Microstructural assessment of additive-manufactured Inconel 718 samples subjected to heat treatments for enhanced mechanical properties

Microstructural assessment of additive-manufactured Inconel 718 samples subjected to heat treatments for enhanced mechanical properties | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Research Article Microstructural assessment of additive-manufactured Inconel 718 samples subjected to heat treatments for enhanced mechanical properties Thiago Roberto Felisardo Cavalcante, Douglas Giovanni Bon, Fábio Edson Mariani, and 7 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-6602510/v1 This work is licensed under a CC BY 4.0 License Status: Published Journal Publication published 05 Aug, 2025 Read the published version in Progress in Additive Manufacturing → Version 1 posted You are reading this latest preprint version Abstract Directed energy deposition-laser beam (DED-LB), a laser additive manufacturing (AM) process, has attracted significant attention as a potential alternative to conventional manufacturing methods, due to its high deposition efficiency, flexibility, and precision. Despite these advantages, components produced by DED-LB often face critical challenges, including residual stresses, micro-segregation, and the formation of non-equilibrium phases due to rapid cooling during the AM process. These issues are particularly critical for Ni-based superalloys such as Inconel 718 (IN718), widely used in the aerospace, energy, and marine industries for their excellent high-temperature strength and corrosion resistance. The mechanical performance of IN718 primarily depends on precipitation hardening via γ' and γ'' phases. In contrast, the formation of deleterious phases, such as δ and Laves, can severely impair performance by depleting key alloying elements and increasing brittleness. Thus, heat treatments (HTs) are vital in addressing these challenges by reducing micro-segregation, homogenizing elemental distribution, and promoting the precipitation of strengthening phases. Therefore, this study investigates the effects of six distinct heat-treatment routes on the microstructural evolution, hardness, tensile properties, and fracture behavior of DED-LB IN718 samples. The relationship between microstructure and mechanical responses is analyzed and compared to a forged IN718 counterpart. The results offer valuable insights for optimizing heat-treatment strategies to improve the structural integrity and mechanical reliability of DED-LB-fabricated IN718 components. IN718 heat treatment microstructure additive manufacturing mechanical properties Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Figure 7 Figure 8 Figure 9 Figure 10 Figure 11 Figure 12 1. Introduction Recently, laser additive manufacturing (AM) techniques have been extensively researched and are in a state where they could be considered potential candidates to replace traditional manufacturing methods when needed [ 1 ]. Laser beam directed energy deposition (LB-DED) is a type of AM that deposits thin layers of metal powder and melts them using a high-energy laser beam; furthermore, it can deposit metallic materials and functional coatings. LB-DED has the advantages of high deposition efficiency, flexibility, and precision [ 2 ]. However, due to the rapid cooling of the process, residual stresses, micro-segregation, and the formation of non-equilibrium phases are a common feature in as-built components [ 3 ]. Ni-based superalloys, such as Inconel 718 (IN718), are heat-treatable, hot corrosion resistant, and preferred for the fabrication of aerospace engines, turbine blades, injectors, and many other high temperatures and/or high-pressure components [ 2 ], [ 4 ]. This alloy, specifically, is mainly strengthened by precipitation hardening, via the precipitation of γ’ and γ’’ phases [ 5 ]. Moreover, the solid solution hardening effects of some elements in the γ matrix, such as Nb and Mo, can be additional factors to improve the mechanical properties. The presence of micron/nano-sized γ’ and γ’’ phases controls the mechanical properties of the IN718 alloy. On the other side, deleterious phases, such as δ and Laves phases, can degrade the materials' properties, either due to the incoherent nature of the phase with the matrix, or the inhibition of γ’ or γ’’ precipitation [ 6 ]. The micro-segregation of elements and the presence of deleterious phases, typical in AM components, make the elemental distribution uneven, which also impacts the precipitation of γ’ and γ’’ phases [ 7 ]. Laves phase forms in the interdendritic regions and is known to deteriorate mechanical properties such as strength, ductility, fatigue, and creep, since it depletes essential precipitation elements, besides facilitating crack initiation and propagation due to its brittle nature [ 7 ], [ 8 ]. Applying heat treatments (HT) and controlling the solidification conditions during the deposition are two ways to reduce the micro-segregation and formation of the Laves phase. It has been reported that Nb segregation and Laves phase formation can be reduced by increasing the cooling rate from the liquid material [ 9 ]. However, controlling the solidification conditions is rarely achievable due to the complicated solidification of the AM processes [ 8 ]. As such, post-fabrication HTs are vital to achieve excellent mechanical performance and meet service requirements through the elemental distribution homogenization and precipitation of the strengthening phases [ 10 ]. The precipitation hardening heat treatment involves the dissolution of the Laves phase, since it releases Nb and Ti back to the matrix [ 11 ]. Generally, a two-step direct aging (DA) and solubilization (S) are performed, and, in some cases, a homogenization step is also applied. The solubilization typically occurs in the range of 950 to 1100 o C to dissolve the Laves phase and release Nb and Ti into the matrix; direct aging usually ranges from 600 to 800 o C and results in the formation of γ’ and γ’’ precipitates. Furthermore, homogenization heat-treatments (H) can also eliminate Laves and microsegregation, enhancing the homogeneous precipitation of strengthening phases [ 12 ]. Several studies have been conducted regarding the influence of HTs on AM IN718. Qi et al. [ 12 ] reported that the microstructure and tensile properties of IN718 manufactured by DED-LB with DA achieved high tensile strength values with low ductility. Huang et al. [ 13 ] showed that, for alloys fabricated via PBF-LB process, different solution treatment temperatures, ranging from 980 to 1280 oC for 1h, implied different Laves phase dissolution degrees. Burad et al. [ 7 ] investigated the effect of S and S + DA HTs and concluded that solution treatment at 980 o C, to dissolve the Laves phase and microsegregation spots, led to a good precipitation quality. Zhai et al. [ 14 ] found that a simple aging process could precipitate an amount of γ’ and γ’’ phases capable of enhancing the mechanical properties of IN718, with weak microsegregation and fine Laves phase still present. However, it could not change the typical columnar dendrite structure. Zhao et al. [ 15 ] demonstrated that a homogenization temperature of 1180 o C can eliminate residual stress and promote recrystallization of IN718 manufactured by LB-PBF. Regarding the Laves phase, Sui et al. have studied its dissolution [ 16 ], who investigated the effect of a 1050 o C solubilization with DA on LB-DED IN718 samples, showing a large improvement in high-temperature mechanical behavior. Xu et al. [ 10 ] studied the microstructural evolution and the mechanical performance of AMed IN718 under three HTs: DA, S + DA, and H + S + DA, and concluded that different as-built microstructures require different HTs; a similar study was also performed by Yu et al. [ 17 ], who focused on the fracture behavior of the LB-DED IN718 deposited material. Furthermore, as-built microstructures, according to Jang et al. [ 18 ] studied the precipitation kinetics of secondary phases formed by the heat accumulation during the IN718 DED deposition. This study investigates the effects of six different heat treatment strategies on the microstructural evolution, hardness, tensile properties, and fracture behavior of IN718 samples produced by DED-LB. Special focus is placed on the influence of these treatments on the precipitation of strengthening phases and the dissolution of deleterious phases such as the Laves phase. The mechanical performance of the treated AMed samples is critically compared to that of their forged counterparts to assess the effectiveness of post-processing in mitigating typical additive manufacturing defects. The findings aim to support better heat-treatment routes for AMed IN718 alloys, promoting improved structural integrity and mechanical reliability for high-performance applications. 2. Experimental procedure 2.1 Materials For this study, gas-atomized IN718 powder, produced by Carpenter Additive Inc., was used. The powder's chemical composition, expressed in weight percentage (wt.%), was as follows: Ni 52.30, Cr 19.72, Nb + Ta 4.96, Mo 3.30, Co 0.53, Ti 0.39, Si 0.09, Mn 0.04, C 0.01, P 0.01, with Fe in balance. 2.2 Materials processing The depositions were conducted on AISI 1020 steel with a thickness of 5 mm, which was end-milled and cleaned before the experiments. A BeAM Inc. Modulo 250 5-axis Directed Energy Deposition-Laser Beam (DED-LB) system was utilized, equipped with a 1 kW continuous-wave ytterbium fiber laser (Model YLR-1000-MM-WC-Y14, IPG Photonics Inc.), operating at a wavelength of 1070 nm and a focused spot size of 0.8 mm. The laser was mounted on a coaxial deposition head, and the system incorporated a dual powder feeder for material deposition. The parameters utilized for this work were a laser power of 500 W with a crosshead speed of 2000 mm/min and a feed rate of 5 g/min. These parameters were determined in a previous study, which evaluated various combinations of power, scanning speed, and powder feed rate to achieve high-integrity single tracks [ 19 ]. Detailed machine parameters, process schematics, and deposition parameters can be found in previous studies [blinded]. All heat treatments (HT) were carried out in a Jung-brand muffle furnace (LF07013 model). The six HT routes employed in this study were designed to evaluate the effects of homogenization, solubilization, and aging on the microstructure and mechanical properties of the material. Th different kind of heat treatments used are: (i) homogenization (H), consisted of heating the samples to 1100°C for 1.5 h, followed by water quenching; (ii) solubilization (S), where the material is isothermally heat treated at 1000°C for 1.0 h, also followed by water quenching; and (iii) double aging (DA), consisted on an aging at 720°C for 8 h, followed by controlled cooling at 50°C/h down to 620°C, where the samples were held for an additional 8 h, and then air-cooled to room temperature. In addition to these isolated heat treatments, three combinations were also evaluated: homogenization followed by double aging (HDA), solubilization followed by double aging (SDA), and homogenization followed by solubilization and double aging (HSDA). To assess the porosity level after manufacturing, a Bruker Skyscan 1272 microtomograph was utilized. It features an 11 MP detector (4032 × 2688 pixels) and operates at 10 W. The sample, measuring 2 mm × 3 mm × 8 mm, was scanned at a resolution of 2 µm with a step size of 0.3°. The 3D model was reconstructed and rendered using the Dragonfly™ 2022.2 software developed by Comet Group. Volumetric porosity was then calculated by applying color thresholding to distinguish between the pores and the solid material. 2.3 Microstructural characterization The samples underwent standard preparation for metallographic evaluation, involving abrasive cutting, manual grinding with sandpapers, and sequential surface polishing using 1.00 and 0.25 µm diamond suspensions. Final etching was done using waterless Kalling’s reagent (100 ml ethanol, 100 ml HCl, 5 g CuCl₂) to reveal the microstructure. Cross-sectional microstructure observations were conducted with an Olympus LEXT 4100 3D laser microscope and a FEI® FE-50 Scanning Electron Microscope (SEM). Microhardness was assessed using a Buehler Model 1600 6300 tester by ASTM E384 [ 21 ], employing a 0.5 N load for 15 s, Fig. 1 d. Further microstructural details were examined via EBSD using a Quanta 650 FEG-SEM and high-speed acquisition EBSD system. Data were analyzed in TSL OIM and MTEX, with EBSD scans performed over 2000 µm × 1500 µm (step size 1.5 µm) and 500 µm × 500 µm (step size 0.5 µm). Maps achieved ≥ 90% indexing; non-indexed points were corrected using the grain dilation correction method, while grains with less than 2 pixels were excluded. GND densities were obtained via the Kernel Average Misorientation (KAM) and the Weighted Burgers Vector (WBV) [ 22 ] approaches. Boundaries were classified by misorientation: > 15° (high-angle grain boundaries—HAGB), 5–15° (low-angle grain boundaries—LAGB), and 2–5° (subgrain boundaries—SGB). Synchrotron energy-dispersive X-ray diffraction (EDXRD) was carried out at P61 High Energy wiggler beamline/LVP, installed at Petra III, Hamburg, Germany. Transmission-mode scans employed 10 wigglers (30–200 keV range), with spectra collected along 10 mm line scans at 0.2 mm steps. Two HPGe detectors positioned at 2θCh0 = 7.557° and 2θCh1 = 7.594° captured data from a 0.1 × 0.1 × ≈3 mm² gauge volume. Exposure time per spectrum was 1 s. Raw data were processed using the in-house Python scripts. 2.4 Mechanical Testing Digital Image Correlation (DIC) tensile testing was performed using a Deben Microtest Tensile Stage on dog-bone samples (gauge length: 4 mm × 1 mm × 1 mm) at 0.4 mm/min (Fig. 1 a and b). Load, displacement, and axial strain were captured via a Balser camera (2048 × 2048 pixels at 75 fps) using 35-, 50-, and 75-mm lenses. DIC analyses were done using Ncorr to generate strain maps and profiles. Conventional tensile tests followed ASTM E8-24 [ 23 ], performed under displacement control (1 mm/min) on a 100 kN MTS Landmark N servo-hydraulic system. Specimens, extracted horizontally, are illustrated in Fig. 1 c. 3. Results and discussion 3.1 Microstructural characterization 3.1.1 Microstructure of as-built samples Figure 2 presents the porosity analysis for the as-built condition, conducted using X-ray diffraction microtomography. This analysis estimated a volumetric porosity of approximately 0.37%, with pore sizes ranging from 13 µm to 150 µm. The pores exhibited irregular morphologies aligned with the building direction, as evidenced by the different views of the analyzed solid shown in Fig. 1 (a). The porosity size calculations resulted in an average pore size of 64 ± 32 µm and a sphericity of 0.6 ± 0.15 for the as-built condition, as illustrated in Fig. 2 (b). According to Dass et al. [ 24 ], spherical pores (those smaller than 50 µm) and irregularly shaped pores (those larger than 50 µm) are linked to gas entrapment and lack of fusion, respectively. The calculated volumetric porosity of 0.37% falls within the low porosity range and can be achieved under optimal manufacturing conditions in DED. These findings indicate that the processing parameters defined with an energy density of 104 J/mm³ result in high-quality material. In our previous study [blinded], the microstructural analysis of the as-built IN718 samples revealed predominantly dendritic, a characteristic commonly observed in additively manufactured components, especially those produced by DED-LB [ 25 ]. These structures consist of elongated columnar grains with finer, equiaxed grains in a bimodal manner. This dual morphology is typical of rapid solidification conditions inherent to the process. The dendritic morphology and potential strings of the Laves phase were more pronounced in the interdendritic regions. Additionally, structures resembling δ phase and MC carbides were found, which align with features typically formed due to microsegregation of alloying elements during solidification. These findings are consistent with literature reports of dendritic cellular structures resulting from fast thermal cycles [ 26 ]. The formation of δ and Laves phases, driven by microsegregation, is a well-documented phenomenon in as-built IN718. These phases are detrimental as they deplete the matrix of Nb, an essential element for the precipitation of γ’ and γ’’, the primary strengthening phases [ 27 ]. Figure 3 (a) shows the band contrast map superimposed with grain boundaries as black lines for the as-built microstructure, i.e., before the heat treatments. It illustrates a microstructure composed of grains with varying morphologies and sizes, ranging from 3 µm to 200 µm. The grains are predominantly oriented along the building direction, leading to an average grain size of 80 ± 50 µm (see Fig. 3 (b)). Regarding grain boundary characteristics, the distributions of misorientation angles and axes reveal a random orientation for LAGBs, comprising 21% of the total boundaries. In contrast, HAGBs show a transition toward the \(\:⟨001⟩\) direction at misorientation angles greater than 30˚ (refer to Fig. 3 (c)). The significant prevalence of LAGB in the as-built condition is attributed to the high residual stress generated during the manufacturing process. These stresses result in a high density of dislocations caused by the compression and tensile residual stresses within the melt pools [ 28 ], [ 29 ]. 3.1.2 Microstructure of heat-treated samples EBSD assemblies extracted from the heat-treated samples can be seen in Figs. 4 and 5 . It can be seen from Figs. 4 a, d, and g that the DA and both solubilized conditions, S and SDA, still present columnar grain morphology, similar to as-built samples. On the other hand, the homogenized conditions, namely H, HDA, and HSDA (Fig. 5 a, d, and g, respectively), present annealing twins and recrystallized grains. It can be noticed that the misorientation features change from solubilization and aging only treatments when compared to homogenization treatments, Figs. 4 and 5 b), e), and h). The average misorientations of the calculated were approximately 30.6°, 28.5°, and 28.2°, with a low-angle grain boundary (LAGB) percentage of 19.5, 28.4, and 32.4% for DA, S, and SDA samples, respectively. When compared to the homogenized counterparts, it can be noticed that the H, HDA, and HSDA present higher values of misorientation angles, 41.3°, 47.4°, and 48.2°, respectively, and lower quantities of LAGBs: 19.0, 7.8, and 9.6% for H, HDA, and HSDA, respectively. The difference in the average misorientation and percentage of LAGBs can be explained by the recrystallization process promoted by the homogenization treatment. The quantity of LAGBs is generally lower in metallic materials that have undergone recrystallization. This occurs because recrystallization involves the formation of new deformation-free grains with HAGBs, which replace the deformed structure and the LAGBs present in the material before the process [ 30 ]. During recrystallization, the energy stored in the material due to plastic deformation is released by forming new grains. These grains tend to exhibit more random crystallographic orientations, resulting in HAGBs. LAGBs, which are characteristic of substructures within deformed grains, are either eliminated or significantly reduced during this process [ 30 ]. This change is evident in the distribution of misorientation axes, where the average misorientation increases for the homogenized materials. This results in a significant shift in the texture of the grain boundaries due to the formation of \(\:\left\{111\right\}⟨111⟩\) recrystallization twins, which occur at a misorientation of 60°. The evolution in grain morphology and grain boundary orientation transits from microstructures primarily characterized by low-angle grain boundaries (LAGB) in the solubilized conditions to those dominated by recrystallization twins after the homogenization heat treatment. This remarkable versatility of the alloy in producing a wide range of properties is a testament to its potential to create different mechanical properties. Furthermore, it can be noticed when comparing DA with the other heat treatments that there is a microtexture change, deviating from the almost random orientation shown by the DA sample to a more oriented grain boundaries, as can be seen by the Mackenzie curves in Figs. 4 and 5 b, e, and h plots. From a grain size perspective, the DA, S, and SDA samples exhibit similar characteristics, featuring coarse columnar grains. In contrast, the homogenized sample reveals recrystallized equiaxed and bimodal grains within the observed field, with grain sizes ranging mostly from approximately 10 to 240 µm for the H and HDA samples, with a slight increase in grain size for the HSDA sample, with a grain size ranging from approximately 10 to 290 µm. The recrystallization may be attributed to the non-uniform distribution of residual stresses in the as-fabricated samples. According to Humphreys et al. [ 30 ], the static recrystallization temperature of wrought IN718 is approximately 1020°C, which is lower than the 1100°C used in the homogenization treatments of this study. According to the literature, DED processed IN718 exhibits a higher recrystallization temperature than its wrought counterpart, primarily due to the lower residual strain and the high supersaturation of solute atoms in the as-built IN718 produced by DED. In general, wrought IN718 undergoes significant plastic deformation during forging, leading to a considerable accumulation of residual strain energy. As a result, the driving force for recrystallization in wrought IN718 is greater than that in DED-produced IN718, as reported elsewhere [ 31 ]. Furthermore, the fast-cooling rates inherent to the DED-LB process significantly enhance the supersaturated solid solubility of alloying elements in the γ matrix, contributing to increased recrystallization temperature [ 32 ]. Heat treatments significantly influence the microstructure of this alloy, as reflected in the distribution, orientation, and density of dislocations. Figure 6 displays geometrically necessary dislocations (GND) maps corresponding to the various heat treatments. These types of dislocations are linked to the curvature of the grains, which is influenced by the manufacturing process—whether through plastic deformation or the extreme additive manufacturing conditions, where significant residual stresses coming from high cooling rates and repetitive fusion between adjacent layers are present [ 33 ], [ 34 ], [ 35 ]. In this context, two behaviors are observed following the various heat treatments: first, heat treatments involving a single stage do not impact the GNDs magnitude compared to the as-built condition (see Figs. 6 a-d). In contrast, heat treatments that include multiple thermal cycles reduce the magnitude of dislocations (see Figs. 6 e-g). The significant decrease in GNDs observed in the SDA, HAD, and HSDA conditions aligns well with the previously described microstructural changes, where grain size and grain boundary misorientations underwent substantial alterations due to temperature-induced recrystallization. The analysis of dislocation orientations, represented by the weighted Burgers vector orientation triangle, indicates that dislocations are primarily aligned between the \(\:\left[001\right]\) and \(\:\left[101\right]\) directions. In the case of Ni, which has an FCC crystal structure, close brace open angle bracket 110, along the open brace 111, close brace open angle bracket 110, along the \(\:\left\{111\right\}⟨110⟩\) slip systems. Thus, most dislocations correspond to the ½ \(\:\left[101\right]\:\) type, as observed in the as-built and homogenized conditions. However, different heat treatments show a more significant variability in dislocation orientations. This scattering can be attributed to the formation of second phases, such as precipitates or carbides, which generate dislocation tangles and distortions in the lattice. Figure 7 presents the XRD spectra of both as-built and heat-treated samples. In Fig. 7 a, two distinct phases—γ and δ—are identified, along with fluorescence peaks attributed to W, In, and Pb emissions, likely originating from the shielding walls and devices within the synchrotron hutch. Notably, only the H and S heat treatments resulted in a single-phase structure, whereas the as-built condition and all other heat treatment combinations led to the formation of the δ phase. It should be noted that the XRD analysis could not detect the presence of γ’ and γ” phases, which is already expected due to their expected low volume fraction, which ranges from less than 1% to approximately 11%, depending on the employed HT conditions [ 36 ], small size, coherence with the matrix, and superimposition of diffraction peaks [ 37 ]. The as-built IN718 sample exhibited a higher relative intensity for the {200}γ peak than the {111}γ peak. This phenomenon is associated with the preferential growth along the orientation, which is influenced by the significant thermal gradient present along the build direction [ 38 ], [ 39 ]. The XRD patterns of the HT samples revealed subtle differences among {111}γ and {200}γ peaks. Additionally, minor peaks were observed between the {111}γ and {200}γ peaks, indicating the formation of the δ-phase [ 40 ], [ 41 ]. In contrast, the HSDA sample exhibited a {111}γ peak with higher intensity than the {200}γ peak, probably due to the grain growth or recrystallization during the HT at 1100°C [ 42 ]. Furthermore, the XRD spectra of the DA samples suggests the development of a more random orientation, akin to that observed in Mackenzie curve from Fig. 3 c. Moreover, the XRD spectrum of the other HT samples also suggests a more randomized crystallographic orientation, as the relative intensity of matrix peaks follows the calculated values. Detailed diffraction patterns in the vicinity of the {200}γ peaks for all samples are also presented in Fig. 7 . The evolution of the γ{200} peak with the complexity of the heat treatment, i.e., following the sequence of conditions: as-built → DA → SDA → HDA → HSDA, is almost none, indicated by the dashed red line, except the HSDA HT, which shifts to lower interplanar spacing. It indicates a reduction in the lattice parameter of the γ matrix. This shift is primarily attributed to the precipitation of the γ″ phase, which reduces the concentration of dissolved Nb in the γ matrix. However, distinguishing their contributions in the XRD pattern proved challenging due to the issues already mentioned [ 43 ]. Figure 8 shows thermodynamic calculations considering either the non-equilibrium (Scheil-Gulliver) or the thermodynamic equilibrium (phase diagram) for IN718 alloy. δ, γ’, γ” were considered one phase in the calculations due to their similar chemical composition, and they are called Ni₃(Nb, Ti). The solidification path predicted for the present composition of IN718 alloy is as follows: Liquid (L)→ L + γ → L + γ + NbC → L + γ + NbC + Laves + Ni 3 (Nb, Ti) → L + γ + NbC + Laves + Ni 3 (Nb, Ti) + σ. The non-equilibrium thermodynamic conditions inherent to additive manufacturing promote the formation of metastable phases. The Scheil–Gulliver solidification model accurately predicts transition temperatures, revealing Laves phase formation between 1180–1185°C. Below 1180°C, chemical element stoichiometry enables Ni 3 Nb phase development. Microsegregation introduces complexity, with low-solubility elements like Nb, Mo, Ti, and C concentrating in interdendritic regions and forming MC carbides and Laves phases. This segregation can induce variations in precipitation temperatures, highlighting the intricate microstructural evolution during DED-LB processing. On the other hand, phase equilibrium calculations show an insignificant presence of NbC, besides a considerable presence of σ phase at low temperature, which continues to dissolve until 925°C, and Ni 3 (Nb, Ti), achieving its complete dissolution at 1030°C. The phase equilibrium calculation at each heat treatment temperature condition is depicted in Table 1 . Table 1 Phase volume percent at heat treatment temperatures under equilibrium conditions. Temperature (°C) Phases (in volume% %) NbC γ Ni 3 (Nb,Ti) σ 620 0.1 73.8 14.8 11.2 720 0.1 77.3 14.0 8.6 1000 0.1 96.9 2.9 - 1100 0.1 99.9 - - Intermetallic phase formation in DED-LB manufacturing of IN718 alloy emerges from complex interactions between kinetic processes, thermal history, and elemental availability. The alloy demonstrates coexistence of stable (δ) and metastable (γ″) phases with Ni 3 Nb stoichiometry, with phase prevalence critically dependent on temperature gradients and elemental distribution [ 36 ], [ 44 ]. Temperatures exceeding 600°C trigger MC carbide precipitation, predominantly NbC, with carbide growth dependent upon Nb concentration in δ and γ″ phases. The metastable γ″ phase is the primary strengthening precipitate, significantly enhancing mechanical properties. However, prolonged high-temperature exposure induces transformation to stable δ phase, potentially compromising material strength [ 36 ], [ 44 ]. Figure 9 shows the SEM and EDS analysis for all the aged samples. Microstructural comparisons of as-built and DA samples revealed minimal morphological distinctions through SEM micrographic analysis. Characteristic precipitates, specifically Laves phase, were consistently observed in both conditions, suggesting that the applied aging temperatures were insufficient to induce complete dissolution of the Laves phase, Fig. 9 a. Figure 9 b depicts the SDA sample microstructure, showing that the cellular microstructure persisted. It can be seen that localized grain growth occurs compared to the as-built condition. From the EDS analysis, shown as a table on the side of the micrographs, the possible identification of those phases as δ precipitates and carbides can be seen. Homogenization treatments, namely HDA and HSDA, Figs. 9 c and 9 d, respectively, effectively eliminated melt pool boundaries and elemental segregation. Significant microstructural transformations were observed, including extensive recrystallization and grain growth, with grain sizes exceeding 200 µm in a bimodal fashion, as shown in Fig. 5 . Annealing twins were detected, and grain boundaries exhibited small white particles (Fig. 9 c), tentatively identified as carbides; Fig. 9 d also shows the presence of carbides and possibly δ precipitates inside the grains. 3.2 Mechanical Properties Hardness test results for the DED-LB and the forged materials are shown in Fig. 10 . The DA heat treatment enhanced the material's mechanical strength by approximately 50% compared to non-aged samples, reaching values of 475.7 HV, 503.5 HV, 506.1 HV, and 483.47 HV for DA, HDA, SDA, and HSDA, respectively. The authors believe that it occurs due to the precipitation of the strengthening phases of the IN718 alloy, i.e., γ’ and γ’’, which could not be detected by the XRD analysis. Furthermore, it can be noticed that the H and S heat treatments could solubilize the alloy, reducing its hardness values. Although similar to the forged IN718, the hardness values are slightly superior to the DED-LB values. Forged materials generally exhibit higher hardness than additively manufactured materials due to differences in microstructure, defect density, and residual stress distribution. Forging induces severe plastic deformation, leading to dynamic recrystallization, grain refinement, and increased dislocation density, which enhances hardness through the Hall-Petch effect [ 45 ]. In contrast, AM processes often result in columnar and dendritic microstructures with lower dislocation density due to rapid solidification and lack of plastic deformation. Additionally, AM components tend to exhibit higher porosity and process-induced defects, reducing mechanical performance compared to the dense and homogeneous microstructure of forged materials [ 46 ]. Residual compressive stresses in forged materials further contribute to their superior hardness, whereas AM components often require post-processing heat treatments to mitigate tensile residual stresses and improve mechanical properties [ 46 ]. While AM advancements are narrowing this gap, forged materials generally maintain superior hardness due to their refined microstructure and defect-free nature. This effect can be seen for the as-built and AMS 5662 forged conditions in Fig. 10 . Figure 11 a presents the engineering stress-strain curves for all analyzed conditions. Upon initial observation, there is a wide range of combinations of strength and ductility, with yield strengths varying from 300 MPa to 1200 MPa and fracture strains ranging from 0.5 to 0.35. Single-stage heat treatments, such as homogenization and solubilization, result in the highest ductility. In contrast, thermal cycles that involve aging heat treatments, such as DA, SDA, HDA, and HSDA, yield greater strengths than the as-built condition. The strain hardening rate curves demonstrate that ductility is enhanced following the S and H heat treatments. In contrast, the aged conditions exhibit reduced ductility, which leads to localized plastic instability during early deformations, as shown in Fig. 11 b. This loss of ductility in the aged conditions can be attributed to the large number of precipitates at the grain boundaries, as confirmed in Fig. 9 . The formation of precipitates at grain boundaries can increase brittleness, particularly when the particles are incoherent and large. Studies have shown that the coarsening of γ’ precipitates after heat treatment follows the Ostwald ripening mechanism. This process results in a loss of ductility due to strain concentration caused by the incoherent precipitates localized at the grain boundaries [ 47 ], [ 48 ]. The strain profiles observed during tensile tests at various deformations reveal different behaviors depending on the applied heat treatment. After the HDA heat treatment, the strain distribution along the calibrated length shows significant heterogeneity from the start until the onset of plastic instability (see Fig. 11 c). In the as-built condition, the strain profiles exhibit multiple peaks, which indicate strain localization at different points along the sample length, thereby delaying the formation of plastic instability (see Fig. 11 d). In contrast, the homogenized condition displays strain profiles with less heterogeneity than the ones mentioned earlier. These profiles resemble a plateau, facilitating the storage of larger and more uniform deformations before the plastic instability (Fig. 11 e). The behaviors discussed are closely related to the various microstructures described. For instance, the early onset of strain localization in the HDA can be attributed to the brittleness caused by the coarse precipitates at the grain boundaries. In contrast, the multiple peaks observed in the as-built condition can be linked to the residual stresses arising from the manufacturing process, which are repeated between adjacent melt pools. In the homogenized condition, the more uniform strain distribution is due to the absence of precipitates and the larger, more equiaxed grain size resulting from the recrystallized microstructure, which allows for greater dislocation motion. Therefore, a positive synergy between strength and ductility in the studied alloy is achieved by an adequate distribution of coherent and intragranular γ’ precipitates, which impart strength while controlling the number of precipitates at the grain boundaries to avoid ductility loss. Figure 12 presents the room-temperature mechanical properties of the DED-LB processed IN718 alloy in both the as-built state and after six distinct heat-treatment conditions. The yield strength values for the as-built, DA, H, S, HDA, SDA, and HSDA conditions were determined to be approximately 638 MPa, 1195 MPa, 360 MPa, 412 MPa, 1037 MPa, 1189 MPa, and 1046 MPa, respectively. Additionally, the elongation values for these conditions were measured as ≈ 28.5%, 8.0%, 49.3%, 47.1%, 17.8%, 12.1%, and 16.7%, respectively. To characterize the plastic deformation behavior observed during tensile testing, the Hollomon equation (σ = K⋅ε n , where σ represents the true stress, K is the strength coefficient, ε denotes the true strain, and n is the strain-hardening exponent) was employed. The calculated n values for the DA, H, S, HDA, SDA, and HSDA samples were approximately 0.072, 0.281, 0.25, 0.085, 0.06, and 0.082, respectively. These results indicate that the H and S conditions exhibit the highest uniform plastic deformation capacity, whereas the aged conditions display increased hardness and reduced deformability. The elastic modulus (E) values were determined to be ≈ 203 GPa for the as-built sample, 211 GPa, 197 GPa, 203 GPa, 214 GPa, 210 GPa, and 220 GPa for the DA, H, S, HDA, SDA, and HSDA conditions, respectively. This suggests that the HSDA sample demonstrates the greatest resistance to elastic deformation. The presence of γ’/γ” strengthening phases, which develop during the DA process, contributes to an increase in tensile strength, explaining why the homogenized and solubilized conditions, lacking these phases, exhibit comparatively lower strength. Among the evaluated heat treatment conditions, direct aging results in the highest tensile strength; however, it also leads to the lowest ductility. In contrast, solution annealing provides slightly higher tensile strength than homogenization solution aging, though at the cost of reduced ductility, while maintaining comparable overall tensile properties [ 49 ]. 4. Conclusions The study investigated the effects of six different heat-treatment routes on the microstructural evolution, hardness, tensile properties, and fracture behavior of DED-LB IN718 samples. The relationship between microstructures and mechanical responses was discussed and compared to the forged IN718 counterpart. The conclusions reached are: 1. The results showed that different heat treatment strategies can tailor the microstructure and mechanical properties of DED-LB IN718 samples. 2. Aging treatments (DA, SDA, HDA, and HSDA) increased hardness due to the precipitation of strengthening phases. 3. Homogenization treatments (H, HDA, and HSDA) promoted recrystallization and grain growth, leading to a more uniform strain distribution and enhanced ductility. 4. The homogenized condition displayed strain profiles with less heterogeneity, facilitating the storage of larger and more uniform deformations before the plastic instability. 5. A balance between strength and ductility was achieved through the adequate distribution of coherent and intragranular γ’ precipitates, while controlling the number of precipitates at the grain boundaries to avoid ductility loss. Declarations Acknowledgments BLINDED MANUSCRIPT Funding BLINDED MANUSCRIPT Data availability . BLINDED MANUSCRIPT Conflict of interest statement. 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. Author Contribution T. R. F. Cavalcante was responsible for writing the manuscript, editing the text, and preparing all figures. D. G. Bon conducted the mechanical testing and contributed to the analysis and interpretation of mechanical behavior. J. A. Muñoz performed the digital image correlation (DIC) analysis and contributed to data interpretation. G. G. Ribamar, J. P. Oliveira, and A. B. Pereira carried out the simulation work and participated in the X-ray diffraction (XRD) analysis. F. E. Mariani contributed to the experimental design and supervised the additive manufacturing process. J. C. Muñoz, R. T. Coelho and J. A. A. Diaz provided project supervision, critical review of the manuscript, and guidance throughout the research. All authors discussed the results, contributed to the final version of the manuscript, and approved its submission. Acknowledgement AcknowledgmentsThe support of the Center for Research and Innovation in Materials and Structures (CEPIMATE) is deeply appreciated. This research used facilities of the Brazilian Nanotechnology National Laboratory (LNNano), part of the Brazilian Centre for Research in Energy and Materials (CNPEM), a private non-profit organization under the supervision of the Brazilian Ministry for Science, Technology, and Innovations (MCTI). The Electron Microscopy Laboratory staff are acknowledged for their assistance during the experiments (proposal SEM-FIB-C1-20233619). We acknowledge DESY (Hamburg, Germany), a member of the Helmholtz Association HGF, for providing the experimental facilities. Parts of this research were carried out at PETRA III (proposal I-20230101 EC), and we would like to thank Dr. Guilherme Abreu Faria and Dr. Marc-André Nielsen for their assistance in using beamline P61A. Funding Funded by Fundação de Amparo à Pesquisa do Estado de São Paulo – FAPESP, grant No. 2020/09079-2. This study was partly financed by the Conselho Nacional de Desenvolvimento Científico e Tecnológico – CNPq, grant No. 306960/2021–4. Thiago Roberto Felisardo Cavalcante recognizes the financial support through the Ph.D. scholarship from the Coordenação de Aperfeiçoamento de Pessoal de Nível Superior – Brasil (CAPES) – Finance Code 001 and CAPES PrInt, process 88887.886782/2023-00. JPO acknowledges funding by national funds from FCT – Fundação para a Ciência e a Tecnologia, I.P., in the scope of the projects LA/P/0037/2020, UIDP/50025/2020, and UIDB/50025/2020 of the Associate Laboratory Institute of Nanostructures, Nanomodelling and Nanofabrication – i3N. The present study was in part developed in the scope of the Project “Agenda ILLIANCE” [C644919832-00000035 | Project nº 46], financed by PRR – Plano de Recuperação e Resiliência under the Next Generation EU from the European Union.Data availability. The data supporting this study’s findings are available from the corresponding author, J. A. Avila, upon reasonable request. 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Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-6602510","acceptedTermsAndConditions":true,"allowDirectSubmit":true,"archivedVersions":[],"articleType":"Research Article","associatedPublications":[],"authors":[{"id":461272154,"identity":"244a338f-c059-48e2-aea8-1121149b041a","order_by":0,"name":"Thiago Roberto Felisardo Cavalcante","email":"","orcid":"","institution":"Universidade de São Paulo","correspondingAuthor":false,"prefix":"","firstName":"Thiago","middleName":"Roberto Felisardo","lastName":"Cavalcante","suffix":""},{"id":461272155,"identity":"54a0543a-e05d-47d0-9880-f719c5beeff3","order_by":1,"name":"Douglas Giovanni Bon","email":"","orcid":"","institution":"Universidade de 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tests\u003c/p\u003e","description":"","filename":"1.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/2d128f58fe3ca2e4421bd237.jpg"},{"id":83437231,"identity":"397e27dd-4054-4a1d-a2ac-0e71b9fa078f","added_by":"auto","created_at":"2025-05-26 08:40:59","extension":"jpg","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":80692,"visible":true,"origin":"","legend":"\u003cp\u003e(a) Porosity distribution in the volume analyzed and (b) porosity and sphericity distributions for the as-built condition.\u003c/p\u003e","description":"","filename":"2.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/65e36127d6648676b06b1233.jpg"},{"id":83437234,"identity":"e2b1b964-51a2-4e61-9dff-6623861b9a43","added_by":"auto","created_at":"2025-05-26 08:40:59","extension":"jpg","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":125526,"visible":true,"origin":"","legend":"\u003cp\u003ea) General view of as-built microstructure; b) Grain size distribution of as-built sample, and c) misorientation angles of as-built samples.\u003c/p\u003e","description":"","filename":"3.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/e50fa203371370355af732d2.jpg"},{"id":83437443,"identity":"8a271674-07f7-42df-8a67-87e188e27b7c","added_by":"auto","created_at":"2025-05-26 08:49:00","extension":"jpg","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":218359,"visible":true,"origin":"","legend":"\u003cp\u003eEBSD results for: DA sample in the first line, S sample in the second line, and SDA condition in the third line. The first column shows the IPF map superimposed with the band contrast map. The second column shows semiquantitative and quantitative misorientation analysis. The third column shows semiquantitative grain area distribution.\u003c/p\u003e","description":"","filename":"4.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/376d2b91a6b2379440229758.jpg"},{"id":83437439,"identity":"1d4b3e7f-03e4-4933-bde3-e8f689271c99","added_by":"auto","created_at":"2025-05-26 08:49:00","extension":"jpg","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":224460,"visible":true,"origin":"","legend":"\u003cp\u003eEBSD results for: H sample in the first line, HDA sample in the second line, and HSDA condition in the third line. The first column shows the IPF map superimposed with band contrast map; the second column shows semiquantitative and quantitative misorientation analysis; the third column shows semiquantitative grain area distribution.\u003c/p\u003e","description":"","filename":"5.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/eb458040a72bddaf1fd5fe3d.jpg"},{"id":83437442,"identity":"a86d6d4c-62f2-4b3f-beb3-5dcc5458050d","added_by":"auto","created_at":"2025-05-26 08:49:00","extension":"jpg","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":196519,"visible":true,"origin":"","legend":"\u003cp\u003eGeometrically necessary dislocations (GNDs) maps and the weighted Burgers vector (WBV) representation for the different analyzed conditions. a) AB, b) DA, c) S, d) H, e) HDA, f) HSDA, and g) GND distributions.\u003c/p\u003e","description":"","filename":"6.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/1d0b4c56a4eebab9ea020b4e.jpg"},{"id":83437237,"identity":"dd14f68f-84f7-4330-8485-68bf5da255b6","added_by":"auto","created_at":"2025-05-26 08:40:59","extension":"jpg","order_by":7,"title":"Figure 7","display":"","copyAsset":false,"role":"figure","size":64940,"visible":true,"origin":"","legend":"\u003cp\u003eEnergy-dispersive X-ray diffraction pattern of the IN718 fabricated via DED-LB in the as-built and heat-treated conditions, along with an in-depth view of the {200}γ peak.\u003c/p\u003e","description":"","filename":"7.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/77e6667e7d8af0634feaa43f.jpg"},{"id":83437247,"identity":"5a4eb034-1771-4777-a66d-dad32b57ccb1","added_by":"auto","created_at":"2025-05-26 08:41:00","extension":"jpg","order_by":8,"title":"Figure 8","display":"","copyAsset":false,"role":"figure","size":58355,"visible":true,"origin":"","legend":"\u003cp\u003ea) Solidification path of IN718 alloy, predicted by Thermo-Calc software using Scheil-Gulliver solidification mode; b) Phase volume fraction of the DED-LB IN718 phases.\u003c/p\u003e","description":"","filename":"8.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/8ae842f6392e51a4f0ac0f85.jpg"},{"id":83437441,"identity":"8650e052-d1bd-4f55-b9f1-703bd082f5cd","added_by":"auto","created_at":"2025-05-26 08:49:00","extension":"jpg","order_by":9,"title":"Figure 9","display":"","copyAsset":false,"role":"figure","size":99492,"visible":true,"origin":"","legend":"\u003cp\u003eScanning electron microscopy images, EDS measurements, and EBSD band contrast maps with HAGB as black lines for a) DA, b) SDA, c) HDA, and d) HSDA HT conditions.\u003c/p\u003e","description":"","filename":"9.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/750ceb5d0c1c56856d3db6f7.jpg"},{"id":83438174,"identity":"63467b0f-95f9-43a1-b9b2-3b2096ae2569","added_by":"auto","created_at":"2025-05-26 08:57:00","extension":"jpg","order_by":10,"title":"Figure 10","display":"","copyAsset":false,"role":"figure","size":46027,"visible":true,"origin":"","legend":"\u003cp\u003eVickers Hardness results for the AB, forged, and HT conditions.\u003c/p\u003e","description":"","filename":"10.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/681324195d0cf6e472a838f1.jpg"},{"id":83438173,"identity":"63745233-f216-40b4-bdc2-738cb91e562a","added_by":"auto","created_at":"2025-05-26 08:57:00","extension":"jpg","order_by":11,"title":"Figure 11","display":"","copyAsset":false,"role":"figure","size":116868,"visible":true,"origin":"","legend":"\u003cp\u003ea) Engineering stress-strain curves for all HT conditions; b) Strain hardening rate curves for all HT conditions; c) Strain hardening of the HDA condition, d) Strain hardening of the AB condition and e) Strain hardening of the H condition.\u003c/p\u003e","description":"","filename":"11.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/48471ed1df029a1a49a0b7c0.jpg"},{"id":83437444,"identity":"5b9bfd8a-2f05-479d-b6da-13e6115d368c","added_by":"auto","created_at":"2025-05-26 08:49:00","extension":"jpg","order_by":12,"title":"Figure 12","display":"","copyAsset":false,"role":"figure","size":32729,"visible":true,"origin":"","legend":"\u003cp\u003eTensile test results for the heat-treated DED-LB IN718 samples.\u003c/p\u003e","description":"","filename":"12.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/1d1f071fb90f656284121214.jpg"},{"id":88814246,"identity":"fd50794c-ca9e-4287-af84-98e4dd877241","added_by":"auto","created_at":"2025-08-11 16:09:02","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":2142189,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/31693f3a-4baf-4877-ba03-0b33fb5bf569.pdf"},{"id":83437242,"identity":"285955b8-bcc4-4cf2-bd4d-c55d5bf885f9","added_by":"auto","created_at":"2025-05-26 08:41:00","extension":"jpg","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":60683,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eGraphical Abstract\u003c/strong\u003e\u003c/p\u003e","description":"","filename":"GraphicalAbstract.jpg","url":"https://assets-eu.researchsquare.com/files/rs-6602510/v1/b90f2087956208eb3fe8d9db.jpg"}],"financialInterests":"No competing interests reported.","formattedTitle":"Microstructural assessment of additive-manufactured Inconel 718 samples subjected to heat treatments for enhanced mechanical properties","fulltext":[{"header":"1. Introduction","content":"\u003cp\u003eRecently, laser additive manufacturing (AM) techniques have been extensively researched and are in a state where they could be considered potential candidates to replace traditional manufacturing methods when needed [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e]. Laser beam directed energy deposition (LB-DED) is a type of AM that deposits thin layers of metal powder and melts them using a high-energy laser beam; furthermore, it can deposit metallic materials and functional coatings. LB-DED has the advantages of high deposition efficiency, flexibility, and precision [\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e]. However, due to the rapid cooling of the process, residual stresses, micro-segregation, and the formation of non-equilibrium phases are a common feature in as-built components [\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eNi-based superalloys, such as Inconel 718 (IN718), are heat-treatable, hot corrosion resistant, and preferred for the fabrication of aerospace engines, turbine blades, injectors, and many other high temperatures and/or high-pressure components [\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e], [\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e]. This alloy, specifically, is mainly strengthened by precipitation hardening, via the precipitation of γ\u0026rsquo; and γ\u0026rsquo;\u0026rsquo; phases [\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e]. Moreover, the solid solution hardening effects of some elements in the γ matrix, such as Nb and Mo, can be additional factors to improve the mechanical properties. The presence of micron/nano-sized γ\u0026rsquo; and γ\u0026rsquo;\u0026rsquo; phases controls the mechanical properties of the IN718 alloy. On the other side, deleterious phases, such as δ and Laves phases, can degrade the materials' properties, either due to the incoherent nature of the phase with the matrix, or the inhibition of γ\u0026rsquo; or γ\u0026rsquo;\u0026rsquo; precipitation [\u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eThe micro-segregation of elements and the presence of deleterious phases, typical in AM components, make the elemental distribution uneven, which also impacts the precipitation of γ\u0026rsquo; and γ\u0026rsquo;\u0026rsquo; phases [\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e]. Laves phase forms in the interdendritic regions and is known to deteriorate mechanical properties such as strength, ductility, fatigue, and creep, since it depletes essential precipitation elements, besides facilitating crack initiation and propagation due to its brittle nature [\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e], [\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e]. Applying heat treatments (HT) and controlling the solidification conditions during the deposition are two ways to reduce the micro-segregation and formation of the Laves phase. It has been reported that Nb segregation and Laves phase formation can be reduced by increasing the cooling rate from the liquid material [\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e]. However, controlling the solidification conditions is rarely achievable due to the complicated solidification of the AM processes [\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e]. As such, post-fabrication HTs are vital to achieve excellent mechanical performance and meet service requirements through the elemental distribution homogenization and precipitation of the strengthening phases [\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eThe precipitation hardening heat treatment involves the dissolution of the Laves phase, since it releases Nb and Ti back to the matrix [\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e]. Generally, a two-step direct aging (DA) and solubilization (S) are performed, and, in some cases, a homogenization step is also applied. The solubilization typically occurs in the range of 950 to 1100 \u003csup\u003eo\u003c/sup\u003eC to dissolve the Laves phase and release Nb and Ti into the matrix; direct aging usually ranges from 600 to 800 \u003csup\u003eo\u003c/sup\u003eC and results in the formation of γ\u0026rsquo; and γ\u0026rsquo;\u0026rsquo; precipitates. Furthermore, homogenization heat-treatments (H) can also eliminate Laves and microsegregation, enhancing the homogeneous precipitation of strengthening phases [\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eSeveral studies have been conducted regarding the influence of HTs on AM IN718. Qi et al. [\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e] reported that the microstructure and tensile properties of IN718 manufactured by DED-LB with DA achieved high tensile strength values with low ductility. Huang et al. [\u003cspan citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e] showed that, for alloys fabricated via PBF-LB process, different solution treatment temperatures, ranging from 980 to 1280 oC for 1h, implied different Laves phase dissolution degrees. Burad et al. [\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e] investigated the effect of S and S\u0026thinsp;+\u0026thinsp;DA HTs and concluded that solution treatment at 980 \u003csup\u003eo\u003c/sup\u003eC, to dissolve the Laves phase and microsegregation spots, led to a good precipitation quality. Zhai et al. [\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e] found that a simple aging process could precipitate an amount of γ\u0026rsquo; and γ\u0026rsquo;\u0026rsquo; phases capable of enhancing the mechanical properties of IN718, with weak microsegregation and fine Laves phase still present. However, it could not change the typical columnar dendrite structure. Zhao et al. [\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e] demonstrated that a homogenization temperature of 1180 \u003csup\u003eo\u003c/sup\u003eC can eliminate residual stress and promote recrystallization of IN718 manufactured by LB-PBF. Regarding the Laves phase, Sui et al. have studied its dissolution [\u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e], who investigated the effect of a 1050 \u003csup\u003eo\u003c/sup\u003eC solubilization with DA on LB-DED IN718 samples, showing a large improvement in high-temperature mechanical behavior. Xu et al. [\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e] studied the microstructural evolution and the mechanical performance of AMed IN718 under three HTs: DA, S\u0026thinsp;+\u0026thinsp;DA, and H\u0026thinsp;+\u0026thinsp;S\u0026thinsp;+\u0026thinsp;DA, and concluded that different as-built microstructures require different HTs; a similar study was also performed by Yu et al. [\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e], who focused on the fracture behavior of the LB-DED IN718 deposited material. Furthermore, as-built microstructures, according to Jang et al. [\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e] studied the precipitation kinetics of secondary phases formed by the heat accumulation during the IN718 DED deposition.\u003c/p\u003e \u003cp\u003eThis study investigates the effects of six different heat treatment strategies on the microstructural evolution, hardness, tensile properties, and fracture behavior of IN718 samples produced by DED-LB. Special focus is placed on the influence of these treatments on the precipitation of strengthening phases and the dissolution of deleterious phases such as the Laves phase. The mechanical performance of the treated AMed samples is critically compared to that of their forged counterparts to assess the effectiveness of post-processing in mitigating typical additive manufacturing defects. The findings aim to support better heat-treatment routes for AMed IN718 alloys, promoting improved structural integrity and mechanical reliability for high-performance applications.\u003c/p\u003e"},{"header":"2. Experimental procedure","content":"\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e \u003ch2\u003e2.1 Materials\u003c/h2\u003e \u003cp\u003eFor this study, gas-atomized IN718 powder, produced by Carpenter Additive Inc., was used. The powder's chemical composition, expressed in weight percentage (wt.%), was as follows: Ni 52.30, Cr 19.72, Nb\u0026thinsp;+\u0026thinsp;Ta 4.96, Mo 3.30, Co 0.53, Ti 0.39, Si 0.09, Mn 0.04, C 0.01, P 0.01, with Fe in balance.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec4\" class=\"Section2\"\u003e \u003ch2\u003e2.2 Materials processing\u003c/h2\u003e \u003cp\u003eThe depositions were conducted on AISI 1020 steel with a thickness of 5 mm, which was end-milled and cleaned before the experiments. A BeAM Inc. Modulo 250 5-axis Directed Energy Deposition-Laser Beam (DED-LB) system was utilized, equipped with a 1 kW continuous-wave ytterbium fiber laser (Model YLR-1000-MM-WC-Y14, IPG Photonics Inc.), operating at a wavelength of 1070 nm and a focused spot size of 0.8 mm. The laser was mounted on a coaxial deposition head, and the system incorporated a dual powder feeder for material deposition. The parameters utilized for this work were a laser power of 500 W with a crosshead speed of 2000 mm/min and a feed rate of 5 g/min. These parameters were determined in a previous study, which evaluated various combinations of power, scanning speed, and powder feed rate to achieve high-integrity single tracks [\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e]. Detailed machine parameters, process schematics, and deposition parameters can be found in previous studies [blinded].\u003c/p\u003e \u003cp\u003eAll heat treatments (HT) were carried out in a Jung-brand muffle furnace (LF07013 model). The six HT routes employed in this study were designed to evaluate the effects of homogenization, solubilization, and aging on the microstructure and mechanical properties of the material. Th different kind of heat treatments used are: (i) homogenization (H), consisted of heating the samples to 1100\u0026deg;C for 1.5 h, followed by water quenching; (ii) solubilization (S), where the material is isothermally heat treated at 1000\u0026deg;C for 1.0 h, also followed by water quenching; and (iii) double aging (DA), consisted on an aging at 720\u0026deg;C for 8 h, followed by controlled cooling at 50\u0026deg;C/h down to 620\u0026deg;C, where the samples were held for an additional 8 h, and then air-cooled to room temperature. In addition to these isolated heat treatments, three combinations were also evaluated: homogenization followed by double aging (HDA), solubilization followed by double aging (SDA), and homogenization followed by solubilization and double aging (HSDA).\u003c/p\u003e \u003cp\u003eTo assess the porosity level after manufacturing, a Bruker Skyscan 1272 microtomograph was utilized. It features an 11 MP detector (4032 \u0026times; 2688 pixels) and operates at 10 W. The sample, measuring 2 mm \u0026times; 3 mm \u0026times; 8 mm, was scanned at a resolution of 2 \u0026micro;m with a step size of 0.3\u0026deg;. The 3D model was reconstructed and rendered using the Dragonfly\u0026trade; 2022.2 software developed by Comet Group. Volumetric porosity was then calculated by applying color thresholding to distinguish between the pores and the solid material.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec5\" class=\"Section2\"\u003e \u003ch2\u003e2.3 Microstructural characterization\u003c/h2\u003e \u003cp\u003eThe samples underwent standard preparation for metallographic evaluation, involving abrasive cutting, manual grinding with sandpapers, and sequential surface polishing using 1.00 and 0.25 \u0026micro;m diamond suspensions. Final etching was done using waterless Kalling\u0026rsquo;s reagent (100 ml ethanol, 100 ml HCl, 5 g CuCl₂) to reveal the microstructure. Cross-sectional microstructure observations were conducted with an Olympus LEXT 4100 3D laser microscope and a FEI\u0026reg; FE-50 Scanning Electron Microscope (SEM). Microhardness was assessed using a Buehler Model 1600 6300 tester by ASTM E384 [\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e], employing a 0.5 N load for 15 s, Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ed.\u003c/p\u003e \u003cp\u003eFurther microstructural details were examined via EBSD using a Quanta 650 FEG-SEM and high-speed acquisition EBSD system. Data were analyzed in TSL OIM and MTEX, with EBSD scans performed over 2000 \u0026micro;m \u0026times; 1500 \u0026micro;m (step size 1.5 \u0026micro;m) and 500 \u0026micro;m \u0026times; 500 \u0026micro;m (step size 0.5 \u0026micro;m). Maps achieved\u0026thinsp;\u0026ge;\u0026thinsp;90% indexing; non-indexed points were corrected using the grain dilation correction method, while grains with less than 2 pixels were excluded. GND densities were obtained via the Kernel Average Misorientation (KAM) and the Weighted Burgers Vector (WBV) [\u003cspan citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e] approaches. Boundaries were classified by misorientation: \u0026gt; 15\u0026deg; (high-angle grain boundaries\u0026mdash;HAGB), 5\u0026ndash;15\u0026deg; (low-angle grain boundaries\u0026mdash;LAGB), and 2\u0026ndash;5\u0026deg; (subgrain boundaries\u0026mdash;SGB).\u003c/p\u003e \u003cp\u003eSynchrotron energy-dispersive X-ray diffraction (EDXRD) was carried out at P61 High Energy wiggler beamline/LVP, installed at Petra III, Hamburg, Germany. Transmission-mode scans employed 10 wigglers (30\u0026ndash;200 keV range), with spectra collected along 10 mm line scans at 0.2 mm steps. Two HPGe detectors positioned at 2θCh0\u0026thinsp;=\u0026thinsp;7.557\u0026deg; and 2θCh1\u0026thinsp;=\u0026thinsp;7.594\u0026deg; captured data from a 0.1 \u0026times; 0.1 \u0026times; \u0026asymp;3 mm\u0026sup2; gauge volume. Exposure time per spectrum was 1 s. Raw data were processed using the in-house Python scripts.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec6\" class=\"Section2\"\u003e \u003ch2\u003e2.4 Mechanical Testing\u003c/h2\u003e \u003cp\u003eDigital Image Correlation (DIC) tensile testing was performed using a Deben Microtest Tensile Stage on dog-bone samples (gauge length: 4 mm \u0026times; 1 mm \u0026times; 1 mm) at 0.4 mm/min (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea and b). Load, displacement, and axial strain were captured via a Balser camera (2048 \u0026times; 2048 pixels at 75 fps) using 35-, 50-, and 75-mm lenses. DIC analyses were done using Ncorr to generate strain maps and profiles. Conventional tensile tests followed ASTM E8-24 [\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e], performed under displacement control (1 mm/min) on a 100 kN MTS Landmark N servo-hydraulic system. Specimens, extracted horizontally, are illustrated in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e"},{"header":"3. Results and discussion","content":"\u003cdiv id=\"Sec8\" class=\"Section2\"\u003e \u003ch2\u003e3.1 Microstructural characterization\u003c/h2\u003e \u003cdiv id=\"Sec9\" class=\"Section3\"\u003e \u003ch2\u003e3.1.1 Microstructure of as-built samples\u003c/h2\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e presents the porosity analysis for the as-built condition, conducted using X-ray diffraction microtomography. This analysis estimated a volumetric porosity of approximately 0.37%, with pore sizes ranging from 13 \u0026micro;m to 150 \u0026micro;m. The pores exhibited irregular morphologies aligned with the building direction, as evidenced by the different views of the analyzed solid shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e(a). The porosity size calculations resulted in an average pore size of 64\u0026thinsp;\u0026plusmn;\u0026thinsp;32 \u0026micro;m and a sphericity of 0.6\u0026thinsp;\u0026plusmn;\u0026thinsp;0.15 for the as-built condition, as illustrated in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e(b). According to Dass et al. [\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e], spherical pores (those smaller than 50 \u0026micro;m) and irregularly shaped pores (those larger than 50 \u0026micro;m) are linked to gas entrapment and lack of fusion, respectively. The calculated volumetric porosity of 0.37% falls within the low porosity range and can be achieved under optimal manufacturing conditions in DED. These findings indicate that the processing parameters defined with an energy density of 104 J/mm\u0026sup3; result in high-quality material.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eIn our previous study [blinded], the microstructural analysis of the as-built IN718 samples revealed predominantly dendritic, a characteristic commonly observed in additively manufactured components, especially those produced by DED-LB [\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e]. These structures consist of elongated columnar grains with finer, equiaxed grains in a bimodal manner. This dual morphology is typical of rapid solidification conditions inherent to the process. The dendritic morphology and potential strings of the Laves phase were more pronounced in the interdendritic regions. Additionally, structures resembling δ phase and MC carbides were found, which align with features typically formed due to microsegregation of alloying elements during solidification. These findings are consistent with literature reports of dendritic cellular structures resulting from fast thermal cycles [\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e]. The formation of δ and Laves phases, driven by microsegregation, is a well-documented phenomenon in as-built IN718. These phases are detrimental as they deplete the matrix of Nb, an essential element for the precipitation of γ\u0026rsquo; and γ\u0026rsquo;\u0026rsquo;, the primary strengthening phases [\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e]. Figure\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e(a) shows the band contrast map superimposed with grain boundaries as black lines for the as-built microstructure, i.e., before the heat treatments. It illustrates a microstructure composed of grains with varying morphologies and sizes, ranging from 3 \u0026micro;m to 200 \u0026micro;m. The grains are predominantly oriented along the building direction, leading to an average grain size of 80\u0026thinsp;\u0026plusmn;\u0026thinsp;50 \u0026micro;m (see Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e(b)).\u003c/p\u003e \u003cp\u003eRegarding grain boundary characteristics, the distributions of misorientation angles and axes reveal a random orientation for LAGBs, comprising 21% of the total boundaries. In contrast, HAGBs show a transition toward the \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\u0026lang;001\u0026rang;\\)\u003c/span\u003e\u003c/span\u003e direction at misorientation angles greater than 30˚ (refer to Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e(c)). The significant prevalence of LAGB in the as-built condition is attributed to the high residual stress generated during the manufacturing process. These stresses result in a high density of dislocations caused by the compression and tensile residual stresses within the melt pools [\u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e], [\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec10\" class=\"Section3\"\u003e \u003ch2\u003e3.1.2 Microstructure of heat-treated samples\u003c/h2\u003e \u003cp\u003eEBSD assemblies extracted from the heat-treated samples can be seen in Figs.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e and \u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e. It can be seen from Figs.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea, d, and g that the DA and both solubilized conditions, S and SDA, still present columnar grain morphology, similar to as-built samples. On the other hand, the homogenized conditions, namely H, HDA, and HSDA (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ea, d, and g, respectively), present annealing twins and recrystallized grains. It can be noticed that the misorientation features change from solubilization and aging only treatments when compared to homogenization treatments, Figs.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e and \u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003eb), e), and h). The average misorientations of the calculated were approximately 30.6\u0026deg;, 28.5\u0026deg;, and 28.2\u0026deg;, with a low-angle grain boundary (LAGB) percentage of 19.5, 28.4, and 32.4% for DA, S, and SDA samples, respectively. When compared to the homogenized counterparts, it can be noticed that the H, HDA, and HSDA present higher values of misorientation angles, 41.3\u0026deg;, 47.4\u0026deg;, and 48.2\u0026deg;, respectively, and lower quantities of LAGBs: 19.0, 7.8, and 9.6% for H, HDA, and HSDA, respectively.\u003c/p\u003e \u003cp\u003eThe difference in the average misorientation and percentage of LAGBs can be explained by the recrystallization process promoted by the homogenization treatment. The quantity of LAGBs is generally lower in metallic materials that have undergone recrystallization. This occurs because recrystallization involves the formation of new deformation-free grains with HAGBs, which replace the deformed structure and the LAGBs present in the material before the process [\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e]. During recrystallization, the energy stored in the material due to plastic deformation is released by forming new grains. These grains tend to exhibit more random crystallographic orientations, resulting in HAGBs. LAGBs, which are characteristic of substructures within deformed grains, are either eliminated or significantly reduced during this process [\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e]. This change is evident in the distribution of misorientation axes, where the average misorientation increases for the homogenized materials. This results in a significant shift in the texture of the grain boundaries due to the formation of \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\left\\{111\\right\\}\u0026lang;111\u0026rang;\\)\u003c/span\u003e\u003c/span\u003e recrystallization twins, which occur at a misorientation of 60\u0026deg;.\u003c/p\u003e \u003cp\u003eThe evolution in grain morphology and grain boundary orientation transits from microstructures primarily characterized by low-angle grain boundaries (LAGB) in the solubilized conditions to those dominated by recrystallization twins after the homogenization heat treatment. This remarkable versatility of the alloy in producing a wide range of properties is a testament to its potential to create different mechanical properties. Furthermore, it can be noticed when comparing DA with the other heat treatments that there is a microtexture change, deviating from the almost random orientation shown by the DA sample to a more oriented grain boundaries, as can be seen by the Mackenzie curves in Figs.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e and \u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003eb, e, and h plots.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFrom a grain size perspective, the DA, S, and SDA samples exhibit similar characteristics, featuring coarse columnar grains. In contrast, the homogenized sample reveals recrystallized equiaxed and bimodal grains within the observed field, with grain sizes ranging mostly from approximately 10 to 240 \u0026micro;m for the H and HDA samples, with a slight increase in grain size for the HSDA sample, with a grain size ranging from approximately 10 to 290 \u0026micro;m. The recrystallization may be attributed to the non-uniform distribution of residual stresses in the as-fabricated samples. According to Humphreys et al. [\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e], the static recrystallization temperature of wrought IN718 is approximately 1020\u0026deg;C, which is lower than the 1100\u0026deg;C used in the homogenization treatments of this study. According to the literature, DED processed IN718 exhibits a higher recrystallization temperature than its wrought counterpart, primarily due to the lower residual strain and the high supersaturation of solute atoms in the as-built IN718 produced by DED. In general, wrought IN718 undergoes significant plastic deformation during forging, leading to a considerable accumulation of residual strain energy. As a result, the driving force for recrystallization in wrought IN718 is greater than that in DED-produced IN718, as reported elsewhere [\u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e]. Furthermore, the fast-cooling rates inherent to the DED-LB process significantly enhance the supersaturated solid solubility of alloying elements in the γ matrix, contributing to increased recrystallization temperature [\u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eHeat treatments significantly influence the microstructure of this alloy, as reflected in the distribution, orientation, and density of dislocations. Figure\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e displays geometrically necessary dislocations (GND) maps corresponding to the various heat treatments. These types of dislocations are linked to the curvature of the grains, which is influenced by the manufacturing process\u0026mdash;whether through plastic deformation or the extreme additive manufacturing conditions, where significant residual stresses coming from high cooling rates and repetitive fusion between adjacent layers are present [\u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e], [\u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e], [\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e]. In this context, two behaviors are observed following the various heat treatments: first, heat treatments involving a single stage do not impact the GNDs magnitude compared to the as-built condition (see Figs.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ea-d). In contrast, heat treatments that include multiple thermal cycles reduce the magnitude of dislocations (see Figs.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ee-g). The significant decrease in GNDs observed in the SDA, HAD, and HSDA conditions aligns well with the previously described microstructural changes, where grain size and grain boundary misorientations underwent substantial alterations due to temperature-induced recrystallization.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe analysis of dislocation orientations, represented by the weighted Burgers vector orientation triangle, indicates that dislocations are primarily aligned between the \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\left[001\\right]\\)\u003c/span\u003e\u003c/span\u003e and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\left[101\\right]\\)\u003c/span\u003e\u003c/span\u003e directions. In the case of Ni, which has an FCC crystal structure, close brace open angle bracket 110, along the open brace 111, close brace open angle bracket 110, along the \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\left\\{111\\right\\}\u0026lang;110\u0026rang;\\)\u003c/span\u003e\u003c/span\u003e slip systems. Thus, most dislocations correspond to the \u0026frac12; \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\left[101\\right]\\:\\)\u003c/span\u003e\u003c/span\u003etype, as observed in the as-built and homogenized conditions. However, different heat treatments show a more significant variability in dislocation orientations. This scattering can be attributed to the formation of second phases, such as precipitates or carbides, which generate dislocation tangles and distortions in the lattice.\u003c/p\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e presents the XRD spectra of both as-built and heat-treated samples. In Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ea, two distinct phases\u0026mdash;γ and δ\u0026mdash;are identified, along with fluorescence peaks attributed to W, In, and Pb emissions, likely originating from the shielding walls and devices within the synchrotron hutch. Notably, only the H and S heat treatments resulted in a single-phase structure, whereas the as-built condition and all other heat treatment combinations led to the formation of the δ phase. It should be noted that the XRD analysis could not detect the presence of γ\u0026rsquo; and γ\u0026rdquo; phases, which is already expected due to their expected low volume fraction, which ranges from less than 1% to approximately 11%, depending on the employed HT conditions [\u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e], small size, coherence with the matrix, and superimposition of diffraction peaks [\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eThe as-built IN718 sample exhibited a higher relative intensity for the {200}γ peak than the {111}γ peak. This phenomenon is associated with the preferential growth along the \u0026lt;\u0026thinsp;001\u0026thinsp;\u0026gt;\u0026thinsp;orientation, which is influenced by the significant thermal gradient present along the build direction [\u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e], [\u003cspan citationid=\"CR39\" class=\"CitationRef\"\u003e39\u003c/span\u003e]. The XRD patterns of the HT samples revealed subtle differences among {111}γ and {200}γ peaks. Additionally, minor peaks were observed between the {111}γ and {200}γ peaks, indicating the formation of the δ-phase [\u003cspan citationid=\"CR40\" class=\"CitationRef\"\u003e40\u003c/span\u003e], [\u003cspan citationid=\"CR41\" class=\"CitationRef\"\u003e41\u003c/span\u003e]. In contrast, the HSDA sample exhibited a {111}γ peak with higher intensity than the {200}γ peak, probably due to the grain growth or recrystallization during the HT at 1100\u0026deg;C [\u003cspan citationid=\"CR42\" class=\"CitationRef\"\u003e42\u003c/span\u003e]. Furthermore, the XRD spectra of the DA samples suggests the development of a more random orientation, akin to that observed in Mackenzie curve from Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ec. Moreover, the XRD spectrum of the other HT samples also suggests a more randomized crystallographic orientation, as the relative intensity of matrix peaks follows the calculated values.\u003c/p\u003e \u003cp\u003eDetailed diffraction patterns in the vicinity of the {200}γ peaks for all samples are also presented in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e. The evolution of the γ{200} peak with the complexity of the heat treatment, i.e., following the sequence of conditions: as-built \u0026rarr; DA \u0026rarr; SDA \u0026rarr; HDA \u0026rarr; HSDA, is almost none, indicated by the dashed red line, except the HSDA HT, which shifts to lower interplanar spacing. It indicates a reduction in the lattice parameter of the γ matrix. This shift is primarily attributed to the precipitation of the γ\u0026Prime; phase, which reduces the concentration of dissolved Nb in the γ matrix. However, distinguishing their contributions in the XRD pattern proved challenging due to the issues already mentioned [\u003cspan citationid=\"CR43\" class=\"CitationRef\"\u003e43\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure 8 shows thermodynamic calculations considering either the non-equilibrium (Scheil-Gulliver) or the thermodynamic equilibrium (phase diagram) for IN718 alloy. δ, γ\u0026rsquo;, γ\u0026rdquo; were considered one phase in the calculations due to their similar chemical composition, and they are called Ni₃(Nb, Ti). The solidification path predicted for the present composition of IN718 alloy is as follows:\u003c/p\u003e \u003cp\u003eLiquid (L)\u0026rarr; L\u0026thinsp;+\u0026thinsp;γ \u0026rarr; L\u0026thinsp;+\u0026thinsp;γ\u0026thinsp;+\u0026thinsp;NbC \u0026rarr; L\u0026thinsp;+\u0026thinsp;γ\u0026thinsp;+\u0026thinsp;NbC\u0026thinsp;+\u0026thinsp;Laves\u0026thinsp;+\u0026thinsp;Ni\u003csub\u003e3\u003c/sub\u003e(Nb, Ti) \u0026rarr; L\u0026thinsp;+\u0026thinsp;γ\u0026thinsp;+\u0026thinsp;NbC\u0026thinsp;+\u0026thinsp;Laves\u0026thinsp;+\u0026thinsp;Ni\u003csub\u003e3\u003c/sub\u003e(Nb, Ti) + σ.\u003c/p\u003e \u003cp\u003eThe non-equilibrium thermodynamic conditions inherent to additive manufacturing promote the formation of metastable phases. The Scheil\u0026ndash;Gulliver solidification model accurately predicts transition temperatures, revealing Laves phase formation between 1180\u0026ndash;1185\u0026deg;C. Below 1180\u0026deg;C, chemical element stoichiometry enables Ni\u003csub\u003e3\u003c/sub\u003eNb phase development. Microsegregation introduces complexity, with low-solubility elements like Nb, Mo, Ti, and C concentrating in interdendritic regions and forming MC carbides and Laves phases. This segregation can induce variations in precipitation temperatures, highlighting the intricate microstructural evolution during DED-LB processing.\u003c/p\u003e \u003cp\u003eOn the other hand, phase equilibrium calculations show an insignificant presence of NbC, besides a considerable presence of σ phase at low temperature, which continues to dissolve until 925\u0026deg;C, and Ni\u003csub\u003e3\u003c/sub\u003e(Nb, Ti), achieving its complete dissolution at 1030\u0026deg;C. The phase equilibrium calculation at each heat treatment temperature condition is depicted in Table\u0026nbsp;\u003cspan refid=\"Tab1\" class=\"InternalRef\"\u003e1\u003c/span\u003e.\u003c/p\u003e \u003cp\u003e \u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab1\" border=\"1\"\u003e \u003ccaption language=\"En\"\u003e \u003cdiv class=\"CaptionNumber\"\u003eTable 1\u003c/div\u003e \u003cdiv class=\"CaptionContent\"\u003e \u003cp\u003ePhase volume percent at heat treatment temperatures under equilibrium conditions.\u003c/p\u003e \u003c/div\u003e \u003c/caption\u003e \u003ccolgroup cols=\"5\"\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c3\" colnum=\"3\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c4\" colnum=\"4\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c5\" colnum=\"5\"\u003e\u003c/div\u003e \u003cthead\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c1\" morerows=\"1\" rowspan=\"2\"\u003e \u003cp\u003eTemperature (\u0026deg;C)\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colspan=\"4\" nameend=\"c5\" namest=\"c2\"\u003e \u003cp\u003ePhases (in volume% %)\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c2\"\u003e \u003cp\u003eNbC\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c3\"\u003e \u003cp\u003eγ\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c4\"\u003e \u003cp\u003eNi\u003csub\u003e3\u003c/sub\u003e(Nb,Ti)\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c5\"\u003e \u003cp\u003eσ\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003c/thead\u003e \u003ctbody\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e\u003cb\u003e620\u003c/b\u003e\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e0.1\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e73.8\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c4\"\u003e \u003cp\u003e14.8\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c5\"\u003e \u003cp\u003e11.2\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e\u003cb\u003e720\u003c/b\u003e\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e0.1\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e77.3\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c4\"\u003e \u003cp\u003e14.0\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c5\"\u003e \u003cp\u003e8.6\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e\u003cb\u003e1000\u003c/b\u003e\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e0.1\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e96.9\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c4\"\u003e \u003cp\u003e2.9\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c5\"\u003e \u003cp\u003e-\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e\u003cb\u003e1100\u003c/b\u003e\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e0.1\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e99.9\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c4\"\u003e \u003cp\u003e-\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c5\"\u003e \u003cp\u003e-\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003c/tbody\u003e \u003c/colgroup\u003e \u003c/table\u003e\u003c/div\u003e \u003c/p\u003e \u003cp\u003eIntermetallic phase formation in DED-LB manufacturing of IN718 alloy emerges from complex interactions between kinetic processes, thermal history, and elemental availability. The alloy demonstrates coexistence of stable (δ) and metastable (γ\u0026Prime;) phases with Ni\u003csub\u003e3\u003c/sub\u003eNb stoichiometry, with phase prevalence critically dependent on temperature gradients and elemental distribution [\u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e], [\u003cspan citationid=\"CR44\" class=\"CitationRef\"\u003e44\u003c/span\u003e]. Temperatures exceeding 600\u0026deg;C trigger MC carbide precipitation, predominantly NbC, with carbide growth dependent upon Nb concentration in δ and γ\u0026Prime; phases. The metastable γ\u0026Prime; phase is the primary strengthening precipitate, significantly enhancing mechanical properties. However, prolonged high-temperature exposure induces transformation to stable δ phase, potentially compromising material strength [\u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e], [\u003cspan citationid=\"CR44\" class=\"CitationRef\"\u003e44\u003c/span\u003e].\u003c/p\u003e\u003cp\u003eFigure \u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e9\u003c/span\u003e shows the SEM and EDS analysis for all the aged samples. Microstructural comparisons of as-built and DA samples revealed minimal morphological distinctions through SEM micrographic analysis. Characteristic precipitates, specifically Laves phase, were consistently observed in both conditions, suggesting that the applied aging temperatures were insufficient to induce complete dissolution of the Laves phase, Fig.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e9\u003c/span\u003ea. Figure\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e9\u003c/span\u003eb depicts the SDA sample microstructure, showing that the cellular microstructure persisted. It can be seen that localized grain growth occurs compared to the as-built condition. From the EDS analysis, shown as a table on the side of the micrographs, the possible identification of those phases as δ precipitates and carbides can be seen. Homogenization treatments, namely HDA and HSDA, Figs.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e9\u003c/span\u003ec and \u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e9\u003c/span\u003ed, respectively, effectively eliminated melt pool boundaries and elemental segregation. Significant microstructural transformations were observed, including extensive recrystallization and grain growth, with grain sizes exceeding 200 \u0026micro;m in a bimodal fashion, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e. Annealing twins were detected, and grain boundaries exhibited small white particles (Fig.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e9\u003c/span\u003ec), tentatively identified as carbides; Fig.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e9\u003c/span\u003ed also shows the presence of carbides and possibly δ precipitates inside the grains.\u003c/p\u003e \u003c/div\u003e \u003c/div\u003e \u003cdiv id=\"Sec11\" class=\"Section2\"\u003e \u003ch2\u003e3.2 Mechanical Properties\u003c/h2\u003e \u003cp\u003eHardness test results for the DED-LB and the forged materials are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e10\u003c/span\u003e. The DA heat treatment enhanced the material's mechanical strength by approximately 50% compared to non-aged samples, reaching values of 475.7 HV, 503.5 HV, 506.1 HV, and 483.47 HV for DA, HDA, SDA, and HSDA, respectively. The authors believe that it occurs due to the precipitation of the strengthening phases of the IN718 alloy, i.e., γ\u0026rsquo; and γ\u0026rsquo;\u0026rsquo;, which could not be detected by the XRD analysis. Furthermore, it can be noticed that the H and S heat treatments could solubilize the alloy, reducing its hardness values. Although similar to the forged IN718, the hardness values are slightly superior to the DED-LB values. Forged materials generally exhibit higher hardness than additively manufactured materials due to differences in microstructure, defect density, and residual stress distribution. Forging induces severe plastic deformation, leading to dynamic recrystallization, grain refinement, and increased dislocation density, which enhances hardness through the Hall-Petch effect [\u003cspan citationid=\"CR45\" class=\"CitationRef\"\u003e45\u003c/span\u003e]. In contrast, AM processes often result in columnar and dendritic microstructures with lower dislocation density due to rapid solidification and lack of plastic deformation. Additionally, AM components tend to exhibit higher porosity and process-induced defects, reducing mechanical performance compared to the dense and homogeneous microstructure of forged materials [\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eResidual compressive stresses in forged materials further contribute to their superior hardness, whereas AM components often require post-processing heat treatments to mitigate tensile residual stresses and improve mechanical properties [\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e]. While AM advancements are narrowing this gap, forged materials generally maintain superior hardness due to their refined microstructure and defect-free nature. This effect can be seen for the as-built and AMS 5662 forged conditions in Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e10\u003c/span\u003e.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e11\u003c/span\u003ea presents the engineering stress-strain curves for all analyzed conditions. Upon initial observation, there is a wide range of combinations of strength and ductility, with yield strengths varying from 300 MPa to 1200 MPa and fracture strains ranging from 0.5 to 0.35. Single-stage heat treatments, such as homogenization and solubilization, result in the highest ductility. In contrast, thermal cycles that involve aging heat treatments, such as DA, SDA, HDA, and HSDA, yield greater strengths than the as-built condition.\u003c/p\u003e \u003cp\u003eThe strain hardening rate curves demonstrate that ductility is enhanced following the S and H heat treatments. In contrast, the aged conditions exhibit reduced ductility, which leads to localized plastic instability during early deformations, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e11\u003c/span\u003eb. This loss of ductility in the aged conditions can be attributed to the large number of precipitates at the grain boundaries, as confirmed in Fig.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e9\u003c/span\u003e. The formation of precipitates at grain boundaries can increase brittleness, particularly when the particles are incoherent and large. Studies have shown that the coarsening of γ\u0026rsquo; precipitates after heat treatment follows the Ostwald ripening mechanism. This process results in a loss of ductility due to strain concentration caused by the incoherent precipitates localized at the grain boundaries [\u003cspan citationid=\"CR47\" class=\"CitationRef\"\u003e47\u003c/span\u003e], [\u003cspan citationid=\"CR48\" class=\"CitationRef\"\u003e48\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eThe strain profiles observed during tensile tests at various deformations reveal different behaviors depending on the applied heat treatment. After the HDA heat treatment, the strain distribution along the calibrated length shows significant heterogeneity from the start until the onset of plastic instability (see Fig.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e11\u003c/span\u003ec). In the as-built condition, the strain profiles exhibit multiple peaks, which indicate strain localization at different points along the sample length, thereby delaying the formation of plastic instability (see Fig.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e11\u003c/span\u003ed). In contrast, the homogenized condition displays strain profiles with less heterogeneity than the ones mentioned earlier. These profiles resemble a plateau, facilitating the storage of larger and more uniform deformations before the plastic instability (Fig.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e11\u003c/span\u003ee). The behaviors discussed are closely related to the various microstructures described. For instance, the early onset of strain localization in the HDA can be attributed to the brittleness caused by the coarse precipitates at the grain boundaries. In contrast, the multiple peaks observed in the as-built condition can be linked to the residual stresses arising from the manufacturing process, which are repeated between adjacent melt pools. In the homogenized condition, the more uniform strain distribution is due to the absence of precipitates and the larger, more equiaxed grain size resulting from the recrystallized microstructure, which allows for greater dislocation motion.\u003c/p\u003e \u003cp\u003eTherefore, a positive synergy between strength and ductility in the studied alloy is achieved by an adequate distribution of coherent and intragranular γ\u0026rsquo; precipitates, which impart strength while controlling the number of precipitates at the grain boundaries to avoid ductility loss.\u003c/p\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e12\u003c/span\u003e presents the room-temperature mechanical properties of the DED-LB processed IN718 alloy in both the as-built state and after six distinct heat-treatment conditions. The yield strength values for the as-built, DA, H, S, HDA, SDA, and HSDA conditions were determined to be approximately 638 MPa, 1195 MPa, 360 MPa, 412 MPa, 1037 MPa, 1189 MPa, and 1046 MPa, respectively. Additionally, the elongation values for these conditions were measured as \u0026asymp;\u0026thinsp;28.5%, 8.0%, 49.3%, 47.1%, 17.8%, 12.1%, and 16.7%, respectively. To characterize the plastic deformation behavior observed during tensile testing, the Hollomon equation (σ\u0026thinsp;=\u0026thinsp;K\u0026sdot;ε\u003csup\u003en\u003c/sup\u003e, where σ represents the true stress, K is the strength coefficient, ε denotes the true strain, and n is the strain-hardening exponent) was employed. The calculated n values for the DA, H, S, HDA, SDA, and HSDA samples were approximately 0.072, 0.281, 0.25, 0.085, 0.06, and 0.082, respectively. These results indicate that the H and S conditions exhibit the highest uniform plastic deformation capacity, whereas the aged conditions display increased hardness and reduced deformability. The elastic modulus (E) values were determined to be \u0026asymp;\u0026thinsp;203 GPa for the as-built sample, 211 GPa, 197 GPa, 203 GPa, 214 GPa, 210 GPa, and 220 GPa for the DA, H, S, HDA, SDA, and HSDA conditions, respectively. This suggests that the HSDA sample demonstrates the greatest resistance to elastic deformation.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe presence of γ\u0026rsquo;/γ\u0026rdquo; strengthening phases, which develop during the DA process, contributes to an increase in tensile strength, explaining why the homogenized and solubilized conditions, lacking these phases, exhibit comparatively lower strength. Among the evaluated heat treatment conditions, direct aging results in the highest tensile strength; however, it also leads to the lowest ductility. In contrast, solution annealing provides slightly higher tensile strength than homogenization solution aging, though at the cost of reduced ductility, while maintaining comparable overall tensile properties [\u003cspan citationid=\"CR49\" class=\"CitationRef\"\u003e49\u003c/span\u003e].\u003c/p\u003e \u003c/div\u003e"},{"header":"4. Conclusions","content":"\u003cp\u003eThe study investigated the effects of six different heat-treatment routes on the microstructural evolution, hardness, tensile properties, and fracture behavior of DED-LB IN718 samples. The relationship between microstructures and mechanical responses was discussed and compared to the forged IN718 counterpart.\u003c/p\u003e\n\u003cp\u003eThe conclusions reached are:\u003c/p\u003e\n\u003cp\u003e1. The results showed that different heat treatment strategies can tailor the microstructure and mechanical properties of DED-LB IN718 samples. \u0026nbsp;\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e2.\u0026nbsp; Aging treatments (DA, SDA, HDA, and HSDA) increased hardness due to the precipitation of strengthening phases. \u0026nbsp;\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e3.\u0026nbsp; Homogenization treatments (H, HDA, and HSDA) promoted recrystallization and grain growth, leading to a more uniform strain distribution and enhanced ductility. \u0026nbsp; \u0026nbsp;\u003c/p\u003e\n\u003cp\u003e4.\u0026nbsp; The homogenized condition displayed strain profiles with less heterogeneity, facilitating the storage of larger and more uniform deformations before the plastic instability. \u0026nbsp;\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e5. A balance between strength and ductility was achieved through the adequate distribution of coherent and intragranular \u0026gamma;\u0026rsquo; precipitates, while controlling the number of precipitates at the grain boundaries to avoid ductility loss.\u0026nbsp;\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eAcknowledgments\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eBLINDED MANUSCRIPT\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eFunding\u0026nbsp;\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eBLINDED MANUSCRIPT\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eData availability\u003c/strong\u003e. BLINDED MANUSCRIPT\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eConflict of interest statement.\u003c/strong\u003e 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.\u003c/p\u003e\u003ch2\u003eAuthor Contribution\u003c/h2\u003e\u003cp\u003eT. R. F. Cavalcante was responsible for writing the manuscript, editing the text, and preparing all figures. D. G. Bon conducted the mechanical testing and contributed to the analysis and interpretation of mechanical behavior. J. A. Mu\u0026ntilde;oz performed the digital image correlation (DIC) analysis and contributed to data interpretation. G. G. Ribamar, J. P. Oliveira, and A. B. Pereira carried out the simulation work and participated in the X-ray diffraction (XRD) analysis. F. E. Mariani contributed to the experimental design and supervised the additive manufacturing process. J. C. Mu\u0026ntilde;oz, R. T. Coelho and J. A. A. Diaz provided project supervision, critical review of the manuscript, and guidance throughout the research. All authors discussed the results, contributed to the final version of the manuscript, and approved its submission.\u003c/p\u003e\u003ch2\u003eAcknowledgement\u003c/h2\u003e\u003cp\u003eAcknowledgmentsThe support of the Center for Research and Innovation in Materials and Structures (CEPIMATE) is deeply appreciated. This research used facilities of the Brazilian Nanotechnology National Laboratory (LNNano), part of the Brazilian Centre for Research in Energy and Materials (CNPEM), a private non-profit organization under the supervision of the Brazilian Ministry for Science, Technology, and Innovations (MCTI). The Electron Microscopy Laboratory staff are acknowledged for their assistance during the experiments (proposal SEM-FIB-C1-20233619). We acknowledge DESY (Hamburg, Germany), a member of the Helmholtz Association HGF, for providing the experimental facilities. Parts of this research were carried out at PETRA III (proposal I-20230101 EC), and we would like to thank Dr. Guilherme Abreu Faria and Dr. Marc-Andr\u0026eacute; Nielsen for their assistance in using beamline P61A. Funding Funded by Funda\u0026ccedil;\u0026atilde;o de Amparo \u0026agrave; Pesquisa do Estado de S\u0026atilde;o Paulo \u0026ndash; FAPESP, grant No. 2020/09079-2. This study was partly financed by the Conselho Nacional de Desenvolvimento Cient\u0026iacute;fico e Tecnol\u0026oacute;gico \u0026ndash; CNPq, grant No. 306960/2021\u0026ndash;4. Thiago Roberto Felisardo Cavalcante recognizes the financial support through the Ph.D. scholarship from the Coordena\u0026ccedil;\u0026atilde;o de Aperfei\u0026ccedil;oamento de Pessoal de N\u0026iacute;vel Superior \u0026ndash; Brasil (CAPES) \u0026ndash; Finance Code 001 and CAPES PrInt, process 88887.886782/2023-00. JPO acknowledges funding by national funds from FCT \u0026ndash; Funda\u0026ccedil;\u0026atilde;o para a Ci\u0026ecirc;ncia e a Tecnologia, I.P., in the scope of the projects LA/P/0037/2020, UIDP/50025/2020, and UIDB/50025/2020 of the Associate Laboratory Institute of Nanostructures, Nanomodelling and Nanofabrication \u0026ndash; i3N. The present study was in part developed in the scope of the Project \u0026ldquo;Agenda ILLIANCE\u0026rdquo; [C644919832-00000035 | Project n\u0026ordm; 46], financed by PRR \u0026ndash; Plano de Recupera\u0026ccedil;\u0026atilde;o e Resili\u0026ecirc;ncia under the Next Generation EU from the European Union.Data availability. The data supporting this study\u0026rsquo;s findings are available from the corresponding author, J. A. 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J Manuf Mater Process 8(4). \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003e10.3390/jmmp8040135\u003c/span\u003e\u003cspan address=\"10.3390/jmmp8040135\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eDeng D, Peng RL, Brodin H, Moverare J (Jan. 2018) Microstructure and mechanical properties of Inconel 718 produced by selective laser melting: Sample orientation dependence and effects of post heat treatments. Mater Sci Engineering: A 713:294\u0026ndash;306. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003e10.1016/j.msea.2017.12.043\u003c/span\u003e\u003cspan address=\"10.1016/j.msea.2017.12.043\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":false,"hideJournal":true,"highlight":"","institution":"","isAcceptedByJournal":true,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true},"keywords":"IN718, heat treatment, microstructure, additive manufacturing, mechanical properties","lastPublishedDoi":"10.21203/rs.3.rs-6602510/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-6602510/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eDirected energy deposition-laser beam (DED-LB), a laser additive manufacturing (AM) process, has attracted significant attention as a potential alternative to conventional manufacturing methods, due to its high deposition efficiency, flexibility, and precision. Despite these advantages, components produced by DED-LB often face critical challenges, including residual stresses, micro-segregation, and the formation of non-equilibrium phases due to rapid cooling during the AM process. These issues are particularly critical for Ni-based superalloys such as Inconel 718 (IN718), widely used in the aerospace, energy, and marine industries for their excellent high-temperature strength and corrosion resistance. The mechanical performance of IN718 primarily depends on precipitation hardening via γ' and γ'' phases. In contrast, the formation of deleterious phases, such as δ and Laves, can severely impair performance by depleting key alloying elements and increasing brittleness. Thus, heat treatments (HTs) are vital in addressing these challenges by reducing micro-segregation, homogenizing elemental distribution, and promoting the precipitation of strengthening phases. Therefore, this study investigates the effects of six distinct heat-treatment routes on the microstructural evolution, hardness, tensile properties, and fracture behavior of DED-LB IN718 samples. The relationship between microstructure and mechanical responses is analyzed and compared to a forged IN718 counterpart. The results offer valuable insights for optimizing heat-treatment strategies to improve the structural integrity and mechanical reliability of DED-LB-fabricated IN718 components.\u003c/p\u003e","manuscriptTitle":"Microstructural assessment of additive-manufactured Inconel 718 samples subjected to heat treatments for enhanced mechanical properties","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-05-26 08:40:55","doi":"10.21203/rs.3.rs-6602510/v1","editorialEvents":[{"type":"communityComments","content":0}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true}}],"origin":"","ownerIdentity":"98c82716-01c6-4af3-9e97-91d3b693f02c","owner":[],"postedDate":"May 26th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"posted","subjectAreas":[],"tags":[],"updatedAt":"2025-08-11T16:05:33+00:00","versionOfRecord":{"articleIdentity":"rs-6602510","link":"https://doi.org/10.1007/s40964-025-01279-y","journal":{"identity":"progress-in-additive-manufacturing","isVorOnly":false,"title":"Progress in Additive Manufacturing"},"publishedOn":"2025-08-05 15:57:36","publishedOnDateReadable":"August 5th, 2025"},"versionCreatedAt":"2025-05-26 08:40:55","video":"","vorDoi":"10.1007/s40964-025-01279-y","vorDoiUrl":"https://doi.org/10.1007/s40964-025-01279-y","workflowStages":[]},"version":"v1","identity":"rs-6602510","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-6602510","identity":"rs-6602510","version":["v1"]},"buildId":"XKTyCvWXoU3ODBz1xrDgd","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}