the fracture surface morphology of the AA specimen exhibits characteristics akin to columnar grain structures at the crack origin site (a). Such features are exclusively observed on the fracture crack surfaces of specimens following heat treatment. This can be attributed to two primary factors: The material’s bimodal microstructure, characterized by a varied grain size distribution, results in an inhomogeneous stress distribution. The equiaxed grain areas, reinforced by the pinning effect of Al₃Sc particles at the grain boundaries, demonstrate greater strength than the columnar grain regions. Consequently, crack initiation is more prone to occur in the columnar grain areas. The existence of larger grains within the bimodal microstructure contributes to a reduced crack propagation velocity, which in turn, augments the material’s fracture toughness. Figure 12c,d, representing the fracture surfaces at temperatures of 100 degrees and 200 degrees, respectively, corroborate the presence of this phenomenon.

Ductile fractures are generally characterized by three distinct regions: the fiber zone, which is also known as the goose down zone, the radiation zone, and the shear lip zone. These zones are the key components of fracture morphology [37]. As depicted in Figure 13a,c, the fiber zone of the tensile fracture is examined across a range of temperatures, showing an increase in irregularity with rising temperature. The crack propagation within this zone is slower in comparison to other areas. The fracture surface features a jagged pattern punctuated by holes, which at a microscopic level, is composed of many fibrous ‘mini peaks’, each with a 45° inclination to the tensile axis. This inclination indicates that the fiber zone is a result of the progressive enlargement and merging of microcracks during plastic deformation [38]. The dimples on the fracture surface are a consequence of dislocation slip, with the second-phase particles acting as the main impetus. Figure 13b,d show the dimple morphology on the fracture surfaces at 100 °C and 200 °C, respectively. The dense distribution of dimples indicates that the material exhibits high ductility. The dimples at 200 °C are larger than those at 100 °C, suggesting that the specimens underwent greater plastic deformation before fracture. This observation is consistent with the measured elongation results. The size and shape of these dimples are correlated with the material’s properties, including its strength, the dimensions, form, and dispersion of the second-phase particles, as well as the stress conditions, encompassing both the nature and magnitude of the stress applied [39]. Figure 9 displays the secondary phase particles located at the grain boundaries in the equiaxed zone, as detected by EDS. The elemental analysis confirms their enrichment with Scandium (Sc) and Zirconium (Zr). Drawing from Wang’s characterization [40], these particles are identified as the primary Al₃(Sc, Zr) phase, which is directly responsible for the dimple formation observed in the tensile fracture surfaces of Al-Mg-Sc-Zr alloy systems. In alignment with dislocation theory [41], it is around these Al₃(Sc, Zr) particles that dislocation loops are formed, contributing to the material’s ductile fracture behavior. As depicted in Figure 14, in the absence of external forces, dislocation loops are in equilibrium, held in check by the repulsion from Al₃(Sc, Zr) particles and the stress of dislocation stacking. When an external force is sufficiently large, dislocations around the Al₃(Sc, Zr) particles are set into motion once more, moving toward the particles. The formation of microvoids occurs when the elastic strain energy that has built up in front of the dislocation loop reaches a level that can overcome the interfacial bonding forces between the Al₃(Sc, Zr) particles and the aluminum matrix. The formation of microvoids significantly alleviates the repulsive forces on dislocations, leading to a surge of dislocations toward the microvoids under external force, which in turn, drives their enlargement. At this juncture, the dislocation sources that were quiescent behind the dislocation loop become re-energized as the constraint from dislocation accumulation is lifted, initiating new dislocation loops that migrate incessantly toward the microvoids. This results in a swift and unstable enlargement and coalescence of the microvoids. Dislocations, capable of amassing on diverse slip planes, may induce the formation of microvoids through their movement along one or several planes. Concurrently, dislocations on other slip planes are drawn to the microvoids, further promoting their expansion [42].

4. Conclusions

The microstructure, room-temperature and high-temperature tensile properties and fracture morphology of Al-Mg-Sc-Zr samples prepared by laser powder bed fusion were studied. The main results are as follows:

• The microstructure is characterized by neatly aligned melt pools, with the boundaries marked by fine equiaxed grains and the interiors by larger equiaxed grains, creating a distinct bimodal distribution. EBSD analysis indicates an even dislocation distribution across the alloy, devoid of any textural preference.
• Room-temperature tensile tests demonstrated that the heat-treated specimens achieved ultimate tensile strengths up to 560.6 MPa, coupled with an elongation of 11.1%. Tensile strengths for both untreated and treated samples were modestly higher in the H direction than in the Z direction, with the discrepancy being less than 5%. A progressive decrease in strength with increasing test temperature was observed, with values at 100 °C being 435.1 MPa and 25.8%, at 150 °C being 269.4 MPa and 20.1%, at 200 °C being 102.8 MPa and 47.9%, and at 250 °C being 54.0 MPa and 72.2%, respectively. This trend signifies a trade-off between strength and ductility with rising temperature, with the material exhibiting greater plasticity at elevated temperatures.
• Fracture surface analyses from tensile tests across the temperature range disclosed an abundance of dimples, indicative of ductile fracture behavior. The dimples are mainly attributed to the presence of the second-phase Al₃(Sc, Zr) particles. Upon the application of substantial external forces, dislocations surrounding these particles become active again, leading to fracture when the forces surpass the interfacial bond strength between the particles and the aluminum matrix.