# Feasibility study of advanced manufacturing processes: Integrating LPBF and LMD for Inconel 718

Pinku Yadav, Olivier Rigo, Corinne Arvieu, Eric Lacoste

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Pinku Yadav, Olivier Rigo, Corinne Arvieu, Eric Lacoste. Feasibility study of advanced manufacturing processes: Integrating LPBF and LMD for Inconel 718. Journal of Advanced Joining Processes, 2025, 11, pp.100296. 10.1016/j.jajp.2025.100296. hal-05039563

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# Feasibility study of advanced manufacturing processes: Integrating LPBF and LMD for Inconel 718

Pinku Yadav<sup>a,b</sup>, Olivier Rigo<sup>b</sup>, Corinne Arvieu<sup>a</sup>, Eric Lacoste<sup>a,*</sup>

## ARTICLE INFO

**Keywords:**

Hybrid Manufacturing

Laser Melt Deposition

Laser Powder Bed Fusion

Micro Tensile Tests

## ABSTRACT

Laser hybrid manufacturing combines Laser Powder Bed Fusion (LPBF) and Laser Melt Deposition (LMD) to overcome LPBF's size constraints and LMD's lower geometric precision. This study explores the feasibility of hybrid LPBF-LMD processing for Inconel 718, focusing on interface properties and mechanical performance. Hybrid samples were first fabricated using LPBF, followed by LMD, with LMD process parameters optimized using a second-order parabolic model. Two LPBF variants as-built and solution-annealed were evaluated to assess their influence on interface characteristics. Microstructural analysis revealed a fine-grained LPBF region and a coarser LMD region with distinct texture, both demonstrating defect-free metallurgical bonding. Microhardness measurements showed a gradient at the LPBF interface, increasing from 346 ± 20 HV at the build plate to 410 ± 18 HV, influenced by solidification and thermal gradients. The LMD region exhibited a lower hardness of 314 ± 12 HV, correlating with its coarser microstructure. Tensile tests showed that as-built LPBF-LMD samples had higher elongation (26.76 ± 2 %) compared to solution-annealed samples (8.29 ± 2 %), with the LPBF region contributing more to ductility. These findings provide key insights into optimizing hybrid LPBF-LMD processing for high-performance components, enabling improved repair strategies and multifunctional part design in aerospace, energy, and other critical applications.

## Introduction

In recent years, additive manufacturing (AM) has gained significant popularity across various industries due to its ability to overcome traditional geometrical constraints. According to the Wohlers Report, the adoption of AM technologies in industrial production environments is on the rise (Wohlers, 2019). Among the most prominent laser-based AM technologies are Laser Powder Bed Fusion (LPBF) or Selective Laser Melting (SLM) and Laser Melt Deposition (LMD), also known as Direct Energy Deposition (DED) or laser cladding. However, these technologies have their limitations: LPBF is typically constrained by build size and has longer production times due to lower build rates, while LMD is limited in terms of geometric complexity and the precision of produced parts (Buchbinder et al., 2011).

To capitalize on the strengths of different AM technologies, combining them in a "hybrid manufacturing" approach is increasingly seen as a promising solution. For instance, LMD can achieve higher deposition rates of up to 1000 cm³/h, while LPBF can produce fine structures with feature resolutions as small as 100–200 µm. By

integrating these technologies, it is possible to mitigate the limitations each poses individually. Recently demonstrated the potential of combining LMD and LPBF processes to manufacture turbine blades, where fine structures (cooling channels) were produced using LPBF and the fir-tree root was added using the LMD process, reducing the production time from 13.85 hr to 5.35 hr (Graf et al., 2015). This not only improves production time but also allows LMD to be used for repairing expensive LPBF-manufactured parts or adding additional functionalities to existing LPBF components (Petrat et al., 2016).

Hybrid manufacturing could be crucial in achieving higher production rates while maintaining fine detail in parts. However, the rapid melting and solidification cycles inherent in these processes can lead to significant challenges, such as high residual stresses, dislocation generation, and material segregation. Additionally, AM processes are often associated with undesirable phase evolution and defects (e.g., pores) that can negatively impact part performance (Wang et al., 2018), (Godec et al., 2020). Therefore, it is essential to control and thoroughly study these challenges within hybrid manufacturing, especially when dealing with complex multi-system environments. The combination of LPBF and

<footnote><sup>a</sup> CNRS, Arts et Metiers Institute of Technology, University of Bordeaux, Bordeaux INP, INRAE, I2M Bordeaux F-33400 Talence, France</footnote>

<footnote><sup>b</sup> SIRRIS, Rue du Bois Saint-Jean 12 4102 Seraing, Belgium</footnote>

<footnote>* Corresponding author.
E-mail address: Eric.lacoste@u-bordeaux.fr (E. Lacoste).</footnote>
<table><caption>**Table 1**
Composition of Inconel 718 alloy powder used for LMD and LPBF printing.</caption><thead><tr><th></th><th>Ni</th><th>Cr</th><th>Fe</th><th>Nb+Ta</th><th>Mo</th><th>Ti</th><th>Al</th></tr></thead><tbody><tr><td>Powder for LMD</td><td>52.8</td><td>18.5</td><td>18</td><td>4.8</td><td>3.5</td><td>0.75</td><td>0.3</td></tr><tr><td>Powder for LPBF</td><td>50–55</td><td>17–21</td><td>Bal.</td><td>4.75–5.5</td><td>2.8–3.3</td><td>0.75–1.15</td><td>0.30–0.70</td></tr></tbody></table>

<table><caption>**Table 2**
Process parameters used for LPBF printing.</caption><thead><tr><th>Power (W)</th><th>Scanning Speed (mm/s)</th><th>Hatch Distance (mm)</th><th>Layer thickness (mm)</th><th>Scanning Strategy</th><th>Rotation (°)</th></tr></thead><tbody><tr><td>212.5</td><td>850</td><td>0.09</td><td>0.06</td><td>10 mm Stripes</td><td>67</td></tr></tbody></table>

DED processes is gaining increasing attention from both the research community and industry. Early applications, particularly in the 2010s, explored this hybrid approach for Ti6Al4V alloys in medical and aerospace fields. Liu et al. (Microstructure and mechanical properties of LMD-SLM 2024) demonstrated that while LPBF offers superior material properties and geometric precision, integrating DED enables the fabrication of larger structures. Their study confirmed a strong metallurgical bond at the interface, with no significant defects such as pores or cracks. Another key application of LPBF-DED hybridization is repairing LPBF components, as investigated by Oh et al. (Oh et al., 2019). Their research on 316L stainless steel revealed that DED process parameters and deposition depth play a crucial role in interface bonding. However, mechanical testing showed that repaired samples had lower tensile strength and elongation compared to fully LPBF-manufactured parts, which was attributed to suboptimal interface bonding. Further studies reinforced these findings. Qin et al. (Qin et al., 2019) confirmed that the DED region remains the limiting factor in mechanical performance, even after post-processing heat treatments. Graf et al. (Graf et al., 2013) highlighted the efficiency benefits of hybrid manufacturing, showing that combining LPBF and DED can reduce turbine blade production time by over 50 percent compared to conventional methods. Despite these advancements, the role of the LPBF-DED interface in determining overall mechanical behavior remains insufficiently explored, necessitating further investigation into its impact on strength and reliability.

Inconel 718, a Ni-Cr-based superalloy, is extensively used in high-temperature applications such as aviation and aerospace due to its superior thermal and mechanical properties under extreme conditions (Mei et al., 2016). The alloy's strength is derived from precipitation hardening by γ'' (Ni₃Nb, D022) and γ' (Ni₃Al/Ti, L12) phases, and it also contains the δ phase (Ni₃Nb, D0a) and secondary carbides like M₂₃C₆ and M₆C, (Alam et al., 2010; Kuo et al., 2009; Niang et al., 2010), Extensive research on wrought Inconel 718 has highlighted the importance of δ-phase precipitation during hot deformation, which improves fatigue properties by refining grain structure (Kañetas et al., 2020). This phase also strengthens grain boundaries, enhancing crack resistance during fatigue (Hu et al., 2016; Pedroso et al., 2023). In contrast, AM Inconel 718, which is not intended for post-processing via hot deformation, is influenced by the high cooling rates of the process that induce dislocation substructures, potentially triggering recrystallisation during subsequent annealing (Tucho et al., 2017). The parameters of AM and heat treatments such as homogenization or solution annealing play a critical role in the microstructure evolution and mechanical properties of AM Inconel 718 parts. A significant challenge in AM Inconel 718 is the segregation of Nb, Ti, and Mo, leading to the formation of Laves phases ((Ni, Cr, Fe)₂(Nb, Mo, Ti), D024), which are brittle and can negatively impact mechanical properties by acting as crack initiation sites, especially when forming long-chained particles in interdendritic regions or cell boundaries (Dinda et al., 2012). This makes solution annealing prior to ageing an essential step in processing AM Inconel 718 to mitigate these detrimental effects.

To the best of the authors' knowledge, no significant study has systematically examined the mechanical behavior of hybrid LPBF-LMD Inconel 718 components. Existing literature suggests that the LMD region typically has inferior mechanical properties due to its coarser microstructure, which limits the ultimate strength and ductility of hybrid parts. However, beyond ultimate strength, the key factors influencing the yield stress, interface bonding, and deformation behavior in hybrid LPBF-LMD structures remain unidentified. Moreover, while prior studies have investigated LPBF-DED hybridization in other alloys, they have not examined the impact of post-processing conditions on interface integrity and mechanical performance, particularly for Inconel 718. Understanding these factors is crucial for optimizing hybrid manufacturing strategies and ensuring structural reliability.

This study aims to fill this gap by investigating the feasibility of hybrid LPBF-LMD processing for Inconel 718. Hybrid samples were fabricated using LPBF followed by LMD, with two LPBF conditions—as-built and solution-annealed examined to assess their influence on interface properties. This approach is particularly relevant for repairing LPBF-manufactured parts, making it essential to evaluate both conditions. Microstructural analysis confirms that the LPBF region exhibits a fine grain structure, while the LMD region has coarser grains, influencing mechanical performance. In-situ SEM tensile tests and micro-hardness measurements reveal a hardness gradient at the interface, with LPBF contributing more to elongation. These findings provide key insights into interface integrity, mechanical performance, and post-processing optimization, ensuring the reliability of hybrid LPBF-LMD components for industrial applications.

## Experimental section

### Materials

Inconel 718 (chemical composition tabulated in Table 1) powder was used as filler for LMD printing on LPBF Inconel 718 samples which also acts as a substrate. The Inconel 718 powder with a particle size distribution between 45–150 µm was supplied by Flame Spray Technologies (FST, Duiven, the Netherlands). Similarly, for LPBF samples, the gas-atomized Inconel 718 powder with particle size distribution 45–106 µm was supplied by the carpenter additive (UK). The chemical composition of the LPBF Inconel 718 powder is also summarized in the Table 1. The powder was kept in an airtight container before printing to avoid humidification.

### Experimental setup

#### Laser powder bed fusion

LPBF samples were fabricated using RenAM 500Q from Renishaw, Germany. The machine is equipped with 4 continuous Yb- fiber lasers achieving a maximum power of 500 W and a laser spot diameter of 85 µm. Also, it has a building platform size of 250*250*300 mm³. Before printing, the building plate was preheated to a temperature of 170 °C, and an oxygen level of 200 ppm was achieved. Argon was used as an inert gas. The optimized parameters are summarized in the Table 2 (Design space for Inconel 718 material was found internally, the results are not part of this study.).

#### Laser melt deposition (LMD)

LMD printing was conducted using an IBARMIA ZVH45/1600 Add+Process hybrid machine, which integrates 5-axis milling and
<figure><img src="image_3.png" /><figcaption>Fig. 1. (a) Experimental setup used for LMD printing, (b) LPBF-LMD Hybrid printed parts.</figcaption></figure>

<table><caption>**Table 3**
Process parameters used for LPBF printing.</caption><thead><tr><td>Process Parameter</td><td>Level</td></tr></thead><tbody><tr><td>Laser Power (W)</td><td>900, 650, 400</td></tr><tr><td>Feed rate (mm/min)</td><td>700, 500, 300</td></tr><tr><td>Powder mass flow (g/min)</td><td>8, 6, 4</td></tr><tr><td>Type of coaxial nozzle</td><td>4-stream Discrete</td></tr></tbody></table>

turning with LMD technology. The system is equipped with a Precitec YC52 cladding head featuring 125 mm collimating optics and 250 mm focusing optics. Powder deposition was facilitated by a Sulzer Metco TWIN-10-C Powder Feeder, and the laser source was a Yb-Fiber Rofin FL030 with a 3 kW output and a continuous wavelength of 1.07 µm. A 4-stream coaxial discrete nozzle from Precitec was used for material delivery. The LPBF samples were secured using cylindrical fixtures, as shown in Fig. 1. High-purity argon (99.999 % purity, 5.0 grade) was used as shielding with a flow rate of 21 L/min to ensure an inert atmosphere during the process.

## Parameter optimization for LMD

Single clad tracks of 60 mm were printed and analysed to find the best parameters for obtaining defect-free desirable morphology of the tracks for LMD printing. The process parameters used for optimization are tabulated in the Table 3.

An average of four different cross-sections perpendicular to the laser scanning direction was selected for optimization. The clad tracks' first and last 10 mm were avoided due to instabilities subjected at the start and stop conditions. The cross-sections were analyzed for porosity, and clad morphology, i.e., height (h), width (w), area (A), dilution depth (p), and dilution area (Da), as shown in Fig. 2. The samples were mirror polished and chemically etched using Kallings 2 etchant for 15 s. The cross-section was measured using the Motic SMZ-143 microscopy (Motic, Hong Kong, China) and Clemex Captiva® software (Clemex, Longueuil, QC, Canada).

## Hybrid samples and heat treatment

To study the feasibility of advanced manufacturing using LMD and LPBF AM processes, cylindrical Inconel 718 samples with a diameter of 16 mm and length of 55 mm were fabricated using the LPBF process. The LPBF samples were then used as a substrate for further LMD printing of Inconel 718 samples. Cylinders with the same geometry as LPBF were fabricated using LMD on LPBF samples. Fig. 3 shows the schematic and actual samples combining LPBF and LMD printing.

The optimal annealing conditions involved a 2 h hold at 1150 °C in a pre-heated furnace, followed by cooling to 700 °C for 12 hours under an argon atmosphere (Yadav et al., 2023). Two different geometries for mechanical characterization samples were machined from Ø12 mm cylinders using electric discharge machining (EDM): (1) flat, dog

<figure><img src="image_4.png" /><figcaption>Fig. 2. Optical micrograph of clad and schematic of clad characteristics (H.A.Z denoted Heat Affected Zone).</figcaption></figure>
<figure><img src="image_5.png" /><figcaption>Fig. 3. Schematic for the printed samples and actual samples of LMD printing along with machined tensile samples.</figcaption></figure>

bone-shaped tensile samples with a gauge length of 20 mm and width of 5 mm for micro in-situ tensile tests, and (2) flat rectangular samples with a width of 5 mm and length of 20 mm for micro-Vickers hardness tests. For all specimens, the LPBF-LMD interface was positioned at the centre of the test section. In this study, the regions of the hybrid samples produced by the LPBF process and the LMD process are referred to as the "LPBF zone" and "LMD zone," respectively.

## Microstructural investigations

The microstructure was examined using optical microscopy and scanning electron microscopy (SEM). The samples were mechanically polished to a mirror finish, followed by etching for 10 s with an Aqua Regia solution, consisting of 30 vol% HNO₃ and 70 vol% HCl. Electron backscatter diffraction (EBSD) analysis was performed using a JSM-7100 TTLS LV SEM, equipped with an HKL EBSD system.

For grain boundary analysis, misorientations greater than 10° were considered, with a minimal grain area defined as 10 pixels (equivalent to 20 µm² or a 2.5 µm diameter). These criteria were based on the guidelines established by Brandon (Brandon, 1966) for face-centred cubic (FCC) alloys. Grain size analysis was based on the equivalent grain diameter, which represents the diameter of a disk with the same area as the grain.

Twin boundaries were identified by a misorientation of 60° ± 5°, twisted about the (111) plane, in accordance with the original work by Brandon (Brandon, 1966). These twin boundaries were excluded from the grain size calculations.

## Mechanical characterizations

### Vickers micro-hardness (mHV)

Indentation matrices of 3 columns (37 measurement points) spaced out by 1 mm were performed on as-built and heat-treated hybrid samples. The first column is indented 1 mm away from the sample's extremity to avoid side effects disruption. The matrix is manually situated so that the interface between LPBF and LMD is at mid-height. The tested area represents a rectangle of 4.5 mm in height by 5 mm in length, corresponding to half a cylindrical sample of 12 mm in diameter. The load applied on the 136° pyramid was 2 kg, and maintained 10 s for each indent. The Vickers hardness (HV2) is then deduced from the diagonals.

### In-situ tensile tests

The tests were executed on a stroke-controlled Kammrath-Weiss 5 kN module from Eden Instrument installed in a Zeiss SEM. The in situ tensile tests were performed with a constant stroke rate of 2 mm/s. During loading, automated image capture was carried out at the frequency of 10 Hz. In addition, tests were occasionally paused at constant load levels to take further images manually at various magnifications. For better identification of the microstructure, samples were manually polished and slightly etched with Aqua Regia solution for 10 s before the tensile tests.

For each condition, two specimens were subjected to tensile testing to calculate the average mechanical properties of IN718 alloy. Before the tensile test, the size of specimens was measured again in order to confirm the effective cross-section, due to the reduction effect of the polishing.

## Results and discussions

### LMD process

#### Parameter optimization for LMD process

The average four values of the cross-sections of all the clads were used to study the single-clad characteristics (w, h, A, p). Each process parameter (i.e., Laser Power, Feed rate, and Mass flow rate) affects the clad characteristics. Therefore, it is crucial to study the correlation between the process parameter and clad characteristics. Based on the process optimization results, principal component analysis (PCA) was used to study the correlation, as shown in Fig. 4. For a detailed explanation of Principal Component Analysis and its application in material characterization, refer to (Hotelling, 1933). Here values >0 show the positive correlation, and values <0 represent the negative correlation among the variables. The intensity of the color is related to the scale. Notably, the clad height shows a positive correlation with the powder mass flow rate and a strong negative correlation with the feed rate. It is important to clarify that Area_AR refers to the aspect ratio of the clad area, calculated as Width/Depth, while Dilution_AR pertains to the aspect ratio of the dilution area, determined as Width/Dilution Depth.

This relationship is also evident in the experimental value plots, as illustrated in Fig. 5. The figure displays the clad height, width, and aspect ratio (Area_AR) in relation to Laser Power, Feed Rate, and Mass Flow Rate. It can be observed that Laser Power had no significant impact on the height, although the width increased with higher laser power. As
<figure><img src="image_6.png" /><figcaption>**Fig. 4.** PCA correlation matrix between process parameters and clad characteristics.</figcaption></figure>

anticipated, both height and width decreased with an increase in feed rate.

The optimal conditions for LMD printing were selected as follows:

1. Firstly, the dilution depth should be higher than the clad height, as it will ensure deeper penetration of the beneath layer to create defect-free and proper adhesion.

2. Secondly, the aspect ratio should be between 3–5, ensuring good overlap between the tracks as reported in the previous study (Ramiro-Castro et al., 2020; Ramiro et al., 2020).

3. Thirdly, a mass flow rate of 6.7 g/min was chosen intuitively as a slower mass flow rate leads to longer printing time, whereas a higher mass flow rate could lead to more thick tracks.

Therefore, a laser power of 650 W, a feed rate of 500 mm/min, and a mass flow rate of 6.7 g/min were chosen for this study.

## Layer height optimization

The layer height uniformity in the LMD process is less consistent compared to other AM techniques such as LPBF or Electron Beam Melting (EBM), making it essential to calculate the correct layer height to reduce the risk of porosity formation between tracks. Therefore, to optimize the layer height theoretically, there are models already reported in the literature (P. Ramiro et al., 2020). A simple second-order parabola model was used, utilizing the single track's clad characteristics as inputs to determine the height and width for future overlapping clad tracks. The model showed a high level of accuracy, as reported in (P. Ramiro et al., 2020). The overall methodology and equations are discussed (P. Ramiro et al., 2020).

The height and width for the first 20 tracks were calculated using a 45% overlap, as shown in Fig. 6. As can be seen, for the first 4 tracks, the height and width are not stable (related to the LMD process itself). Therefore, the height and width of the 5th track were used while printing.

## Microstructure

The optical micrographs in Fig. 7 illustrate the interface between two scenarios: AB LPBF-LMD and HT LPBF-LMD. In the as-built (AB) LPBF part (Fig. 7a), the microstructure is finer, characterized by smaller melt

pools, while the LMD region exhibits a coarser microstructure with wider melt pools. After heat treatment (HT), the LPBF section no longer shows melt pool lines, as these features have been obliterated by the heat treatment process. However, the formation of twins is clearly visible. The bond area and heat-affected zone (HAZ) are distinctly marked in Fig. 7. It is also noteworthy that the LMD region contains more spherical defects, predominantly gas porosity, compared to the LPBF-processed material. In both cases, good metallurgical bonding is noticed.

Fig. 8 presents SEM images of the As-built LPBF, HT LPBF, LMD, and bond area, highlighting the distinct microstructural features in each region. The As-built (AB) LPBF (Fig. 8b) reveals a fine cellular dendritic structure with a few visible second-phase particles along the cell walls. This cellular dendritic structure is associated with dislocation formation due to rapid solidification, a phenomenon well-documented inadditively manufactured Inconel 718. It has been noted that these cells are actually tubes, depending on their orientation, and correspond to nanoscale solidification dendrites formed during rapid cooling (Godec et al., 2020; Teixeira et al., 2021).

After solution annealing at 1050 °C, the typical cell structure in the LPBF parts is no longer visible. Instead, the microstructure shows globular secondary particles (δ phase) along with some annealed twin boundaries (Godec et al., 2020; Teixeira et al., 2021) (Fig. 9). The electron image of the as-received LMD part reveals a rough dendritic structure with coarser second-phase particles in the interdendritic regions (Fig. 10). Furthermore, the electron image of the LMD part shows clear evidence of segregation along the interdendritic areas, which appear as topographical variations due to the etching effect (Fig. 10).

As reported in previous studies, both LPBF and LMD parts exhibit fine Laves phases and oxide particles, with the oxides typically presenting a circular shape. Similar oxides were observed in this study, identified as Al₂O₃ particles formed due to aluminium's high affinity for oxygen (Yu et al., 2018), likely originating from the feedstock powder. Despite both additive manufacturing processes occurring in an inert atmosphere, oxygen infiltrates the melt through a thin oxide layer on the surface of each powder particle (Godec et al., 2020; Hybrid additive manufacturing of Inconel 2025). The Laves phase is characterized by a higher concentration of Nb and slightly more Mo than the γ-matrix.

At the interface between the LPBF and LMD parts, the LMD region shows a reduced number of second-phase particles (Fig. 11) compared to the bulk of the LMD part. The electron micrograph of the bond area reveals the absence of a cellular structure (Fig. 11). This lack of cell structure at the bond is attributed to the material undergoing remelting and recrystallization, leading to reduced segregation. However, some small islands of segregation remain because the process did not eliminate segregation.

Fig. 12 displays the EBSD and texture analysis of the As-built (AB) LPBF, heat-treated (HT) LPBF, and LMD samples. The AB LPBF sample shows a columnar grain structure aligned along the build direction, with an average grain size of 14.4 ± 3 µm. After heat treatment, the grains in the LPBF sample become much larger and equiaxed, with an average grain size of 21.1 ± 2 µm. This increase in grain size is due to solution annealing at 1050 °C, which promotes grain growth and the dissolution of smaller grains. Similar to the AB LPBF, the LMD sample exhibits elongated columnar grains along the build direction, but with a significantly larger average grain size of 56.8 µm—approximately four times that of the AB LPBF sample. This difference is attributed to the slower cooling rates in the LMD process compared to the LPBF process.

The corresponding Inverse Pole Figures (IPF) reveal that the AB LPBF sample has a strong texture in the (001) direction, parallel to the build direction. Following solution annealing, the columnar grains in the LPBF part disappear due to recrystallization, resulting in a random texture and the formation of equiaxed grains, as shown in Fig. 12b. In contrast, the LMD sample displays a strong texture in the (101) direction, parallel to the transverse direction, with a texture intensity of 3.7.

Fig. 13 also presents a high-magnification view of the interface for two scenarios: AB LPBF-LMD and HT LPBF-LMD. In both cases, epitaxial
<figure><img src="image_7.png" /><figcaption>Fig. 5. Plots of clad width, height, and area aspect ratio with respect to process parameters: laser power, feed rate and mass flow rate. [Note: Please zoom for better clarity].</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>Fig. 6. Clad height (left) and width (right) for an overlap of 45 % predicted by second order parabola model.</figcaption></figure>

grain growth across the interface is observed. The EBSD analysis on the LMD side of the bond area reveals a mixture of columnar dendrites and smaller equiaxed grains at the interface. Additionally, the columnar grains are oriented perpendicular to the plane of the bond. Near the

bond area on the LMD side, the EBSD analysis reveals periodically occurring regions of smaller grains (Fig. 13).

Fig. 14 illustrates the misorientation angle distribution for the AB-LPBF, HT-LPBF, and LMD samples. The AB-LPBF and LMD samples
<figure><img src="image_9.png" /><figcaption>Fig. 7. Optical micrographs (OM) illustrating different regions of the LPBF-LMD processed sample: (a) As-built LPBF-LMD interface, (b) HT LPBF-LMD interface, (c) As-printed LPBF region, (d) As-printed LMD region, and (e) HT LPBF region.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>Fig. 8. (a) SEM image of the as-printed LPBF region, showcasing the overall microstructure. (b) Zoomed-in SEM image of the same region, providing a detailed view of the finer microstructural features (Melt pool boundaries marked with dotted lines).</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>Fig. 9. (a) SEM image of the heat-treated (HT) LPBF region, highlighting the overall microstructure. (b, c) Magnified views showing secondary particles along with twins within the HT LPBF region.</figcaption></figure>
<figure><img src="image_12.png" /><figcaption>**Fig. 10.** (a) SEM image of the as-built LMD region, illustrating the overall microstructure. (b) Zoomed-in SEM image of the same region, providing a closer look at the finer microstructural features.</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>**Fig. 11.** SEM images of the LPBF-LMD interfaces: (a, b) As-built LPBF-LMD interface showcasing the transition region, and (c, d) Heat-treated (HT) LPBF-LMD interface highlighting the microstructural changes post heat treatment.</figcaption></figure>

both exhibit a significant amount of subgrain boundaries (misorientation angles <5°), accounting for 43% and 41% of the boundaries, respectively. Inside the grains, a misorientation was observed due to the stress that occurred during the LPBF process, leading to slightly rotated dendrites inside the grains (Hybrid additive manufacturing of Inconel 2025), (Godec et al., 2020). Additionally, the LMD samples display a lower proportion of low-angle grain boundaries (misorientation angles <15°) compared to the AB-LPBF sample. This suggests that the LPBF process, with its rapid solidification, induces a higher density of dislocations, leading to the formation of a cellular structure. In contrast, the LMD sample, which experiences slower cooling, has a lower dislocation density.

Furthermore, the misorientation data for the heat-treated (HT) LPBF sample show a reduction in the number of subgrain boundaries following solution annealing, indicating a decrease in dislocation

density due to the heat treatment process (Fig. 14). However, there is a slight but noticeable increase in the number of annealing twins, as depicted in the graph. This observation aligns with the well-documented behavior of Inconel 718, which is known to develop annealing twins during solutionizing heat treatment.

### *In-situ micro tensile tests*

The plasticity behavior of the interface in both scenarios was investigated using in-situ SEM tensile tests. Figs. 15 and 16 present the stress-strain curves for the samples during these tests. The results show that the AB LPBF-LMD samples exhibited better elongation compared to the HT LPBF-LMD samples, with elongations of 26.76 ± 2% and 8.29 ± 2%, respectively. For the AB LPBF-LMD samples, the average contributions to elongation from the LPBF and LMD regions were 63.5% and
<figure><img src="image_14.png" /><figcaption>**Fig. 12.** EBSD and IPF images of (a) As-built LPBF, (b) Heat-treated (HT) LPBF and (c) LMD sample highlighting the microstructural evolution and texture changes induced by heat treatment.</figcaption></figure>
<figure><img src="image_15.png" /><figcaption>**Fig. 13.** EBSD IPF images of the LPBF-LMD interfaces: (a) As-built LPBF-LMD interface illustrating the grain orientation and texture across the transition region, and (b) Heat-treated (HT) LPBF-LMD interface showcasing the microstructural changes induced by the post-processing heat treatment.</figcaption></figure>

<figure><img src="image_16.png" /><figcaption>**Fig. 14.** Misorientation angle distribution for different processing conditions.</figcaption></figure>

36.5 %, respectively. In contrast, for the HT LPBF-LMD samples, the contributions were 25 % from the LPBF region and 75 % from the LMD region. Despite the differences in elongation, the ultimate tensile strength (UTS) for AB LPBF-LMD samples is 833 ± 8 MPa, while for HT LPBF-LMD samples, it is 825 ± 10 MPa, showing no significant variation. However, the yield strength (YS, at 0.2 % strain) differs, with AB LPBF-LMD samples exhibiting a YS of 676 ± 10 MPa and HT LPBF-LMD samples a YS of 629 ± 8 MPa. The fracture for AB LPBF-LMD samples occurred in the LPBF part, whereas for HT LPBF-LMD samples, the failure occurred in the LMD region due to the presence of porosity.

The observed behaviours are qualitatively linked to the strain distribution. Fig. 15 illustrates the different deformation stages for the AB LPBF-LMD sample. Up to approximately 1.5 % total strain, the microstructure remains largely unchanged. Subsequently, surface deformations begin to appear in the grains of the LMD region. The grains start to tilt, revealing what appear to be boundaries between the deposition layers. Failure in the LMD zone initiates between these layers. Notably, the crack does not initiate from the side of the elongated grains but rather in areas where grain boundaries are likely perpendicular to the loading direction, suggesting an intergranular failure mechanism. The interface between the LPBF and LMD regions, marked in orange, shows no significant deformation and does not contribute to the failure. Occasionally, slip bands (highlighted in red arrow) are observed on a smaller scale. These bands cross the dendritic structures without deviation but change direction from one grain to another, indicating the influence of crystallographic texture on the deformation behaviour. Surface deformations are also noted in the LPBF region towards the end of the test, although they are less pronounced. Fewer slip bands are

observed in the as-built LPBF part, while they are more prominent in the annealed LPBF samples (Fig. 16).

### Micro hardness

Micro-hardness (HV2) measurements taken at the interface of the hybrid samples are shown in Fig. 17. A clear difference in hardness is observed between the as-built LPBF and LMD microstructures. The as-built LPBF exhibits a Vickers micro-hardness of 346 ± 20 HV at the build plate, which increases gradually to 410 ± 18 HV as we move away from the build plate. This gradient in hardness is likely due to the layer-wise solidification and thermal gradients during the LPBF process, leading to a finer microstructure further from the build plate. In contrast, the LMD region shows a lower hardness of 314 ± 12 HV, with larger grains, dendrites, and a lower dislocation density compared to the finer microstructure of the as-built LPBF. Interestingly, the hardness of the heat-treated LPBF sample remains constant across the entire height of the part, indicating that the heat treatment has homogenized the microstructure. Further studies are needed to fully understand the underlying mechanisms of this behaviour.

### Conclusions

The feasibility of hybrid manufacturing was explored by integrating the Laser Powder Bed Fusion (LPBF) and Laser Metal Deposition (LMD) processes for Inconel 718. This study focused on both as-built and solution heat-treated states (1150 °C for 2 h). The following conclusions emerged from our mechanical and microstructural analyses:

* Hybrid manufacturing that combines two additive manufacturing processes presents a promising approach for producing complex, high-performance parts across diverse applications. Defect-free metallurgical bonding was successfully achieved in both LPBF states without the need for additional surface modifications.

* The two processes yield distinct microstructures: the LPBF part displays a fine, columnar grain structure with a well-defined dislocation cell substructure, whereas the LMD part reveals a coarser microstructure with grain sizes three times larger than those seen in the LPBF part. Notably, discontinuous Laves phases and oxides were identified in the LMD section.

* Annealing at 1050 °C for 2 hours effectively dissolved the Laves phases in the LPBF part, resulting in the formation of globular delta phases at the grain boundaries.

* The mechanical properties of the hybrid samples were predominantly governed by the LMD portion, which served as a weak point for crack initiation, likely due to the higher gas porosity observed in this section.
<figure><img src="image_17.png" /><figcaption>Fig. 15. Stress-strain curve for the as-built LPBF-LMD part, illustrating the mechanical response under tensile loading. Insets show SEM micrographs taken at different locations along the interface, highlighting the microstructural features contributing to the mechanical behavior. Red dashed lines indicate the grains and red arrows shows the direction of slip bands.</figcaption></figure>

<figure><img src="image_18.png" /><figcaption>Fig. 16. Stress-strain curve for the heat-treated (HT) LPBF-LMD part, depicting the mechanical response after thermal processing. Insets display SEM micrographs from various locations along the interface, revealing the microstructural evolution and its influence on the material's mechanical behavior. Red dashed lines indicate the grains and red arrows shows the direction of slip bands.</figcaption></figure>

* The LMD part exhibited lower microhardness compared to both the as-built and solution heat-treated LPBF parts. In the heat-treated LPBF part, hardness decreased linearly as the distance from the interface increased, a phenomenon warranting further investigation.

* Elongation results showed that the as-built LPBF-LMD samples exhibited better elongation compared to the heat-treated LPBF-LMD

samples, with the LPBF region contributing more to elongation than the LMD region.

Overall, hybrid manufacturing offers a valuable method for producing complex parts at elevated production rates while preserving finer details. This approach effectively combines the strengths of both LPBF
<figure><img src="image_19.png" /><figcaption>**Fig. 17.** Microhardness profiles across the interface for LPBF-LMD parts: As-built (AB) LPBF-LMD, showing the hardness variation due to the as-fabricated microstructure, and Heat-treated (HT) LPBF-LMD, highlighting the changes in hardness distribution resulting from thermal processing. Measurement points from 1 to 19 indicates the LMD part and 20 to 37 points are on LPBF part.</figcaption></figure>

and LMD processes to address their individual limitations. Furthermore, it holds promise for the repair and modification of existing AM-produced parts, although additional research is essential for a comprehensive understanding of bond behavior at the microstructural level.

## Funding

This work was conducted as part of the "ENABLE" project funded by the European Union's Marie Skłodowska-Curie Actions (MSCA) Innovative Training Networks (ITN) H2020-MSCA-ITN-2017 under the grant agreement Number 764979.

## CRediT authorship contribution statement

**Pinku Yadav:** Writing - review & editing, Writing - original draft, Project administration, Methodology, Investigation, Conceptualization.
**Olivier Rigo:** Supervision, Software, Resources, Formal analysis.
**Corinne Arvieu:** Supervision, Resources, Data curation.
**Eric Lacoste:** Supervision, Project administration, Funding acquisition.

## Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Acknowledgments

The authors would like to thank the ADD team at Lortek IK4, Spain for providing valuable suggestions and access to LMD equipment.

## Data availability

Data will be made available on request.

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