# A high-strength TiB2 -modified Al-Si-Mg-Zr alloy fabricated by laser powder-bed fusion

Yaoxiang Geng, Keying Lv, Chunfeng Zai, Zhijie Zhang, Anil Kunwar

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Yaoxiang Geng, Keying Lv, Chunfeng Zai, Zhijie Zhang, Anil Kunwar. A high-strength TiB2 - modified Al-Si-Mg-Zr alloy fabricated by laser powder-bed fusion. Acta Metallurgica Sinica, 2025, 10.1007/s40195-025-01825-1. hal-04993333

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Submitted on 17 Mar 2025

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**Abstract:** To increase the strength of the laser powder-bed fusion (LPBF) Al-Si-based aluminum alloy, TiB₂ ceramic particles were selected to be mixed with high-Mg content Al-Si-Mg-Zr powder, and then a novel TiB₂/Al-Si-Mg-Zr composite was fabricated using LPBF. The results indicated that a dense sample with a maximum relative density of 99.85% could be obtained by adjusting the LPBF process parameters. Incorporating TiB₂ nanoparticles enhanced the powder's laser absorption rate, thereby raising the alloy's intrinsic heat treatment temperature and consequently facilitating the precipitation of Si and β" nanoparticles in the α-Al cells. Moreover, the rapid cooling process during LPBF resulted in numerous alloying elements with low-stacking fault energy dissolving in the α-Al matrix, thus promoting the formation of the 9R phase. After a 48 h direct aging treatment at 150 °C, the strength of the alloy slightly increased due to the increase of nanoprecipitates. Both yield strength and ultimate tensile strength of the LPBF TiB₂/Al-Si-Mg-Zr alloy were significantly higher than that of other LPBF TiB₂-modified aluminum alloys with external addition.

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**Keywords:** Laser powder-bed fusion, aluminum alloy, TiB₂ ceramic particle, 9R phase, mechanical properties

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# 1. Introduction

Laser powder-bed fusion (LPBF) is an advanced additive manufacturing (AM) technology that enables efficient and rapid production of complex and precise parts.

Using LPBF, it is possible to effectively reduce the weight of parts and enhance their performance by employing structure topology optimization and integrating multi-part production [1-6]. Aluminum alloys are extensively utilized in various industries to manufacture key components. However, the current application of LPBF for aluminum alloys primarily focuses on near-eutectic Al-Si, and Al-Si-Mg alloys [7].

Unfortunately, the strength of these alloys is lower than that of traditional high-strength aluminum alloys, thereby imposing restrictions on their application.

Ceramic nanoparticles, such as TiB₂, are commonly incorporated to improve the strength of the LPBF Al-Si-based aluminum alloys. Li et al. [8] investigated the microstructural evolution and mechanical performance of the LPBF in-situ 11.6 wt% TiB₂/AlSi10Mg composite. The results showed that the alloy exhibited fully fine-equiaxed grains and contained well-dispersed nano-TiB₂ particles within the grains and rod-like nano-Si precipitates within the cells. Both nano-TiB₂ particles and nano-Si precipitates showed a highly coherent interface with the α-Al matrix. The alloy exhibited remarkable mechanical properties, including a yield strength (YS) of about 400 MPa, an ultimate tensile strength (UTS) of approximately 530 MPa, and an elongation of approaching 15.5%. Feng et al. [9] fabricated an in-suit 6.5 wt.% TiB₂/AlSi10Mg composite through LPBF. The alloy exhibited a finely equiaxed microstructure without noticeable anisotropy, as well as achieved an exceptional combination of high tensile strength and elongation with a UTS of 536.9 MPa, a YS of 332.3, and an elongation of 16.5%. The excellent mechanical properties can be
attributed to the systematic optimization of the processing parameters and the strong
work-hardening capability of the alloy. Xi et al. [10] mixed TiB₂ and AlSi10Mg
powders at different concentrations and systematically studied the effects of different
TiB₂ content on the manufacturing quality, microstructure, and mechanical properties of
the LPBF TiB₂/AlSi10Mg alloy. The result indicated that the relative density showed a
visible decrease with increased TiB₂ addition. However, the grains of the alloy became
refined with the increasing TiB₂ content. The UTS, YS, and elongation reached a
maximum value of 444 MPa, 283 MPa, and 4.2%, respectively, at a TiB₂ fraction of 2
wt.%. Both in-situ synthesis and external addition of TiB₂ nanoparticles are effective
methods for improving the mechanical properties of LPBF Al-Si-Mg alloys. However,
it has been noticed that the strength and plasticity of the LPBF TiB₂-modified Al-Si-Mg
alloy prepared through the external addition are significantly lower than those fabricated
by in-situ synthesis, which restricts the industrial application of these alloys.

Enhancing the mechanical properties of the matrix material is an effective
approach to improving the mechanical properties of Al-based composite materials [11,
12]. Recently, Geng et al. [13, 14] discovered that increasing the Mg content in LPBF
Al-Si-Mg alloys can improve their mechanical properties effectively. When the content
of Mg was increased to1.5 wt.%, the YS, UTS, and elongation of the new LPBF high-
Mg content Al-Si-Mg alloy reached 295 ± 6 MPa, 476 ± 3 MPa, and 8.4% ± 0.6%,
respectively. The strength of the alloy was much higher than that of traditional LPBF
Al-Si-Mg alloy [1]. Moreover, it has been demonstrated that the addition of Zr alloying
elements to this alloy could further enhance its strength and plasticity, with YS of 343 ±
3 MPa, UTS of 485 ± 4 MPa, and elongation of 10.2% ± 0.2% [15]. In this letter, 2 wt.
% TiB₂ particles were mixed with high-Mg content Al-Si-Mg-Zr powder, and then a
novel TiB₂/Al-Si-Mg-Zr composite was fabricated using LPBF. The densification and the influence of aging treatments on the microstructure and mechanical properties of the alloy were systematically studied.

## 2. Experimental

### 2.1. Powder material

The Al-Si-Mg-Zr alloy powder with sizes ranging from 15 µm to 53 µm was prepared using a vacuum induction N₂ gas atomization process. The chemical composition of the powder was tested by an inductively coupled plasma atomic scanning spectrometer (ICP, PE OPTIMA 7300DV). The composition of the powder is listed in Table 1. The raw materials were made of high-purity aluminum ingot (99.99 wt %), pure magnesium (99.9 wt%), pure silicon (99.99 wt%), and Al-5Zr master alloy (99.5 wt%). The average size of the powder was measured to be 36.9 µm. The TiB₂/Al-Si-Mg-Zr composite powder with 2 wt% TiB₂ and 98 wt.% Al-Si-Mg-Zr was prepared using mechanical mixing at 200 rpm for 12 h. The size of TiB₂ particles ranged from 50 nm to 500 nm. The scanning electron microscopy (SEM) image of the TiB₂/Al-Si-Mg-Zr mixed powder and its corresponding energy dispersion spectrometer (EDS) mapping are presented in Fig. 1. The composite powder displayed a nearly spherical shape. The TiB₂ nanoparticles were mainly attached to the surface of Al-Si-Mg-Zr powder, with a portion of them exhibiting agglomeration. This is a typical phenomenon when mixing the nanoparticles with other powders [16].
<figure><img src="image_2.png" /><figcaption>Fig. 1. (a) SEM surface morphology and (b-f) corresponding EDS mapping of the TiB₂/Al-Si-Mg-Zr composite powder.</figcaption></figure>

<table><caption>Table 1 Chemical composition of the Al-Si-Mg-Zr powder (wt.%).</caption><thead><tr><th rowspan="2">Powder</th><th colspan="5">Chemical composition</th></tr><tr><th>Si</th><th>Mg</th><th>Zr</th><th>Fe</th><th>Al</th></tr></thead><tbody><tr><td>Al-Si-Mg-Zr</td><td>7.43</td><td>1.57</td><td>0.37</td><td>0.08</td><td>Bal.</td></tr></tbody></table>
## 2.2. Specimens preparation

The specimens were produced using an EP M260 machine (E-Plus-3D, China) with dimensions of 55 mm (length) × 15 mm (width) × 15 mm (height). The process parameters and scanning strategy are presented in Table 2 and Fig. 2a, respectively. Fig. 2b shows the morphology of the LPBF specimens fabricated under different process parameters. The upper surface of the specimens was flat, without defects such as deformation and cracking. To regulate the mechanical properties, the as-built specimens were aged at 100 °C–500 °C for 2 h and 150 °C for 2 h–48 h, respectively. After the aging treatment, the specimens were cooled in air.

<table><caption>Table 2 LPBF process parameters of the specimens.</caption><thead><tr><td>Parameters</td><td>Value</td></tr></thead><tbody><tr><td>Laser power</td><td>250 W, 350 W</td></tr><tr><td>Laser beam diameter</td><td>100 µm</td></tr><tr><td>Powder layer thickness</td><td>30 µm</td></tr><tr><td>Hatch spacing</td><td>100 µm</td></tr><tr><td>Laser scanning speed</td><td>800 mm/s, 900 mm/s, 1000 mm/s, 1100 mm/s, 1200 mm/s</td></tr><tr><td>Angle of rotation</td><td>67°</td></tr></tbody></table>

<figure><img src="image_3.png" /><figcaption>Fig. 2. (a) The schematic of the scanning strategy and (b) the morphology of the LPBF TiB$_{2}$/Al–Si–Mg–Zr composite specimens.</figcaption></figure>
## 1.2.3. Density, microstructure, and mechanical tests

The mass density of the specimens was determined via the Archimedes method. A LEXTOLS4000 laser scanning confocal microscopy (LSCM) was applied to measure the roughness ($R_a$) of the top surface of the specimens. Microstructure analysis of the LPBF specimens was performed using optical microscopy (OM, Axio Vert. A1), SEM (SM-6480), and transmission electron microscopy (TEM, JEM-2010F). An EDS was employed to conduct a component analysis of the specimens. Before observing the microstructure using OM and SEM, the samples were subjected to chemical etching with Keller's reagent (1 mL HF + 1.5 mL HCl + 2.5 mL $HNO_3$ + 95 mL distilled water) for approximately 15 s. TEM samples were prepared using twin-jet electropolishing and subsequent ion-beam thinning. Grain size and crystal orientation were analyzed using electron backscatter diffraction (EBSD, OXFORD, Nordlys Nano). A Vickers hardness tester (HRS-150) was utilized to measure the microhardness of the samples. The loading and dwell time were set to 300 g and 15 s, respectively. The tensile properties of the alloy were tested using a UTM5105 universal testing machine at a constant strain rate of 1 mm/min, following the ASTM E8-04 standard. All microstructure characterization and mechanical property tests were conducted on the specimens fabricated with a laser power of 250 W and a laser scanning speed of 1000 mm/s. All the tensile test specimens were perpendicular to the building direction. Three specimens were tested under each condition, and the strength and elongation values were averaged.

## 2.3.3. Results and discussion

### 2.3.1. Processability

The variations in relative density and volumetric energy density (VED) of the TiB$_2$/Al-Si-Mg-Zr samples fabricated at different process parameters are illustrated in
Fig. 3. The samples produced at a laser power of 250 W demonstrated good processability, with relative densities ranging from 99.44% (900 mm/s) to 99.85% (1000 mm/s). However, when the laser power was increased to 350 W and the laser scanning speed was 800 mm/s and 900 mm/s, the relative densities of the samples were less than 99%. As the laser scanning speed exceeded 1000 mm/s, the relative densities of the samples began to rebound and exceeded 99.3% (Fig. 3a). The VED plays an important role in the formation of porosity during the laser melting process [17, 18]. VED can be expressed as $VED = P/(v \cdot h \cdot t)$, where P is the laser power (W), v is the scanning speed (mm/s), h is the hatch distance (mm), and t is the layer thickness (mm). According to the calculation results, when the VED was 129.63 and 145.83 J/mm³, the relative densities of the samples were below 99.0%, as shown in Fig. 3b.

<figure><img src="image_4.png" /><figcaption>Fig. 3. The change of (a) relative density and (b) VED of the TiB₂/Al-Si-Mg-Zr composite fabricated under different process parameters.</figcaption></figure>

To analyze the reasons for the variations in the relative density of the samples, the SEM topography, the $R_a$ of the upper surface, and the OM microstructure of the LPBF TiB₂/Al-Si-Mg-Zr samples were examined, as depicted in Fig. 4. The SEM result indicated that the samples revealed a flat surface, except for splashes with varying sizes
1 that formed as droplets expelled from the melt pool during LPBF (Figs. 4a and 4d) [19].
2 The $R_a$ values of the samples fabricated at different process parameters ranged from
3 17.66 µm to 29.67 µm (Figs. 4b and 4e). The samples fabricated at 350 W exhibited a
4 lower $R_a$ due to the higher energy input [3, 20]. Generally, a rough surface is not
5 beneficial for increasing the processability of the LPBF alloy, as it increases the
6 instability of powder spreading and melt pool during the LPBF process [21]. In the
7 present study, the upper surface of the samples obtained at 250 W was relatively rough,
8 but their relative densities were higher. This indicates that surface roughness is not the
9 primary factor in influencing the relative density of the present LPBF TiB$_2$/Al-Si-Mg-
10 Zr alloy. The OM images showed that only metallurgical pores were observed in the
11 samples, without any keyholes and cracks (Figs. 4c and 4f). At higher laser power (350
12 W) and lower scanning speed (800-900 mm/s), the excessive input of laser energy, i.e.,
13 high VED, resulted in the intense vaporization of volatile elements (especially for Mg)
14 and decomposition of vapor, which led to the formation of numerous large-sized
15 metallurgical pores during the LPBF process and consequently reduced the relative
16 density of the samples (Fig. 3) [22]. In contrast, the LPBF Al-Si-Mg-Zr matrix alloy
17 demonstrated exceptional process stability across various parameters [15]. This can be
18 attributed to the incorporation of TiB$_2$ into the alloy, which enhances the powder's laser
19 absorption rate [8], thus increasing the temperature of the melt pool, and ultimately
20 promoting the formation of larger metallurgical pores during the LPBF process at a
21 higher VED. To obtain a high density in the LPBF TiB$_2$/Al-Si-Mg-Zr alloy, the VED
22 value should not surpass 116.67 J/mm$^3$ (Fig. 3b).
<figure><img src="image_5.png" /><figcaption>Fig. 4. (a, d) SEM and (b, e) LSCM upper surface topography and (c, f) OM images of the TiB₂/Al-Si-Mg-Zr alloy fabricated at various process parameters.</figcaption></figure>

## 3.2 Microstructure

### 3.2.1. As-built alloy

Figs. 5a and 5b show EBSD orientation maps of the LPBF TiB₂/Al-Si-Mg-Zr alloy obtained from the vertical and horizontal planes. On the vertical plane, an equiaxed-columnar bimodal grain structure was revealed with equiaxed grains dominating (Fig. 5a). On the horizontal plane, the alloy showed equiaxed grains (Fig. 5b). During the LPBF process, the presence of TiB₂ nanoparticles and the precipitation of the primary (Al, Si)₃Zr phase served as nucleation sites for α-Al grains due to its low
mismatch degree with the aluminum matrix, resulting in a strong grain refining effect in the alloy [23]. The combined grain refinement effect of multiple types of nanoparticles led to a higher proportion of equiaxed grains in the current alloy to compare the LPBF Al-Si-Mg-Zr matrix alloy [8, 24]. The pole figures inset in Fig. 5a and Fig. 5b indicate that the α-Al grains of the alloy displayed a weak <100> texture, with maximum texture densities of 2.09 and 1.86 at the vertical and horizontal planes, respectively, which was weaker than that of the LPBF Al-Si-Mg-Zr matrix alloy (~3.1) [15]. In the LPBF aluminum alloys, the preferential orientation of the <100> texture of α-Al mainly derived from the columnar grain due to the directional solidification of the material inside the molten pool in the absence of nucleation points during LPBF [25]. In the present LPBF TiB₂/Al-Si-Mg-Zr alloy, the incorporation of TiB₂ nanoparticles effectively reduced the proportion of columnar crystals, thereby reducing the <100> texture of α-Al. The SEM image of the alloy revealed a distinctive cellular substructure, comprising fine and slightly coarser cell regions separated by a heat-affected zone (HAZ) (Fig. 5c), reminiscent of the characteristics observed in other Al-Si-based alloys produced via LPBF [26-28]. Moreover, the alloy showed a uniform distribution of larger particles with several micrometers in size. Further EDS analysis revealed that these particles were the segregation of Ti element, indicating the presence of TiB₂, as shown in Fig. 6. The high-magnification SEM image revealed a higher concentration of nanoparticles dispersion within the cellular α-Al matrix (Fig. 5d), in contrast to the LPBF Al-Si-Mg-Zr matrix alloy [15]. This result indicates that adding TiB₂ promotes the precipitation of the nanoparticles in the α-Al matrix. The research conducted by Cheng et al. [29] evidenced that adding 3.0 wt.% TiB₂ to Al-Si-Mg powder notably reduced the laser reflectivity of the powder to just 40% of that displayed by Al-Si-Mg
1 powder alone. The high laser energy absorption rate of the TiB₂/Al-Si-Mg-Zr powder
2 leads to an obvious rise in the molten pool temperature during the LPBF process. This
3 elevated temperature, conversely, raises the inherent heat treatment temperature of the
4 material that has been deposited, which promotes the precipitation of nanoparticles in
5 the α-Al cells [2].

<figure><img src="image_6.png" /><figcaption>Fig. 5. (a, b) EBSD orientation maps obtained from (a) parallel and (b) perpendicular to the building directions, and (c) low and (d) high magnification (originated from the fine region) SEM images of the LPBF TiB₂/Al-Si-Mg-Zr samples obtained from parallel to the building direction.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>Fig. 6. (a) SEM image and (c-f) corresponding EDS mapping of the LPBF TiB<sub>2</sub>/Al-Si-Mg-Zr alloy, (b) Al, (c) Si, (d) Mg, (e) Zr, and (f) Ti.</figcaption></figure>

Fig. 7 shows the TEM bright field (BF) image and corresponding EDS mapping of the LPBF TiB<sub>2</sub>/Al-Si-Mg-Zr sample. The results indicated that the elements of Si, Mg, Zr, Ti, and Fe mainly accumulated in the cell boundary of the alloy. The Zr element primarily diffused to the TiB<sub>2</sub> particles, forming the (Ti, Zr)B<sub>2</sub> phase [30]. Moreover, the segregation of Si elements and the co-segregation of Si and Zr elements were observed within the cellular α-Al, suggesting the formation of the Si nanoparticles with axiolitic shape and (Al, Si)<sub>3</sub>Zr particles with bulk shape [8, 23]. Further TEM analysis confirmed the presence of the β'' and 9R long period stacked ordered (LPSO) phases within the α-Al cell (Fig. 8). The 9R phase is characterized by nine atomic layers oriented {111}. It exhibits a high density of stacking faults (SFs). SFs are known to enhance alloys' mechanical properties significantly [31]. However, forming the 9R phase in pure aluminum poses challenges due to its high stacking fault energy (SFE) [32]. Previous studies have shown that incorporating alloying elements such as Fe, Ti, Si, Mg, and Zr can effectively reduce the SFE of pure aluminum [33]. However, conventional
processing techniques face challenges in forming the 9R phase in aluminum alloys due to these alloying elements' low equilibrium solid solubility. The LPBF process, characterized by rapid cooling (cooling rate ~ $10^5-10^6$ K/s) [34], allows alloys with supersaturated solid solutions that exceed the solid solution limit of alloying elements. In the present study, the LPBF TiB$_2$/Al-Si-Mg-Zr alloy exhibits a high concentration of Mg, Zr, Si, Ti, and impurity Fe elements dissolving, which plays a crucial role in reducing the SFE to facilitate the formation of the 9R phase [35]. The formation mechanism of the 9R phase in the aluminum alloy has been widely investigated and disclosed in the previous nanometals [36, 37].

<figure><img src="image_8.png" /><figcaption>Fig. 7. (a) TEM-BF image and (b-g) corresponding EDS mapping of the LPBF TiB$_2$/Al-Si-Mg-Zr sample, (b) Si, (c) Mg, (d) Zr, (e) Ti, and (f) Fe.</figcaption></figure>
<figure><img src="image_9.png" /></figure>

Fig. 8. (a, b) TEM-HR images, and (c) inverse FFT obtained from the FFT of the inset in (b). FFT image originated from (a) indicated the presence of β" phase.

Inverse FFT analysis revealed that the 9R structure consists of three Shockley partials. When viewed along the [110] direction, the stacking sequence of the 9R phase is observed to follow the ...ABC/BCA/CAB/...

### 3.2.2. Aged alloy

Fig. 9 presents the SEM images of the LPBF TiB₂/Al-Si-Mg-Zr alloy after aging at various temperatures for 2 h. The microstructure did not show a significant difference from SEM images between the as-built and 150 °C aged alloys (Fig. 5d and Fig. 9a). After 200 °C aging treatment, the number and size of nanoparticles embedded in α-Al cells began to increase (Fig. 9b). When the aging temperature increased to 250 °C, the Si nanoparticles grew further, and the Si-rich cell boundaries began to decompose and became discontinuous (Fig. 9c) [38]. At 400 °C, Si-rich cell boundaries were completely decomposed, and spherical particles with a size of several hundred nanometers were uniformly distributed in the aluminum matrix (Fig. 9e). Further increased the aging temperature to 500 °C, the Si particles grew larger than 1 µm. Those results were similar to those observed in other LPBF Al-Si-based aluminum alloys [38, 39]. The SEM images revealed no obvious change in the microstructure after direct aging treatment at 150 °C for less than 24 h (Figs. 10a-10c). Nevertheless, when the aging time increased
1 to 48 h, the number of nanoparticles embedded in the α-Al matrix slightly increased, as

2 shown in Fig. 10d.

<figure><img src="image_10.png" /><figcaption>Fig. 9. SEM images of the LPBF TiB<sub>2</sub>/Al-Si-Mg-Zr alloy after aging at (a) 150 °C,</figcaption></figure>

(b) 200 °C, (c) 250 °C, (d) 300 °C, (e) 400 °C, and (f) 500 °C for 2 h.
1

<figure><img src="image_11.png" /></figure>

2

3 **Fig. 10.** SEM images of the LPBF TiB<sub>2</sub>/Al-Si-Mg-Zr alloy after aging at 150 °C for

4 (a) 4 h, (b) 12 h, (c) 24 h, and (d) 48 h.

5

6 **3.3. Mechanical property**

7 Fig. 11 depicts the variations in the Vickers hardness of the LPBF TiB<sub>2</sub>/Al-Si-Mg-
8 Zr alloy after direct aging treatment at different conditions. As illustrated in Fig. 11a,
9 with the increase of the aging temperature, the Vickers hardness of the samples initially
10 experienced a slight decline, reducing from 146 ± 9 HV in the as-built state to 137 ± 3
11 HV at 200 °C. When the aging temperature rose to 250 °C, the hardness of the sample
12 rapidly dropped to 102 ± 1 HV and subsequently decreased gradually with the further
13 increase in the aging temperature. When the aging temperature reached 500 °C, the
14 hardness of the sample reduced to 61 ± 2 HV. The reduction in hardness at a
15 temperature below 200 °C is mainly because of the release of internal stress and the

1

2
1 coarsening of nanoparticles in the sample. The rapid decrease in hardness observed as
2 the temperature rises above 200 °C can be attributed to the decomposition of the Si-rich
3 cell boundaries [28]. At 150 °C, as the aging time increases, the hardness of the sample
4 remained relatively stable, fluctuating between 140 HV (48 h) and 149 HV (12 h), as
5 shown in Fig. 11b.

<figure><img src="image_12.png" /><figcaption>Fig. 11. The variations of Vickers hardness of the LPBF TiB<sub>2</sub>/Al-Si-Mg-Zr alloy after aging (a) at different temperatures for 2 h and (b) at 150 °C for different times.</figcaption></figure>

Fig. 12a presents the engineering stress-strain curves of the as-built and aged LPBF TiB<sub>2</sub>/Al-Si-Mg-Zr alloy. For comparison, the stress-strain curves of the as-built and 300 °C-2 h aged LPBF Al-Si-Mg-Zr matrix alloy are also presented in Fig. 12a. The values of the YS, UTS, and elongation obtained from the stress-strain curves are summarized in Table 3. For the LPBF TiB<sub>2</sub>/Al-Si-Mg-Zr alloy, the YS, UTS, and elongation in its as-built state were 408±9 MPa, 514±8 MPa, and 4.7±1.2%, respectively. The YS and UYS were much higher than those of the LPBF matrix alloy (Table 3). At 150 °C, when the aging time was less than 24 h, there were no significant changes in the strength and elongation of the alloys. As the aging time increased to 48 h, the YS of the alloy rose to 419 ± 3 MPa. When the aging temperature exceeded 250 °C,
1 the strength of the alloy decreased rapidly, while the elongation increased effectively.
2 After 300 °C-2 h aging treatment, the YS and UTS decreased to 182±1 and 279 ± 1
3 MPa, respectively, with an elongation of 13.5 ± 4%, which were lower than those of
4 LPBF Al-Si-Mg-Zr alloy. The above results also indicate that the incorporation of
5 highly stable ceramic particles cannot improve the thermal stability of LPBF Al-Si-Mg-
6 based aluminum alloys. Fig. 12b and 12c compare the YS and UTS of the present alloy
7 with other LPBF aluminum alloys modified by the external addition of TiB₂ particles.
8 The YS and UTS of the as-built and 150 °C aging treatment LPBF TiB₂/Al-Si-Mg-Zr
9 alloy were much higher than those of comparable LPBF alloys while maintaining a high
10 elongation [10, 40-44].

<figure><img src="image_13.png" /><figcaption>Fig. 12. (a) Tensile stress-strain curves, and comparison of the (b) YS and (c) UTS of the LPBF TiB₂/Al-Si-Mg-Zr alloy with other previously reported external addition TiB₂-modified aluminum alloys fabricated using LPBF.</figcaption></figure>

<table><caption>Table 3 Mechanical properties of as-built and aging LPBF TiB₂/Al-Si-Mg-Zr and Al-Si-Mg-Zr alloys.</caption><thead><tr><th>Alloy</th><th>Condition</th><th>YS (MPa)</th><th>UTS (MPa)</th><th>Elongation<br>(%)</th></tr></thead><tbody><tr><td rowspan="3">TiB<sub>2</sub>/Al-Si-Mg-Zr</td><td>As-built</td><td>408±9</td><td>514±8</td><td>4.7±1.2</td></tr><tr><td>150 °C-12 h</td><td>406±10</td><td>518±6</td><td>4.8±1.1</td></tr><tr><td>150 °C-24 h</td><td>403±11</td><td>520±12</td><td>4.7±1.1</td></tr></tbody></table>
<table><tbody><tr><td></td><td>150 °C-48 h</td><td>419±3</td><td>517±10</td><td>4.4±1.2</td></tr><tr><td></td><td>250 °C-2 h</td><td>221±10</td><td>345±7</td><td>8.3±0.7</td></tr><tr><td></td><td>300 °C-2 h</td><td>182±1</td><td>279±1</td><td>13.5±4</td></tr><tr><td>Al-Si-Mg-Zr</td><td>As-built</td><td>343±3</td><td>485±4</td><td>10.2±0.2</td></tr><tr><td></td><td>300 °C-2 h</td><td>198±1</td><td>305±6</td><td>28.3±3.3</td></tr></tbody></table>

1

2 Fig. 13 shows the SEM fracture morphology images of the as-built and aged LPBF

3 TiB₂/Al-Si-Mg-Zr alloy. In the as-built sample, the coexistence of tearing ridges and

4 small dimples suggests a mixed ductile and brittle failure (Fig. 13a) [45]. Moreover, the

5 part of the fracture surface exhibited flat, columnar grain-like features (as depicted in

6 Fig. a1), similar to those found in the LPBF Al-Mg-Sc-Zr alloy [46]. This observation

7 indicates that the columnar grains within the molten pool are prone to intergranular

8 fracturing. For the 150 °C-24 h aged sample, there was no obvious difference when

9 compared to the as-built sample (Fig. 13b). When the sample was aged at 300 °C for 2

10 h, the equiaxed dimples were observed across the whole fracture surface, as shown in

11 Fig. 13c, indicating micro-void coalescence and deformation by slip. This fracture

12 feature indicates a typical ductile failure [45].

13

14

15

16

17

18

19

20

1

2
<figure><img src="image_14.png" /><figcaption>2 Fig. 13. SEM low- and high-magnification fracture morphology images of the LPBF</figcaption></figure>

<caption>3 TiB₂/Al-Si-Mg-Zr composite at different heat treatment conditions, (a) as-built, (b)</caption>

<caption>4 150 °C-24 h, and (c) 300 °C-2 h.</caption>

5 The high strength of the as-built alloy can be primarily attributed to the
precipitation of the high number of nanoparticles within the cellular α-Al induced by
high inherent heat treatment temperature during the LPBF process (Fig. 5d) and the
existence of TiB₂ nanoparticles. Extensive research has shown that Si, TiB₂, β", and
(Al,Si)₃Zr, nanoprecipitates are coherent or semi-coherent with the α-Al matrix in the
LPBF Al–Si–Mg-based aluminum alloys [8, 13, 47]. These nanoparticles effectively
impede dislocation movement by exploiting the differences in bulk modulus
differentials between the precipitates and the surrounding Al matrix, leading to a notable
increase in alloy strength [48]. Besides, the formation of the 9R phase in the current
alloy has the potential to distort the crystal lattice at the interfaces and prevent the
emission of dislocations. Moreover, molecular dynamics (MD) simulations have shown
that the interfacial transition boundaries (ITBs) between aluminum and the 9R phase act
as strong barriers to the movement of dislocations, allowing for a stepwise movement of
dislocations along the ITBs [49, 50]. These effects can effectively enhance the strength
and plasticity of the alloy. Recently, Li et al. [47, 49] observed the 9R phase in the
LPBF Al–Mg–Zr and Al–Mg–Si–Sc–Zr alloys. The mechanical property tests
conducted in these alloys have shown that the 9R phase can effectively enhance the
strength and plasticity of these alloys. The 9R phase would be the key factor in
achieving high strength and maintaining high elongation in the present LPBFTiB₂/Al–
Si–Mg–Zr alloy. After undergoing a long-term aging treatment at 150 °C, the YS of the
alloy is slightly increased due to the increase in the number of nanoparticles in the α-Al
matrix (Fig. 10d). Nevertheless, since the majority of the nanoparticles have already
precipitated during the LPBF process, the strength increment of the aging alloy is
limited (Table 1), which is in contrast to a significant increase in strength observed in
the LPBF Al–Si–Mg alloy [13, 38]. When the aging temperature increased to 250 °C
and 300 °C, the strength of the alloy was significantly reduced. The possible reasons can be ascribed to the following aspects. Firstly, the broken or disappearance of Si-rich cell boundaries results in the elimination of the contribution of boundary strengthening. Kim et al. [51] reported the stress reached in the cell boundaries is more than twice that in the α-Al phase during loading, which imparts high mechanical strength of the alloy. However, the load-bearing effect of the Si particles that were decomposed from Si-rich cell boundaries after high-temperature aging is inferior to that of the continuous Si-rich cell boundaries reinforcement [52]. Secondly, the desolvation of Si and Mg in the α-Al matrix decreases the solid solution strengthening. Thirdly, the incoherency and coarsening of Si and β" nanoparticles reduce the precipitation strengthening of the alloy.

## 4. Conclusion

In summary, a novel external addition TiB₂-modified Al-Si-Mg-Zr composite was fabricated using LPBF. The influences of process parameters on the processability and the effects of aging treatment on the microstructure and mechanical properties of the LPBF alloy were systematically studied. The primary conclusions can be summarized as follows.

(1) The incorporation of TiB₂ nanoparticles increased the laser absorption rate of the powder, which promoted the formation of numerous large-sized metallurgical pores at high VED during the LPBF process. The alloy fabricated at a low laser power of 250 W demonstrated good processability with a maximum relative density of 99.85%.

(2) The high laser absorption rate of the powder led to an increased intrinsic temperature of the samples and thus promoted the precipitation of various nano-strengthening phases within the cellular α-Al. Especially, the 9R-LPSO phase was observed in a bulk Al-Si-Mg-based aluminum alloy, which might play a vital role in
1 achieving a harmonious balance between the strength and plasticity of aluminum alloys.

2 (3) After direct aging treatment at 150 °C for 48 h, the YS of the alloy slightly

3 increased to 419±3 MPa due to the increase in the number of nanoparticles in the α-Al

4 matrix. However, when the aging temperature exceeded 250 °C, the strength of the alloy

5 decreased rapidly because of the decomposition of the Si-rich cell boundary and the

6 weakening of solid solution and precipitation strengthening.

7 (4) Both YS and UTS of the as-built and 150 °C aging treatment TiB₂/Al-Si-Mg-

8 Zr alloy were much higher than those of the LPBF external addition TiB₂-modified

9 aluminum alloys while maintaining a high elongation, indicating broad application

10 prospects in the industry.

11 Acknowledgments

12 This work was supported by the National Natural Science Foundation of China

13 (Nos. 51801079 and 52001140) and the National Science Centre, Poland (Narodowe

14 Centrum Nauki) (No. UMO-2021/42/E/ST5/00339).

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