Review on laser directed energy deposited aluminum alloys

Lightweight aluminum (Al) alloys have been widely used in frontier fields like aerospace and automotive industries, which attracts great interest in additive manufacturing (AM) to process high-value Al parts. As a mainstream AM technique, laser-directed energy deposition (LDED) shows good scalability to meet the requirements for large-format component manufacturing and repair. However, LDED Al alloys are highly challenging due to their inherent poor printability (e.g. low laser absorption, high oxidation sensitivity and cracking tendency). To further promote the development of LDED high-performance Al alloys, this review offers a deep understanding of the challenges and strategies to improve printability in LDED Al alloys. The porosity, cracking, distortion, inclusions, element evaporation and resultant inferior mechanical properties (worse than laser powder bed fusion) are the key challenges in LDED Al alloys. Processing parameter optimizations, in-situ alloy design, reinforcing particle addition and field assistance are the efficient approaches to improving the printability and performance of LDED Al alloys. The underlying correlations between processes, alloy innovation, characteristic microstructures, and achievable performances in LDED Al alloys are discussed. The benchmark mechanical properties and primary strengthening mechanism of LDED Al alloys are summarized. This review aims to provide a critical and in-depth evaluation of current progress in LDED Al alloys. Future opportunities and perspectives in LDED high-performance Al alloys are also outlined.

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Introduction
With the rapid progress of modern manufacturing industry, metal additive manufacturing (AM) technologies are booming in high-end manufacturing fields such as electric vehicles, aerospace and biomedicine [1][2][3][4].The global metal AM market is expected to reach about 8 billion USD by 2025, with a compound annual growth rate of 34.9% [5].This fantastic development trend owes to four core advantages of metallic AM over traditional manufacturing: (i) high freedom in structural complexity; (ii) high utilization rate of feedstocks; (iii) achieving rapid prototyping; (iv) capability to integrate an assembly part into a single part.The aerospace industry is one of the leading drivers of AM [6][7][8][9].With the demand to be lightweight, the proportion of Al alloy applied in aerospace structural parts and automotive fields is increasing continuously due to its excellent strength-ductility synergy, wear resistance, and thermodynamic stability [10][11][12].In recent years, many high-strength Al alloys dedicated to AM have been developed and put into practical application [13][14][15].For instance, Scalmalloy ® , developed by Airbus, has been employed to fabricate a bionic partition in the A320 aircraft cabin (figure 1(a)) [16].Similarly, the high-strength 7A77 Al alloy developed by Houston Research Laboratory has been used in an aerial topological gyroid heat exchanger [17].
The European space telecommunications company Thales Alenia Space, in collaboration with 3D printing services company Poly-Shape, has printed a large AlSi7Mg antenna support for satellites with a mass of only 1.13 kg [18].As shown in figure 1(b) [16], the AM-ed topologically optimized AlSi10 Mg antenna bracket for Sentinel satellites was fabricated by laser powder bed fusion (LPBF), whose weight was only half that of the conventionally fabricated counterparts.Figure 1(c) [19] shows an LPBF-ed Al6061-RAM2 heat sink.
Mercedes has successfully printed high-strength Al thermostat covers for old trucks with high-density comparable to traditional die-cast Al parts (figure 1(d)) [19].Based on these successful application examples, research on the integrated control of the 'Material-Microstructure-Performance' of AM Al alloy has become a hotspot in recent years to meet the stringent high-performance and high-reliability requirements in engineering applications.Laser AM (LAM) and wire arc AM (WAAM) are commonly used advanced metal AM technologies.Compared with WAAM, LAM has some general advantages over LDED in the following aspects: (i) the fine laser beam can process a tiny feature to achieve high precision and high-resolution manufacturing; (ii) the high energy density of the laser beam allows the metal material to melt and solidify rapidly, thus achieving grain refinement and a higher material strength; (iii) less material waste, due to less materials removal during machining; (iv) more flexibility in manufacturing complex structures and geometries [20][21][22][23].Laser-directed energy deposition (LDED) and LPBF are the two typical branches of LAM, using different manners of material delivery [19].LPBF technology uses the laser to selectively melt the powder bed to form a thin layer, repeatedly, in a layer-by-layer manner to form three-dimensional parts, as shown in figure 2(a) [24].In contrast, the powder feeds to the melt pool directly in LDED (figure 2(b)) [25].As shown in figure 2(c), compared with LPBF, LDED possesses unique advantages in the following aspects: (i) a magnitude higher deposition rate [26,27]; (ii) large-formation component manufacturing; (iii) local repair of high-value parts [28]; (iv) high flexibility in preparing gradient multi-materials [29]; and (v) higher flexibility when integrating with other technique (such as machining and rolling) [30].
Driven by these advantages, LDED Al alloys have important research value and development potential, which enable component consolidation, increasing device integrity and reducing assembly difficulty and production costs effectively.The research on the formability, microstructure evolution and mechanical properties of LDED Al alloy has attracted much attention in recent years and is expected to advance the aerospace and automobile manufacturing field.
Nevertheless, the poor printability of Al is related to the inherent characteristics of Al alloy and the LAM process (e.g.high thermal conductivity and high laser reflectivity, strong oxidizing sensitivity, high hydrophilicity, poor powder fluidity, high thermal expansivity and high cooling rate [31][32][33]).The near-eutectic Al-Si alloy is relatively mature and reliable for LAM processing due to its outstanding printability [34][35][36].However, high-strength Al alloys such as Al-Cu, Al-Mg and Al-Zn alloys still face significant challenges in LAM, owing to high oxidation sensitivity, porosities, and cracking [20,[37][38][39][40].So far, the trade-off between printability and the mechanical properties of LPBF Al alloy has been successfully broken through by some strategies such as the addition of secondary particles (e.g.TiC [41], TiH 2 [42], SiC p [24], TiB 2 [43]) or alloying elements (Er [44], Y [45], Ta [46], Sc [47], Zr [48], etc).However, there is a lack of reports about LDED Al alloys with high strength due to the significant technological differences between LDED and LPBF in terms of laser energy density, deposition environment, thermal histories, etc.
In recent years, there have been several reviews on the AM of Al alloys, but most of them focused on the LPBF Al alloys [49][50][51] while lacking an elaborated review on LDED Al alloys.The findings in LPBF Al alloys (e.g.materials composition, laser processing parameters, microstructure evolutions, and achievable mechanical properties) cannot transfer to LDED directly.Hence, there is a need for a focused review of LDED Al alloys to understand the challenges, progress and future trends in LDED Al alloys.This work could provide frontier and inspiring information for AM Al alloys to further extend their industrial applications.

High thermal conductivity and laser reflectivity.
As shown in figure 3(a) [58], when the laser irradiates the melt pool, the laser energy is partially absorbed by the material.
Therefore, the laser absorption rate and thermal conductivity of the material itself directly affect the energy efficiency and material deposition rate.Notably, the surface reflectivity of Al can be over 90%, and the thermal conductivity of solid Al is nearly 240 W mK −1 , which is five times that of steel (45 W mK −1 ) [51].As a result, the LDED process for Al alloys usually requires higher laser power, which inevitably leads to the vaporization of alloying elements with low boiling points, such as Mg and Zn.Additionally, the high thermal conductivity of Al alloys associated with the high cooling rate of the LDED process will lead to a high cracking tendency in the material.

High oxidation sensitivity and hydrophilicity.
Generally, Al powders have a dense alumina layer on the surface, which may further reduce their laser absorption rate, and the reduction degree depends on the thickness of the alumina layer [40,59].The presence of an oxide layer may also influence the heat capture of inflight powders and melt pools.The dense alumina layer can impair the bonding between new depositing powder with the solidified material, significantly reducing the build efficiency and density.Powder particles that are not well bonded will bounce off and fall over the solidified melt pool, seriously affecting the process efficiency and causing workpiece pollution.However, Ghasemi et al [60] found that the oxide levels in LPBF processed pure Al, AlSi12, and AlSi10Mg are all reduced compared with the raw powders, suggesting a higher oxygen release than oxygen involvement from oxidation.However, this has not been substantiated in LDED Al alloys.The LPBF process is under inert gas protection with low oxygen, while the LDED process is in the ambient environment with shielding argon gas only.
In addition to oxygen, the solubility of hydrogen in molten Al reaches 0.7 ml•100 g −1 .When the melt pool temperature drops to the liquidus temperature, the hydrogen solubility drops to 0.04 ml•100 g −1 , and most of the hydrogen will form hydrogen bubbles, which will evolve into hydrogen pores with the solidification of the melt pool.Therefore, the LDED Al alloys favor a high-purity protective atmosphere (e.g.Ar) or a high vacuum environment during deposition.

Poor powder and melt fluidity.
The poor powder fluidity may cause the aggregation and inconsistent feeding rate of powders, making powders easily deviate from the instructions of the computer, due to mechanical scraper and Van der Waals forces and static electricity between powders during the deposition process, affecting the design accuracy [49,61].Powder size and morphology have a great effect on fluidity.The larger and more spheroidal powders have better fluidity than irregular or angular smaller ones [62].Regarding melt fluidity, cast Al-Si alloys with eutectic systems are generally better than wrought Al alloys [61].The melt fluidity is positively correlated with the energy input, and good melt fluidity may promote the filling of shrinkage pores during solidification and reduce defects [63].In general, adding Si contributes to improving melt fluidity, thus suppressing hot cracking [64].In addition, Al alloy powder has a high humidity requirement.If the powder is damp, sending the powder out of the powder hopper is difficult, leading to nozzle blockage [59].

High thermal expansivity and large solidification range.
The high thermal expansivity and large solidification range of Al alloy could induce a large amount of residual stress at a high cooling rate and tend to cause solidification cracks due to solidification shrinkage [65,66].Overall, Al-Si alloys' solidification ranges and thermal expansivity are relatively low, as shown in figures 3(b)-(d) [34,[67][68][69][70][71].Therefore, most of the research on LDED Al alloys is around Al-Si alloys, especially the traditionally cast AlSi10Mg and AlSi7Mg alloys, whose printability and castability are excellent.However, the room-temperature strength of as-deposited and heat-treated Al-Si alloys is not high enough to meet industrial requirements, and their poor deformability also limits diversified applications [54,67,68,72].Wrought Al alloys, such as Al-Cu, Al-Mg and Al-Zn alloys, possess a wider solidification temperature range, higher crack sensitivity, high probability of alloying elements evaporating, and inferior printability, but their excellent mechanical properties offer them great research value.

Technical challenges
LPBF and LDED are the mainstream techniques for LAM Al alloys.However, due to the difference in deposition environments, thermal histories, cooling rates and solidification modes, LDED Al alloys encounter more challenges in achieving fine grain, uniform element distributions and high mechanical properties than LPBF.

Solidification behavior.
The solidification behavior of materials after laser processing is illustrated in figure 4(a).The relationship between solidification rate (R) and laser scan speed (v) is [73]: where dT dt denotes the cooling rate, and G and α represent the thermal gradient and the angle between the laser scan direction and the solidification direction, respectively.Generally, the v value of LPBF (typical 500-2 000 mm s −1 ) is much higher than that in LDED (typical 10-40 mm s −1 ).At the same time, the G value in LDED is much lower than LPBF [74].Hence, the cooling rate in LPBF is much higher (1-3 orders of magnitude) than LDED (table 1).The relationship between the dendrite arm spacing (λ) of the solidified microstructures and the cooling rate is as follows [75]: Therefore, the dendrite spacing is negatively correlated with the cooling rate.Thus, the microstructures in LPBF are much finer than those of LDED, as the low magnification stereomicroscope photos and the electron backscattered diffraction (EBSD) inverse pole figure (IPF) maps shown in figures 4(b)-(e) and table 1 [31,47,76].The coarser microstructures lead to the poorer contribution of grain boundary  strengthening in the components fabricated by LDED compared with those fabricated by LPBF.Moreover, Gong et al [76] reported that low-angle grain boundaries (LAGBs, misorientation is less than 15 • ) occupy a larger proportion in LDED samples.In contrast, high-angle grain boundaries (HAGBs) dominate LPBF samples (table 1).

Deposition environment.
Compared with LPBF, LDED usually utilizes higher laser power and larger laser beam size to achieve higher build efficiency.However, the ensuing disadvantage is a greater melting loss of alloying elements and more defects, which impairs the mechanical properties [6].In addition, LPBF is generally carried out in an inert enclosed environment, which can avoid high-temperature oxidation in the AM process to a large extent [79,80].Besides, the laser power in LDED is usually higher than in LPBF; thus, the alloying elements with low boiling points are more likely to evaporate and form porosities, as compared in figures 4(f) and (g) [78].Due to these differences between LPBF and LDED, the mechanical properties of LDED-processed Al alloys are generally inferior to those produced by LPBF when both achieved relative densities higher than 99%, as shown in figure 4(h) [78].

Powder feeding.
Low-density Al powders are more vulnerable to dust splash caused by laser bombardment during LAM.Unlike LPBF, which lays powder in advance, LDED's powder feeding is dynamic, which makes it easier to cause defects in components and dust in the cavity, affecting the printing environment [81].In addition, due to the poor fluidity and high hydrophilicity of Al powders, it is difficult to achieve continuous and stable powder feeding in the LDED process [59,63], which could increase the defects and induce poor surface quality of the component.

Challenges in process optimization
The LDED of Al alloys is a non-equilibrium solidification process with high cooling rates and complex repeated thermal cycles.These LDED process characteristics and the Al alloys' inherent disadvantages make Al alloys prone to metallurgical defects in LDED, including porosities, cracks, inclusions, evaporation of alloying elements, residual stresses and distortions.This section will review the formation mechanism of metallurgical defects in LDED Al alloys.
Generally, high laser reflective Al powders require high laser energy input to achieve a complete melting of materials and well-built components.The keyhole porosities resulting from the excessive energy input are relatively large due to the vaporization of alloying elements and the entrapment of inert gases.In contrast, the porosities caused by insufficient energy input are usually caused by low energy power, fast laser scanning speed, and the prevented laser energy absorption by the oxidation films on the powders [85].This kind of porosity is generally irregular in shape and at the interlayer or at intermelt pools.The porosities resulting from the reduced hydrogen solubility pertain to metallurgical defects and are usually small.It is generally believed that increasing the scan speed can increase the solidification speed of the melt pool; thus, lowering the chance of forming pores at the solid-liquid interface.In general, bubbles or unfused particles are dragged from the bottom of the melt pool [86] as driven by Marangoni flow due to temperature gradients within the melt pool [87].The greater the temperature gradient is, the stronger the Marangoni convection, making bubbles more likely to be trapped in the melt pool [68].
As for the Al-Zn, Al-Cu, and Al-Mg alloys, alloying elements with low boiling points, such as Zn and Mg, may quickly vaporize under the high energy input during the LDED process [39,57,82].Svetlizky et al [39] have found that Al5083 powders acquired by gas atomization were transformed into Al5754 in composition due to the selective vaporization of elements during the LDED process.The powder mass flow rate and laser scan speed affected the evaporation degree of alloying elements to some extent [88,89].The higher powder mass flow rate increases the flux of powder feed gas, which brings more oxygen to the melt pool from the air.However, the increase of powder mass flow would reduce the temperature of the melt pool, which helped to prevent the vaporization of alloying elements.Singh et al [82] utilized Ni (with low reflectivity and photon absorption) to coat 7050 Al powders for LDED use.The Ni coating effectively reduced the lack of fusion due to increased laser absorption.However, Ni tended to segregate at grain boundaries and form brittle Al 3 Ni [90], which requires subsequent plastic deformation to refine and homogenize the microstructures.
The formation of porosities in Al alloys decreases the strength [91] and fatigue performance [92,93] of the component since porosities can be stress concentrators under loading [94].Therefore, the inherent characteristics of Al alloys, processing technologies and feedstock compositions are the key factors affecting the porosities of LDED Al alloys.Optimizing the processing parameters, improving the feedstock quality, and post-processing are the mainstream means of reducing the porosities of LDED Al alloys, which will be summarized and discussed in the following sections.

2.3.2.
Cracks.The crack is a common defect observed in the LDED Al alloys and is directly related to the physical and metallurgical characteristics of Al alloys [95].High crack susceptibility is generally caused by volume solidification shrinkage and thermal shrinkage during the rapid heating and cooling process in the LDED process [96][97][98][99].There are four kinds of cracking mechanisms in LDED Al alloys, i.e. solidification cracks, liquation cracks, second phase induced cracks, and segregation induced liquation [96,100,101].
First, excessive energy input may generate tensile stresses on the solidified regions due to diverse thermal shrinkage degrees of the solidified and unsolidified regions, leading to solidification cracks.Once the residual liquid that flows into the interdendritic region is insufficient to compensate for the shrinkage discrepancy and thermal strain, the voids between solidifying grains may develop into solidification cracks [100,102], as shown in figures 5(a)-(d), which are common defects in LDED high strength Al alloys.
Second, preferentially solidified layers are reheated during subsequent deposition, partially melting the eutectic phase with a low melting point when the temperature exceeds the eutectic temperature.The molten eutectic phase can form the liquid film along the grain boundary.It is easy for this to evolve into a liquation crack when under residual tensile stress, resulting in the delamination of the melting zone, as shown in figures 5(e)-(f) [40,96,104].
Third, the large-size second phases are another possible inducement of crack initiation in LDED Al alloys.It has been proposed that the second phase with quasi-crystalline structures may be formed due to the rapid cooling of LAM [105].Zhai et al [106] found that although the generation of the quasi-crystalline phase is a good way to strengthen Al alloys, they show brittle and hard features at low temperatures.More second phases with quasi-crystalline structures may result in cracking under hot stress and microscopic stress concentration around sharp corners with a regular angular shape.Moreover, Su et al [96] found that insufficient energy input caused by low laser power may generate microscopic stress around the second phases, resulting in crack initiation; thus, increasing the laser power can effectively inhibit crack initiation, as seen in figure 5(g).
Fourth, the segregation may induce a change of grain boundary energy (γ GB ) and surface energy (γ S ), which also affects the cracking tendency, as illustrated in figure 5(h) [96,103].As for the embrittlement segregation, the stability of the free surface tends to develop toward more stable conditions, thus reducing the γ S value.Meanwhile, the γ GB value is increased.This may increase the supercooling required for the bridging of the adjacent solidification boundaries.When γ GB > γ S , a repulsive type of interaction between the solidifying boundaries is promoted.Thus, the liquid on the interface is beneficial to the separation of grain boundaries.The grain bridging is postponed to the later phases of solidification; therefore, the repulsive contact would increase the susceptibility to hot cracking.As for the strengthening segregation, the γ S increases, and the γ GB may reduce, reducing the supercooling required for the bridging of the adjacent solidification boundaries.When γ GB < γ S , an attractive interaction between the adjacent solidifying boundaries is energetically favored.Thus, grain bridging may occur earlier at lower supercooling, which reduces the hot cracking susceptibility [103].The segregation of different elements has a fundamental effect on grain boundary interaction.For example, segregation of Mg is more likely to lead to grain boundary embrittlement.In contrast, the segregation of Cu may enhance grain boundary cohesion due to the generation of Cu-Al bonds across the grain boundaries [107].
Optimizing processing strategy to reduce the temperature gradient and solidification rate is usually utilized to reduce the cracking of LDED Al alloy, such as preheating the substrate, optimizing energy density, and alternating scan strategies [88,99].Also, enhancing the fluidity of the melt and refining the grain sizes by blending reinforcement additives inside are effective means to restrain crack initiation in LDED Al alloys [108,109].

Residual stresses and distortions.
Residual stress and distortions induced by high thermal gradient and comprehensive thermal history are also commonly observed in LDED Al parts.Mercelis and Kruth [110] have proposed a temperature gradient model, which reveals the mechanisms that form residual stresses during AM, as shown in figure 6(a) [111].In the laser heating process, thermal expansion will be suppressed by surrounding low-temperature parts; thus, the heated part would be subjected to compressive residual stress.After the laser source is displaced, the previously heated zone would have shrunk.However, suppressed by the pregenerated plastic strain during the heating phase, tensile residual stress is formed.The residual stresses are easy to accumulate at the pore boundaries, as shown in figure 6(b) [112].These may cause a series of issues: the distortion of the workpiece, the delamination of the workpiece and the substrate, the premature failure of the workpiece, cracking, and mechanical anisotropy [97,[113][114][115].The main strategies to mitigate residual stress and distortions of LDED Al alloys are as follows: (i) Transform the microstructures from columnar grains to equiaxed grains [33,116].Due to the vertically extending heat flow, which is opposite to the growth direction of grains, the microstructures of LDED Al alloys usually consist of coarse and epitaxial columnar grains [117].As shown in figure 6(c), the thermal gradient needed to form columnar structures is much higher than for equiaxed structures [33], and the columnar structures usually lead to low residual stress resistance and high crack sensitivity.Generally, adding ceramic particles or in-situ forming intermetallics to stimulate heterogeneous nucleation is the typical approach to promote columnar to equiaxed transition (CET) [31,[118][119][120][121].(ii) Utilize preheating treatment or innovative laser scan paths [99,122].The preheating of the substrate can significantly decrease the thermal gradient during deposition, thus reducing the accumulated distortion [123].The substrate preheating can achieve better residual stress reduction.A higher preheating temperature usually corresponds to lower peak residual stress [122].Moreover, different scan paths would also affect substrate distortion.Generally, a spiral scan path may acquire a more uniform temperature field compared with transverse and longitudinal scan paths [122].(iii) Post-heat treatment and in-situ treatment.Hot isostatic, solid solution, and artificial aging treatment are commonly used, which can effectively optimize the microstructures and reduce defect sizes [6,113].(iv) Surface treatment.Laser shock peening (LSP) is a commonly used surface treatment technique for Al alloys.The plasma shock wave generated by the nanosecond laser pulse irradiates the surface of the material to produce a certain degree of local plastic deformation on the surface of the sample, which is an encouraging and potential technique to improve the quality of LDED Al alloys and promote the healing of defects [113,124].

Inclusions.
The inclusions in LDED Al alloys are mainly introduced by oxygen and hydrogen, and the natural aging of Al powder during storage is also a possible trigger.In general, powder aging has no effect on powder size and morphology [125,126].However, oxidation inclusions will form during natural powder aging, affecting the powder's laser absorption efficiency (figure 7(a)), subsequent wettability and shape of the melt pool (figures 7(b) and (c)).The number of inclusions increases with increasing ambient temperature, humidity and aging time [59,127], as shown in figure 7(d).
As the oxygen content in the powder increases, the surface tension of the melt pool decreases, resulting in a significant decrease in the height of the track and an increase in the width [59], as shown in figure 7(e).Based on previous investigations, hydrogen levels in the tens of ppm can lead to the formation of pores [128,129].Hydride inclusion will increase the porosity of the deposited materials by two to three times, which will seriously affect the strength and fatigue life of the deposited components [59].Normally, the pores caused by hydrogen inclusion are mainly located on the periphery of a single track, while large holes are formed due to the agglomeration of small holes (figures 7(d)-(f)) [59].Due to the strong oxygen affinity, the Al powder surface may also possess either an alumina layer or an aluminum hydroxide layer.Note that at 700 • C, the contact angle between solid alumina and molten Al is greater than 90 • [130], and this poor wettability will cause great difficulty for powder to enter the melt pool.In addition, the aluminum hydroxide may break down at high temperatures, creating detrimental water vapor on the surface of the powder [131].Once the escape velocity of the pores cannot keep up with the cooling rate of the melt pool, shell pores will form inside the solidified material, as shown in figure 7(f).
Due to the existence of inclusions, the LDED process of naturally aged Al powders requires a higher energy input.As such, a corresponding parameter optimization should be conducted to address the passive influences of inclusions.

Selective evaporation of alloying elements.
As LDED of Al requires a high energy input, the selective evaporation of alloying elements of the melt pool is a common phenomenon, especially for elements with low boiling points, such as Al-Mg alloys and Al-Zn alloys [37,120].As shown in figure 8(a), Muneoka et al [132] have investigated the six stages of evaporative release: (i) forming plasma, (ii) expansion of the cavitation bubble, (iii) expansion of the double-layer cavitation bubble, (iv) bubble shrinkage, (v) bubble stasis, and (vi) bubble burst.Ibrahimkutty et al [133] characterized the ascending process of the plasma plume  The evolution of the bubble burst process from single laser shots at a fixed delay of 320 µs [133].Reproduced from [133].CC BY 4.0.
inside the cavitation bubble using x-rays.During this process, the evaporation products were released into the air, as shown in figure 8(b).
It has been found that the elements' evaporation degree is closely related to the energy density [134,135].Svetlizky et al [39] found that a high scan speed is beneficial to obtaining low energy density and reducing element loss in LDED Al 5083 alloys, as listed in table 2. Interestingly, the element evaporation also strongly correlates with the powder feed rate (PFR).Following the data listed in tables 2 and 3, the Mg loss is varied on the PFR under the same scan speed and laser energy density.A higher PFR leads to more intensive Mg evaporation, and this is because the higher PFR promotes the reaction of oxyphilic elements (e.g.Mg) with oxygen [136].It was supposed that the gas-dynamic phenomena [137] should also be responsible for the influence of the PFR on the element evaporation.
Nevertheless, increasing the PFR can also cause a reduction in the melt-pool temperature, which in turn weakens the element evaporation.Thus, the definite reasons for the element evaporation in LDED Al alloys deserve further exploration.

Process parameters improvement
The high laser reflectance of Al alloy greatly limits the LDED processing window.It is a tricky task to produce high-quality Al alloy components with LDED technology, as there are a massive number of parameters related to the deposition process, as presented in figure 9 [50].These parameters will affect the thermal histories, solidification behaviours and the subsequent densification.

Laser-related parameters.
The varying laser-related parameters have an immediate impact on the thermal input.As Singh et al found [33], the surface energy density is positively associated with the laser power and is negatively associated with the scan speed and laser spot size, as shown in figures 10(a)-(c).In addition, the laser power affects the momentum of the powders.Da Silva et al [81] quantified AlSi10Mg powder's interaction with different laser powers through a high-speed imaging camera and found that the powder density in the deposition state gets higher with the increase of laser power, as shown in figure 10(d).When the laser power increased from 0 kW to 6 kW, the peak density of small-size, medium-size and large-size powder particles increased by 27.8%, 21.2% and 52.8%, respectively [81].It was indicated that the appropriate increase in laser power could partly reduce the powder trajectory deviation.Moreover, the laser scan speed determines the grain growth rate to some extent, which can refer to the above formula (1).Generally, the G/R and G•R values determine the solidification structure morphology and size, respectively [138].High scan speeds can achieve a high growth rate, promoting microstructure refinement, as shown in figure 10(e) [138].Additionally, the scan speed also significantly influences the morphology and size of intermetallic phases.Ramakrishnan and Dinda [138] found that the size of Al-W intermetallic compounds was refined with the increase in scan speed, which is advantageous for improving mechanical properties due to the suppressed growth from a high cooling rate.
In terms of improving the printability and formability of LDED Al alloy, the adjustment of laser-related parameters will also have a beneficial effect.It has been proved that the density of the deposited part is mainly influenced by the laser energy density [141].Insufficient energy densities prevent laser energy from fully melting the powder, resulting in a lack of fusion or poor interlayer bonding.Conversely, extremely high energy density could cause the turbulent flow in the melt pool to absorb shielding gas, cause element evaporation, and form pores. Keyhole pores are the typical defects caused by excessive energy input.Liu et al indicated that laser power and scan speed are the two major factors that determine the relative density by using the signal-to-noise (S/N) ratio for determining relative density, as shown in figure 10(f) [139].It can be seen from figure 10(g) that the energy density below 100 J•mm −3 corresponds to a lower relative density due to incomplete fusion.When the energy density reaches 125 J mm −3 , the relative density of LDED AlSi10Mg can reach 99% or higher.However, suppose the energy density is further increased, it may cause the molten pool turbulence to absorb too much heat and cause the material to evaporate, thus reducing the relative density.Moreover, the scan speed had a more significant effect on the relative density at lower laser power, while the laser power showed a more substantial impact on the relative density at higher scan speed [139].Regarding laser type, lasers with a wavelength between 980 nm and 1080 nm are commonly used in LDED [142].The absorption rate of Al alloy of the traditional 1064 nm infrared laser is only about 7% [143].Thus, Al powders require a high power density to be molten.Moreover, the liquid Al has much higher absorptivity than the solid Al, especially with keyhole formation.Thus, it requires high laser energy to melt the powders, which easily induces pores during deposition.Furthermore, the repeated reflection of the infrared laser in the pore could cause droplet splash, affecting the LDED process stability.It is worth mentioning that implementing a blue laser with a wavelength of 450 nm is a good way to enhance the laser absorption rate to 14.5% due to the shorter wavelength [144], which is about twice as much as the infrared laser.Compared with the infrared laser, the temperature and ionization of the plasma under the blue laser are lower, which will significantly reduce the laser scattering and the production of droplet splash, as shown in figures 10(h)-(k).

Deposition-related parameters.
Regarding deposition-related parameters, the PFR affects the continuity of the track and the regularity of the track geometry.A proper increase in deposition rate is conducive to forming more even tracks.The PFR also directly determines the layer thickness and indirectly influences the solidification conditions and internal heat cycle in LDED.Higher deposition rates correspond to a higher layer thickness.Li et al [145] found that the formation of equiaxed crystals can be promoted by appropriately reducing the layer thickness in LDED Al-5Si-1Cu-Mg alloy, as shown in figures 11(a) and (b).The flow rate of shielding gas affects the quality of LDED Al alloy components from the following aspects: (1) It may affect the flow rate of Al powders, and the higher flow rate of shielding gas would increase the amount of powder in the melt pool.
(2) It would affect the speed of Al powder.The acceleration of the shielding gas flow rate would make Al powder more likely to rebound [146].(3) It dramatically affects the oxidation degree of Al powder during the deposition process.(4) It affects the porosity of LDED Al alloys.Excessive shielding gas flow rate may lead to melt pool disturbance, which can easily cause holes during the rapid solidification process.This phenomenon is more pronounced in Al alloys due to their higher solidification rate and lower density of melt pool.Therefore, moderate shielding gas flow rate is generally selected to minimize the melt pool disturbance on the premise of ensuring the low oxidation rate of Al powder, so as to reduce the generation of pores.
Additionally, related work has explored many novel scan strategies that are innovative in space and time dimensions in recent years to improve the printability and formability of Al alloys in LDED [147,148].Gu et al [147] investigated the differences in the microstructure of Al 2024 alloy deposited by different scan strategies, i.e. unidirectional deposition and bidirectional deposition.The bidirectional deposition presented a bidirectional cubic texture, while unidirectional deposition exhibited a unidirectional strongly solidified fiber texture, as shown in figures 11(c) and (d), indicating that the bidirectional deposition strategy is more conducive to reducing the residual stress generated during LDED.Li et al [148] investigated interlayer pause strategies to eliminate hot cracks.Increasing the interlayer pause time can reduce the heat accumulation during LDED, effectively refining the columnar grains and promoting the CET, as shown in figures 11(e)-(i).It is advantageous to eliminate hot cracks along the columnar grain boundaries and reduce porosities.However, excessive extension of interlayer retention may result in a lack of interlayer fusion due to too little heat accumulation.
Recently, optimization strategies of preheating or watercooling substrates have been tried to reduce metallurgical defects in LDED Al alloys [149][150][151].Bhagavatam et al [149] found that preheating the substrate to 260 • C reduced the residual stress and eliminated the solidification cracks in LDED 7075 Al alloy.It was found that a preheated substrate reduced the cooling rate of the molten pool, giving enough time for the bubbles produced by the evaporation of Zn and Mg to escape from the molten pool, thus significantly reducing the porosity (figures 11(j) and (k)).Wang et al [150] compared the microstructure configurations of LDED Al-Mg-Sc-Zr components corresponding to air-cooled and water-cooled substrates (figures 11(l) and (m)).The microstructures of the LDED part with water-cooled substrate exhibited duplex grain structures, while the microstructures with air-cooled substrate were equiaxed.This can be ascribed to the increased average cooling rate of the fusion boundary under water cooling conditions, which inhibited the precipitation of primary Al 3 (Sc, Zr).Also, for this reason, more secondary Al 3 (Sc, Zr) precipitated in water-cooled samples during the aging process, which increased the strengthening effect of precipitation strengthening.Thus, the yield strength (YS) of water-cooled samples reached twice that of air-cooled samples.Guo et al [151] found that the microstructure of 5052 Al alloy components prepared by underwater wire-feeding laser deposition was finer than that of in-air wire-feeding laser deposition, due to the faster cooling rate under water.Meanwhile, the loss of Mg was less, and thus the hardness of components prepared under water was higher.

Machine learning assisted process optimization.
Selecting an appropriate processing window is crucial for processing difficult LDED Al alloys.However, the processing window optimization process requires high time costs for Al alloys with poor printability.In recent years, the rapid development of machine learning provided a shortcut for optimizing processing windows in metal AM, which can effectively facilitate faster selection of the most appropriate process parameters [152][153][154].The leading machine learning algorithm is supervised learning, which obtains the optimal model through training from existing data according to the known relationship between the input and output results.The main processes include data acquisition, feature treatment, machine learning, deep learning models, result training and testing, as shown in figure 12. Caiazzo and Caggiano [155] proposed a machine learning approach based on artificial neural networks to accurately estimate the correct laser power, scan speed and powder feeding rate to achieve a specified geometry of LDED Al 2024 alloy, whose mean absolute percentage errors was as low as 2.0%, 5.8%, and 5.5%, respectively.

In-situ alloying
In-situ alloying of elements in Al alloys during LDED can affect the solidification behavior, grain growth, and precipitation [158].For instance, adding Sc and Zr has attracted wide attention due to their excellent effects on changing the solidification range, refining microstructures, and precipitate strengthening [31,150,[159][160][161][162].The maximum solid solubility of Sc and Zr in Al is only 0.35 wt.% and 0.26 wt.%, respectively [163,164], but rapid solidification can increase the amount of Sc and Zr in the Al matrix by several orders of magnitude [164,165].Xiao et al [160] found that in-situ formed primary Al 3 (Sc, Zr) particles can precipitate in the melt pool inner region and melt pool boundary of as-deposited Al-Mn-Mg-Sc-Zr alloy.Wang et al [31] found that the fraction of in-situ formed particles at the fusion boundary was approximately twice that of the interior of the melt pool, and the overall content of Al 3 (Sc, Zr) particles increased with the increase of Sc/Zr content, as shown in figures 13(a)-(d).Furthermore, the critical cooling rate for the nucleation of Al 3 (Sc, Zr) was studied based on an equilibrium phase diagram (figure 13(e)).As shown in figure 13(f), with the increasing of Sc/Zr contents from 0.37 wt.% to 0.51 and 0.71 wt.%, the critical cooling rates for suppressing the nucleation of the Al 3 (Sc, Zr) phase were approximately increased from 1.1 × 10 2 K•s −1 -2.8 × 10 3 K•s −1 and 3 × 10 4 K•s −1 , corresponding to figures 13(a)-(c) respectively.It can be postulated that the increased Sc/Zr contents can promote the nucleation of Al 3 (Sc, Zr) particles.
As shown in figure 13(g), the value of the ratio of thermal gradient and growth rate (G/R) usually determines the grain size, and the lower G/R value facilitates the formation of equiaxed microstructures.LDED Al alloy has a high thermal gradient, usually leading to columnar grain formation.However, under heterogeneous nucleation on primary Al 3 (Sc,Zr) particles, G/R can be expressed as [166]: where a and n are material-dependent constants, φ denotes the volume fraction of equiaxed grains, and N 0 represents the nucleation density.Based on the low lattice misfit between Al 3 (Sc, Zr) particles and α-Al ( [ 112] Al // [ 112] Al3(Sc,Zr) and (110) Al // (110) Al3(Sc,Zr) [31]), the value of N 0 increases, which offers sufficient nucleation sites to realize CET, although the value of G/R is large.It was found that the columnar grains in molten pools gradually undergo a greater degree of CET with the increase of Sc and Zr content, as shown in figure 13(h).In addition, the secondary Al 3 Sc formed during the artificial aging process also has a highly efficient strengthening effect [167].Interestingly, it was found that Zr can significantly retard the coarsening of the precipitated phase and diffuse into Al 3 Sc to form secondary Al 3 (Sc, Zr) with a core-shell structure [168].Li et al [64] also found that Zr, as a surfaceactive element, can effectively reduce the solid-liquid interfacial energy and surface tension of Al alloys, thereby reducing their viscosity and improving the inherent poor fluidity of the Al melt pool.Therefore, Sc and Zr elements are very suitable for improving the printability and performance of LDED Al alloys.In addition, the repeated thermal cycle also promotes in situ aging and coarsening of precipitated phases, thereby affecting the mechanical performances [169,170].
In addition to the Sc and Zr mentioned above, alloying of Si is also usually used to reduce the high hot cracking susceptibility of Al-Zn-Mg-Cu and Al-Cu LDED components, since Si induces the formation of a low melting point Al-Si eutectic phase, the molten Al-Si eutectic phase can backfill cracks during deposition.Li et al [64] found that the addition of 2.9 wt.% Si to Al-5.0Zn-2.0Mg-1.5Cualloy significantly improved its printability in LDED.The cracks were completely eliminated, and the mean grain size was refined from 106.9 µm to 68.6 µm.The highest texture intensity was also decreased from 5.283 to 3.561 mrd.Moreover, Ce has been found to significantly improve the high-temperature stability of LDED Al alloys in recent years.Plotkowski et al [171] found that the coarsening of microstructure was not observed in the heat-affected zone and heat-treated LDED Al-12Ce alloy, which can be ascribed to the excellent coarsening-resistant ability of Al 11 Ce 3 that was formed during the rapid solidification process of LAM [172].In addition, Yang et al [158] investigated the influence of alloying of Ni, Mn, and Mg on the microstructure of Al-Si-Cu alloy.The results showed that adding Mg promoted the transformation of α-Al grains from coarse columnar to fine equiaxed shape and hindered the transformation of δ phase to γ phase (figures 13(i)-(k)).Meanwhile, Mg alloying could also facilitate the formation of the Al 15 Mn 3 Si 2 phase with better mechanical properties.Due to the efficient strengthening effect of Mn-rich phases and Ni-rich phases (figures 13(l)-(n)), LDED Al-18Si-10Cu-10Ni-5 Mg-5Mn alloy exhibited the highest wear resistance at 350 • C.

Reinforcement particle addition
The minor addition of reinforcement particles (mostly ceramics like TiC [53,173], TiB 2 [52,174], TiCN [65], and SiC [175]) into Al alloys during laser-based AM can endow Al alloys with improved laser absorption rate, high specific strength, high hardness, and good thermostability, etc.The mainstream reinforcement particles can be divided into ex-situ ones and in-situ ones.

Ex-situ reinforcement particles.
The preparation methods of ex-situ particle reinforced LDED Al alloy can be mainly divided into two categories: (i) One is to add the reinforced particles into the Al melt during the smelting period, which is conducted to sufficient dispersion.Then, the particle-reinforced Al alloy powder of the desired diameter can be obtained after postprocessing (such as melt filtration and gas atomization [39,176]).The molten salt assisted method can effectively introduce particles into Al melt.However, molten salt contamination will be observed when the volume fraction of reinforced particles exceeds 10 vol% [177].It also directly limits the application of particle-reinforced Al alloy powder prepared by the gas atomization method.Lin et al [176] successfully prepared 35vol%TiC/Al powder by melt filtration.Most of the TiC nanoparticles would gather on the droplet's surface, as shown in figures 14(a)-(c).During the LDED process, the reinforced particles on the surface of Al powder can effectively absorb the laser beam and boost the melting of the Al.Moreover, the nanoparticles in LDED samples were evenly dispersed throughout the matrix without forming clusters, as shown in figure 14(d) [176].This method is more suitable for high reinforcement particle content.(ii) The other method is to mechanically mix Al powders with reinforced particle powders by a powder mixer like a powder blender or ball milling, which is more flexible and widely used.After mixing, the reinforcement particles would adhere to the surface of Al powder, as shown in figures 14(e)-(g).Li et al [34] found that wet ball milling with anhydrous ethanol as the solvent improved the combination of reinforcement particles with Al powders and the homogeneous distribution of particles.Ethanol can increase the wettability between the reinforced particles and the Al powders, to effectively alleviate particle aggregation, and help to achieve a uniform distribution of reinforcement particles in the Al matrix after LDED, as shown in figures 14(h)-(j).
Generally, in particle-reinforced LDED Al alloys, the interfacial bonding and wettability between reinforcement particles and the Al matrix are critical [20].The TiC/Al interface and TiB 2 /Al interface exhibited in figures 15(a) and (b) [176,178] can facilitate the transfer of loads [179].It is worth mentioning that some ex-situ particles play a refining role by forming a transition layer.For example, the Al-Ti layer formed in LDED TiB 2p /Al at the TiB 2 /Al interface [20], indirectly promoting α-Al nucleation, as shown in figures 15(c)-(e).Interestingly, the Al-Ti layers were only observed during LDED but were challenging to form during LPBF.This is because the cooling rate of LDED is 50-100 times lower than LPBF's, which is conducive to the full interaction between TiB 2 and Al melt, thus promoting the formation of an Al-Ti layer with high nucleation promotion efficiency [20].However, detrimental interfacial reactions may also occur after the addition of reinforcement particles.Mi et al [180] found that SiC may decompose into Si atoms and C atoms during deposition when the temperature is higher than the dissolution temperature of SiC.When Si atoms were exposed to liquid Al, the dissolution velocity of the SiC surface was accelerated.Thus, unevenly distributed lamellar Al 4 C 3 was more likely to form on the corresponding location, which was harmful to the tensile strength.
In addition, ex-situ reinforcement particles play an important role in refining grains, which benefits reducing crack sensitivity and improving the formability of Al alloys in the process of LDED.On the one hand, more nucleation sites are provided; on the other hand, reinforcement particles act as growth blockers and are nailed to grain boundaries to impede grain growth [182,183].Chen et al [38] found that 0.  to increase the laser energy absorption.Under heat conduction, the absorbed energy can be transferred to the surrounding powders.Due to the increased laser absorption after adding ex-situ reinforcement particles, the formability of LDED Al alloys is improved prominently.Li et al [181] found that the number of pores significantly decreased with the TiC content increasing from 0.5 wt.% to 2.0 wt.%, as shown in figure 15(i).
Secondary particles can also promote internal convections in the molten pool.Marangoni convection exists in the zone of low surface tension to high surface tension in the melt pool, and its intensity can be expressed as [184]: where M a denotes the intensity of dimensionless Marangoni, L, ∆γ, µ, and v represent the melt pool length, the difference in surface tension, dynamic viscosity, and kinematic viscosity, respectively.Since the reinforcement particles can assist the Al alloy melt pool to absorb more energy, this decreases the dynamic viscosity and increases the intensity of Marangoni, as shown in figure 15(j).The resulting convection inside the melt pool can be enhanced, which is equivalent to stirring, thus improving the homogenization of the LDED microstructures of Al alloys [38].Gonzalez et al [101] found that the successful CET in AA7075 alloys reinforced by TiC remarkably weakened the residual stress inside the LDED sample, as shown in figures 15(k) and (l).Moreover, some studies confirmed that Al composites reinforced by two or more hybrid ceramics seem more effective in improving formability and printability than a single reinforcement [185,186].
The reinforcement effect of the ex-situ particles is generally considered as 'ex-situ' enhancement [187], whose load transfer strengthening effect inside the composites was found diminished due to the poor wettability and differences in the thermophysical properties of the ceramic/metal interface [188].Also, the mechanical mixing method to add ex-situ reinforcement particles was found challenging to realize the uniform dispersion of particles [189], and the ball milling method easily destroys the structural integrity of reinforcement particles [190].In recent years, some studies have verified that in-situ reinforcement particles can also effectively improve the printability of Al alloys in LAM [41,43,191,192].This section will review the effect of in-situ reinforcement particles on the LDED Al alloys, except where the alloying elements form intermetallic compounds or eutectic phases with the Al matrix.
In general, there are two main in-situ synthesizing patterns, i.e. the reinforcement particles are in-situ synthesized during the smelting before LAM [43] or during the LAM process [192].In terms of the former, Sun et al [43] used the in-situ mixed salt method to fabricate as-cast Al-Cu-Mg composite reinforced by TiB 2 nanoparticles, then the vacuum gas atomization method was used to fabricate Al-Cu-Mg powders with in-situ TiB 2 nanoparticles, as shown in figure 16(a).It was found that the powders had a cellular-dendritic microstructure without the internal porosity, and the in-situ TiB 2 particles can distribute homogeneously without obvious agglomeration, as shown in figure 16(b).As for the latter, Yi et al [192] induced a combustion reaction between Ti and B 4 C to generate TiB 2 and TiC aided by the LAM.In this work, the Ti powders and B 4 C powders were ball milled first, and then the collosol was uniformly sprayed onto the powder mixture to make Ti powders coated with B 4 C (B 4 C@Ti).Finally, the B 4 C@Ti powders were tightly packed around the AlSi10Mg powders after ball milling, as shown in figure 16(c).TiB 2 and TiC particles were in-situ generated under high-energy laser induction, as shown in figure 16(d).The analysis showed that the contents of TiB 2 , TiC and residual B 4 C in LAM AlSi10Mg composites were 1.46 wt.%, 0.63 wt.% and 1.91 wt.%, respectively.Moreover, the in-situ TiC and TiB 2 particles were formed in the elemental transition zone surrounding the residual B 4 C particles (figures 16(e)-(h)), where B 4 C particles were dissolved and formed an enrichment zone of B and C elements in the liquid phase of AlSi10Mg.Compared with the ex-situ particles, in-situ ones are not damaged and have higher chemical stability; thus, the laser absorption rate can be improved more significantly.Sun et al [191] found that the laser reflectivity of 2024 Al alloy powders was reduced with the increasing contents of in-situ synthesized TiB 2 nanoparticles, which was conducive to acquiring the combination of near-full densification and high forming quality.However, excessive in-situ particles may lead to excessive energy absorption and overburning of the powder (figures 16(i) and (j)).
Until now, most studies about the in-situ reinforcement particles applied in LAM of Al alloys focused on LPBF, while using in-situ reinforcement particles to LDED in Al alloys has rarely been reported.Based on the conclusion that in-situ particles significantly improve the printability of Al powders, LDED Al alloys containing in-situ particles are very promising and valuable for research.

Field-assisted LDED of Al alloys
Field-assisted LDED technologies have gradually emerged for processing Al alloys, which offer good capabilities in effectively reducing defects and residual stresses, refining microstructures and improving the performance of Al alloys [66,74,113,193,194].
Some studies modified the scan path to improve the melt pool fluidity, reduce element segregation, and homogenize the structure in LDED Al alloy.Cui et al [194] fabricated LDED 2319 Al alloy with low porosity and high mechanical properties by using the circular oscillation mode of the wire-feeding AM equipment with an oscillating laser (figure 17(a)).The oscillation of the laser was equivalent to exerting a stirring effect on the melt pool to keep it in a laminar flow mode, which was conducive to stabilizing the keyhole and reducing the porosity (figures 17(b) and (c)).To further enhance the surface quality and mechanical properties of LDED Al alloys, Zhou et al [113] reported LSP assisted LDED of AlSi10Mg, as shown in figure 17(d), whose maximal pressure of the shock wave can attain the gigapascal level.After in-situ LSP treatment, the trade-off in strength and ductility was overcome by microstructure refinement, dislocation strengthening, and compressive residual stress.Meanwhile, the surface roughnesses of S a and S z were significantly reduced by 24.3% and 21.4%, respectively.As shown in figure 17(e), the pore shape was transformed from spherical to oval after the introduction of the shock wave; both the width and length of the pores were reduced observably.It can be demonstrated that severe plastic deformation can heal the defects to some degree.In recent years, the laser-arc hybrid was found to be an effective method to combine a high-power laser and adaptable arc, which was beneficial to realize a high build rate and low defect rate [195,196].Miao et al [193] and Wu et al [66] reported the Al-Si and Al-Cu alloys fabricated by laser-arc hybrid AM, which combines the advantages of WAAM and LAM, with the ascendancy of low price, high deposition rate and high cooling rate over individual LDED technologies, as shown in figure 17(f), acquiring fine mechanical properties.As shown in figures 17(g)-(i), the combination of laser and arc promoted a strong Marangoni convection inside the keyhole, which was conducive to dendrite fracture [66].The dendrite fragments were found to serve as the nucleation sites and accelerated microstructure refinement.Moreover, LDED assisted by a magnetic field was an effective method to accelerate melt pool convection and improve the material formability [197,198], whose experimental setup can be referred to in figures 17(j) and (k).Schneider et al [199] found that the pore quantity in the LDED AlMg3 alloy was reduced under the influence of a magnetic field (figure 17(l)).This can be attributed to the pore escaping from the melt pool and passing through the oxide layer [199].As a surface enhancement strategy, LDED assisted with deep surface rolling has great tolerance to residual stress relaxation at high temperatures [200,201].As shown in figures 17(m)-(o), the deep surface rolling technology applies high-pressure fluid to make the rolling ball float in a socket and roll freely along the part surface; thus, the near-surface zones undergo plastic deformation.As the neutron diffraction residual stresses analysis results obtained by Zhuang et al [201] in Al 7075 alloy clad by Al-12Si alloy shown in figure 17(p), deep surface rolling can generate large compressive residual stresses that can reach considerable depths below the laser clad surface, which was beneficial to enhance fatigue resistance.Zhang et al [202] investigated the ultrasonic-assisted LDED of the 4047Al alloy (figures 17(q) and (r)).A dense sample with the density of 99.1% was obtained by optimizing the processing parameters, which was almost equivalent to that of a cast sample.In addition, the α-Al grains and Si phases can be markedly refined compared with cast samples, resulting in higher mechanical properties.

Grain morphology and intermetallic distribution in as-built state
The as-built microstructure of LDED Al alloy is affected by a high thermal gradient and repeated thermal cycle, which leads to very different microstructures from conventional cast parts.Pan et al [32] particularly compared the grain size, morphology and texture of AA3104 alloy processed by LDED with that of direct casting.The microstructures of the LDED sample  were columnar with a large grain size of 0.5-1 mm and intense texture, which was different from the equiaxed grains in directly cast samples with a mean size of ∼95 µm (figures 18(a) and (d)).In addition, LAGBs were the majority in the LDED sample, whose contribution to element diffusion, precipitation, and mechanical property strengthening was weak, as shown in figures 18(b) and (e).Most LAGBs originated from forming substructures at high cooling rates during LDED.
As shown in figures 18(c) and (f), intermetallic compound particles precipitated at grain boundaries with lower cooling rates in directly cast AA3104 samples.In comparison, non-equilibrium solidification results in chemical inhomogeneity in LDED samples.High-temperature thermal cycling dissolved Mg 2 Si, increasing the concentration of Si in the matrix and promoting the Al 6 (Fe, Mn) to AlFeMnSi phase transformation.Rapid cooling also provided a higher degree of supersaturation and subcooling for the α phase precipitation, resulting in the area fraction of the α phases in the LDED sample being three times that in the direct casting sample.
In addition, due to the different thermal histories among the top, middle and bottom regions along the build direction, the actual layer thickness and the morphology of intermetallics are also different in other areas of the LDED Al alloys [203].In LDED Al-Cr-Fe alloy, the bottom layer thickness was 210 µm, while the top layer thickness increased to 380 µm due to heat accumulation [203].As the number of layers increased and the temperature of the LDED component increased, the volume of the melt pool would increase, allowing more powders to enter the melt pool.Similarly, the complex thermal history within a single melt pool led to the hierarchical size of intermetallics.In general, the cooling rate of the inner melt pool is much greater than that in the melt pool boundary, and the growth time of the primary intermetallics is limited in the inner melt pool.However, the intermetallics at the edge of the melt pool were heated repeatedly during layerby-layer deposition, which may grow up and coarsen easily, as observed for the Al 13 Fe 4 in LDED-processed Al-Cr-Fe alloy (figures 18(g)-(k)) [203].
In LDED of Al-Cu and Al-Zn alloys, the wide solidification interval generally results in a high proportion of columnar grains [38,88,121].In general, only the bottom and top parts are equiaxed, and the rest of the middle portion of the component shows columnar grains.The region near the substrate will likely form equiaxed grains because of the higher cooling rate.The cooling rate in the middle part is much lower than that at the bottom; thus, the size of columnar grains at the middle and top region is usually larger, as shown in figures 19(a)-(f).The critical thermal gradient condition for complete columnar grain growth can be expressed as [166]: where N 0 , ∆T N , and ∆T C represent the number of heterogeneous nuclei per unit volume, the supercooling for heterogeneous nucleation, and the supercooling of the columnar dendrite tip, respectively.However, the temperature gradient can reach a critical state at the top of the melt pool.Thus, equiaxed grains would form when the maximum grain growth undercooling is greater than the nucleation undercooling.If the depth of the remelting layer exceeds the equiaxed zone at the top of the previous deposition layer, columnar grains could be induced.The microstructure of LDED Al-Si alloy differs from that of other Al alloys due to its low melting point, narrow solidification temperature range and the Al-Si eutectic reaction.As shown in figures 19(g) and (h), the longitudinal section of LDED AlSi10Mg alloy has three typical regions, namely, the fine columnar region with developed secondary dendrite arms, the coarse columnar region and a thin heat-affected zone at the melt pool boundary.The cross-section of columnar grains has symmetrical equiaxed structures and continuous eutectic structures between dendrites (figures 19(i) and (j)).As observed from the Al-Si equilibrium phase diagram in figure 19(k), during the cooling in LDED, Al-Si alloy underwent the first isomorphic reaction to form α-Al and then the Al-Si eutectic phase through the eutectic reaction.The value of the cooling rate (G × R) determines the solidification mode and grain morphology, while the G/R value determines the grain size.Shi et al [204] obtained the values of G • R and G/R at different positions of the melt pool by finite element model using MATLAB with the results shown in figure 19(l).From P1 to P9, the thermal gradient decreased from 4.9 × 10 5 to 3.9 × 10 5 K s −1 , while the solidification rate increased from 4.2 × 10 −3 to 1.6 × 10 −2 m s −1 .Therefore, the value of G • R increased gradually from the bottom to the top of the melt pool, while the value of G/R decreased gradually, as shown in figure 19(m), directly leading to the planar-cellulardendritic transition from the bottom of the melt pool to the top.The similarity of the primary dendrite arm spacing and the perturbation wavelength of the instability of planar growth resulted in continuous epitaxial columnar crystals [205].The difference in secondary dendrite arm spacing between coarse and fine-grain zones can also be ascribed to different cooling rates [206].
Moreover, adding Sc or Zr can generate a layered structure with alternating columnar and equiaxed grains, as shown in figures 19(n)-(p) [31,33,112,150,207].Even more critical, columnar grains in LDED Al alloys can be eliminated by tuning Sc/Zr content (figures 13(h)-(m)).According to the reported temperature-time evolution at different pool depths in LDED Al-Mg-Sc-Zr alloys shown in figure 19(q) [31], the initial melt, whose peak temperature of the melting boundary was lower than the liquidus temperature of Al 3 Sc, would consist of supersaturated liquid and residual Al 3 Sc.Incorporated with the lower cooling rate in this zone, high-density primary Al 3 Sc would precipitate at the fusion boundary during resolidification (figures 13(a)-(c)) and effectively promote the heterogeneous nucleation of α-Al, forming a fine equiaxed grain region.Different peak temperatures lead to the difference in Al 3 Sc density at different depths of the melt pool, which is an important reason why the columnar and equiaxed grains are alternating, as shown in figures 19(n)-(p).
In addition, the shape and size of melt pools may vary from the bottom to the top in LDED Al alloy components, which significantly affects the direction of grain growth.The optimal grain growth direction is along the melt pool heat dissipation's fastest path [147].Due to the different curvature of the melt pool in different regions, there is a slight difference in the direction of grain growth (figures 11(c) and (d)).

Intrinsic heat treatment (IHT)
The LDED process of layer-by-layer deposition has a unique thermal cycle, which can trigger the IHT in the deposited materials, thus making possible the elimination of the subsequent complicated heat treatment process [208].This characteristic is particularly prominent in Al-Cu, Al-Li, Al-Sc-Zr, Al-Zn-Mg, and Al-Si alloys [37,150,[209][210][211].The IHT effect in LDED Al alloy has the following influence: (i) there may be differences in the type and size of precipitates along the build direction due to complex thermal history; (ii) precipitate coarsening; (iii) promote the precipitation of coherent nano-precipitates which play a crucial role in strengthening the material; (iv) the nano-scale precipitates that are precipitated during IHT inhibit grain growth; and (v) spheroidizing Si phases in Al-Si alloys.

Temperature-time sensitivity of precipitates.
Chen et al [209] proved the temperature-time sensitivity of the precipitates in different regions of LDED 2A97 Al-Li alloy.The component would experience a repeated thermal cycle due to a layerwise deposition, affecting the elements' distribution, precipitation behavior, and morphology of non-equilibrium solidified microstructures.Generally, in the deposition process, the bottom region is most affected by the heat cycles and stays in the high-temperature state for the longest time.The heat accumulated and the IHT effect during deposition had little impact on the top area.The T 2 -Al 6 CuLi 3 and T B -Al 7 Cu 4 Li phases exist in each region, whose content in the bottom region was In addition, precipitate coarsening occurring in the IHT process is one of the critical factors causing the unsatisfactory mechanical properties of LDED Al alloy.Kürnsteiner et al [211] found that 10 23 m −3 Al 3 (Sc, Zr) particles could be precipitated during IHT in Scalmalloy ® alloys, but the precipitates at the bottom layer were coarsened significantly.However, after adding 0.4 wt.% Zr, the excess Zr element was solidly dissolved in the Al matrix and formed Zr-rich shells, which greatly impeded the coarsening of secondary Al 3 (Sc, Zr) precipitates during IHT, as shown in figure 20(e).This finding has significant guiding potential for improving the mechanical properties of Al-Sc-Zr alloy.

Precipitation of coherent nano-precipitates.
Furthermore, the shearable nano-precipitates in Al alloys had remarkable strengthening and toughening effects [52,212,213].Li et al [210] found that the IHT during the LDED process can also facilitate the precipitation of shearable nano-precipitates in Al alloys.In LDED Al-5.32Si-1.19Cu-0.46Mg-0.09Fe alloy, the π-Fe phases, Si phases, θ-Al 2 Cu phases and Q-Al 5 Si 6 Mg 8 Cu 2 phases were formed in the process of rapid solidification.They remained stable during the thermal cycle, whose peak temperature was above the liquidus temperature of Al-5.32Si-1.19Cu-0.46Mg-0.09Fe alloy.When the peak temperature of the thermal cycle was between the Al-Si eutectic transition temperature and the liquidus temperature, the Si, π-Fe and θ-Al 2 Cu phases aggregated and coarsened.As the thermal cycle progresses, when the peak temperature was lower than the Al-Si eutectic transition temperature, the θ-Al 2 Cu and Q-Al 5 Si 6 Mg 8 Cu 2 phases dissolved, and the π-Fe phases transformed into β-Fe phases.When the peak temperature was further reduced, a large amount of Q ′ phases that were coherent with α-Al ((0001 [ 110] Al [214]) precipitated, and the phase composition inside this region tended to be stable.Therefore, many Q ′ phases existed in the bottom region, as shown in figures 20(f)-(h).In Al-Zn-Mg-Cu alloys, Li et al [37] found that the MgZn 2 phases were in a metastable state under the high cooling rate of LDED.In addition, due to the addition of Si, nano-scale Mg 2 Si phases were formed during IHT, and the nailing effect of nano-Mg 2 Si and nano-MgZn 2 phases significantly refined the microstructure of the as-built samples.

Spheroidization of Si phases.
Beyond that, IHT may also cause the spherization of Si phases, which is beneficial for reducing the stress concentration in Al-Si alloys.As shown in figure 20(i), (∼90 ± 2.0)% Si phases were spheroidized during IHT of LDED AlSi10Mg alloy [204], and the rest of the Si phases kept the non-sphere shapes, which was related to the diffusion and cluster of Si elements, as well as the nucleation and growth of Si phases during the thermal cycle.After the spheroidized Si phase with a critical radius is precipitated during IHT, the spheroidized Si phase may dissolve into the matrix again when the thermal-cycling temperature is higher than the equilibrium temperature of the Al/Si phase.In contrast, the nucleation, growth, and coarsening of the Si phases may be promoted when the thermal-cycling temperature is lower than the equilibrium temperature of Al/Si phase [204].
By customizing the IHT process during the intermittent deposition strategy through machine learning and integrating it synchronously with the 3D geometry, 4D printing with high energy efficiency and sustainability can be achieved.This idea has attracted the attention of some researchers in the field of steel, and in-depth explorations have been carried out [208].

Post heat treatment
The LDED process usually results in uneven microstructures and stress accumulation inside Al alloy components.Postheat treatment is generally necessary to optimize the microstructure to achieve higher mechanical properties.The roles of post-heat treatment on microstructure evolution of LDED Al alloys mainly include the following aspects: (i) to alleviate element enrichment in as-deposited samples and promote the uniform distribution of elements; (ii) weakening residual stresses; (iii) homogenizing the misorientation of grain boundaries; (iv) spheroidizing the Si phase; and (v) facilitating precipitationhardening.
Li et al [36] found that the element enrichment phenomenon was severe in the as-deposited Al-5Si-1Cu-Mg alloy, and many plate intermetallic phases were formed.After solution treatment at 500 • C for 6 h, Cu and Mg elements were evenly distributed, and many Al 2 Cu phases were dissolved in the matrix, as shown in figures 21(a) and (b).Also, as mentioned in section 4.2, many LAGBs generally exist in asdeposited LDED Al alloys, which is unfavorable to mechanical properties.It was found that the misorientation angle distribution of grain boundaries was more evenly distributed after post-heat treatment, due to adequate recrystallization diffusion time, as shown in figures 21(c)-(h) [96].For LDED Al-Si alloys, another essential purpose of post heat treatment is to spheroidize Si phases, which results from the thermallyactivated growth along the steadiest plane with the minimum free energy [215].The spheroidized Si phase can be more evenly dispersed in the Al matrix and better hinder crack initiation and propagation [216].Moreover, it was found that the spheroidization degree of Si phases may increase with the increase of solution temperature, but extended solution time may lead to the re-coarsening of the spheroidized Si phases, as shown in figures 21(i)-(o) [36,67].In addition, an important reason that LDED Al alloys require post-heat treatment is to induce phase transformation to generate precipitates with a significant strengthening effect, or to promote further the precipitation of second phases that precipitate in the process of IHT, improving the mechanical properties.The reported phase transitions during the post-heat treatment of several LDED Al alloys have been listed in table 4.

Tensile properties of LDED-processed Al alloys
The tensile property of LDED-processed Al alloys is one of the most intuitive ways to evaluate their potential engineering application.As outlined above, the interplays among materials composition design and modification (e.g.particle reinforcement), process parameters optimization, and deployment of appropriate heat treatment profiles are critical to achieving good mechanical properties.The room-temperature tensile properties of different series of LDED-processed Al alloys in as-built and heat-treated conditions are summarized in figures 22(a)-(d) and table 5.

Mechanical properties of Al-Si alloys.
For the Al-Si alloys, it is easier to generate cellular structures (∼5 µm) with fine Si particles distributed at cellular boundaries due to the high cooling rate of LDED (figure 19(j)).Thus, the mechanical properties of LDED Al-Si alloys are often higher than those of the traditional cast.AlSi10Mg alloy with good printability is the most widely used Al alloy in LDED, mainly attributed to its suitable solidification interval and low hot cracking tendency.The ultimate tensile strength (UTS) of AlSi10Mg alloy can exceed 250 MPa, but the elongation (EL) hardly exceeds 8% [67,76,113,204,217].Zhou et al [113] achieved excellent UTS of 423 MPa along with EL of 6.4% in LDED AlSi10Mg through layer-by-layer LSP during LDED, showing a higher strength-plasticity synergy than that of LPBFprocessed AlSi10Mg alloy.This is because the LSP treatment significantly promoted grain refinement, accelerated dislocation proliferation, and introduced compressive residual stress strengthening.In Al-Si alloy, the α-Al grain nucleates around the Si particles.With the increasing concentration of Si in the liquid phase, the alloy composition develops toward Al-Si eutectic, which contributes a lot to the strength of Al-Si alloy [219,220].The mechanical properties of Al-Si alloys can be significantly improved by optimizing the morphology and size of the Al-Si eutectic phase [204].In AlSi7Mg alloy and Al-5Si-1Cu-Mg alloy, the UTS (150-250 MPa) decreased due to the decrease in Si content [34,36,145].Note that the UTS of Al-5Si-1Cu-Mg alloy after T6 treatment could increase to 304-416 MPa while maintaining EL above 10% [36].The spheroidized Si phases after heat treatment with dimensions within 1-5 µm can further improve the properties of the asbuilt Al-Si alloy.The aggregation, diffusion, and Ostwald ripening of Si phases in the heat treatment process make the Si phase's distribution more uniform, reducing the mechanical anisotropy of the as-built sample.During the tensile process, a tensile fracture often starts from the brittle eutectic Si.The Si particles could be torn into smaller Si fragments, thus separating the two adjacent grains.It was proposed that increasing solution temperature could significantly reduce the proportion of torn Si particles and significantly enhance the mechanical properties of LDED Al-Si alloys [36].[39].Sc, Zr, and reinforcement particles are often used to improve the printability and formability of LDED Al-Mg alloys, and the subsequent strength-plasticity synergy in as-built conditions can be improved significantly [150,159,183,207], which can achieve the highest level of strength in the wide series  [20,38,119,147,194,209]; Al-Mn alloy [160]; Al-Si alloy [34-36, 59, 67, 68, 76, 113, 145, 204, 217]; Al-Mg alloy [39,120,150,159,183,207]; Al-Mg-Si alloy [218]; Al-Zn alloy [37,64,78,88,96,112,181], Al-Cr alloy [203]), and (b) post-heat-treated (Al-Cu alloy [38,194], Al-Mn alloy [160], Al-Si alloy [36], Al-Mg alloy [120,150], Al-Mg-Si alloy [148,218], and Al-Zn alloy [37,64,96,112,181]).
of LDED Al alloys (figures 22(a) and (b)).Hua et al [159] explored LDED Al-9 Mg alloy with high Mg content, which achieved a higher Mg solid solubility, grain refinement (2-30 µm equiaxed grains) and a weak texture.Under this condition, the UTS at the as-built state reached 400 MPa, the highest among the LDED of Al-Mg alloys.Moreover, the LDED of Al-Mg-Sc-Zr alloy assisted with water cooling of the substrate is a valuable approach toward acquiring excellent strength-plasticity synergy.Wang et al [150] found that the effective heat release through water cooling of the substrate effectively suppressed the precipitation and coarsening of primary and in-situ secondary Al 3 (Sc, Zr), which contributed to the formation of heterogeneous microstructures and promoted the precipitation of secondary Al 3 (Sc, Zr) during artificial aging.Thus, outstanding strength-plasticity synergy in the as-build state can realize further strengthening after direct aging.Analogously, Xiao et al [160] added Sc and Zr in LDED Al-Mn-Mg alloy, and both the as-built and heat treated samples acquired excellent strength-plasticity synergy, with the UTS reaching 519 MPa after heat treatment.Moreover, some investigations studied the optimization of Al-Mg alloys by adding Sc/Zr and Ti-based reinforcement together [120,183].The reinforcing particles can help Sc/Zr to refine the microstructures and decrease the porosities further.Interestingly, it was proposed that the extra reinforcement particles and Al 3 (Sc, Zr) may compete in the priority and effectiveness of microstructure refinement [120,221].On the one hand, the prevenient phases, as a small fraction of all the possible active particles for the heterogeneous nucleation, may prevent other grain refiners from obtaining the undercooling needed for nucleation by releasing latent heat during the phase growth [222].On the other hand, fine reinforcement particles may aggregate into clusters and play a dominant role in stimulating heterogeneous nucleation before precipitating Al 3 (Sc, Zr).In addition, as for Ti-based reinforcement particles, Zhao et al [120] found that Ti and Sc tend to segregate on the precipitate surface.Thus, three negative effects on mechanical properties may be caused by: (i) reducing the effectiveness of heterogeneous nucleation cores; (ii) generating harmful needlelike intermetallics (e.g.(Ti, Zr) 5 Si 3 , Al 3 Ti, and Al 4 Ti, etc [183,223]); and (iii) co-poisoning between Zr and Ti [120,224].It was found that although the microstructures of LDED AA 5024 + 3 wt.%TiC exhibited fully fine equiaxed grains, they were coarser than those of AA 5024.Moreover, the addition of 3 wt.%TiC reduced the UTS in heat-treated AA 5024 alloy from 347 to 335 MPa along with the YS reduction from 296 to 265 MPa.

Al alloy References Alloy composition
As-built   [69,225,226].For this reason, the content of Zn in the feedstock of LDED Al-Zn alloys reported so far can only reach ∼6 wt.% [112].Deng et al [112] confirmed the alloy composition of Al-Zn-Mg-Sc alloy with the lowest hot-cracking sensitivity in LDED by adopting the model proposed by Kou [227] based on the liquid film theory.Cu and Mg elements segregate easily at a high cooling rate and form a large number of brittle intergranular compounds (e.g.Al 2 CuMg), which have a high possibility of causing solidification cracking and poor mechanical properties, and the strengthening effect of Zn and Mg elements are also limited [37].Methods to improve the printability of Al-Zn alloys include: (i) Alloying anti-thermal cracking alloy elements [37,64,112].Due to the effective filling effect of Al-Si eutectic on cracks, Si element can be applied to eliminate cracks in LDED Al-Zn alloys [37,64].Sc and Zr elements are also commonly used to promote the formation of equiaxed grain structures in LDED Al-Zn alloys, reducing hot-cracking sensitivity and promoting precipitation strengthening [64,112].Li et al [64] reported that the strength-plasticity synergy of Al-5.10Zn-1.90Mg-1.47Cu-2.90Sialloy increased with the addition of Zr, but the addition level had a maximum limit value.Excessive Zr would make it difficult to reach the critical undercooling degree in the low solute density region and challenging to promote new grain nucleation.(ii) Addition of reinforcement particles [37,181] with low thermal conductivity.ZrO 2 [37] and TiC [181] have been reported to strengthen LDED Al-Zn-Mg-Cu-Si alloys.ZrO 2 and TiC can increase laser absorption of Al powders, interact with the dislocations and effectively hinder the dislocation motion, thus enhancing the relative density and strength of LDED-processed Al alloys.However, it is easy to induce the initiation and propagation of cracks around the particles; therefore, a high content of reinforcement particles may reduce the plasticity dramatically [181,183].(iii) Optimization of laser processing parameters and deposition strategies [88].Wang et al [88] found that the porosity was negatively correlated with the specific energy (the ratio of laser power to wire feeding mass per unit time) and k value (the ratio of wire feeding rate to laser scan speed) in LDED AA7075 alloy.When the laser power decreases with the increase of layers, it can avoid heat accumulation and is beneficial to achieving high strength-plasticity synergy in LDED Al-Zn alloys (UTS of 402 MPa, and EL of 9.2%).However, it is still difficult to eliminate the large pores and cracks in the samples due to interlayer cooling between adjacent deposition layers.(iv) Optimization of heat treatment profile [96].Heat treatment with multiple solutions plus multiple aging treatments is more significant in promoting the precipitation of the dominant strengthening phases η ′ -MgZn 2 , θ ′ -Al 2 Cu, and S-Al 2 CuMg phases, thus it is more conducive to the improvement of the mechanical properties of LDED Al-Zn alloys [96,112].According to existing reports, the UTS of LDED Al-Zn alloys can stably reach 300-400 MPa after heat treatment, as plotted in figure 22(d) [37,64,96,112].

Mechanical properties of Al-Cu alloys.
Like Al-Zn alloys, Al-Cu alloys are also subjected to high thermal cracking sensitivity, resulting in limited tensile properties.The UTS of as-built LDED Al-Cu alloys can only reach 200-300 MPa, which is still far from the commercial Al-Cu alloys processed by rolling and extrusion (typical 400-500 MPa) [53,[228][229][230].TiB 2 has been reported several times to reinforce LDED 2024 Al alloys.The TiB 2 size was noted to tremendously influence the strengthening effects [20,38,119].Chen et al [38] found that large TiB 2 particles with an average size of 2-10 µm enhanced the strength of LDED 2024 Al alloy, while the ductility was reduced due to the coarse TiB 2 particles and network-like θ-Al 2 Cu phases distributed on the grain boundaries acted as the crack sources.However, Wang et al [20] and Wen et al [119] found that TiB 2 particles with a mean size of ∼500 nm can effectively enhance the strengthplasticity synergy of LDED 2024 Al alloys.The solid solution heat treatment can dissolve θ-Al 2 Cu phases and precipitate T-Al 20 Cu 2 Mn 3 phases, improving grain bonding and strengthplasticity synergy.The aging process promotes the precipitation of the semi-coherent θ ′ phase, further increasing strength.But aging may cause the regeneration of Al 2 Cu phases, which is unfavorable to the tensile properties [38,194].The UTS and EL of LDED Al-Cu alloy after heat treatment are in the range of 350-450 MPa and 8%-20%, respectively [38,194], which is outstanding for LDED Al alloys after heat treatment, as shown in figure 22(d).

Mechanical properties of Al-Mg-Si alloys.
Al-Mg-Si alloys also have high thermal cracking sensitivity, which is unsuitable for AM [231].Due to the high difficulty in preparing its LDED components, there is relatively little research on the LDED of Al-Mg-Si alloys.Lee et al [218] found that the UTS and EL of as-built LDED 6061 Al alloy were only 76.8 MPa and 2.1%, respectively.The strengthening of B 4 C particles can enhance the UTS and EL to 270.6 MPa and 6.4%, respectively, by effectively alleviating the formation of cracks, pores, and balling during LDED.Li et al [148] successfully eliminated thermal cracks by effectively releasing thermal stress and interrupting the continuous growth of the intergranular liquid film through an interlayer pause.The tensile properties of T6-treated LDED 6061 Al alloys are even better than those of wrought 6061 Al alloys.

Mechanical anisotropy of LDED Al alloys
Due to the melt-pool-by-melt-pool and layer-by-layer deposition manner in LDED, the materials will undergo different thermal histories and solidification behavior in different regions.This results in different grain morphologies and intensity of texture, which may directly lead to the anisotropy of mechanical properties.It was indicated that the microstructures in the longitudinal section had intense <100> texture.The grains tended to grow along [001] orientation [112,194,204], as shown in figure 23(a).Unlike the longitudinal section, the microstructures in the cross-section are close to equiaxed crystals (figure 23(b)).In addition, Hua et al [159] found  that the porosity in the longitudinal section and horizontal section was also different.The longitudinal section's porosity was higher, leading to lower elongation (figures 23(c)-(h) and table 6).In Al-Si alloys, the distribution of Si phases has a preference orientation.Li et al [145] found that when the tensile test was along the longitudinal direction, microcracks in adjacent Si phases were more likely to join, leading to fracture due to the small spacing between Si phases along the crack growth direction.This made the sample alloys possess better ductility in the horizontal direction, as per the diagram shown in figures 23(i) and (j).In addition, Li et al [145] also proposed that reducing the layer thickness contributed to the formation of equiaxed grains and weakened the orientation preference of Si phases, thus reducing the anisotropy in mechanical properties, as shown in table 6.Generally, the tensile properties along the horizontal direction are better than those along the longitudinal direction due to better metallurgical bonding and even grain size [159].

Grain boundary strengthening.
Grain boundary strengthening mainly results from grain boundary obstruction to dislocation movement, which can be quantified by the Hall-Petch formula [160,232]: where σ 0 represents the friction stress of pure Al that equals to 20 MPa, k is the Hall-Petch coefficient (0.17 MPa•m 1 2 ), and d denotes the average grain diameter.When the microstructure is refined, the grain boundary density will increase, which can better hinder the dislocation motion and improve the YS.Wang et al [20] reported that 3 wt.%TiB 2 nanoparticles refined the microstructures of LDED Al-Cu alloys from 419.1 µm to 12.3 µm, contributing 18.9 MPa of grain boundary strengthening to the YS increment.Moreover, it is worth mentioning that heterostructured microstructures are typical in LDED Al alloys [31,112].The contribution of the heterostructured grains to the YS of LPBF AlZnMgCuScZr alloy was estimated according to the optimized Hall-Petch relationship as follows [233]: where V UFG and V (MEG+CCG) represent the volume fraction of the ultrafine grains and the medium-sized equiaxed grains and coarse columnar grains, respectively; d 1 and d 2 denote the average grain size of the ultrafine grains and the medium-sized equiaxed grains and coarse columnar grains, respectively.

Solid solution strengthening.
Solid solution strengthening is induced by the strain field, which can hinder dislocation movement due to the difference in atomic radius and shear modulus between alloying elements and the Al matrix, which can be represented as [234]: where H i is the solid solution strengthening constant of different alloy elements, C i is the alloy element concentration (at%), and n i is the power law coefficient of different alloy elements.However, it should be noted that, due to the aging precipitation process of IHT and post-heat treatment, some alloying elements exist in the form of a precipitated phase.Thus, the value of solid solution strengthening is difficult to determine, but the calculated lower limit value obtained through Thermo-Calc calculation and the upper limit value obtained according to formula (8) can be given [233].

Precipitation strengthening.
First, for the precipitates or reinforced particles that are larger than 8 nm [235,236], the mechanism between them and dislocations is mainly Orowan strengthening, which may form a dislocation loop around the precipitates or reinforcement particles and hinder dislocation motion under the action of the stress field.The contribution value of Orowan strengthening to the YS can be estimated by [235,237]: where M, b and G denote the Taylor factor, burgers vector, and shear modulus of Al, which equals 3.06, 0.286 nm, and 25.4 GPa, respectively.λ represents the spacing of particles, which can be estimated by formula (9), where f and r denote the particle volume fraction and particle average radius, respectively.Second, there may be a strain field close to the precipitates and reinforcement particles when the interface between particles and matrix is coherent, offering an interaction effect with dislocations.The contribution part of the interaction effect is called coherency strengthening and can be calculated by [238]: ( rf 0.5Gb where α ε equals 2.6, which is a constant for FCC metal.ε represents the lattice misfit between the particles and the Al matrix.Xiao et al [160] reported that the estimated ∆σ Orowan and ∆σ CS of the secondary Al 3 (Sc, Zr) in the Al-Mn-Mg-Sc-Zr alloy was 175 and 68 MPa, respectively, and the total contribution of precipitate strengthening to the YS was 243 MPa, which was the highest among all the strengthening mechanisms.Third, if the particles possess a small enough size and can be cut directly by the moving dislocations, the antiphase boundary produced by the cutting process can lead to ordering strengthening, which can be calculated according to [239]: where γ APB is the energy of the antiphase boundary of the cut precipitates.The ordering strengthening is the dominant strengthening mechanism of the precipitates with a mean size of ∼2 nm [240].Fourth, when the dislocation interacts with the precipitates, the energy of the dislocation changes.The modulus strengthening caused by the difference in shear modulus can be calculated by the following formula [64,241]: where ∆G denotes the difference of shearing modulus between precipitates and Al matrix, and m equals 0.85.Sun et al [242] reported the dislocation shearing mechanism of Al 2 CuLi (T1) phases in LDED Al-Cu-Li alloys, contributing greatly to the excellent high-temperature mechanical properties and work-hardening capacity.

Dislocation strengthening.
The difference in thermal expansion coefficient between the Al matrix and precipitates or reinforcement particles can generate a large number of dislocations, due to the complex thermal cycle during the LDED process of Al alloy, and the interaction between dislocations may generate a local stress field, whose contribution to YS can be expressed as [243,244]: where β is the dislocation strengthening factor, equal to 0.16 for Al, and ρ CTE represents the dislocation density.Wang et al [20] investigated the mechanical properties of LDED Al-Cu alloys strengthened by TiB 2 particles.They estimated that the vast difference in the coefficient of thermal expansion between the Al-Cu matrix and TiB 2 (the thermal expansion coefficient of TiB 2 and Al is 7.8 × 10 −6 and 24 × 10 −6 K −1 , respectively [245]) contributed 21.3 MPa to the YS.
5.3.5.Load transfer strengthening.Concerning the LDED particle reinforced Al alloys, the good interface bonding between reinforcement particles and Al matrix can effectively assist the Al matrix to bear load, thus increasing YS.It can be calculated as follows [246]: where σ i is the interface bonding between reinforcement particles and the Al matrix.It was reported that the strong interface bonding between TiB 2 and Al in LDED 3 wt.%TiB 2 /2024 Al alloy contributed 27.9 MPa to the YS [20].
5.3.6.Other strengthening mechanisms.Zhou et al [113] reported compressive residual stress strengthening in LDED AlSi10Mg alloy treated by in-situ LSP.The in-situ LSP treatment benefited the healing of pores and decreased the stress intensity factor [247].The substantial contribution value of compressive residual stress strengthening to the YS can be estimated by [248]: where σ t and ε p denote the tensile residual stress before the treatment and surface plastic strain, respectively.v represents the Lamé constant, which equals 0.32.L p and r p are the affected depth and pulse laser radius, respectively.Based on the calculation result by Zhou et al [113], the contribution of compressive residual stress strengthening to YS was 83 MPa, which was higher than 10 MPa of grain boundary strengthening and 17 MPa of dislocation strengthening in their study.So far, numerous studies have calculated the specific contribution of different strengthening mechanisms to the YS in different Al alloys.In Al-Si alloys, Si phases play a significant role in mechanical property strengthening, especially Si precipitate strengthening and load transfer strengthening.The calculation by Shi et al [204] showed that the contribution value of Si precipitate strengthening and load transfer strengthening in LDED AlSi10Mg alloy was approximately 82.4 and 51.3 MPa, which dominated 41.4% and 25.8% in the overall estimated YS.In Al-Mg alloys, the solid solution strengthening of the Mg element is dominating.Wang et al [150] found that the solid solution strengthening of Mg contributed 19.5% to the calculated YS of direct-aged LDED Al-3.31Mg-0.52Sc-0.22Zralloy with substrate water cooling considering Mg evaporation loss.Notably, the coherency strengthening of Al 3 (Sc, Zr) contributed the most to the YS, whose proportion reached 48.0%.The remarkable coherencystrengthening effect of Al 3 (Sc,Zr) has also been found in Al-Zn alloys and Al-Mn alloys [64,160].In Al-Zn alloys, Al-Cu alloys, and Al-Mg-Si alloys, the precipitate strengthening is the primary strengthening mechanism.Liu et al [249] compared the mechanical properties of Al-7.75Zn-1.43Mg-2.33Cu alloy fabricated by WAAM and laser-arc hybrid AM.It was found that the Orowan strengthening of η and η ′ phases dominated 89.8% of the overall grain boundary strengthening, solid solution strengthening, and Orowan strengthening in terms of the YS enhancement of laser-arc hybrid AM compared with wire arc manufacturing.Chen et al [38] investigated the influence of Al 2 Cu precipitate evolution on the mechanical properties of heat-treated LDED TiB 2 /Al 2024 composites.During aging, Cu elements gradually segregated and tended to form coherent θ ′′ phases, semi-coherent θ ′ phases and completely non-coherent θ phases.Moreover, the content of the θ ′ -Al 2 Cu phase increased with the increasing aging time.It was indicated that the YS of TiB 2 /Al 2024 composites was markedly improved from 120.65 MPa at solid solution state to 245.73 MPa after aging for 15 h.In Al-Mg-Si alloy, Li et al [218] found that the precipitation process during artificial aging was prominently facilitated due to the formation of ultra-high supersaturated solid solution in LDED 6061 Al samples.As a result, the number density of nanoscale precipitates (β ′′ -Mg 5 Si 6 and Q ′ -AlCuMgSi) in the as-built samples was much higher than that in the traditional cast and wrought 6061 Al alloys, which was conducive to realize more significant coherency strengthening of nanoscale precipitates and better strength-plasticity synergy in T6-treated LDED 6061 Al alloy.

Reasons for mechanical property difference between LDED and LPBF
LDED and LPBF are the two mainstream LAM processes, and some studies specified the difference between LPBF and LDED-processed Al alloys [76,78,207].The mechanical performances of Al alloys processed by LDED are generally inferior to those processed by LPBF.The underlying reasons could be associated with the following aspects: (i) the melting loss of alloying element; (ii) feedstock oxidation; (iii) grain size; (iv) element distribution; (v) heterostructured grains; (vi) dislocation density; and (vii) intermetallic distribution.The difference in grain size, alloying element melting loss, and feedstock oxidation of LPBF and LDED samples has been described in section 2.2.Other differences will be discussed in this section.
(i) The higher cooling rate of LPBF increases the solute trapping capacity of the Al matrix, which facilitates the uniform distribution of solute atoms (figure 24(a)).In contrast, the solute elements are inclined to distribute near the dendrite boundaries in LDED Al alloy (figure 24(b)).This is because a lower cooling rate increased the solute distribution time while reducing the solubility [76].Therefore, Al alloys processed by LPBF show better solution strengthening than LDED.(ii) Heterostructured grains and intermetallic distribution.
Wang et al [207] studied the LDED and LPBF of Al-Mg-Sc-Zr alloys; completely equiaxed grain structures with an average grain size of 8 µm were formed in LDED sample, while heterogeneous grain sizes composing refined cellular and coarse columnar grains were created in the LPBF sample.This heterogeneous microstructure in LPBF samples is mainly due to the different precipitation behavior of the primary Al 3 (Sc, Zr) phase at the fusion boundary and inner melt pool.In contrast, the primary Al 3 (Sc, Zr) phase distributed throughout the entire melt pool in LDED samples achieves a more uniform grain size.Note that the back stress hardening caused by strain partitioning between equiaxed grains bands and columnar grains improved the strain hardening ability and strengthductility synergy of the LPBF sample.Moreover, the strain hardening rate of LPBF samples can be divided into three stages, which were different from the trend of continuous decline of strain hardening rate in LDED samples, as shown in figure 24(a).At low strain, the strain hardening rate continued to decline and was lower than that of LDED samples, and then in the second stage, the strain hardening rate started to rise and then maintained a constant value, and then decreased in the third stage.The larger strain hardening rate of the LPBF sample at the high strain stage was beneficial to delaying the local deformation and neck shrinkage of LPBF samples during the tensile process.As listed in table 7, the YS of the LPBF sample was approximately 2.8 times that of the LDED sample while keeping similar elongation.Zhu et al [233] also reported the heterostructured grains in LPBF AlZnMgCuScZr alloy, the volume fractions of the ultrafine grains and the mediumsized equiaxed grains and coarse columnar grains were 47% and 53%, respectively.The estimated contribution from grain boundary strengthening was about 141 MPa, which accounted for 23.8% of the contribution value of all strengthening mechanisms to YS. (iii) Dislocation density.The proliferation capacity of dislocation is normally correlative with grain size and nonuniformity.Grain refinement would promote dislocation accumulation at grain boundaries and increase the dynamic recovery rate.The heterostructured grain is conducive to generating strain gradients during deformation and geometrically necessary dislocations [250].
According to Wang et al's study [207,251], the alternate coarse-grain and fine-grain distributions in the LPBF sample could effectively distribute the strain distribution and promote the generation of geometrically necessary dislocations, thereby enabling the back stress strengthening effect in the material to obtain a good strength-ductility synergy.
Intriguingly, Gong et al [76] processed AlSi10Mg alloy by combining LPBF and LDED and found that the microhardness of the LPBF zone was higher than the LDED zone (figure 24(b)).Under the influence of heat input of the LDED process, spheroidization and dendrite disappearance of Si phases occurred in the LPBF region below the interface.In addition, an equiaxed grain region with a height of 80 µm was formed near the interface, as shown in figure 24(c).The tensile properties showed that the fracture occurred in the LDED region.

Summary
This review focused on research progress in LDED Al alloys with an elaborated overview of the challenges, strategies and opportunities.The challenges associated with the intrinsic characters of Al alloys and manufacturing processes were summarized, and potential strategies to improve the printability were elucidated (figure 25).The correlation between microstructures and mechanical properties was discussed, and benchmarking mechanical properties of a wide range of Al alloys were established.
From the perspective of materials, the inherent characteristics of Al alloys (i.e.high laser reflectivity, high oxidation sensitivity and hydrophilicity, poor fluidity and large solidification interval) are the fundamentals leading to poor printability in LDED.The technical differences between LDED and LPBF in terms of the print environment, thermal history, and cooling rates also raised more challenges for the application and extension of LDED Al alloys.The challenges in process optimizations are the generation of porosity, cracks, residual stress, inclusions, selective evaporation of alloying elements, and coarse columnar grains during LDED Al alloys, which could lead to building failure and poor mechanical properties.Materials innovations by in-situ alloying and reinforcing particle addition, process parameters optimization, and auxiliary field deployment are the key strategic approaches to improve the material printability, reduce defects, and enhance resultant mechanical performance.Generally, optimizing processing windows is the most straightforward approach when depositing a commercial Al alloy with relatively good printability.Apart from optimizing the laser energy input (e.g.laser power and scan speed), the multidirectional deposition, interlayer pause, substrate preheating or cooling could affect the printability and performance.
The microstructures of LDED-processed Al alloys can be affected by process parameters, auxiliary fields, alloying elements and reinforcement particles, and the intrinsic and post heat treatments.The formation of refined equiaxed grains and heterostructured grains due to different thermal histories and LDED parameter control could strengthen the materials.Intrinsic and post heat treatments can optimize the phase composition, homogenize element distribution, reduce residual stress and spheroidize the Si phases to enhance the mechanical properties.in-situ alloying is vital in changing the solidification range, refining microstructures, and facilitating precipitate via heterogeneous nucleation and second-phase strengthening.At present, Sc and Zr are the most commonly used alloying elements, demonstrating good capability in tailoring printability, microstructure and mechanical properties.Adding reinforcement particles (e.g.TiC, SiC, TiB, and TiB 2 ) effectively compensates for the high laser reflectivity and increases the laser absorptivity and mechanical performance.It is worth noting that ex-situ reinforcement particle addition is more convenient, and the particle content is easy to control.In comparison, in-situ formation of reinforcement particles in the Al matrix can obtain uniform particle distributions with good particle/matrix wettability and bonding strength.The LDED Al alloys with auxiliary fields like mechanical deformation, magnetic field and acoustic fields can affect the melt pool dynamics and grain recrystallizations, which also provides possibilities for tuning the microstructure of Al alloys.
The mechanical property comparison shows that the asbuilt Al-Mg alloys exhibit relatively good tensile strengthductility synergy, but as-built Al-Zn alloys and Al-Cu alloys show unsatisfactory performances due to the high probability of alloying element evaporation and high thermal cracking sensitivity.Nevertheless, post-heat-treated Al-Zn alloys and Al-Cu alloys exhibit greater potential in achieving excellent mechanical properties (figure 22).The underlying strengthening mechanisms in LDED-processed Al alloys include grain boundary strengthening, solid solution strengthening, precipitate strengthening, dislocation strengthening, and other strengthening mechanisms, such as residual stress strengthening.Furthermore, the mechanical properties gaps between LDED and LPBF Al alloys are discussed in this review.The reasons for poor mechanical properties in LDED-processed Al alloys generally can be ascribed to oxidation, severe element segregation and evaporation, coarse microstructures, and a lower dislocation due to a lower cooling rate.Notably, the literature related to the fatigue properties and high-temperature tensile properties of LDED Al alloys is still minimal, leaving a knowledge gap in objectively assessing the potential practical application of LDED Al alloys.

Perspective
There is an increasing demand by the leading aerospace companies (e.g.Rolls Royce and Boeing) for high-strength and high-value Al alloy components to be deployed in aerospace and in maintenance, repair and overhaul.To fill the gap between R&D and industrial applications, as summarized in figure 26), future innovations in materials, process, defects monitoring and microstructural control for LDED Al alloys are critical to further improve the printability and performance of Al alloys and drive certification and applications.

Materials customization.
The current commercial Al alloys for LDED mostly inherit compositions designed for traditional manufacturing routes, which may suffer from poor printability.New Al materials customization needs to address the technical features of LDED.For instance, higher Zn and Mg content in feedstock may be recommended to compensate for the elements' evaporation like Mg.Also, as the Al element shows high oxygen affinity, the deoxidizing element addition to react with oxygen could be a good approach to reduce the oxidation of Al.As for the Al-Zn and Al-Cu alloys, the relationship between hot cracking sensitivity and alloy composition can be predicted by physical hot cracking models, which can be used as a basis for more accurate alloy composition design.Machine learning can be utilized effectively to assist the composition design and the customization of Al alloys for the LDED process, which saves the cost and time of trial-and-error processes [257,258].

Process innovation.
The optimal design of process parameters and the prediction of the process window can be realized by a proper algorithm [259,260].Based on some practical model training, such as cutting-edge convolutional neural networks and machine learning models, the probability of defect occurrence under specific parameters can be accurately predicted.The prediction accuracy of some current models can reach more than 90% [152,252].In addition, it is worthwhile to develop new energy sources for LDED Al alloys.Compared with traditional infrared wavelength lasers,  [6,30,138,139,[252][253][254][255][256].Reprinted from [138], © 2019 Elsevier B.V. All rights reserved; Reprinted from [139], © 2018 Elsevier Ltd.All rights reserved; Reproduced from [30].CC BY 4.0; Reproduced from [254].CC BY 4.0; Reprinted from [252], © 2023 Elsevier Ltd.All rights reserved; Reprinted from [253], © 2023 Elsevier Ltd.All rights reserved; Reprinted from [255], © 2022 Central South University.Publishing services by Elsevier B.V. on behalf of KeAi Communications Co. Ltd; Reprinted from [256], © 2021 Elsevier B.V. All rights reserved; Reprinted from [6], © 2021 Elsevier Ltd.All rights reserved.
Al alloys have a higher absorption rate for green or blue lasers [140,261], which can improve the density of LDED Al alloys and realize the manufacture of near-net-shape structures.In the case of alloying element evaporation and powder contamination, the feedstock composition can be monitored and compensated in real time.Moreover, the inert environment LDED system protected by inert gas or even a vacuum environment LDED system are also reasonable approaches for LDED Al alloys.

Microstructure control.
The precipitates, fine equiaxed grains and fine-coarse heterogeneous grains have demonstrated good capability in strengthening Al alloys.Diverse effective heat treatment strategies can be tried in the future to better optimize intermetallic phases, Si phases, and precipitate behavior, such as ingeniously utilizing IHT and multi-step aging.There is also a need to understand the formation and control mechanisms of heterogeneous grains (with fine and coarse grains at different regions) to further enhance the bulk mechanical properties of Al alloys.In addition, multiple energy field-assisted designs, such as mechanical field, thermal field, acoustic field, and magnetic field can be attempted to overcome the limitations and improve the printability and formability of LDED Al alloys [193,194,199,201].At present, there is only a limited amount of literature investigating field-assisted LDED Al alloys, especially ultrasound-assisted, electron beam-assisted and magneticassisted LDED Al alloys.

Defect and surface quality monitoring.
Due to the inherent characteristics of Al alloys (such as high thermal conductivity and laser reflectance, high oxidation sensitivity and hydrophilicity, poor fluidity, high coefficient of thermal expansion and wide solidification interval) and the technical challenges of LDED, it is generally difficult to obtain high density and defect-free LDED Al alloys.In industrial applications, controlling the dimensional accuracy and surface quality of LDED Al alloys has also been challenging.Online defect detection can be popularized in the case of stringent requirements for the quality control of LDED-fabricated parts.Through the extraction and visualization of signal features, the corresponding defects and dimensional errors can be detected, classified, and predicted by the training model.in-situ monitoring by visual and acoustic signals are the main methods for online defect and dimension error prediction [262,263].Vision-based monitoring has somewhat accomplished industrial readiness, but disadvantages lie in the high time and economic cost.The real-time monitoring of LDED based on acoustics requires high sensor resolution and may need further efforts to reduce interference by background noise.Multi-sensor monitoring and data fusion by combining multiple signals (e.g.visual, thermal, and acoustic signals) could enhance the accuracy and reliability of in-situ defect detection.
Moreover, applying additive and subtractive hybrid systems can instantly remove the defects detected or allow instant surface treatment and machining during real-time monitoring.Also, the deposition process's temperature distribution, residual stress, etc., can be predicted and optimized by applying numerical simulation and modeling, thereby providing practical guidance for improving the surface quality and dimensional accuracy of LDED Al alloy components [264,265].On this basis, defect elimination, high dimensional accuracy, and high surface quality can be successfully achieved in LDED Al alloys to fabricate high-quality parts.

Certification and application.
High-value parts repairing and large-format components manufacturing are two critical strategic applications of LDED Al alloy.These highvalue and large format Al alloy components can be a wing spar, bracket, nozzle liner, nuclear thermal propulsion chamber, powerhead shell, etc.For the sake of safety, the certification of precision aviation LDED Al parts is quite strict, which requires the highest level of quality and repeatability.LDED is a new technique compared with traditional casting and plastic forming, implying that the LDED-processed Al parts need to undergo rigorous validation and certification.New qualification standards and standard operating procedures may be needed for benchmarking and referencing.Realizing the digitalization and intelligent manufacturing of aviation parts could also reduce operator variation and enhance process stability and repeatability.New digitalized processing chains for LDED higher performance Al alloy components will be the future R&D forefront to improve production efficiency since the current LDED process needs multiple workforce interventions.
In summary, the exploration of LDED high-performance Al alloys is still an ongoing topic that attracts many worldwide researchers.How to address the inherent disadvantages of Al alloys in LDED and give full play to their advantages by combining material science and innovative processes could still be the main direction for future research.Microstructural engineering strategies to obtain favorable microstructures and mechanical properties should also be highlighted for LDED high-performance Al alloys.Process stability and repeatability are critical for the certification and application of high-value and large-format Al components produced by LDED.

Figure 2 .
Figure 2. Process schematics and comprehensive comparison of LPBF and LDED.(a) and (b) Schematic diagram of LPBF and LDED, respectively.(c) Radar map illustrating comparison between LPBF and LDED.

Figure 3 .
Figure 3. Laser absorption rate and solidification behavior of Al alloys.(a) Laser absorption rates of typical metals, and the temperature versus solid fraction in (b) Al-Zn alloy and AlSi10Mg alloy, (c) Al-Mg alloy, and (d) Al-Cu alloy.Reproduced from [58].CC BY 4.0; Reproduced from [69], with permission from Springer Nature; Reprinted from [70], © 2016 Elsevier B.V. All rights reserved; Reprinted from [71], © 2020 Elsevier B.V. All rights reserved.

Figure 5 .
Figure 5. Cracking and formation mechanisms in LDED Al alloys.(a) Scanning electron microscopy (SEM) of microcracks in LDED 7075 Al alloys.(b) SEM of the red frame in (a) exhibiting the crack surface.(c) SEM of the green frame in (a) revealing insufficient liquid filling during solidification.(d) SEM of blue frame in (c) exhibiting grain boundary intermetallic particles.Reprinted from [100], © 2021 Published by Elsevier Ltd on behalf of The Society of Manufacturing Engineers.(e) Liquation cracks of LDED 7075 Al alloys produced on the cold plate [101].(f) Liquation cracks of LDED 7075 Al alloys produced on the preheated plate.Reprinted from [101], © 2022 The Authors.Published by Elsevier B.V. (g) Cross-section observations of LDED produced 7075 Al alloys under different powers of 200 W, 240 W, and 280 W, respectively.Reprinted from [96], © 2022 Elsevier B.V. All rights reserved.(h) An illustration depicting possible effects of solute segregation on cracking behaviours of LDED of Al alloys.Reprinted from [103], © 2021 Acta Materialia Inc. Published by Elsevier Ltd.All rights reserved.

Figure 7 .
Figure 7. Inclusions in LDED Al alloys.(a) The relationship between the laser absorbance of the Al powder and the illumination wavelength.(b) Light optical microscope map of single tracks built with virgin powders.(c) Light optical microscope map of single tracks built with aged powders.(d) The relationship between oxygen and hydrogen contents in Al powder with aging time.(e) 1-dimensional measurements of the tracks extracted from light optical microscope maps of the single tracks.Reproduced from [59].CC BY 4.0.(f) Schematic diagram of the formation of pores by inclusion in LDED of Al.Reproduced from [89].CC BY 4.0.

Figure 8 .
Figure 8. Selective evaporation of alloying elements in LDED Al alloys.(a) The six stages of evaporative release.Reproduced from [132].© IOP Publishing Ltd.All rights reserved.(b)The evolution of the bubble burst process from single laser shots at a fixed delay of 320 µs[133].Reproduced from[133].CC BY 4.0.

Figure 10 .
Figure 10.Influence of laser-related parameters on microstructures of LDED Al alloys.(a)-(d) Computed thermal cycle for different (a) laser power, (b) deposition speed, and (c) laser spot size.Reprinted from [33], © 2021 Elsevier Ltd.All rights reserved.(d) AlSi10Mg powder density maps obtained from high-speed imaging.Reproduced from [81].CC BY 4.0.(e) A graph representing the relationship between temperature gradient and growth rate in LDED of Al alloys.Reprinted from [138], © 2019 Elsevier B.V. All rights reserved.(f) Main effect diagram of S/N ratio-relative density in LDED of AlSi10Mg alloys.(g) Correlation between relative density and energy density of LDED AlSi10Mg alloy.Reprinted from [139], © 2018 Elsevier Ltd.All rights reserved.(h) Energy density of a Gaussian infrared laser beam.(i) Melt pool generated by the infrared laser.(j) Energy density of a flat-top blue laser beam.(k) Melt pool generated by the blue laser.Reprinted from [140], © 2023 The Author(s).Published by Elsevier B.V.

Figure 12 .
Figure 12.Machine learning processes in LDED Al alloys.(a) Data acquisition of machine learning.(b) Feature treatment of machine learning.(c) Proposed in-situ defect detection framework through acoustic signal processing and deep learning.Reprinted from [152], © 2023 Elsevier B.V. All rights reserved; Reprinted from [157], © 2020 Elsevier Inc.(d) Optimizing and verification of machine learning.Reprinted from [156], © 2020 The Society of Manufacturing Engineers.Published by Elsevier Ltd.All rights reserved.

Figure 14 .
Figure 14.The characterization of powders with ex-situ reinforcement particles in LDED Al alloys.(a) SEM of powders of TiC-reinforced Al alloys.(b) Detailed view of powders of TiC-reinforced Al alloys.(c) Enlarged view of (b).(d) 52 • tilted SEM cross-section of the LDED samples fabricated by the powders of TiC-reinforced Al alloys shown in (a).Reproduced from [176].CC BY 4.0.(e)-(g) EDS mapping results of Al and Ti of AA7075 powders with TiB 2 reinforcement particles.(h) SEM of 4 wt.%TiB 2 /AA7075 LDED component.(i) and (j) Electron probe wavelength-dispersive spectroscopy (EPWDS) result of Ti and B in 4 wt.%TiB 2 /AA7075 LDED component.Reprinted from [121], © 2018 Elsevier Ltd and Techna Group S.r.l.All rights reserved.
5 wt% ex-situ TiB 2 particles refined the LDED 2024 Al alloy from 431.53 µm to 114.76 µm, refining by about 26.6%, as shown in figures 15(f) and (g).Moreover, the addition of TiC can effectively improve the laser absorption rate of Al powder by nearly one order of magnitude, as shown in figure 15(h).The ex-situ particles can cover the surface of the Al powders

Figure 16 .
Figure 16.Role of in-situ reinforcement particles in laser additive manufacturing Al alloys.(a) Schematic diagram of technique route for Al-Cu-Mg powders with in-situ TiB 2 nanoparticles.(b) SEM-BSE, IPF-EBSD map, and corresponding grain boundary map of the cross-section of Al-Cu-Mg powders with in-situ TiB 2 nanoparticles.Reprinted from [43], © 2022 Elsevier B.V. All rights reserved.(c) SEM of the (B 4 C@Ti)/AlSi10Mg composite powder.(d) BSE of the AlSi10Mg matrix composites with B 4 C@Ti after LAM.(e)-(g) TEM of the detailed microstructure near the residual unreacted B 4 C phase in AlSi10Mg matrix composites.(h) High-resolution TEM (HRTEM) at the interface between the in-situ ceramics and Al matrix.Reprinted from [192], © 2023 Elsevier B.V. All rights reserved.(i) The laser reflectivity curves of the 2024 Al powders with different contents of in-situ TiB 2 nanoparticles.(j) The relative density of LAM 2024 Al alloys with different contents of in-situ TiB 2 nanoparticles.Reproduced from [191].CC BY 4.0.

Figure 17 .
Figure 17.Field-assisted LDED of Al alloys.(a) Sketch map of the LDED device with wire oscillating laser.(b) and (c) Effects of laser spot oscillating on pore distribution: (b) t + 0 s; (c) t + 1 s.Reprinted from [194], © 2022 The Authors.Published by Elsevier Ltd on behalf of the Chinese Mechanical Engineering Society (CMES).(d) Schematic diagram of the LSP-assisted LDED.(e) Sketch map of the pore closure mechanism under the shock wave at different maximum pressures.Reprinted from [113], © 2022 The Author(s).Published by Elsevier B.V. (f) Sketch map of the laser-arc hybrid additive manufacturing system.(g)-(i) The grain refinement mechanism map: (g) Melt pool with single TIG.(h) Melt pool with hybrid laser-TIG.(i) Enlargement of block A in (h).Reprinted from [66], © 2019 Elsevier B.V. All rights reserved.(j) Circuit layout for AC magnet.(k) Laser welding system under a electromagnetic field.(l) Pore distribution in as-weld AlMg3 alloy without or with a magnetic field.Reprinted from [199], Copyright © 2013 The Authors.Published by Elsevier B.V. (m) Sketch map of the laser cladding process [201].(n) and (o) Deep surface rolling processing on a laser-clad specimen using a computer numerically controlled machine.(p) Neutron diffraction results of hydrostatic stresses in the laser-clad samples.Reprinted from [201], Copyright © 2014 Elsevier B.V. All rights reserved.(q) Sketch map of the ultrasonic vibration-assisted LDED process of 4047 Al alloy.(r) Optical images of the 4047 Al alloy fabricated by ultrasonic vibration-assisted LDED under different remelting laser powers.Reproduced from [202].CC BY 4.0.

Figure 18 .
Figure 18.The typical microstructures in as-built LDED Al alloys.(a)-(f) Microstructure of LDED and direct-cast AA3104 alloy: (a) and (d) IPF map.(b) and (e) Grain boundary misorientation image.(c) and (f) optical and SEM maps.Reprinted from [32], © 2021 The Authors.Published by Elsevier Ltd. (g)-(i) OM and (l) and (m) SEM images of the microstructure of Al-Cr-Fe alloy: (g) overview perspective of a single melt pool.(h) Enlarged view of coarse intermetallic region (Region I).(i) Enlarged view of coarse laths.(j) Enlarged view of fine intermetallic region (Region II).(k) Enlarged view of fine particles.Reprinted from [203], © 2020 Elsevier B.V. All rights reserved.

Figure 19 .
Figure 19.The typical grain morphologies in LDED of different Al alloys.(a)-(e) OM images from the top area to the bottom area of LDED Al 2024.Reprinted from [147], © 2018 Elsevier Ltd.All rights reserved.(f) Microstructure formation mechanism of LDED Al 2024 alloy.Reprinted from [38], © 2020 The Author(s).Published by Elsevier B.V. (g)-(j) Microstructures of as-built LDED AlSi10Mg alloy: (g) OM image of the longitudinal section.(h) Enlarged view of the yellow box in (g).(i) OM map of the cross-sectional morphology of dendrites, (j) SEM image of the cross-sectional morphology of dendrites.(k) Al-Si binary equilibrium phase diagram calculated by Thermo-Calc ® software applying the TCBIN database.(l) A schematic diagram of the quasi-steady melt pool: P1 to P9 are on the central axis of the melting temperature isothermal surface.(m) The variation tendency of the G•R and G/R.Reprinted from [204], © 2023 Elsevier B.V. All rights reserved.(n)-(p) IPF maps of mixed grains at the top, middle, and bottom regions of LDED Al-0.5Sc-0.5Zralloy.Reprinted from [33], © 2021 Elsevier Ltd.All rights reserved.(q) Temperature-time dependency chart of the thirteen diverse points at the central axis of the melt pool from the bottom region to the top region.Reprinted from [31], Copyright © 2021, Elsevier.

Figure 20 .
Figure 20.Effect of intrinsic heat treatment on the precipitates in LDED Al alloys.(a) Mechanism diagram of regional differences in thermal cycling and precipitated phases of 2A97 Al-Li alloy during the LDED process.(b)-(d) Bright-field TEM images and SAD patterns of different regions of LDED 2A97 Al-Li alloys along the building direction.Reprinted from [209], © 2021 Elsevier Ltd.All rights reserved.(e) Atom probe tomography measurements taken from the bottom of the LDED Scalmalloy® alloy and AlSc1.0Zr0.4alloy.Reprinted from [211], © 2019 Elsevier B.V. All rights reserved.(f) and (g) TEM and HRTEM images of nanosized Q ′ phase in LDED Al-5.32Si-1.19Cu-0.46Mg-0.09Fe alloy, respectively.(h) Comparison of XRD patterns for the top and bottom region of LDED Al-5.32Si-1.19Cu-0.46Mg-0.09Fe samples.Reprinted from [210], © 2017 Elsevier B.V. All rights reserved.(i) Bright-field TEM image and sphere and non-sphere Si phases of the LDED AlSi10Mg alloy.Reprinted from [204], © 2023 Elsevier B.V. All rights reserved.

Figure 23 .
Figure 23.Understanding the mechanical anisotropy in LDED Al alloys.(a) and (b) EBSD pole figure (PF) and IPF of LDED Al-Zn-Mg-Sc alloy along the longitudinal section and cross-section, respectively.Reprinted from [112], © 2022 Published by Elsevier Inc. (c)-(h) Tensile fracture morphologies of LDED Al-Mg-Sc-Zr alloys at the longitudinal section and cross-section.Reprinted from [159], © 2022 Published by Elsevier Ltd on behalf of the Chinese Mechanical Engineering Society (CMES).(i) and (j) Simplified diagram of the fracture process of the columnar grain zone along the build direction and perpendicular to the build direction.Reprinted from [145], © 2019 Elsevier B.V. All rights reserved.

Table 1 .
Comparison of cooling rates in melt pools during LPBF and LDED of AlSi10Mg alloys.Reprinted from [76], © 2022 Elsevier B.V. All rights reserved.

Table 2 .
[39]er feed rate (PFR), scan speed and laser energy density for LDED 5083 Al powders with a constant laser power of 600 W[39].