Multimaterial laser powder bed fusion: multifunctionality and application – part 2

Multifunctionality plays an important role for many metallic components. In the energy sector, for example, producing conformal cooling channels (CCCs) for molds and dies require efficient heat extraction under clamp loads (Kariminejad et al., 2024; Lamarche-Gagnon et al., 2024; Wagner and Nóbrega, 2025). Similarly, for compact recuperator heat exchangers for aerospace with microchannels, high heat transfer per unit volume is vital while withstanding pressure differentials and the structural loads of the parts (Careri et al., 2023; An et al., 2024; Li et al., 2025). Biomedical implants should have functionalities such as osseointegration and chemical resistance in addition to high mechanical strength ( S et al., 2024 ; El-Bassyouni, Mouneir and El-Shamy, 2025; Srinivasan et al., 2025). Multifunctional electronic circuits may require electrical conductivity within a structural matrix (Li et al., 2017; Tan and Low, 2018; Malikėnaitė et al., 2025).

Traditionally, these parts are produced with multistep milling and machining, injection molding, forging and assembly (Imgrund et al., 2007; Krimpenis and Iordanidis, 2023). To enable further complexity, additive manufacturing such as Powder Bed Fusion (PBF) can be used as part of a hybrid process for production (Wang et al., 2023). Recently, by developing multi powder deposition systems within some laser powder bed fusion (LPBF) machines, multimaterial LPBF enables production of multifunctional parts using properties of different metallic materials and alloys (Mehrpouya et al., 2022; Meyer et al., 2023; Schmidt, Jensch and Härtel, 2023; Griffis et al., 2025). Incorporating the multimaterial concept into high resolution LPBF systems, compared to, for example, directed energy deposition (DED) systems, enables manufacture of complex multimaterial, multifunctional structural designs using methods like topology optimization and generative designs. The comparative benefits of PBF over other multimaterial technologies (DED and Binder Jetting) are detailed in Table 1 of Part 1 Moutablaleh et al., (2025).

Few reviews are published on multimaterial LPBF techniques (MM-LPBF) and their interface properties so far, including the Part 1 of this review (Moutablaleh et al., 2025; Mehrpouya et al., 2022; Wang et al., 2023; Guan and Wang, 2023; Yao et al., 2023; Kavousi Sisi et al., 2024; Zainelabdeen et al., 2024; Griffis et al., 2025). Although these reviews are informative, they lack the end-user perspective on the applications. This review paper studies the interface properties of the produced metallic pairs based on the added functionalities. Therefore, multifunctionality is the focus of this paper. It also includes previous knowledge on the same pairs but different manufacturing processes in the discussion. Each alloying system is discussed based on five functionalities of (1) strength and ductility, (2) wear resistance, (3) thermal and electrical conductivity, (4) corrosion resistance and (5) shape memory and superelasticity, which can be added locally to single material LPBF. The discussion for the pairs is based on the end-user application perspective to help designers select proper materials for their particular needs and functional requirements.

This manuscript represents the second part of a comprehensive review of multimaterial-LPBF (MM-LPBF). While Part 1 established the fundamental feasibility and metallurgical integrity of these structures, Part 2 transitions to an applied, end-user perspective. The initial study categorized current powder deposition systems, including 2D layer-on-layer, 3D suction-based and ultrasonic vibrating nozzle techniques, while evaluating their respective limitations regarding powder contamination and resolution. A critical finding in Part 1 was that multimaterial bonding success is primarily dictated by the “mismatch” in intrinsic material properties, such as melting points, thermal conductivity and coefficients of thermal expansion (CTE). These disparities often lead to nonuniform heat flow and the formation of brittle intermetallic phases (IMCs) at the diffusion zone. (Moutablaleh et al., 2025)

Building upon this foundational understanding of interface physics and the mitigation frameworks (such as CALPHAD simulations and compatible interlayers) proposed in Part 1, this paper (Part 2) categorizes the existing literature based on functional requirements. By focusing on five key functionalities of strength/ductility, wear resistance, conductivity, corrosion resistance and shape memory, this work aims to guide designers in selecting proper material pairs and processing parameters for specific industrial applications.

The literature for this two-part review was gathered through a systematic and structured approach using the Scopus and Web of Science databases. The search strategy was divided into two primary sections to capture the technological, metallurgical and functional aspects of the field. Section 1 targeted Multimaterial Metal Additive Manufacturing, using search terms of “multi material metal AM/3D printing,” “bi-metal AM,” “metallic pair AM,” “multi material LPBF,” and “metallic pair LPBF”. Section 2 focused on single material LPBF Foundations, subdivided into metallurgical aspects (e.g. “thermal expansion LPBF,” “melt pool LPBF” and “intermetallic LPBF”) and functional applications (e.g. “wear resistance LPBF,” “corrosion resistance LPBF” and “shape memory LPBF”). To bridge the gap in specific multimaterial knowledge, we also evaluated literature from dissimilar metal joining and DED to identify overlapping metallurgical solutions. Following the initial screening to remove duplicate records and irrelevant topics, the selection process resulted in 37 primary papers that specifically described metallic pairs produced by LPBF. The systematic flow of this identification, screening and categorization process is illustrated in Figure 1. These 37 papers form the core of the discussion in Part 2, where they are categorized based on the specific industrial functionalities they address.

It is important to have a clear objective for transitioning from single to multimaterial, as there are additional challenges specific to MM-LPBF. The key motivation is the addition of one or more functionalities that cannot be achieved using a single material. Figure 2 shows an overview of various multimaterial pairs produced by LPBF in the literature, including ferrous, nickel, copper, aluminum and titanium alloys (Liu et al., 2014; Anstaett, 2017; Demir and Previtali, 2017; Wei et al., 2018, 2019; Nadimpalli et al., 2019; Wei et al., 2021; Borisov et al., 2021; Duval-Chaneac et al., 2021; Lesko et al., 2021; Mohd Yusuf et al., 2021; Schneck et al., 2021; Wits and Amsterdam, 2021; Wei et al., 2022; Bareth et al., 2022; Chen et al., 2022; Cunha et al., 2022; Demir et al., 2022; Marques et al., 2022, p. 718; Walker et al., 2022; Bartolomeu et al., 2023). The addition of a second material, similar or dissimilar to the first material, adds functionalities that are not available solely for the first material. Figure 3 summarizes the functionalities reported in the literature. In this study, pairs were categorized according to their added functionality to the first material.

Stainless steel to copper alloys: Certain single materials require high strength and/or ductility for their applications. For example, copper alloys have excellent electrical and thermal conductivity and corrosion resistance but lack strength in low- and high-temperature applications in heat exchangers, electrical components, cryogenic and heat sinks in fusion reactors and rocket engine combustion chambers (Li et al., 2020; Kavousi Sisi et al., 2024). In regions of high local stress, the functionality of copper alloys can be reinforced using stainless steel. In multimaterial DED, large miscibility gaps and lack of intersolubility are reported to cause the two to segregate into discrete droplets rather than form a continuous solid solution. This leads to compositionally inhomogeneous fusion zones (Magnabosco et al., 2006; Chen et al., 2013; Reichardt et al., 2021). Challenges such as intermetallic formation, thermal expansion coefficient differences, thermal conductivity mismatch and melting temperature mismatch have been reported for austenitic stainless steel and copper alloy LPBF pairs (Hu et al., 2022). Unwanted liquid metal embrittlement (LME) cracking can also be observed at iron and copper interfaces. As a case study, 316 L austenitic stainless steel was processed using Cu10Sn (Figure 4a) (Wei et al., 2019; Chen et al., 2022). Despite the sound metallurgical bonding with the formation of iron-rich spherical particles and fine copper-rich particles on the CuSn10 side, and large diffusion zone of 550 μm, brittle intermetallic phases such as Cu8Ni and Cu9NiSn3 are developed. This causes deviations in hardness in the diffusion zone. The higher thermal conductivity of copper alloys promotes epitaxial grain growth across the melt pool boundary at the diffusion zone, resulting in a high ultimate strength of 420 MPa, which is better than that of conventional approaches for processing steel and copper multimaterial alloys. Therefore, by applying better process parameters, it is possible to achieve superior diffusion at stainless steel and copper alloy interfaces. To overcome intermetallic formation, a nickel interlayer has been reported to be compatible with steel and copper with very high solubility with iron and copper in multimaterial DED (Ashley Reichardt and Beese, 2021). This might be the case for MM-LPBF helping to eliminate the formation of detrimental intermetallic phases, for example, by including more nickel concentration in the composition of alloys involved.

Nickel to copper alloys: In rocket engine combustion chambers and fusion reactors, where copper-based components are common, local high-temperature strength can be supplied by incorporating nickel superalloys such as In718, or by nickel-iron alloys such as Invar36 to withstand local stresses (Figure 4b). For MM-LPBF of copper and nickel alloys, differences in their melting points and thermal conductivity are the main challenge, since the risk of intermetallic formation in nickel-copper systems is low (Campbell et al., 2002). Owing to their different melting points, unmelted nickel particles might blend with melted copper, creating residual stresses and a lack of fusion. Differences in the thermal conductivity and laser absorptivity between the two alloys can lead to incomplete melting, with un-melted Invar36 particles remaining owing to their higher thermal conductivity (Wei et al., 2021). These problems stem from unoptimized process parameters for nickel and copper alloys and can be mitigated by proper process parameter design.

Nickel alloys to stainless steel: To add a localized high-temperature strength to stainless steel exhaust components, pressure vessels, steam and gas turbine blades, etc. nickel superalloys such as In718 or HastelloyX can be integrated (Figure 4c) (Wei et al., 2018; Duval-Chaneac et al., 2021; Mohd Yusuf et al., 2021; Wits and Amsterdam, 2021). These alloys are frequently joined via fusion welding; however, they are highly susceptible to solidification cracking. Cracks are reduced by refining the grain boundaries, since alloying elements and impurities segregate into interdendritic regions and causes low-melting liquid films to exist along the grain boundaries (Robinson and Scott, 1980). Laser processing in additive manufacturing results in higher cooling rates and finer microstructures and is an advantage for this refining. The main challenges for MM-LPBF of nickel and austenitic stainless steel systems are brittle intermetallic formation, carbon migration and carbide precipitation. However, their thermophysical properties match well and have good solubility (Kavousi Sisi et al., 2024). For multimaterial DED, Lin et al., (2005) reported the precipitation of the γ’ phase with increasing nickel content, gradually increasing the hardness. In another study on multimaterial DED, 304 L stainless steel is DED processed with In625 nickel alloy as a gradient, and it is reported that Nb and Mo carbides presented in high concentrations at a crack at the locations with an approximately 82 wt% of 304 L alloy (Carroll et al., 2016). These monocarbides were predicted by thermodynamic equilibrium phase fractions along the compositional gradients. Because MM-LPBF has high cooling rates, the formation of these carbides, such as NbC, can be limited; however, further studies are needed on their nonequilibrium phase calculations. Because they have similar thermal expansion coefficients, and their existing elements of Fe, Ni and Cr, have good solubility, they generally form large diffusion zones with few defects. For 316 L and In718 pairs produced by MM-LPBF, only a few porosities and cracks on the 316 L side, and some lack of fusion on the In718 side have been reported. In addition to the optimized process parameters, post-processing techniques such as high-pressure torsion (HPT) have been reported to help reduce these porosities and cracks further (Sahu et al., 2024). However, the risk of Laves intermetallic phases forming in the diffusion zone or on the In718 side is high. For 316 L and HastelloyX pairs produced by MM-LPBF, HastelloyX nickel alloy shows susceptibility to hot cracking in its columnar grains owing to the micro-segregation of Mn and Si that results in the formation of ultra-fine dendritic arms of less than 1 µm in size. These fine precipitates lead to crack nucleation and propagation as the material solidifies (Rankouhi et al., 2022). In addition to micro-segregation, the 316 L tends to grow more equiaxed grains, while the nickel alloy grows more columnar grains. This results in intergranular hot cracking along the columnar grain boundaries. If the formation of brittle intermetallic phases and carbides is controlled, these pairs can be directly bonded to each other owing to their similar thermophysical properties.

Martensitic stainless steel to Austenitic stainless steel: In the power generation industry, such as in ultra-supercritical (USC) power plants, precipitation-hardened martensitic stainless steel is welded to the surface of austenitic stainless steels to improve the local strength and creep performance (Dak and Pandey, 2020). The only problem with austenitic and martensitic stainless steels produced by MM-LPBF is their differences in thermal expansion coefficients (Kavousi Sisi et al., 2024). Examples include the LPBF processing of MS1 martensitic stainless steel with 316 L (Figure 4d) (Nadimpalli et al., 2019) and 15-5PH martensitic stainless steel with 316 L (Liang et al., 2023). The diffusion zones for these multimaterial pairs are relatively large (approximately 100 μm) due to their similar elemental compositions, leading to a higher diffusivity. Some lack of fusion has been reported on the 15-5PH side, which can be mitigated by optimizing the relative energy density (Liang et al., 2023). Therefore, this pair can be directly bonded or with a gradient to mitigate thermal expansion coefficient differences. Studies on other ferritic alloys, such as low-carbon ferritic steels and stainless steels, have been performed in welding and DED processes (Ashley Reichardt and Beese, 2021) but not in MM-LPBF yet.

Cobalt alloys to stainless steel: In tooling and mold-making industries, for example, dies and cutting tools, high-temperature strength is critical. Cobalt alloys, known for their high thermal stability and superior toughness, can be combined with high-strength maraging or precipitation-hardened steels, such as MS1 or 17-4PH (Figure 4e), to provide superior performance (Steponavičiūtė et al., 2022). The advantage of combining cobalt alloys with stainless steel for high-temperature applications is their matching thermal-physical properties, such as melting point, thermal conductivity, thermal expansion coefficients and negligible presence of secondary phases (Kavousi Sisi et al., 2024). Therefore, this combination resulted in a good metallurgical bond without visible pores or cracks at the interface. Some mixed structures at the interface underwent martensitic transformation from the FCC columnar grains, leading to solute segregation. Some untransformed FCC islands were observed, likely due to the growth orientation before solidification, local stresses from martensitic volume expansion, uneven distribution of alloying elements and the negligible presence of secondary phases. A homogeneous diffusion zone between 180 μm and 200 μm was observed with no detrimental brittle phases. Pasco et al., (2023) applied heat treatment to CoCrMo-MS1 alloys to improve their ductility, relieve residual stress and enhance their structural homogeneity. The results showed successful precipitation and stress relief, with the tensile strength maintained at approximately 1185 MPa before and after treatment. The heat treatment significantly increased the elongation from approximately 14.7% to 27.1%, resulting in a tougher multimaterial structure. In conclusion, this pair can be directly bonded with LPBF with optimized process parameters because of their similar thermophysical properties and lack of detrimental brittle phase formation with their combination.

Cobalt alloys to titanium alloys: For medical applications, cobalt alloys, such as CoCrMo, can be added to titanium alloys, such as Ti6Al4V, for superior local strength. The reported challenges regarding their MM-LPBF processing are the formation of brittle intermetallic phases and thermal expansion coefficient mismatches (Kavousi Sisi et al., 2024). In one study, the CoCrMo-Ti6Al4V pair processed by LPBF exhibited a small diffusion zone of approximately 5 μm with no reported defects at the interface (Figure 4f) (Bartolomeu et al., 2023). This relatively small diffusion zone is similar to the combination of cobalt alloys and stainless steel and can be attributed to the higher mass density of elements in the CoCrMo alloy, leading to more significant differences in dynamic viscosity between the two alloys. This can result in higher differences in the diffusion coefficients. In addition, a lower energy density during the LPBF process could contribute to a reduced diffusion zone. Schober and Peng, (2016) Interestingly, no typical intermetallic phases, such as Ti₂Co, which are usually found in composites of these two alloys, were observed (Bartolomeu et al., 2023). The gradual increase in hardness also indicated the absence of these intermetallic phases. Therefore, rapid solidification in MM-LPBF mitigates the formation of most known brittle intermetallic phases. Therefore, this pair can be bonded with a gradient with optimized process parameters to promote further diffusion at the interface.

Aluminum alloys to Titanium and Copper alloys: In some industries such as aerospace, lightweight components are crucial. Aluminum alloys, with a density of about 2.7 g/cm³, are significantly lighter than other alloys like titanium (4.42 g/cm³), steels (8.05 g/cm³), copper (8.94 g/cm³), nickel (8.9 g/cm³) and cobalt (8.9 g/cm³). This makes aluminum alloys ideal for creating lightweight products. Wu et al., (2022) examined the interfacial characteristics of a AlSi10Mg and Ti6Al4V aluminum-titanium pair produced by MM-LPBF. Brittle intermetallic formation, melting temperature mismatch and thermal expansion coefficient mismatch have been reported (Kavousi Sisi et al., 2024). They found that the transition from the titanium side to the aluminum side begins with the formation of a Ti₂Al layer, followed by TiAl with Ti₅Si₃ nanoparticles, eventually forming rod-like TiAl₃ toward the AlSi10Mg, along with an opposite heat dissipation direction (Figure 5 a-c). As a result, the tensile test showed brittle fracture at the interface, but an ultimate strength of approximately 265 MPa was reported, which is higher than that of traditional aluminum and titanium multimaterials. To solve the problem of interface brittleness (Wu et al., 2022; Zhang et al., 2023) introduced a 1 mm copper interlayer based on thermodynamic calculations and the eutectoid-forming potential in titanium-copper and aluminum-copper binary alloy systems. The copper interlayer prevented the formation of titanium-aluminum intermetallics and suppressed interface cracking, yielding a crack-free interface with α-Al → Al₂Cu → AlCu₂Ti → CuTi₂ → α′-Ti. This structure has lower hardness and higher wettability, contributing to a well-bonded interface. Therefore, introducing a compatible interlayer, such as copper, at the interface of aluminum and titanium alloys is necessary to eliminate the formation of various titanium-aluminum brittle phases.

Sing et al., (2015) reported on AlSi10Mg and C18400 copper alloy interfaces produced by MM-LPBF. Interestingly, Al₂Cu is formed at the interface with low hardness values, showing higher strength and elongation than both base metals, with fractures occurring on the copper side. This indicates excellent bonding between these two alloys. In another study, Demir and Previtali, (2017) processed AlSi12 alloy with pure Fe as a gradient, but reported large cracks and higher hardness at the interface, indicating the formation of brittle intermetallic FeAl phase. Aluminum alloys can be MM-LPBF processed with copper alloys very well. However, they are highly susceptible to the formation of intermetallic phases in titanium alloys and iron. Therefore, a copper interlayer is necessary to mitigate the formation of intermetallic phases.

Techniques such as Chemical Vapor Deposition (CVD), thermal spraying and Physical Vapor Deposition (PVD) are constantly applied to metallic parts to improve their surface properties. Also, extensive research has been conducted on cladding, particularly laser cladding (LC) (Baragetti et al., 2009; An et al., 2020; Xiang et al., 2023). Because of higher stress concentrations created by surface defects, they can nucleate and grow more easily than bulk defects (Rokhlin and Kim, 2003; Li and Zhang, 2011). Surface coatings with wear-resistant materials can mitigate this problem by reducing the impact of surface defects and preventing fatigue and tensile failure, thereby extending the product lifespan. However, there is a limitation in the thickness of the surface-coating techniques. The current coating technologies typically offer only a few hundred micrometers (Prengel, Pfouts and Santhanam, 1998). This can be inadequate for applications that require thicker coatings for greater surface protection and longer lifespan. Therefore, MM-LPBF is an advantage over other coating and cladding techniques because it can offer a higher thickness. Another advantage of the MM-LPBF is its ability to print complex geometries. Other chemical and physical surface coating methods may not be suitable for very complex geometries such as lattice structures. Therefore, MM-LPBF can provide thicker, solid and wear-resistant surface materials without geometrical limitations.

Martensitic stainless steel with austenitic stainless steel: Martensitic stainless steels such as 15-5PH and MS1 have higher wear resistance than austenitic stainless steels owing to the presence of hard martensite phases. This is advantageous for the surface protection of austenitic stainless steels (Nadimpalli et al., 2019; Liang et al., 2023). For MM-LPBF of 316 L and 15-5PH, some lack of fusion was reported on the 15-5PH side (Figure 6a). A strong {101} texture with larger columnar grains was observed on the 316 L side, whereas a lighter {111} texture with smaller equiaxed grains on the 15-5PH side, consistent with their single LPBF processing. No visible cracks or defects are observed. Similarly, processing austenitic SS420 with martensitic MS300 resulted in a higher tensile strength at the interface than the austenitic steel alone, indicating the added strength from the martensitic MS300 (Tan et al., 2020). Chromium-rich particles in SS420 acted as homogeneous nucleators for MS300 grains, harmonizing the mismatch between the two materials. The epitaxial columnar structure continued along the maximum temperature gradient, with some martensite near the interface, thus enhancing the bonding quality (Figure 6b). Therefore, strong interfaces can be created with optimized process parameters.

Cobalt to martensitic stainless steel: Cobalt alloys are known for their excellent wear-resistant properties, and they have melting points, densities and thermal coefficients similar to those of stainless steel (Antony, 1983). When MS1 martensitic stainless steel was processed with Co-37Cr-7Mo and MM-LPBF (Figure 6c) (Pasco et al., 2023), a strong bond was achieved without keyholes or lack of fusion, with elongated columnar grains on the MS1 side, and finer grains on the cobalt side, creating a robust interface. Successful precipitation strengthening and stress relief after heat treatment have been reported to double the elongation value at the interface without reducing the strength. A similar positive outcome was reported for the CoCrMo and 17-4PH pair, which demonstrated good bonding and acceptable elemental diffusion (Figure 6d) (Steponavičiūtė et al., 2022).

Cobalt alloys to titanium alloys: CoCrMo has a wear rate of around 10E⁻⁶mm³/N/m, about 65 times more resistant than Ti6Al4V, with a wear rate of 6.5E⁻⁴mm³/N/m. This makes cobalt-titanium alloys ideal for applications where wear resistance is critical, such as bone implants, knee replacements and hip joints. Bartolomeu et al., (2023) MM-LPBF-processed CoCrMo with Ti6Al4V for hip implants and achieved enhanced wear resistance without forming typical Ti₂Co intermetallic compounds with a diffusion zone of 5 µm (Figure 6e).

In some environments with polluted air, acidic conditions, high salinity, and high temperatures, metals corrode and rapidly fail at a lower lifetime than in normal conditions (Shi, Yang and Liaw, 2017). A conventional technique to prevent corrosion is to add small concentrations of elements such as Cr, Ni and Mo, in which they form protective passive films on the surface and prevent further corrosion of the underlying alloy (Shi, Yang and Liaw, 2017). Another method involves applying protective coatings; however, thickness limitations still exist. Techniques that provide thicker coatings for complex geometries can expand the lifetimes of parts in harsh environments. Thus, MM-LPBF is a good candidate for this purpose.

Pure Ti to titanium alloys: In a living body, corrosion can release harmful substances during biodegradation (Arjunan et al., 2020). ASTM standards F3001-14 and 2923–14 standardized Ti6Al4V for additive manufacturing, but this class of titanium alloys contains vanadium, a toxic impurity that can be released into the body during corrosion (Zhang et al., 2018). Producing entirely pure titanium parts is not feasible for medical implants because of ductility and insufficient strength. To solve both issues, pure titanium can be integrated into Ti6Al4V for greater biocompatibility, where it is locally necessary. Although these alloys belong to the same system, pure titanium has fine martensitic α′ needles, whereas Ti6Al4V has columnar primary β grains with needlelikeneedlelike martensitic α′ phases (Figure 7a). This significantly different microstructure creates a challenging interface (Borisov et al., 2021). Despite not observing cracks or lack of fusion at the interface, CT scans revealed 50 μm residual pores in different areas. These were eliminated using hot isostatic pressing (HIP). After HIP, the microstructure recrystallized with equiaxed α grains in pure Ti and α + β grains in Ti6Al4V. Tensile test failure occurred at the Ti-Ti6Al4V interface, indicating somewhat brittle behavior, presenting challenges for this pair, requiring further study.

Nickel alloys to stainless steel: In highly corrosive environments, such as salt reactors, nickel-based superalloys, such as Inconel HX, can be added to common stainless steel 316 L for high-temperature corrosion resistance (Rankouhi et al., 2022). For welding and DED of nickel-iron dissimilar systems, the migration of various species across the joint interface, the formation of martensite and the precipitation of unwanted carbides have been reported. The strength and ductility analysis indicates that the nickel-iron multimaterial systems are compatible in terms of thermophysical properties and solubility. For MM-LPBF, the HX nickel alloy and 316 L stainless steel pairs exhibited a sound, graded interface with no visible pores or cracks (Rankouhi et al., 2022). The HX side exhibited perpendicular grain boundaries and a cellular/dendritic solidification microstructure with visible Marangoni flow vortices in the diffusion zone. Some submicron-sized defects were identified at the diffusion zone and on the HX side caused by inclusions such as carbides pulled out during sample preparation, and some gas pores at the grain boundaries and within the grains (Figure 7b). No significant defects are observed on the 316 L side. The diffusion zone was more HX-dominant, as indicated by the same inclusion defects and primary dendritic arm spacing (PDAS) size in the diffusion zone. SEM and XRD analyses revealed no detrimental intermetallic compounds at the grain boundaries. Interestingly, the diffusion zone remained consistent at 240 μm regardless of whether HX was on the top or bottom, indicating that gravity and buoyancy forces did not significantly affect melt-pool dynamics. However, EDS observations suggested that the secondary material on top primarily controlled the composition of the diffusion zone. The tensile test revealed that fractures occurred on the softer 316 L side, indicating good mechanical properties at the interface and diffusion zone. Similar corrosion resistance results were obtained for In718 and 316 L corrosion resistance. Figure 7(c) shows the grain orientation map of the pair, with mostly columnar texture in the < 001> direction for both materials, as they share the same face-centered cubic (FCC) crystallographic structure when the building orientation z-axis is aligned with favorable grain growth for this structure (Duval-Chaneac et al., 2021). Therefore, printing thick nickel alloys over complex stainless-steel components can be achieved with optimized VED for both sides to avoid a lack of fusion and to prevent columnar grain growth (Cruz et al., 2023).

Molds, electrical components, cooling chambers and other energy applications often require high electrical and thermal conductivity (Kavousi Sisi et al., 2024). Table 1 shows the electrical and thermal conductivity advantages of silver and copper compared to other metals. Copper’s relatively low strength can however be a disadvantage, leading to its use with stronger alloys (Seo et al., 2012; Ramkumar et al., 2018). MM-LPBF can produce complex copper bimetallic structures that meet both strength and conductivity requirements in electrical panels or cooling components.

Copper alloys to stainless steel: For superior heat transfer properties in manufacturing tools, such as dies and molds, copper cooling channels can be incorporated into stainless steel molds using AM technologies (Li, Syed and Pinkerton, 2006; Vaezi et al., 2013). Reduced design limitations can help these copper cooling channels to have superior heat transfer properties. This can be achieved MM-LPBF processing of steels and copper alloys. The main challenges with iron-copper interfaces are the high copper reflectivity, poor iron and copper solubility, large melting pools, mismatched thermal conductivities and melting temperatures (Ashley Reichardt and Beese, 2021). A gradient of H13 tool steel and pure copper produced using a partitioned powder hopper resulted in porosity and cracks (Beal et al., 2004). However, brittle intermetallic phases have not been reported for iron-copper pairs. In another study, the 316 L stainless steel and CuSn10 copper alloy pair encountered lack of fusion and LME (Chen et al., 2022; Schanz et al., 2022; Wei et al., 2022). Despite the presence of some microcracks on the 316 L side, elemental diffusion at the interface with no brittle intermetallic phase formation has been reported. Spherical iron-rich particles were observed on the CuSn10 side, with copper-rich particles within these spheres. This multimaterial pair achieved a tensile strength of 424 MPa, which was higher than that of conventional steel and copper combinations. Therefore, an optimized process parameter can overcome microcracks and lack of fusion. Similarly, Liu et al., (2014) reported a diffusion zone of 750 μm for a 316 L and C18400 pair with tensile failure at the softer copper side, indicating strong bonding at the interface. Similarly, porosity on the copper side and dispersed cracks near the interface on the 316 L side have been reported (Figure 8a). For the tool steel 1.2709 and copper alloy 2.1293 pair (Anstaett, 2017), fewer cracks were observed when the copper alloy was placed on top because of its ductility and ability to absorb internal stresses (Figure 8b). This trend was confirmed by the same author for tool steel 1.2709 and copper alloy CW106C pair, where cracks appeared when the steel was on top, indicating a better gradient structure when copper was on top. Schneck et al., (2021) manufactured four tool steel 1.2709 injection nozzles with CW106C copper alloy inlets using LPBF. The results showed delamination, cracks, porosity and cross-contamination, suggesting that a better scanning strategy in the transition zone is required for these pairs. Similarly, 316 L and Houvador copper alloy (Fu, Yeong and Chen, 2016) exhibited a low-quality interface with cracks and pores even at high energy densities. Cunha et al., (2022) reported the successful processing of stainless steel SS420 with pure copper, achieving good bonding with minimal pores and cracks in the diffusion zone. In conclusion, steel and copper MM-LPBF parameter optimization can be challenging owing to their differences in thermal conductivity, melting point and laser absorptivity. However, in the positive side, brittle intermetallic phases have not been reported for MM-LPBF interfaces.

Copper alloys to nickel alloys: Aerospace applications where nickel alloys are constantly used, owing to their excellent high-temperature strength, require local electrical and thermal conductivities higher than those of nickel alloys, such as copper alloys. The main reported challenge for nickel-copper interfaces is the mismatch in their thermal properties. Nickel-copper dissimilar joints produced by laser melt deposition (LMD) have been reported to exhibit interfacial solid strength (Razzaq et al., 2024). The In718 nickel alloy and pure copper pair produced by MM-LPBF exhibited some lack of fusion and gas-entrapped pores (Figure 8c) (Marques et al., 2022). Therefore, the nickel-copper interface can be processed relatively easily with optimized process parameters. Iron-nickel Invar36 and CuSn10, however, exhibited LME cracks and unmelted nickel particles, even when applying a higher VED, possibly because of the significant differences in their thermal conductivity, melting points and laser absorptivity (Wei et al., 2021) (see Figure 8d).

Copper alloys to titanium alloys: Local electrical and thermal conductivities can also be integrated into titanium alloys. Although no brittle interface formation has been reported, the melting temperature and thermal expansion coefficient mismatch is an issue for the titanium-copper multimaterial system (Kavousi Sisi et al., 2024). Wei et al. produced a Ti6Al4V and CuSn10 pairs by MM-LPBF in three forms of direct bonding, remelting and gradient interfaces (Wei et al., 2022) (Figure 8e). Complex titanium-copper intermetallic compounds at the interface caused micropores, cracks and delamination in all three cases. These observations suggest that despite high elemental diffusivity of copper alloys, significant differences in their thermal properties with other alloying systems create challenges in achieving defect-free components. Therefore, more precise process optimization is required.

Silver to copper alloys: Sometimes, even higher local thermal and electrical conductivities than copper are needed in the energy industry. The Ag7.5Cu silver alloy has been MM-LPBF with CuSn10 for high thermal conductivity applications (HETCs) (Chen et al., 2023) (Figure 8f). A good Marangoni flow was observed at the CuSn10 and Ag7.5Cu interface when CuSn10 was on top, but unstable keyhole structures occurred when CuSn10 was at the bottom. Therefore, replacing the substrate with a substrate with a higher thermal conductivity can enhance Marangoni convection, strengthen interfacial bonding and reduce defects (Chen et al., 2023).

Shape memory alloys (SMAs) have applications in engineering and medical industries, where shape memory and superelasticity properties are required (Mehrpouya et al., 2024). For example, the extrinsic two-way effect can be used in bearings and gears to adjust the clearance in tapered roller bearings (Predki, Knopik and Bauer, 2008). In medicine, these alloys play a key role in minimally invasive surgery and diagnostic applications, stents, filters and catheter tubes (Es-Souni, Es-Souni and Fischer-Brandies, 2005). By integrating SMAs with other properties such as high strength, wear resistance, corrosion resistance, fatigue resistance and high thermal conductivity, their applications can be broadened. The EBSD and XRD results of LPBF-printed Ni50.83Ti49.17 on a 316 L stainless steel build plate indicate that the phases formed in the diffusion zone depend heavily on the process parameters and energy input, as shown in Figure 9 (Ekoi et al., 2022). For instance, a Volumetric Energy Density (VED) of 50 J/mm³ and 50 J/mm³ produces a mixture of martensitic NiTi phases, while a VED of 66.67 J/mm³ yields an austenitic NiTi [Figure 9(a)-(d)]. The microstructure near the interface displays a complex arrangement of martensite, austenite and iron phases in a layered sandwich pattern along the build direction, which contributes to excellent bonding. In addition, the addition of iron near the interface, as well as typical thermomechanical treatments such as cold working followed by annealing (e.g. at 673 K) which induce Ti₃Ni₄ precipitates, can modify the path of martensitic transformation from simple B2 → B19’ to B2 → R → B19’. Further studies on NiTi-X or other shape memory pairs using new MM-LPBF systems can enhance our understanding of their interface properties, as well as their shape memory, and superelastic effects.

Table 2 is a summary of the added functionalities and materials with potential application areas, challenges and possible solutions discussed in this section.

This two-part review established a comprehensive roadmap for navigating the complexities of Multimaterial Laser Powder Bed Fusion (MM-LPBF). While Part 1 defined the technological constraints of powder deposition and the fundamental physics of the melt pool, Part 2 successfully categorized these findings into a functional framework for the end-user. Based on the synthesis of metallurgical compatibility and functional requirements, a clear decision-support strategy emerges. Successful multimaterial bonding is primarily dictated by a compatibility triad, where designers must prioritize matching melting points, thermal conductivity and CTE. When property mismatches in this triad exceed 20%, the integration of functionally graded transitions or compatible interlayers becomes mandatory to prevent delamination and residual stress buildup. However, the choice between a functionally graded transition and a discrete compatible interlayer depends strictly on the predicted volume and brittleness of intermetallic phases (IMCs). In alloying systems where the risk of forming brittle intermetallics is high, process parameter optimization or gradient transitions may inadvertently broaden the brittle zone; in such cases, only a discrete, compatible interlayer is an effective mitigation strategy.

From an application perspective, the selection of material pairs must align with the specific functional categories identified in this review:

• strength/ductility;

• wear resistance;

• conductivity;

• corrosion resistance; and

• shape memory.

For thermal management, pairing copper alloys with tool steels or aluminum alloys provides optimal heat extraction, though high-reflectivity copper necessitates specialized green or blue lasers for dense bonding. For structural and aerospace applications, titanium and nickel-based pairings (e.g. Ti6Al4V to Inconel or Ti-Al systems) offer a pathway to localized high-temperature resistance and weight reduction, but require careful management of the heat-affected zone to prevent cracking. Similarly, the integration of wear-resistant coatings like Stellites or high-carbide steels onto tougher structural bases is feasible provided energy density is controlled to limit dilution. These decisions are further constrained by hardware; 2D systems remain the standard for high-speed production of simple bimetallic structures, whereas complex, localized material placement requires 3D suction-based or ultrasonic systems.

To move multimaterial LPBF from laboratory research to industrial maturity, several critical areas require investigation. The development of in situ monitoring and closed-loop control systems, such as high-speed pyrometry and acoustic monitoring, is essential to detect interfacial defects and adjust laser parameters dynamically as material composition shifts. In addition, the economic and environmental impact of mixed-powder waste must be addressed. Establishing efficient powder separation and recycling techniques is a vital prerequisite for the sustainable viability of MM-LPBF in high-volume manufacturing. Future research must also focus on the standardization of mechanical testing and computational design. There is currently a significant lack of standardized testing protocols for multimaterial interfaces; developing specific ASTM/ISO standards for shear, fatigue and thermal cycling of bimetallic bonds is necessary for certification in the aerospace and medical sectors. Finally, the field would benefit from advanced Material-by-Design software that can automatically suggest transition gradients or specific interlayer materials based on desired functional outputs and the underlying risk of intermetallic formation, effectively bridging the gap between digital design and physical metallurgical constraints.