Advancements in Laser Powder Bed Fusion of Carbon Nanotubes-Reinforced AlSi10Mg Alloy: A Comprehensive Analysis of Microstructure Evolution, Properties, and Future Prospects

Abstract

Laser powder bed fusion (L-PBF) stands out as a promising approach within the realm of additive manufacturing, particularly for the synthesis of CNT-AlSi10Mg nanocomposites. This review delves into a thorough exploration of the transformation in microstructure, the impact of processing variables, and the physico-mechanical characteristics of CNT-AlSi10Mg nanocomposites crafted via the L-PBF technique. Moreover, it consolidates a substantial corpus of recent research, proffering invaluable insights into optimizing L-PBF parameters to attain the desired microstructures and enhanced properties. The review centers its attention on pivotal facets, including the dispersion and distribution of CNTs, the formation of porosity, and their subsequent influence on wear resistance, electrical and thermal conductivity, tensile strength, thermal expansion, and hardness. In line with a logical progression, this review paper endeavors to illuminate the chemical composition, traits, and phase configuration of AlSi10Mg-based parts fabricated via L-PBF, juxtaposing them with their conventionally manufactured counterparts. Emphasis has been placed on elucidating the connection between the microstructural evolution of these nanocomposites and the resultant physico-mechanical properties. Quantitative data culled from the literature indicate that L-PBF-produced parts exhibit a microhardness of 151 HV, a relative density of 99.7%, an ultimate tensile strength of $70\times {10}^3 \frac{{\mathrm{mm}}^3}{\mathrm{N}.\mathrm{m}}$, and a tensile strength of 756 MPa.

1. Introduction

Aluminum and its alloys have become prevalent in industrial applications thank to some beneficial properties, among which satisfactory thermal and electrical conductivity, excellent strength-to-weight ratio, and high corrosion resistance can be mentioned [1,2,3,4,5]. Numerous production techniques have been developed to facilitate the mass production of aluminum components and parts. Among these techniques, stir-casting, thixomolding processes, and powder metallurgy approaches are prominently recognized as conventional methods [6,7]. Hence, a multitude of opportunities emerges for the application of aluminum alloys within diverse industrial domains, encompassing sectors such as transportation, aerospace, energy, and power [1,2,3,8]. Over the past decade, a new-fashioned additive manufacturing technology has emerged as a promising approach to the development of geometrically complex industrial parts and structures using aluminum alloys [4,9,10,11,12]. This novel iteration of powder bed melting/sintering technologies has facilitated the commercial production of components with enhanced efficiency, diminished inventory costs, and shortened lead times [10,13,14,15]. It is worth noting that various additive manufacturing methods, including heterogeneous multi-wire indirect arc additive manufacturing, are suitable for producing in-situ aluminum-based alloys and composites [16].

In this section, the authors endeavor to elucidate the factors contributing to the widespread adoption of AlSi10Mg alloy, a popular aluminum alloy, and laser powder bed fusion (L-PBF) technology, a prevalent additive manufacturing technique, for the fabrication of industrial components. Moreover, this section aims to provide compelling insights into the question regarding the superior efficiency of L-PBF compared to alternative methods when applied to AlSi10Mg alloy.

1.1. AlSi10Mg Alloy: Composition, Characteristics, and Phase Constituents

The initial question to be addressed concerns the underlying factors that are associated with the widespread acceptance and adoption of the AlSi10Mg alloy within the realm of additive manufacturing methodologies. As previously mentioned, aluminum and its alloys have gained significant popularity in the industrial sector for the production of various parts. However, it is important to note that each type of aluminum alloy possesses distinct advantages and disadvantages. For instance, certain alloys necessitate a sufficiently high temperature for modification through heat treatment, whereas others encounter challenges associated with high viscosity during the casting process, leading to well-recognized production complications [9,14,17,18]. In the midst of these considerations, the AlSi10Mg alloy has emerged as a potential solution to address some of the aforementioned drawbacks. AlSi10Mg is a non-heat treatable alloy with high silicon content whose chemical composition is at that of a near-eutectic, so-called hypoeutectic alloy [19,20,21]. Conforming to the Al-Si binary phase diagram, the AlSi10Mg alloy undergoes a near-eutectic transformation at its specific composition, thereby benefiting from enhanced castability. The occurrence of a eutectic phase transformation in the AlSi10Mg alloy has the potential to reduce the temperature range at which the molten alloy initiates the solidification process during casting methodologies [22,23,24]. This phenomenon contributes to the alloy’s advantageous solidification behavior, facilitating more precise control over the casting process and mitigating potential issues related to inconsistent solidification and undesirable microstructural formations [25,26,27].

The AlSi10Mg alloy derives several advantages from its excellent fluidity and low shrinkage, which are considered crucial criteria in material selection for casting applications. Additionally, the alloy exhibits favorable characteristics, such as corrosion resistance, cost-effectiveness [28,29,30], good machinability, weldability, and high mechanical stability at elevated temperatures [31,32,33]. These attributes further contribute to the desirability of the AlSi10Mg alloy for various industrial applications, particularly in additive manufacturing and fabrication processes.

The composition of the AlSi10Mg alloy primarily consists of two principal constituents: silicon (9–11 wt%) and magnesium (0.25–0.45 wt%) [34]. These elements play a crucial role in imparting the alloy with its distinctive characteristics, including desirable mechanical properties, corrosion resistance, and wear behavior. Notably, the inclusion of a controlled amount of silicon effectively reduces the probability of crack formation and hot tearing during solidification [35]. This is attributed to the dimensional expansion tendency of silicon during the solidification process, which effectively counteracts excessive shrinkage and mitigates the unfavorable effects associated with it [36].

The phase composition of the AlSi10Mg alloy primarily comprises α-Al solid solution, Mg₂Si, and Si phases [37,38], as elucidated by the Al-Si-Mg phase diagram. Among these phases, the α-Al solid solution predominates and holds considerable influence over the mechanical properties of the alloy. The presence of the Mg₂Si phase in the AlSi10Mg alloy plays a significant role in mechanical strengthening mechanisms by forming small particles that effectively impede the movement of dislocations. Moreover, the Si phase contributes to enhancing the wear and corrosion resistance of the alloy [39,40,41,42,43]. Notably, the chemical reaction between silicon and magnesium facilitates the formation of Mg₂Si precipitates, which serve to reinforce the matrix, as well as MgAl₂O₄ particles that reduce the probability of aluminum oxide film formation [44]. This interplay of phases and reactions within the alloy structure contributes to its overall mechanical integrity and improved resistance against wear and corrosion [45,46,47].

1.2. Enhancing Performance through Composite Reinforcements

Despite the enumerated advantages highlighted in the preceding section for the AlSi10Mg alloy, it is crucial to acknowledge that monolithic aluminum alloys do possess certain weaknesses that impose limitations on their utilization across diverse industries [48,49]. These weaknesses encompass relatively low strength and hardness, inadequate wear resistance, and a lower thermal expansion coefficient [50,51,52]. Furthermore, although aluminum alloys like AlSi10Mg are frequently employed in additive manufacturing, they still fall short of meeting the growing requirements for high-strength and tough aluminum alloys [53,54]. These shortcomings underscore the ongoing pursuit for advancements in aluminum alloy development to address these performance limitations and fulfill the escalating demands of diverse industrial applications.

One common approach to address these limitations is the utilization of aluminum-based composites that incorporate uniformly distributed reinforcements. This strategy effectively mitigates the aforementioned shortcomings associated with monolithic aluminum alloys. Extensive research has been conducted to explore the impacts of diverse reinforced particles, such as oxides [55,56,57], carbides [58,59,60,61], nitrides [62,63,64], borides [65,66,67], and various allotropes of carbon [68,69,70], on the microstructure and mechanical properties of AlSi10Mg composites fabricated through an L-PBF approach. The findings from these studies have demonstrated promising outcomes, indicating significant enhancements in the resulting composites’ microstructural characteristics and mechanical properties.

Carbon nanotubes (CNTs) have emerged as one of the most promising reinforcing agents, garnering significant attention from researchers since their discovery by Iijima in 1991 [71]. The unique structural properties of CNTs, such as their high aspect ratio, exceptional mechanical strength, and excellent electrical and thermal conductivity, make them particularly attractive for enhancing the performance of composite materials [72,73]. As a result, extensive investigations have been conducted to explore the potential benefits of incorporating CNTs into various matrices, including the AlSi10Mg alloy, with the aim of improving their mechanical and functional properties.

Therefore, the incorporation of composite reinforcements, specifically CNTs, into AlSi10Mg shows great potential for overcoming the inherent limitations of monolithic aluminum alloys. Extensive research in the field of L-PBF has demonstrated promising outcomes in terms of improved microstructural characteristics and enhanced mechanical properties. This avenue of exploration opens up new possibilities for the development of high-performance aluminum-based composites that can meet the evolving demands of diverse industrial applications.

1.3. L-PBF: Advantages and Applications for AlSi10Mg Alloy Processing

The presence of industrial parts featuring intricate geometries and irregular configurations poses practical challenges for conventional manufacturing technologies, particularly in the case of casting processes. Consequently, a new generation of manufacturing technologies has emerged, known as additive manufacturing. This innovative approach enables the high-quality fabrication of intricate parts within a short timeframe, while also providing the flexibility to customize processing variables [74,75]. By utilizing additive manufacturing techniques, manufacturers can overcome the limitations of traditional methods and achieve efficient production of complex components. L-PBF, also known by other names [12,76,77,78], is a versatile additive manufacturing technique with diverse applications in various industries, including aerospace, automotive, and medical sectors [8,79,80,81]. Indeed, this fabrication method falls under the category of powder-bed fusion techniques. In this process, 3D metal components are constructed by depositing a thin layer of powdered feedstock onto a predetermined substrate. A focused laser beam is then used to selectively melt specific areas of the powder at temperatures above the material’s melting point. Subsequently, the molten material solidifies upon cooling to ambient temperature, forming a solid layer. This layer-by-layer approach is repeated until the desired component is fully formed [74,75,82].

Recently, a considerable body of research has emerged focusing on the utilization of L-PBF for the fabrication of intricate internal and external structures using various aluminum alloys, including Al-Si alloys and, notably, AlSi10Mg. However, a fundamental question arises: What makes L-PBF particularly suitable for manufacturing AlSi10Mg, and what potential benefits does it offer?

Several compelling reasons underlie the selection of L-PBF for this purpose, with the most significant ones outlined below:

• The high silicon content in the AlSi10Mg alloy’s chemical composition enhances its fluidity by promoting efficient heat dissipation during solidification. This improved fluidity reduces the likelihood of microcrack formation and facilitates the healing of any existing cracks [35].
• AlSi10Mg alloy exhibits low dimensional shrinkage and limited residual stress, thereby reducing its susceptibility to cracking during the solidification process. This characteristic is beneficial for achieving successful fabrication through L-PBF, as it minimizes the occurrence of detrimental cracks [83].
• The solidification range of AlSi10Mg alloy is significantly narrower compared to higher-strength aluminum alloys like the 7000 series. This narrower solidification range makes AlSi10Mg more amenable to the L-PBF process [84]. Additionally, L-PBF technology can accommodate high-temperature heat treatment cycles for the metallic powder feedstocks, further expanding its suitability for fabricating this alloy.
• L-PBF provides exceptional precision in fabricating complex shapes and geometries, which is crucial for producing components with intricate designs and features. This aspect is particularly significant for the AlSi10Mg alloy, commonly employed in the aerospace industry, where tight tolerances and intricate parts are required [14,85,86].
• L-PBF is an efficient manufacturing process that minimizes material waste compared to traditional methods. This is especially advantageous for the AlSi10Mg alloy, which is a relatively expensive material [87,88,89].
• L-PBF allows for the customization of parts according to specific requirements, which is vital for the AlSi10Mg alloy due to its versatility and diverse applications [85,90,91].
• L-PBF offers short lead times for part fabrication, which is crucial for industries that demand quick turnaround times. This aspect is particularly relevant for the AlSi10Mg alloy, as it is frequently utilized in the aerospace industry, where timely production is critical [92,93,94].

These reasons collectively support the use of L-PBF as a favorable method for processing AlSi10Mg, enabling precise, efficient, customizable, and timely production of components with complex designs and tight tolerances.

1.4. Evolution of Microstructure in L-PBF-Fabricated AlSi10Mg-Based Parts

In the preceding sections, the advantages of utilizing the L-PBF for AlSi10Mg alloy-based parts were discussed. However, an important question remains unanswered: What is the evolution of the microstructure in L-PBF-fabricated AlSi10Mg alloy-based parts? Understanding the microstructure evolution in these parts is crucial for gaining insights into potential issues and recognizing the phenomena associated with microstructure evolution, particularly in L-PBF-processed AlSi10Mg alloy reinforced with CNTs. To gain a better understanding of the microstructure evolution and phase formation in the AlSi10Mg alloy, it is beneficial to refer to the equilibrium pseudo-binary phase diagram of the Al-Mg₂Si alloy system (Figure 1a) and the binary phase diagram of Al-Si alloys (Figure 1b). According to Figure 1b, a eutectic reaction occurs at approximately 577 °C with a silicon content of around 12.3 wt%. Consequently, during the cooling process of the AlSi10Mg melt (containing 10 wt% silicon) to ambient temperature, primary aluminum particles begin to separate in the initial stage. In the subsequent stage, a eutectic reaction takes place, resulting in the formation of a continuous network of the aluminum-silicon eutectic phase [94,95]. Additionally, considering the thermodynamic perspective, the coexistence of silicon and magnesium can potentially give rise to the formation of metastable intermetallic precipitates known as magnesium silicide (Mg_xSi_y) β-phase [36,96], as illustrated in Figure 1a. These magnesium silicide (Mg₂Si) precipitates can effectively enhance the mechanical strength and lead to a reduction in grain size. Consequently, the resulting microstructure is characterized by the dispersion of Mg₂Si precipitates throughout the matrix, which consists of primary aluminum solid solution islands alongside a continuous network of the aluminum-silicon eutectic phase.

As anticipated based on metallurgical principles, the operational parameters employed in the L-PBF process have a direct influence on the morphology and phase distribution within AlSi10Mg alloy-based parts. Among these parameters, the cooling rate during the solidification of the laser beam-driven molten pool plays a significant role. For instance, Sathishkumar et al. [100] conducted a comparative study by fabricating an industrial AlSi10Mg-based part using both L-PBF and conventional casting technologies. The aim was to investigate the potential impact of processing variables on the microstructure and mechanical properties of the sintered parts. The results indicated that the cast alloy exhibited a dendritic α-Al phase structure enveloped by a eutectic Al/Si phase, whereas the L-PBF-fabricated part displayed a cellular α-Al phase structure bound by a fibrous eutectic Al/Si phase. The observed differences in microstructure between the cast and L-PBF-fabricated AlSi10Mg parts can be attributed to the distinct cooling rates during the respective processes. This mechanism is illustrated in Figure 2. During rapid solidification, α-Al precipitates from the molten alloy, and a eutectic reaction (L→α-Al+Si) occurs at approximately 800 K. As the temperature further decreases, nanometric secondary Si phases precipitate from the supersaturated α-Al and distribute evenly throughout the α-Al matrix. This precipitation continues until the entire melt is fully solidified. Upon complete recrystallization, the resulting microstructure consists of α-Al+Si (eutectic) and Si precipitates (secondary) [101]. Figure 3 displays the thermodynamically feasible phases in the microstructure of L-PBF-processed AlSi10Mg alloy.

2. AlSi10Mg Alloy: Comparing L-PBF with Traditional Fabrication Technologies

L-PBF represents a highly regarded and efficient form of additive manufacturing technology. The process commences after establishing the initial configuration and inputting the suitable process parameters into the 3D printer. To ensure uniformity, a roller on the manufacturing plate is employed to evenly distribute an appropriate quantity of powder throughout the entire procedure. The laser beam scans the powder feedstock in a specific manner, guided by the model input as a Computer-Aided Design (CAD) file. Upon completion of the pattern scanning for the current layer, the powder feed piston undergoes an upward displacement by one step/layer, while the production piston moves downward by one step/layer, thereby generating a new layer. This sequential layering process continues until the final layer is formed and the overall procedure is finalized [102,103,104]. The specific region where the laser beam engages with the powder layer is commonly referred to as the “interaction zone”. The interaction zone in the L-PBF typically consists of three distinct regions: (i) the solidification zone, (ii) the unaffected powdery zone, and (iii) the melt-evaporation zone. This complex composition of the interaction zone sets L-PBF apart from traditional manufacturing methods like casting and other powder metallurgy technologies, necessitating a multi-physics approach for its understanding. In order to achieve the desired final parts through L-PBF, careful adjustments of process variables such as laser power, laser spot diameter, and scanning speed are essential [102,104,105].

An additional differentiation between L-PBF and traditional methods in the fabrication of AlSi10Mg alloy-based components is the challenging nature of optimizing the printing process. This difficulty implies that the occurrence of microstructural and macrostructural defects, as well as the likelihood of their formation, is higher in L-PBF. Consequently, recent investigations focusing on AlSi10Mg alloy fabricated via L-PBF have placed significant emphasis on optimizing process variables [106,107], understanding microstructure development [108,109], exploring the formation mechanisms of pore defects [110,111], and evaluating the effects of post-treatments on the physico-mechanical properties [112,113]. Notably, Everett et al. [114] presented a case study demonstrating that the local geometry, even under constant production parameters, can influence the evolution of microstructure and mechanical characteristics of AlSi10Mg parts produced using L-PBF.

An additional distinction between L-PBF and traditional fabrication techniques is the extensive range of process parameters and the inclusion of variables that are absent in traditional methods. Notably, L-PBF introduces factors such as the structural angle (i.e., the orientation of the printed part relative to the laser beam), anisotropic thermal gradient, and local geometry of the printing part. These factors add complexity to the optimization and control of the L-PBF process, making it more challenging compared to traditional techniques. Illustratively, a case study focusing on AlSi10Mg alloy fabricated via L-PBF revealed anisotropic behavior concerning conductivity, diffusivity, thermal expansion, and a non-uniform distribution of pores at the borders of the melt pool. The root cause of these phenomena was attributed to the structural angle, highlighting its influence on the material properties [115]. Similarly, in another investigation, a statistical survey was conducted to analyze the grain characteristics in L-PBF-printed AlSi10Mg alloy-based components. The results demonstrated that vertically printed samples exhibited 75% columnar grains and 25% equiaxed grains, whereas these percentages decreased to 49% and 51%, respectively, in horizontally printed samples [116].

In contrast to conventional manufacturing procedures, the most significant distinction between L-PBF technology and traditional methods lies in the resulting microstructural features. In traditional AlSi10Mg casting, the slow cooling rate leads to the rapid separation of the solid solution of silicon in aluminum, forming relatively coarse silicon precipitates. Consequently, the microstructure of the cast AlSi10Mg alloy typically comprises coarse silicon particles, dispersed primary α-Al, intermetallic phases such as Mg₂Si and Fe-rich intermetallic particles, and a continuous eutectic structure of aluminum and silicon within an aluminum matrix. It is worth noting that the microstructure can be significantly influenced by post-heat treatment and variations in the cooling cycles [22,117,118]. In contrast, L-PBF facilitates rapid melting and cooling cycles for each specific location of the printed component, resulting in a distinct microstructure with a unique phase arrangement in the AlSi10Mg alloy. This microstructure significantly differs from that obtained through the casting process. Two distinct microstructure features have been observed in L-PBF: (i) a cellular-dendritic structure of α-Al and (ii) a network of the eutectic Si phase surrounding the α-Al phase boundary. The average dimensions of the cellular dendrites of aluminum in L-PBF are comparatively smaller than those developed through casting [37,117,119,120]. According to literature sources, L-PBF promotes the solidification of α-Al in a cellular morphology and enhances the solubility of silicon in the aluminum matrix [120]. The development of the microstructure and formation of thermo-kinetically probable phases in L-PBF-printed AlSi10Mg-based components are highly influenced by each processing parameter, particularly thermal gradients and scanning specifications. As a result, a wide variety of microstructural features and phase configurations can be expected in L-PBF-fabricated AlSi10Mg-based components (see Refs. [121,122,123,124]).

3. L-PBF-Printed CNT-Reinforced AlSi10Mg Composites

Before delving into the examination of how the incorporation of CNTs impacts the physico-mechanical properties of CNT-modified AlSi10Mg nanocomposites, it is beneficial to address the distribution of CNTs within a metallic matrix. As is well known, numerous studies have sought to create CNT-reinforced metal matrix nanocomposites with exceptional physico-mechanical attributes for various functional and structural applications. However, the inclusion of CNTs in metallic matrices is often hindered by two unresolved challenges: (i) the uneven dispersion of CNTs throughout the metallic matrices, and (ii) the limited interfacial adhesion between CNTs and the metallic matrix [125]. Due to their relatively high aspect ratio and remarkable specific surface area, CNTs are more prone to uneven dispersion within metallic systems compared to carbon particles and fibers [126]. To address this issue, a wide range of techniques have been developed to achieve uniform CNT dispersion and enhance the interfacial chemistry of nanotubes for improved properties. For readers seeking a deeper understanding of the practical aspects addressed in this review article, we recommend exploring recent review articles, like Refs. [125,127,128], that delve into this topic. In essence, techniques for CNT dispersion can be broadly categorized using various methods, which can be grouped into three major categories from a phenomenological standpoint: (i) colloidal mixing, involving physical interactions between CNTs and other species in aqueous and inorganic media; (ii) chemical mixing, where CNTs undergo chemical interactions with other components; and (iii) mechanical mixing, where CNTs are physically separated from each other. Importantly, these categories are not mutually exclusive, as a single dispersion method can belong to two or even all three groups simultaneously. For instance, if a chemical reaction occurs during the ultrasonication of CNTs in an aqueous medium, the ultrasonication method, typically associated with colloidal mixing, may also be classified within the chemical mixing category [125,129,130].

3.1. Process Variables

3.1.1. Impact of CNTs on Scanning Track Morphology

The production of components through L-PBF involves a sequential deposition of material, requiring a significant overlap between the scan tracks to achieve highly densified as-printed parts [131,132]. The ability of the printed layers to adhere to each other is greatly influenced by the wetting behavior of the liquid phase. Given the high energy consumption and resulting elevated temperature of the liquid phase in L-PBF, the oxide layer on the molten aluminum deteriorates, leading to a decrease in dynamic viscosity [133,134]. This implies that specific processing conditions that ensure lower dynamic viscosity and improved spreadability of the molten metal on the previously solidified layer will promote better interlayer wettability and a higher degree of microstructure densification. Consequently, the incorporation of CNTs is expected to have an impact on the morphology of the scanning tracks in L-PBF-manufactured AlSi10Mg-based parts.

The presence of CNTs on the surface of AlSi10Mg powders typically involves weak bonding. Under conditions of low scan speed and limited laser power during the printing process, the input energy for the powder feedstock is reduced, resulting in a lower number of melted particles and a smaller overall volume of the molten pool. This occurs because the surrounding powders are not easily drawn into the molten pool, leading to significant fluctuations in track height and non-uniformity. Furthermore, this phenomenon weakens the driving force for the solidification of the molten pool before spreading on the previously solidified layer. Consequently, in CNT-reinforced AlSi10Mg matrix nanocomposites, the height of the molten pool increases as the scan speed becomes higher [135,136].

Figure 4 illustrates the surface morphology of CNT-reinforced AlSi10Mg composites produced at different scan speeds. It can be observed that, when the scan speed is below 1500 mm/s, the surface roughness decreases due to a wider melting path and favorable overlap between scan tracks. The ideal scanning surface is achieved at a scan speed of 1300 mm/s, indicating a balanced relationship between the scan track width and the set hatch. At very low scan speeds of 1100 and 900 mm/s, the input energy is high, leading to the formation of splash droplets (highlighted in red in Figure 4), which can be detrimental to the subsequent layer. Moreover, if the scan speed exceeds 1500 mm/s, the distribution of scan line width becomes inconsistent due to weak overlap between the scan tracks [135]. These gaps between tracks will be transferred to the printed component and manifest as macrostructural defects, incompletely melted particles, and irregular pores [137].

3.1.2. Impact of Scan Speed on the Distribution of CNTs

The scan speed, which determines the energy input, plays a crucial role in achieving a uniform distribution of CNTs in AlSi10Mg alloy during printing. An optimized scan speed exists, below which a certain degree of CNTs agglomeration can occur. At lower scan speeds, the agglomeration degree of CNTs becomes more pronounced. On the other hand, at higher scan speeds, CNTs tend to disperse more uniformly with less clustering. This phenomenon can be explained by the Marangoni convection, which acts as the primary driving force for CNTs’ movement in the aluminum melt [138]. By reducing the scan speed and creating sharp temperature gradients in the melt pool, the Marangoni convection is intensified [139,140]. This stronger convection enhances the flow of CNTs into the melt pool and increases the frequency of CNTs’ transport to the melt surface. Due to the limited wettability between CNTs and the melt, along with the reduction of the driving force for CNTs’ redistribution within the melt pool, nanotube clusters are more likely to form [135].

3.1.3. Optimizing Laser Energy Density for Densification of CNT-Reinforced AlSi10Mg

The densification of parts produced by L-PBF can be influenced by a range of operational parameters. These variables can be combined into a single formula to define the laser energy density (LED) factor. The LED (expressed in J/mm³) represents the effective energy input per unit volume during the L-PBF process and can be calculated using the following equation [141,142]:

$$\mathrm{LED}=\frac{P}{v\times h\times t}$$

(1)

In the equation, P (W) represents the laser power, v (mm/s) denotes the scanning velocity, h (mm) represents the hatch spacing, and t (mm) is the layer thickness. Based on investigations of AlSi10Mg alloy produced via L-PBF, an optimum LED range has been identified, where high relative density and low defect formation can be achieved. Deviations from this optimal energy density can negatively impact the microstructure and mechanical properties, including hardness and tensile strength. Insufficient LED values may lead to the formation of undesirable pores, microstructural defects, and incomplete melting, while excessive LED values can result in the balling effect [101]. In the study conducted by Wang et al. [101], an optimal LED value of 131 J/mm³ was identified for achieving a high relative density in CNT-reinforced AlSi10Mg alloy-based components (refer to Figure 5). The results depicted in Figure 5 demonstrate that any deviation from this optimal LED value can disrupt the densification behavior during L-PBF of CNT-reinforced AlSi10Mg alloy. Specifically, lower LED values promote the formation of semi-spherical pores with larger sizes, while higher LED values lead to the development of micro-cracks with elliptical or triangular morphologies. The observed increase in density with increasing LED is attributed to the improvement in the wetting characteristics of the molten pool. Conversely, the decrease in density at higher LED values is believed to be caused by the occurrence of the balling effect [101,143].

In line with these findings, Yu et al. [144] demonstrated that utilizing a laser power of 450 W and a scan speed of 2 m/s enables the fabrication of a CNT-AlSi10Mg nanocomposite with minimal CNT agglomeration and noticeable absence of micropores or microstructural defects at the interlayer interfaces. Conversely, Du et al. [145] reported that lower laser powers, such as 50 W, do not facilitate the production of defect-free and fully dense CNT-AlSi10Mg binary components. In such cases, the LED is insufficient to completely melt and sinter the powder feedstock, but increasing the LED leads to the attainment of near fully dense samples with a relative density of 97%. However, it is important to note that low LED values may give rise to other disadvantages, including: (i) the formation of low-thickness layering structures due to limited laser penetration depth and shallower melt pools, (ii) the generation of spherical pores due to the entrapment of protective gas within the melt pools, and (iii) the evolution of large irregular voids due to incomplete melting of the powder feedstock and fusion of the previous layer [145].

3.1.4. Influence of CNTs on Surface Quality

The surface quality of components produced through L-PBF is a crucial aspect that directly influences their physico-mechanical properties. Achieving high surface quality in L-PBF-manufactured parts can be challenging due to the complexity and interrelation of the process variables. Surface roughness, in particular, plays a significant role in the reliability and service life of L-PBF-printed components as rough surfaces can initiate cracking.

Accurate quantification of surface quality and understanding its relationship with processing parameters are therefore essential. Researchers, such as Uzan et al. [146], have investigated the impact of surface roughness on the fatigue behavior of L-PBF-fabricated AlSi10Mg alloy. Their findings indicate that the elimination of superficial defects, including roughness, porosity, and un-melted particles, can significantly enhance the fatigue life of the components. However, according to some literature, surface smoothing does not have a notable effect on the tensile strength of L-PBF-printed AlSi10Mg components [147].

Optimizing the surface quality of L-PBF-manufactured parts requires a thorough understanding of how operational parameters contribute to achieving high-quality surfaces. Key variables such as layer thickness, inter-layer overlap rate, laser power, and scan speed play a crucial role in this process. The combination of scan speed and laser power is particularly important in achieving the necessary surface smoothness [148,149]. Studies, such as the one conducted by Yang et al. [150], have investigated the stability of printed tracks and surface quality of L-PBF-printed AlSi10Mg components in relation to scan speed and laser energy density (LED). The stability of deposited tracks initially improves and then decreases as LED increases, while the vertical surface roughness exhibits an opposite trend. At lower LED values, the balling effect can occur, leading to increased surface roughness. This happens because the molten pool does not have sufficient time to fully spread on the previously printed layer due to short solidification time. On the other hand, higher LED values can result in unfavorable phenomena such as metal evaporation and intensified Marangoni convection, causing surface instability and structural fluctuations.

When LED reaches an optimized level, an ideal surface quality can be achieved. At this point, the molten drops have enough time to solidify and spread on the neighboring layer without strong disruptive Marangoni convection. Similar principles apply to L-PBF-printed CNT-reinforced AlSi10Mg matrix nanocomposites, as demonstrated by Zhao et al. [82]. Lower LED values can lead to the formation of surface pores and discontinuous molten tracks, while increasing the laser power can reduce the presence of surface pores and result in a denser and smoother surface. In summary, finding the optimal combination of scan speed and laser power, within an appropriate LED range, is crucial for achieving high surface quality in L-PBF-printed components.

3.2. Microstructure

3.2.1. Powder Feedstock Preparation

Uniform element distribution and high fluidity are widely recognized as key attributes of an optimal composite powder blend for additive manufacturing. Thus far, various techniques have been employed to process composite powder feedstocks, including powder electrostatic self-assembly, ball-milling, and gas/plasma atomization [151]. These techniques have gained popularity due to their effectiveness in achieving the desired properties in composite powders used for L-PBF manufacturing applications.

• Ball milling is a widely utilized technique for processing composite powder feedstocks in L-PBF manufacturing. It involves the application of high-energy mechanical force to combine nanoparticles with metal powder. This process is typically conducted at room temperature and incorporates a process control agent (PCA). The principal mechanism involved is the repetitive deformation and welding of powder particles through collisions with the grinding balls, resulting in the formation of a fine-grained composite powder. By adjusting the grinding parameters, the processed blend can be optimized. However, this technique has a notable drawback: the resulting composite particles tend to experience undesirable superficial oxidation and exhibit reduced fluidity, leading to a loss of sphericity. These characteristics significantly impact the formability and densification behavior of printed products [152,153];
• Gas or plasma atomization is another technique used for the processing of composite powder feedstocks. In this method, the spherical composite powder is generated by a powder atomizer through the combination of reinforced particles with a molten aluminum alloy. This process yields highly uniform spherical composite powder with excellent flowability. However, it is important to note that this technique is relatively time-consuming and costly [154];
• Electrostatic self-assembly is another approach utilized for the blending of composite powders. It involves the even mixing of two dissimilar powders, where one powder carries a positive surface charge and the other powder carries a negative charge. The electrostatic attraction between the particles enables their uniform distribution. While this technique effectively addresses the limitations associated with ball milling, it is a complex process with limited performance. Moreover, it often leads to poor binding between the particles, as achieving consistent electrical charging can be challenging [155].

Based on a literature survey, it is evident that the ball milling process is widely adopted by researchers investigating CNT-AlSi10Mg binary systems in the context of L-PBF. It is worth noting that advanced techniques are often employed prior to ball milling to enhance the inter-particle bonding between CNTs and Al and achieve a uniform dispersion of the constituent phases. Two conventional approaches include the use of catalytic carriers for in-situ CNTs formation and the application of a binder on the metallic particles prior to ball milling. In the first technique, metallic powders or ceramic particles such as SiC and Al₂O₃ serve as catalytic carriers, providing suitable conditions for the in-situ growth of CNTs. If ceramic carriers are employed, the resulting composite will be a hybrid material with two reinforcing agents. However, if CNTs are grown directly on initial Al particles, the composite will consist of a single reinforcement, namely CNTs (refer to Figure 6) [156,157].

Conventional methods for preparing CNT-AlSi10Mg composites include the use of catalytic carriers for in-situ CNT synthesis and the application of a binder during slurry ball milling. In the first method, catalytic carriers such as metallic powders or ceramic particles like SiC and Al₂O₃ are employed, providing a substrate for the in-situ growth of CNTs. When ceramic carriers are used, the resulting composite contains two reinforcing agents. However, if CNTs are grown directly on initial Al particles, the composite consists of a single reinforcement, namely CNTs (refer to Figure 6) [156,157].

In the second technique, entangled CNTs are dispersed in ethanol along with a binder or surface modifier, typically a polymer. This mixture is subjected to ball milling under moderate and non-aggressive conditions. During milling, the AlSi10Mg particles maintain their spherical shape and can effectively adhere to the surfaces of the CNTs. The binder plays a crucial role in enhancing the inter-particle bonding between CNTs and AlSi10Mg. The milling duration determines the volume percentage of CNTs attached to the particle surfaces, while the rotation speed controls the dispersion of the powder particles [82].

For CNT-reinforced AlSi10Mg matrix nanocomposites, research findings have confirmed that shorter milling times result in a more uniform dispersion and improved flowability of the composite feedstock for L-PBF. Conversely, longer milling times and higher rotation speeds lead to the agglomeration of CNTs. The use of a binder, such as the PVP polymer, during slurry ball milling enhances the bonding strength between CNTs and AlSi10Mg. The CNTs become tightly adhered to the surface of the AlSi10Mg particles during this process, and their content and dispersion can be controlled by adjusting the milling duration and rotation speed [82].

In brief, powder feedstock preparation for CNT-reinforced AlSi10Mg matrix nanocomposites in L-PBF manufacturing involves various techniques, such as ball milling, gas or plasma atomization, and electrostatic self-assembly. Among these techniques, ball milling is widely adopted due to its ability to achieve a fine-grained composite powder. However, it can lead to superficial oxidation and reduced fluidity. Advanced techniques, like in-situ CNTs synthesis using catalytic carriers and the use of binders during ball milling, are employed to enhance inter-particle bonding and achieve uniform dispersion. Shorter milling times and the use of binders have been found to result in a more uniform dispersion and improved flowability, while longer milling times and higher rotation speeds can cause CNT agglomeration. By carefully adjusting the milling parameters, the content and dispersion of CNTs in the feedstock can be controlled, leading to optimized powder characteristics for L-PBF manufacturing.

3.2.2. Densification Behavior

When it comes to consolidating metallic alloy powder feedstocks, achieving full densification is generally not a significant challenge, and fully densified components can be successfully fabricated. However, the addition of CNTs to a metallic matrix can have a substantial impact on the relative density and consolidation rate of a two-component system. This is primarily due to the poor wettability between CNTs and the metallic matrix, as well as the tendency of CNTs to aggregate, leading to the formation of pores that impede densification and reduce the relative density of the resulting nanocomposites. Consequently, achieving full density in CNT-reinforced metal matrix nanocomposites is a critical issue that heavily relies on the distribution of CNTs within the matrix [125,127,128].

In the case of L-PBF-printed CNT-reinforced AlSi10Mg matrix composites, the incorporation of a small amount of CNTs into an Al matrix (typically ranging from 0.03 to 0.05 wt%) has minimal impact on the relative density of the L-PBF-manufactured bulk nanocomposites, which can still reach around 97%. However, as the CNT content is further increased up to 0.5%, the relative density gradually decreases, reaching a value as low as 91.23% [158].

3.2.3. Dominant Mechanisms in Microstructure Evolution

The fluid dynamics, chemical concentration, and temperature gradient at the interface between the solid and melt in the L-PBF process give rise to the formation of Marangoni flow and capillary forces. These capillary forces can exert a torque on nanotubes, causing them to undergo changes in their 3D configuration and distribution within the melt pool. Additionally, CNTs have the ability to absorb higher energy from laser irradiation compared to AlSi10Mg, which increases the likelihood of their thermal decomposition. Consequently, it is common to observe instances where CNTs undergo decomposition, evaporation, or structural deterioration following the L-PBF process [75]. A schematic representation of this concept can be seen in Figure 7.

The microstructure evolution of CNT-reinforced AlSi10Mg matrix nanocomposites in the context of L-PBF processing has received limited attention in research studies. However, the findings from the available research can be summarized as follows:

• In many cases, it is challenging to directly observe CNTs in the microstructure of L-PBF-processed samples. However, the presence of elemental carbon, which confirms the existence of CNTs, can be detected through elemental analysis. The elemental carbon is uniformly distributed within the AlSi10Mg matrix [82]. This suggests that CNTs may undergo partial or complete degradation during high-energy laser irradiation. The observed reduction in CNT length compared to the primary nanotubes and the significant decrease in their volume percentage indicate changes that occur during the L-PBF process [82,158];
• Prior to L-PBF, during the ball milling stage, the entangled CNTs become attached to the surfaces of AlSi10Mg particles while maintaining their tubular morphology. However, when exposed to high-energy laser irradiation, the thermal stability of CNTs is compromised due to defects induced by ball milling. This leads to their thermal decomposition into elemental carbon, which then diffuses within the melt pool as a result of Marangoni flow-induced melt vibration. The elemental carbon undergoes a chemical reaction with the Al matrix, leading to the in-situ formation of Al₄C₃. Consequently, the outer layer of the CNTs in contact with the Al matrix consists of Al₄C₃ [134,159];
• During the L-PBF process, a distinctive microstructure develops in CNT-reinforced AlSi10Mg matrix nanocomposites. This microstructure consists of solidified α-Al grains arranged in a cellular morphology, accompanied by a discontinuous network of the silicon phase. The size of the Al cells increases with higher CNT content [160]. Figure 8 depicts the scanning electron microscopy (SEM) image of L-PBF-printed CNT-reinforced AlSi10Mg matrix nanocomposites, clearly identifying both phases, namely eutectic silicon and cellular aluminum [159,161]. The proposed mechanism for the microstructure evolution in these composites is as follows: (i) Initially, the α-Al phase undergoes solidification, leaving residual silicon at the grain boundaries [120]. (ii) Subsequently, the silicon forms a supersaturated Al-Si solid solution, leading to the segregation of coarse elemental silicon particles and Si-Al eutectic sheets at the primary Al grain boundaries. The resulting microstructure, as depicted in Figure 8, exhibits brighter zones corresponding to the eutectic silicon phase segregated from the AlSi10Mg matrix, while darker regions represent the α-Al phase. It is notable that the aluminum regions are surrounded by silicon particles, maintaining their cellular equiaxed grain structure [75];
• The incorporation of CNTs into the AlSi10Mg matrix can induce a transition in the solidification regime, shifting from a dendritic microstructure to a cellular-dendritic microstructure [162,163];
• Elemental silicon is predominantly concentrated along the grain boundaries of the Al grains. It is deposited in the form of thin eutectic sheets with nanoscale thickness along these boundaries. The precipitation of eutectic silicon becomes more challenging with higher cooling rates employed during L-PBF, resulting in smaller generated precipitates. In other words, a higher LED corresponds to a greater likelihood of silicon particle formation in L-PBF [82];
• Increasing the laser power or LED can promote the growth of primary α-Al grains [82].

Some researchers have provided evidence of the presence of Mg₂Si precipitates and in-situ formed Al₄C₃ resulting from the thermal decomposition of CNTs and their subsequent phase transformation with the alloy matrix [82].

These findings provide valuable insights into the microstructural changes and chemical reactions that occur between CNTs and the AlSi10Mg matrix during the L-PBF process.

3.2.4. Potential Phase Transformations

During the L-PBF process of CNT-reinforced AlSi10Mg matrix nanocomposites, several chemical reactions can occur, leading to various phase transformations. The following is a brief summary of these transformations, highlighting their potential impact on the mechanical properties when properly controlled [75,82,134,158,164,165]:

• CNTs exhibit superior thermal conductivity and possess a large specific surface area, enabling them to efficiently absorb laser energy during the L-PBF process. As a consequence, the CNTs undergo thermal decomposition, leading to either the formation of elemental carbon or their evaporation. The extent of this decomposition is influenced by the laser energy density (LED) applied during L-PBF. Elevated LED values enhance the effectiveness of Marangoni flow, a phenomenon associated with vibrational motion within the melt pool. This intensified fluid dynamics promotes the diffusion of atomic carbon within the molten material, consequently increasing the likelihood of nucleation events for the formation of Al₄C₃ precipitates;
• AlSi10Mg is prone to chemical interaction with oxygen impurities, leading to the formation of an oxide film on the surface of the melt pool during the L-PBF process. Consequently, the presence of this oxide film can give rise to the occurrence of small Al₂O₃ precipitates within the aluminum matrix, attributable to the oxidation of the molten material. These precipitates have been found to contribute to the formation of elongated microcracks that propagate toward the surface of the printed component;
• Under sufficiently low cooling rates, the formation of Mg₂Si and Al₄C₃ phases can occur in the AlSi10Mg alloy. The presence of Mg₂Si arises from the silicon element’s supersaturation within the AlSi10Mg matrix, while the formation of Al₄C₃ is a result of a chemical reaction between the decomposed CNTs and the aluminum matrix. It is expected that Al₄C₃ will predominantly develop on the outer surface of CNTs, specifically at the interface between the nanotubes and the metallic matrix. This phenomenon is depicted in Figure 9, where CNTs are encapsulated by a thin layer of the Al₄C₃ phase. From a thermodynamic standpoint, the Al₄C₃ phase can exist within the temperature range of 600 to 1000 °C; however, if the temperature exceeds 1400 °C, it will decompose into elemental aluminum and carbon. Additionally, it should be noted that higher cooling rates in the melt pools, such as those on the order of 10⁴–10⁵ K/s, can diminish the likelihood of Al₄C₃ formation;

• Although some researchers have hypothesized the presence of the Al₉Si phase resulting from the supersaturation of the Al-Si solid solution and the segregation of silicon in the aluminum matrix, no empirical evidence confirming its existence has been reported to date. This is likely due to the low concentrations of Al₉Si in the parts produced by L-PBF. Further investigations are required to validate the presence of the Al₉Si phase and elucidate its formation mechanism in L-PBF-printed components.

In this context, a pivotal question arises concerning the thermochemical stability and likelihood of formation of two phases, Al₄C₃ and Mg₂Si, within CNT-reinforced AlSi10Mg matrix composites fabricated through L-PBF. The primary objective is to determine which of these phases offers superior thermochemical stability and is more probable to be observed in the resulting microstructures. To shed light on this matter, the free Gibbs energy of the respective phases (in kJ/mole) can be employed for analysis.

$${4\mathrm{Al}+3\mathrm{C}=\mathrm{Al}}_4{\mathrm{C}}_3;\begin{array}{cc}& \end{array}\Delta G_T^0=-\mathrm{216,689.64}+143.88T$$

(2)

$${2\mathrm{Mg}+\mathrm{Si}=\mathrm{Mg}}_2\mathrm{Si};\begin{array}{cc}& \end{array}\Delta G_T^0=-\mathrm{79,180.31}+16.34T$$

(3)

Equations (2) and (3) depict the quantitative relationship between Gibbs free energy and temperature range, elucidating the propensity for formation reactions associated with each phase. The Gibbs free energy-temperature correlation is illustrated in Figure 10. Notably, within the temperature range of 298–922 K, a linear relationship is observed between temperature and Gibbs free energy. A larger $\Delta G_T^0$ value corresponds to a reduced driving force for the formation reaction of the given phase and indicates higher thermal instability. As demonstrated in Figure 10, the $\Delta G_T^0$ value for Al₄C₃ is comparatively smaller than that for Mg₂Si, unequivocally suggesting that Al₄C₃ exhibits greater stability and is expected to form earlier than Mg₂Si within the composite microstructure [159].

3.2.5. Development of Defects

A comprehensive analysis of the existing literature reveals that, in the case of CNT-reinforced AlSi10Mg binary composites fabricated through L-PBF, a considerable number of pores and cracks tends to arise, exhibiting spherical or irregular morphologies. The spherical pores have dimensions smaller than 20 microns, whereas the irregular pores exhibit sizes larger than 50 microns. The literature suggests that as the CNT content increases, the number of pores, particularly spherical ones, also increases. This observation is consistent with the reported variations in relative density [158].

Moreover, it has been found that increasing the laser energy density (LED) during the L-PBF process can effectively mitigate the formation of irregular pores and cracks [2]. Several factors have been proposed to explain the occurrence of microstructural pores and defects during L-PBF in the presence of CNTs. These factors include:

• The presence of gases in the vicinity of the melt pool, which can dissolve into the solidifying metal;
• Agglomeration of CNTs and entrapment of gases between the particles;
• Superficial adsorption of gases on the high-specific-area CNTs during consolidation and L-PBF;
• Incomplete filling of gaps during the rapid solidification process;
• Vaporization and decomposition of CNTs due to the interaction with the laser, leading to the generation of gases [82,158,162].

It is important to note that the term “gas” referred to in the aforementioned factors encompasses hydrogen, argon, helium, nitrogen, methane, and traces of carbon monoxide. These gases can originate from various sources, such as the presence of moisture in the L-PBF chamber or the vaporization of CNTs during the printing process [158,162,166].

3.2.6. Effect of CNTs on Grain Structure

Experimental results consistently demonstrate that the addition of CNTs to AlSi10Mg alloy leads to a refinement of the grain structure in the fabricated components. This observation has been widely reported by various researchers in different contexts. For example, Carpenter et al. [167] observed a direct correlation between grain size and the volume percentage of CNTs in CNT–Nickel deposited coatings. Similarly, Liu et al. [168] reported a reduction in grain size from 0.3 to 0.2 microns upon increasing the CNT content from 1.5 to 3 volume percent in a CNT/Al-5Mg binary system. Hence, it is anticipated that the inclusion of CNTs in AlSi10Mg will result in a finer-grain structure. The underlying mechanism responsible for this grain refinement in the CNT-AlSi10Mg binary system involves the alignment of CNTs along the grain boundaries and their effective inhibition of grain growth in the metallic matrix through a pinning effect during the L-PBF process [159,160,162].

4. Physico-Mechanical Properties

4.1. Hardness

The microhardness of CNT-modified AlSi10Mg nanocomposites produced by L-PBF has been studied in the literature. Researchers have investigated this property from two main perspectives: (1) the influence of CNTs on microstructure development, and (2) the control of process parameters, including crystalline texture, grain configuration, and microstructural evolution. As an illustrative example, Figure 11 depicts the impact of varying laser power and scanning speed on the hardness of CNT/AlSi10Mg nanocomposites manufactured using the L-PBF method.

Comparatively, CNT-AlSi10Mg composites generally exhibit higher microhardness than the pristine alloy matrix when L-PBF parameters are optimized to avoid unfavorable microstructural phenomena. For example, Du et al. [162] reported a 33% increase in microhardness by incorporating CNTs into AlSi10Mg alloy and subsequently densifying the material through L-PBF. The microhardness of the pristine matrix improved from approximately 115 to 143 HV. It is worth noting that the laser energy density and scan speed are critical parameters that significantly influence the occurrence of favorable or unfavorable physico-chemical events.

To provide further insight into the possible outcomes, the following empirical results are reported:

• In the presence of CNTs, a chemical reaction occurs with the alloy matrix, resulting in the formation of a thin film of in-situ Al₄C₃ interfacial phase on the outer surfaces of the nanotubes. This thermodynamically stable film plays a crucial role in enhancing load transfer and improving the microhardness of the composite. However, it should be noted that, if the processing conditions, particularly the energy inputs, are set too high, CNTs may undergo thermal decomposition, leading to a decrease in microhardness [82,159];
• With an increase in laser energy density (e.g., from 89 to 131 J/mm³), the microhardness of the material initially improves (e.g., from 115 HV to 145 HV), but then starts to decrease (e.g., 130 HV). The initial improvement can be attributed to the grain refinement effect caused by the presence of nanotubes, which suppresses grain coarsening during the L-PBF process. However, the subsequent decrease in microhardness can be attributed to the change in microstructural features, specifically the transition from equiaxed to columnar grain orientation, which promotes grain growth and reduces the overall hardness of the material [75,169];
• The addition of CNTs to the pristine alloy matrix leads to improved microhardness by inhibiting the atomic diffusion of alloying elements and the subsequent grain coarsening process. Additionally, during the L-PBF process, there is an ample opportunity for the formation of silicon (Si) precipitates, which is facilitated by the L-PBF technique. This results in the presence of cellular equiaxed α-Al and nanometric fibrous eutectic Si in the aluminum (Al) matrix. The Si can exist in the Al matrix as a solid solution, contributing to solid solution strengthening. Therefore, the combined effects of grain refinement strengthening and solid solution strengthening can be considered as two possible mechanisms that enhance the microhardness of CNT-reinforced AlSi10Mg composites [159];
• The increase in laser energy input up to an optimized level can enhance the hardness of the material by promoting the formation of finer silicon (Si) blocky particles within the aluminum (Al) matrix. This is accompanied by the activation of the Orowan strengthening mechanism, which involves the hindrance of dislocation motion by the presence of these particles. Furthermore, the rapid melting and solidification process induced by the high energy input can generate internal stresses, contributing to increased hardness. However, if the energy input exceeds the optimal range, deviations from the desired microhardness can occur. This can be attributed to the formation of microstructural defects, degradation of CNTs, or dissolution of hardening precipitates in the material. These factors can lead to a reduction in the hardness of the composite [82,162,170].

4.2. Tensile Strength

The available data on the tensile properties of CNT-reinforced AlSi10Mg alloy matrices produced using L-PBF are limited and contradictory in nature. Different studies have reported conflicting results, leading to a lack of consensus. For example, Wang et al. [75] have observed that the incorporation of CNTs can improve the tensile strength of the material by 15% (from 356 MPa to 412 MPa). However, this enhancement in strength comes at the expense of reduced ductility, as evidenced by a 22% decrease in elongation at failure (from 5.5% to 4.3%).

In contrast, Thompson et al. [171] have reported a significant decrease in the tensile properties of CNT-reinforced AlSi10Mg composites. They found that the inclusion of CNTs led to a 13.3% reduction in yield strength, a 10.5% decrease in ultimate tensile strength, and a 4.9% decline in elastic modulus. It is important to note that, in the absence of CNT reinforcement, the unreinforced AlSi10Mg material exhibited gradual yielding and subsequent hardening before tensile failure. Several research studies have investigated the strengthening mechanisms in CNT-reinforced AlSi10Mg nanocomposites. It has been suggested that the contribution of dislocation strengthening in these materials is negligible or can be completely ignored. This assertion is supported by microscopic observations, which have revealed the absence of apparent dislocations in the matrix or around the nanotubes. In contrast, the Orowan strengthening and load-bearing strengthening are identified as the two dominant mechanisms activated. However, if CNTs used becomes sufficiently short, their contribution to the strengthening may be ignorable, because they will not be able to form a robust strain field around themselves and induce a great deal of pile-up dislocations. Similarly, if the volume content of nanotubes exceeds a critical level, the generation of pores in the microstructure can deteriorate the tensile mechanical properties [158,159].

Figure 12 provides a visual representation of crack fronts in relation to the melt pool structure. Notably, the melt pool exhibits a stretched configuration, displaying a high aspect ratio near the crack front [74]. Additionally, it allows for the observation of the presence of the CNT phase within the nanocomposites. Figure 12a,b illustrate how the boundaries of the melt pool and the elongated melt pool structure influence the deviation in the propagation path of the cracks. It was found that, when a crack is confined to the core regions of the melt pool, the cellular structure plays a significant role in resisting crack growth [172]. This finding underscores the intriguing relationship between subgrain size, cellular structure, and the mechanical properties of CNT-reinforced AlSi10Mg composites.

These contradictory findings highlight the complexity of the relationship between CNT reinforcement and the mechanical properties of the AlSi10Mg matrix. The investigation of CNT strengthening mechanisms is crucial for achieving optimal tensile properties in L-PBF-printed AlSi10Mg composites. By understanding the underlying mechanisms, researchers aim to enhance the material’s mechanical performance.

CNT Strengthening Mechanisms

Several strengthening mechanisms contribute to the improved performance of CNT-reinforced aluminum matrix composites. This section aims to elucidate these mechanisms and provide a rationale for the suitability of CNTs as strengthening additives in aluminum matrices. The literature extensively discusses various strengthening mechanisms that contribute to the enhanced properties of aluminum matrices reinforced with CNTs. These mechanisms include load transfer [173,174], grain refinement [173,175], solid solution strengthening [176], Orowan strengthening [177,178,179], and dislocation accumulation resulting from thermal mismatch [180,181]. The collective effect of these mechanisms determines the overall strengthening of the alloys or composites. Calculating the total strengthening requires considering the contribution of each mechanism [182,183]. Therefore, understanding the significance of CNTs in influencing the mechanical behavior of aluminum components requires a brief description of these prevalent strengthening mechanisms.

The load transfer mechanism serves as a direct means of strengthening in CNT-reinforced aluminum matrices. It involves the absorption and distribution of applied loads by the CNTs, thereby transferring the load to the aluminum matrix. This mechanism becomes active when larger CNTs generate a significant number of dislocations during plastic deformation, thereby increasing their relative density within the matrix. The resulting entanglement of dislocations contributes to a notable plastic constraint during deformation, ultimately strengthening the aluminum matrix [184,185,186,187]. The dominant role of the load transfer mechanism as a strengthening mechanism in CNT-reinforced aluminum matrix composites remains a subject of debate. According to a general consensus, the contribution of the load transfer mechanism is interactive and dependent on the aspect ratio of the CNTs. Specifically, a higher aspect ratio of CNTs is believed to enhance their effectiveness in strengthening the aluminum matrix through load transfer [183]. Contrary to the prevailing belief, certain studies have reported that the load transfer mechanism may not be the dominant strengthening mechanism in CNT-reinforced aluminum composites. For instance, Xie et al. [183] conducted quantitative analysis and fitting of experimental results, revealing that dispersion strengthening, rather than load transfer, plays a dominant role in enhancing the mechanical properties of these composites.

Another significant factor in the strengthening mechanism of CNT-reinforced aluminum alloys is grain refinement. The presence of CNTs acts as an effective barrier, impeding the normal growth of aluminum grains. As a result, the addition of CNTs to the matrix significantly reduces the grain size and enhances the mechanical properties [182,188]. This phenomenon can be attributed to the activation of the Hall–Petch effect, wherein the mechanical strength increases due to the restriction of dislocation movement by the crystalline misorientation at the grain boundaries [189,190].

The Orowan strengthening mechanism plays a crucial role in CNT-reinforced aluminum matrix composites, operating through the interaction between dislocations and CNTs. This mechanism impedes the motion of dislocations within the metal matrix. When a dislocation propagates through the matrix, it encounters CNTs embedded in its path. These nanotubes act as effective barriers, compelling the dislocations to bend around them. This bending induces a stress and strain field around the CNTs, thereby increasing the energy required for dislocation motion. Consequently, the strength and stiffness of the composite material are enhanced [181,191,192,193]. A clear indication of the Orowan mechanism in CNT-reinforced aluminum matrix composites is the presence of a significant population of dislocations within the aluminum lattice surrounding the CNTs [194].

The thermal mismatch strengthening mechanism observed in CNT-reinforced aluminum matrix nanocomposites can be attributed to the substantial disparity in their coefficients of thermal expansion (CTE). This phenomenon becomes evident when subjecting the composite material to heating or cooling cycles. During these thermal transitions, CNTs and the aluminum matrix experience disparate rates of expansion or contraction, resulting in the development of stress fields surrounding the CNTs within the composite. These induced stress fields play a pivotal role in enhancing the overall strength of the composite system. They achieve this by effectively impeding the movement of dislocations and increasing the energy required for the propagation of microcracks within the aluminum matrix [181,195,196,197]. Notably, the extent of the CTE disparity effect is directly proportional to the aspect ratio of the CNTs. However, it is noteworthy that the strengthening effect originating from thermal mismatch diminishes as the grain size of the aluminum matrix increases [180,181,198].

4.3. Coefficient of Thermal Expansion

The incorporation of CNTs into AlSi10Mg can effectively reduce the coefficient of thermal expansion (CTE) of the resulting nanocomposites, even at low CNT content. CNTs have a CTE close to zero, which contributes to the suppression of dimensional changes in the nanocomposites with temperature variation. Studies on CNT-reinforced AlSi10Mg nanocomposites produced through L-PBF have demonstrated that higher CNT content leads to lower dimensional changes, especially when the temperature exceeds a critical threshold, such as 250 °C. This indicates improved dimensional stability at elevated temperatures. It is worth noting that there exists a temperature range where the CTE of CNT-reinforced AlSi10Mg nanocomposites may significantly decrease. This phenomenon is attributed to the precipitation of Si atoms from the supersaturated Al lattice during heating [160,199].

4.4. Electrical Resistivity

The formation of a percolation network is a crucial factor in determining the electrical resistance of matrices containing conductive additives. CNTs play a vital role in facilitating the formation of percolation networks in both metallic and polymeric materials. Experimental findings have demonstrated that the incorporation of CNTs into AlSi10Mg matrices fabricated using L-PBF under various laser power and scan speed conditions can enhance their electrical conductivity (refer to Figure 13). For example, the electrical conductivity can be improved from 4.42 μΩ.cm to 0.11–0.52 μΩ.cm. This improvement indicates that CNTs have successfully formed a continuous percolation network within the AlSi10Mg matrix, which is attributed to their uniform dispersion and the robust L-PBF fabrication technology employed [82].

4.5. Wear

The incorporation of CNTs in L-PBF-fabricated AlSi10Mg alloy has a significant impact on the wear resistance of the material, primarily due to changes in surface hardness. It has been observed that CNT-reinforced AlSi10Mg nanocomposites exhibit higher coefficients of friction (COF) compared to the bare alloy, mainly attributed to the harder surface of the composites. Conversely, the lower COF values observed in the bare matrix can be attributed to surface softening, leading to the formation of deeper grooves and larger debris.

Moreover, the presence of CNTs can alter the wear mechanisms. In the absence of CNTs, the primary wear mechanism is pure plowing, whereas in CNT-reinforced AlSi10Mg nanocomposites, cracking and plowing become the dominant wear mechanisms [144]. Figure 14 illustrates the wear-induced grooves and edge cracks on the surfaces of both the worn AlSi10Mg alloy and CNT-AlSi10Mg nanocomposites. Examination of the scratching tracks and wear groove profiles reveals that the composite specimens exhibit smoother tracks compared to the unreinforced sample, indicating the formation of more threadlike debris and craters as the dominant wear mechanisms. Additionally, the worn surfaces of CNT-AlSi10Mg nanocomposites are characterized by shallower and narrower wear grooves, as well as a higher cross-sectional area of the wear groove (Figure 14).

Table 1. A concise review of the quantitative information reported by the literature on the properties and operational parameters for CNT-reinforced AlSi10Mg nanocomposites.

Chemical Composition	Reinforcement Content (wt%)	Dispersion Method	Densification Technology	Considerations	Operational Parameters	Measured Properties	Ref.
CNT-AlSi10Mg	1 wt%	Slurry ball milling and planetary ball milling	SLM	First, CNTs were added to the slurry and then ball milled by a planetary ball milling machine.	Ball-to-powder weight ratio: 10:1 Rotation speed: 100, 200 and 300 rpm Laser power: 240, 300, 360 W Laser λ = 1.064 mm Laser spot size: 0.08 mm scanning speed: 550, 650, and 750 mm/s	Reachable relative density: 90–97% Residual stress: 0–124 MPa Vickers Hardness: 123 HV (vs. 95–105 HV for pure alloy) Electrical resistivity: 0.11 to 0.52 μΩ cm vs. 4.42 μΩ cm for pure alloy	[82]
CNT-AlSi10Mg	1 wt%	Ultrasonication and stirring	SLM	A zig-zag scan strategy was used based on which the substrate should be rotated 90 degrees before starting the next layer.	scan speeds: 900, 1100, 1300, 1500, 1700 and 1900 mm/s. laser power: 370 W scan distance: 105 µm layer thickness: 30 µm	Roughness: 7–16 μm Maximum relative density: 98.53% Maximum hardness: 143.33 HV Reachable tensile strength: 499 MPa Maximum elongation: 7.6%	[135]
CNT-AlSi10Mg	0.5 wt%	Planetary ball milling	SLM	To avoid oxidation of the CNTs/AlSi10Mg mixture during SLM, substrate was pre-heated up to 150 °C.	ball-to-powder ratio: 1:1 rotation rate: 200 rpm total milling time: 4 h laser power: ~500 W Laser spot size: 90 μm continuous laser wavelength: 1080 nm scan speed: 9 m/s	Roughness: 2–5.5 Maximum roughness: 9.7 μm maximum hardness: 128 HV (vs. 126.99 HV for pure alloy) Maximum relative density: 99.3% Average yield strength: 380 MPa (vs. 329 MPa) Elongation: 7% (vs. 9%) wear rate about 33% lower than that of pure AlSi10Mg	[144]
CNT-AlSi10Mg	1 wt%	Planetary ball milling	SLM	The composite manufacturing was carried out in the argon atmosphere to prevent oxidation.	ball-to-powder ratio: 1:1 rotation rate: 200 rpm total milling time: 4 h maximum laser power: 400 W Laser λ = 1064 μm Laser spot size: 0.07 mm scan speeds: 400, 600, 800, 1000, and 1200 mm/s	Microhardness: 151.17 HV (vs. 120.15 HV for pure alloy) Average tensile strength: 498.6 MPa (vs. 439.2 MPa for pure alloy) Yield strength: 309.6 MPa vs. 270.7 MPa for pure alloy) Elongation: 10.6% (vs. 7.5% for pure alloy)	[159]
CNT-AlSi10Mg	0.01 wt% 0.05 wt% 0.1 wt% 0.5 wt% 1 wt% 2 wt% 5 wt%	Ultrasonication-assisted colloidal mixing	Directed energy deposition (DED)	AlSi10Mg powder particles benefit from good sphericity and smooth surfaces, with limited number of planetary particles adhered.	the laser power: 1300 W scanning speed: 10 mm/s initial temperature: 90 °C	Microhardness: 88.8–105.8 HV (vs. 87 HV for pure alloy)	[200]
MWCNT-AlSi10Mg	2 wt%	Combination of low-energy wet-milling and high-energy dry-milling	Spark Plasma Sintering (SPS)	The as-received AlSi10Mg/CNTs was added during the dry-milling stage.	ball to powder ratio: 20:1 rotation rate: 200 rpm total milling time: 120 min SPS temperature: 500 °C SPS duration: 10 min SPS pressure: 50 MPa	Compressive yield strength: 211 MPa (vs. 44 MPa for pure alloy) Ultimate strain: 31.5% (vs. 32.5% for pure alloy)	[201]
CNT-AlSi10Mg	1 wt%	Ball milling	SLM	The unique integrity and structure of CNTs are more likely to be damaged during SLM process due to high operational temperature and rapid solidification.	rotation rate: 100 rpm total milling time: 10 h Ball milling under argon gas powder layer thickness: 25 μm scan line hatch spacing: 130 μm laser power: 400 W volumetric laser energy densities: 89, 109, 131, 174 and 200 J/mm3	Density: 2.545–2.585 gr/cm3 Microhardness:121.4–143.7 HV Tensile strength: 412 MPa (vs. 356 MPa) Elongation: 4.3% (vs. 5.5%)	[75]
CNT-AlSi10Mg	unknown	friction stir processing (FSP)	SLM	FSP process was conducted on SLM-fabricated CNT-AlSi10Mg nanocomposite for more uniformilty.	Laser power: 350 W Scanning speed: 1040 mm/s Hatch spacing: 0.17 mm Layer thickness: 50 μm Energy density: 40 J/m3	Hardness after SLM: 98 HV Hardness after FSP and SLM: 115	[170]
CNT-AlSi10Mg	0.2–1.5 wt%	mechanical ball milling	SPS	Before ball milling, CNT were acid-washed in H2SO4	ball to powder ratio: 15:1 rotation rate: 200 rpm total milling time: 6, 7, 8, and 9 h SPS temperature: 540 °C SPS duration: 18 min SPS pressure: 40 MPa	UTS: 337 MPa (vs. 151 MPa) Yield strength: 241 MPa (vs. 82 MPa) Elongation: 1.9% (vs. 9.16%)	[202]
CNT-AlSi10Mg	0.5–2.5 wt%	High-energy ball milling	SPS	A varying thermal cycle was carried out to reach SPS temperature.	ball to powder ratio: 4:1 rotation rate: 600 r/min total milling time: 2 h SPS temperature: 350–550 °C Heating rate: 100–300 °C/min SPS duration: 5 min SPS pressure: 50 MPa	Reachable relative density: 99.7% Maximum hardness: 98 HV Compressive strength: 756 MPa Optimal Wear rate: $70\times {10}^3 \frac{{\mathrm{mm}}^3}{\mathrm{N}.\mathrm{m}}$ (vs. $120\times {10}^3 \frac{{\mathrm{mm}}^3}{\mathrm{N}.\mathrm{m}}$)	[203]

furnishes a succinct synthesis of the quantitative findings documented in recent scholarly literature. It is evident from the table that microhardness, wear rate, Ultimate Tensile Strength (UTS), and relative density stand as pivotal parameters of interest to researchers. Notably, the maximum values for these properties have been documented as 151 HV, $70\times {10}^3 \frac{{\mathrm{mm}}^3}{\mathrm{N}.\mathrm{m}}$, 756 MPa, and 99.7%, respectively. It is imperative to underscore that these values exhibit a substantial dependency on variables such as the content of CNT, the method employed for dispersion, and the specific operational parameters employed in the respective studies.

5. Challenges and Research Opportunities for CNT-AlSi10Mg Composites

The incorporation of CNTs into aluminum-silicon alloy (AlSi10Mg) matrices offers significant potential for enhancing the mechanical, thermal, and electrical properties of the resulting nanocomposites. However, the L-PBF process of CNT-AlSi10Mg composites poses several challenges that need to be addressed for successful implementation. This section aims to explore these challenges and identify potential research opportunities to overcome them. The following is a brief list of the challenges that will be addressed:

• Dispersion and Alignment: Attaining a homogeneous dispersion and preferentially aligned distribution of CNTs within the AlSi10Mg matrix holds paramount importance in achieving exceptional properties. Challenges arise due to the inherent tendency of CNTs to agglomerate, which leads to inadequate dispersion and inefficiency in load transfer. Overcoming this challenge necessitates the development of effective dispersion techniques and strategies to ensure a uniform distribution throughout the matrix. This issue becomes more complex when employing L-PBF as the densification technique, as the sphericity of the powder feedstock assumes significant importance in this method. Any deficiency in this geometric characteristic can result in the formation of porosity and degradation of mechanical properties;
• Interfacial Bonding: The presence of weak interfacial bonding between CNTs and the AlSi10Mg matrix hampers efficient load transfer between the reinforcement and the matrix, thereby compromising the overall mechanical performance. Enhancing the interfacial bonding is crucial to maximize the potential benefits of CNTs in AlSi10Mg composites. Strategies such as surface functionalization and interfacial engineering can be explored to improve interfacial interactions.
• Thermal Stability: The L-PBF process entails high-temperature conditions that have the potential to degrade CNTs, thereby impacting their structural integrity and properties. Preserving the thermal stability of nanotubes during the L-PBF process is crucial to retain their advantageous characteristics. Novel approaches, including the application of protective coatings or the use of hybrid reinforcement systems, can be investigated to enhance the thermal stability of CNTs.

It is evident that the primary avenues for prospective research involve addressing the previously delineated challenges. These include: (1) Investigating novel dispersion techniques aimed at mitigating agglomeration and achieving uniform dispersion of carbon nanotubes (CNTs) within the AlSi10Mg matrix; (2) Exploring diverse methodologies for enhancing the interfacial bonding between CNTs and the AlSi10Mg matrix; and (3) Investigating strategies to augment the thermal stability of CNTs within nanocomposites. Moreover, the authors suggest the following areas as potential directions for future research:

• Porosity and Defect Control: The authors propose the development of techniques aimed at minimizing porosity and defects in nanocomposites consisting of CNTs and an AlSi10Mg matrix densified via L-PBF. This can be achieved through the optimization of process parameters, manipulation of powder feedstock characteristics, or the implementation of post-processing treatments. By carefully fine-tuning process parameters and optimizing the properties of the powder feedstock, researchers can effectively reduce the occurrence of porosity and defects within the fabricated nanocomposites. Furthermore, the application of appropriate post-processing treatments can further enhance the overall mechanical properties of the materials. Through these advancements, the quality and performance of L-PBF-densified CNT-AlSi10Mg nanocomposites can be significantly improved.
• Multi-Scale Modeling: The authors propose the utilization of advanced modeling techniques, such as finite element analysis (FEA) or molecular dynamics (MD) simulations, to acquire a comprehensive understanding of the interactions occurring between CNTs and the AlSi10Mg matrix within nanocomposites. By employing these sophisticated computational tools, researchers can gain valuable insights into the mechanical behavior, stress transfer mechanisms, and failure modes occurring at different length scales. This enhanced understanding enables more precise design and optimization of CNT-reinforced AlSi10Mg composites. Through the application of FEA or MD simulations, researchers can evaluate and predict the performance of these nanocomposites, providing valuable guidance for future design and development efforts.
• Property–Performance Relationships: The authors propose an investigation into the correlation between the microstructure, processing parameters, and resulting properties of nanocomposites composed of CNTs and an AlSi10Mg matrix. This endeavor entails employing comprehensive characterization techniques, including mechanical testing, microstructural analysis, and measurements of electrical and thermal conductivity. By utilizing these techniques, researchers can establish clear relationships between the material composition, processing conditions, and the desired performance attributes. Through systematic analysis, it becomes possible to gain valuable insights into the influence of microstructural features and processing parameters on the mechanical, electrical, and thermal properties of CNT-AlSi10Mg nanocomposites. This knowledge is instrumental in guiding future material design and processing optimizations for enhanced performance.

By addressing these research opportunities, scientists and engineers can further advance the understanding and utilization of CNT-reinforced AlSi10Mg nanocomposites, leading to improved materials with enhanced properties and expanded applications in various industries.

6. Conclusions

In conclusion, this review provides a comprehensive examination of the L-PBF process applied to CNT-AlSi10Mg nanocomposites. The primary focus was on the formation of microstructure and the influence of processing parameters on a range of mechanical properties. The findings presented in this review highlight the considerable potential of L-PBF as a promising manufacturing technique for fabricating CNT-AlSi10Mg nanocomposites with improved properties. The main issues discussed throughout the paper are summarized as follows:

• By exercising meticulous control over laser power, scanning speed, and layer thickness during the L-PBF process, the microstructure of CNT-AlSi10Mg nanocomposites can be finely tuned. The addition of CNTs to the AlSi10Mg matrix brings about notable enhancements in various properties, including wear resistance, electrical and thermal conductivity, tensile strength, thermal expansion characteristics, and hardness. The incorporation of CNTs imparts reinforcing effects, thereby yielding superior mechanical performance when compared to the pure AlSi10Mg alloy.
• In light of recent research endeavors, a conspicuous dearth of data emerges concerning several critical aspects within the realm of CNT-reinforced AlSi10Mg nanocomposites. These encompass the establishment of an optimal CNT content necessary for the formation of an appropriate percolation network within the aluminum matrix, the quantitative evaluation of CNT agglomeration tendencies, the induction of residual stress within the matrix as a consequence of CNT integration, the potential necessity for supplementary heat treatments, and the possible occurrence of undesirable chemical reactions between CNTs and the metallic matrix leading to the consequential diminishment of physico-mechanical properties. These parameters and practical variables offer promising avenues for future research initiatives, warranting focused attention to unravel their intricacies and implications.
• The impact of processing parameters on both the microstructure and properties of CNT-AlSi10Mg nanocomposites was subjected to comprehensive investigation. The findings highlight the pivotal role of optimized processing conditions, particularly laser power and scanning speed, in attaining favorable microstructural characteristics and mechanical properties. Precise adjustment of these parameters enables control over grain size, porosity, and the distribution of CNTs, thereby exerting a significant influence on the overall performance of the nanocomposites.
• Significant attention is directed towards the characterization of physico-mechanical properties, encompassing wear resistance, electrical and thermal conductivities, tensile strength, thermal expansion, and hardness. These properties hold paramount importance across diverse applications such as aerospace, automotive, and electronics industries, where the demand for lightweight materials exhibiting exceptional mechanical and functional characteristics is particularly high. Accurate characterization of these properties enables researchers and engineers to assess the suitability and performance of CNT-AlSi10Mg nanocomposites for specific application requirements.
• Several challenges and research opportunities in the field of CNT-AlSi10Mg nanocomposites have been identified. These include the dispersion and alignment of CNTs, interfacial bonding between CNTs and the matrix, and ensuring the thermal stability of CNTs during the L-PBF process. Further investigation is required to address these challenges effectively. Subsequent research endeavors ought to concentrate on the elucidation of sophisticated dispersion methodologies, interfacial engineering tactics, and fortifying coatings in order to surmount these obstacles and actualize the complete capability of CNT-AlSi10Mg nanocomposites. By addressing these challenges and exploring the research opportunities, the full potential of CNT-AlSi10Mg nanocomposites can be realized, leading to advancements in their mechanical, thermal, and electrical properties.
• This paper illuminates the recent advancements and inherent challenges encountered in the L-PBF process of CNT-AlSi10Mg nanocomposites, facilitating a comprehensive comprehension of the mechanisms governing microstructure formation and mechanical properties. The discoveries put forth in this study hold substantial value as a valuable reference for researchers and engineers who aspire to refine the manufacturing methodologies and enhance the functional characteristics of CNT-AlSi10Mg nanocomposites across various application domains.
• A comprehensive review of quantitative findings from recent literature underscores the significant focus on microhardness, wear rate, ultimate tensile strength (UTS), and relative density as key research parameters. It is observed that the maximum values for these properties reach 151 HV, $70\times {10}^3 \frac{{\mathrm{mm}}^3}{\mathrm{N}.\mathrm{m}}$, 756 MPa, and 99.7%, respectively. Importantly, these values are highly contingent on factors such as CNT content, the dispersion technique employed, and operational parameters.