Build Orientation Effects on Microstructure and Damage Evolution of Selective Laser Melted in718 Alloy

Abstract

This study reveals the synergistic mechanism of build direction and heat treatment on the microstructure, mechanical properties, and damage evolution behavior of selective laser melted (SLM) IN718 alloy. The results demonstrated that the mechanical properties of as-deposited alloys exhibited significant directional dependence: the 0°-oriented specimens showed higher tensile strength of 1001 MPa, while the 90°-oriented specimens exhibited superior plasticity with elongation of 19.9%. After heat treatment, microstructural segregation was eliminated, and δ phases precipitated at grain boundaries. The tensile strength of 0°-oriented specimens increased by 37.4%, surpassing the 26.4% enhancement observed in 90°-oriented specimens. Quantitative characterization of damage evolution behavior via nanoindentation technology indicates that the 90°-oriented specimens exhibit slower elastic modulus degradation in the early stage of damage, with a higher critical damage factor than the 0°-oriented specimens.

Keywords:
Build orientation; Heat treatment; SLM-IN718 alloy; Nanoindentation; Damage evolution

1. INTRODUCTION

IN718 is a nickel-based superalloy, which is widely used in aerospace and other fields due to its excellent fatigue resistance, corrosion resistance, oxidation resistance, and high-temperature mechanical properties¹-⁵. However, traditional processing methods are difficult to meet its application requirements. Selective laser melting (SLM), as a new type of additive manufacturing process, has the advantages of custom design shape and personalized customization, and can complete the design of fine parts that are difficult to achieve by traditional processing methods⁶-⁸.

When using selective laser melting technology to process specimens, the specimens obtained with different build directions have differences in structure and properties⁹,¹⁰. Relevant studies have shown that the build direction has a significant impact on material properties, and the tensile strength of samples built in the Z direction is relatively weak¹¹. The study by Singh et al.¹² has shown that samples with a 0° build direction exhibit higher tensile strength and higher microhardness but poorer ductility. Samples with a 90° build direction have better ductility. Sun et al.¹³ have found that 0° and 45° samples have more defects in the upper region, while 90° samples show a more uniform defect distribution.

In addition, the defects existing in the microstructure and mechanical properties of SLM-formed alloys mainly include pores, cracks, residual stress, and spheroidization. These defects can be improved to a large extent through reasonable heat treatment measures, thereby enhancing the comprehensive performance of the formed parts¹⁴-¹⁷. Some studies¹⁸ have shown that through optimizing the heat treatment process, the Nb-enriched Laves phase is effectively dissolved, promoting the homogenization of the alloy; the yield strength of the alloy after direct aging treatment is 53.9% higher than that of the original state. The study by Wang et al.¹⁹ has shown that after Homogenization + solution + double aging (HAS) treatment, the elongation of Inconel 718 alloy increases from 21.0% to 27.3%, and the tensile strength increases from 946 MPa to 1570 MPa. The studies by scholar Aripin²⁰-²² have indicated that the microstructure, porosity, and mechanical properties of 17-4 PH stainless steel prepared by selective laser melting (SLM) are significantly affected by build direction and heat treatment. It is mainly composed of austenite and martensite phases, with crescent, triangular, and spherical pores. The 0° build direction usually has higher tensile strength due to higher martensite content and lower porosity, while heat treatments such as H750 can further optimize mechanical properties by regulating phase transformation and retaining molten pool boundaries.

The above studies have shown that heat treatment can effectively improve the microstructure and comprehensive mechanical properties of SLM-IN718 alloy, enabling it to have wider applications in practical engineering. However, in practical engineering, the accumulation of internal damage in materials will seriously affect their in-service performance²³,²⁴. Under loading conditions, internal damage in materials will evolve and develop with deformation and eventually lead to specimen failure. Exploring the law of material damage evolution is crucial for evaluating its in-service performance. Some scholars have conducted research on this. Studies have shown that combining digital image correlation (DIC) and direct current potential drop (DCPD) technologies, ductile damage evolution can be measured through continuous uniaxial tensile tests²⁵, and the method used can easily determine the damage threshold, plastic strain, and critical damage value. In addition, through monotonic tensile and fatigue tests, the changes in elastic modulus of materials under different strain and stress conditions are recorded²⁶,²⁷, and then the damage value can be calculated²⁸.

To sum up, although a large number of scholars have studied the structure and properties of SLM-IN718, the impact of heat treatment on the damage evolution behavior of SLM-IN718 alloys with different build directions has not been involved by scholars so far. In this paper, the microstructure and mechanical properties of SLM-IN718 specimens built with two different orientations (0° and 90°) and their heat-treated counterparts are analyzed. Based on nanoindentation technology²⁹,³⁰, the damage evolution behavior of SLM-IN718 alloy during tensile process is quantitatively characterized. On this basis, a damage evolution equation is established, and the critical damage factor is calculated. Finally, the influence of different printing directions on the mechanical properties and tensile damage evolution behavior of SLM-IN718 alloy is discussed in depth. The results of this study can provide a key basis for the performance optimization of SLM-IN718 alloy. The study clarifies how build direction and heat treatment synergistically affect the microstructure, mechanical properties, and damage evolution of SLM-IN718. These findings enable engineers to optimise build orientation such as 0° for strength-critical components or 90° for ductility-demanding applications and tailor heat treatment protocols to suppress damage accumulation in aerospace components. At the same time, optimizing the heat treatment process according to the research results can further improve the alloy performance, reduce component failure caused by damage accumulation, effectively improve the reliability and service life of SLM-IN718 alloy components, reduce engineering costs, and provide strong technical support for the wide application of such alloy components in aerospace and other fields.

2. MATERIALS AND METHODS

Gas-atomized spherical IN718 powder (15–45 μm particle size) was processed using an EOS-M290 3D printer (Figure 1(a)). The chemical composition of the powder is listed in Table 1. The printing parameters were as follows: laser spot diameter of 0.1 mm, laser power of 285 W, scanning speed of 960 mm/s, scanning spacing of 100 μm, layer thickness of 0.04 mm, and interlayer rotation angle of 67°. A 67° interlayer rotation ensures that no repetition occurs until more than 1800 layers are scanned, and the different scanning patterns between adjacent layers help improve the bonding strength between layers, as shown in Figure 1(b). Meanwhile, two build directions (0° and 90°) were adopted. The dimensions of the as-formed specimens are shown in Figure 1(c). To improve surface quality, ensure dimensional accuracy, and thereby reduce interference from factors such as stress concentration and impurities on the test results, the specimens were polished with 400-grit sandpaper before the experiment, and Figure 1(d) shows the polished specimens.

Figure 1
SLM process: (a) EOS-M290 3D metal printer. (b) Schematic diagram of SLM forming process. (c) Specimen size. (d) Specimens of the actual picture.

Table 1
Major elements of SLM-IN718 alloy powders (wt.%).

C	Si	Mn	S	P	Cr	Ni
0.05	0.043	0.03	0.002	0.034	19.01	52.30
Al	Nb	Ti	Cu	B	Mo	Fe
0.57	5.07	1.00	0.02	0.003	3.06	Bal

Heat treatment was performed on the specimens using an SG-GL1200 vacuum tube furnace. The temperature was increased at a rate of 5 °C/min, held at 980 °C for 1 h, and then cooled by air cooling (AC). The names and processes of as-deposited (AD) and heat-treated (HT) specimens are listed in Table 2.

Table 2
Name of specimens and process.

Specimen	AD-0	AD-90	HT-0	HT-90
Build Orientation	0°	90°	0°	90°
Heat Treatment	No	No	980°C/1h/AC	980°C/1h/AC

Prior to microstructure observation, the specimens were first ground with waterproof sandpaper (up to 2000 grit). After fine grinding, they were polished to obtain scratch-free mirror-like specimens. The specimens were etched with a corrosive solution with a ratio of HCl:C₂H₅OH:CuCl₂ = 20 ml:20 ml:1 g for 30 s. The micrographs of 0°-oriented specimens were taken on the XOY plane in Figure 1(b), while those of 90°-oriented specimens were taken on the YOZ plane. A Leica DM 1700 M upright metallographic optical microscope was used to observe the microstructure of the specimens (in accordance with ASTM E3-11³¹). A scanning electron microscope (model Apreo S LoVac) was used to observe the microstructure and capture images of the fracture surfaces of the specimens after tensile testing.

The elastic modulus of the specimens at different tensile stages was measured using a Nano Indenter G200 in-situ nanoindentation tester in accordance with ASTM E2546-15³². Based on the elongation of different specimens obtained from the tensile property tests, 15 stages were set for each specimen. After each tensile stage, the specimen was removed from the tensile testing machine, ground, polished, and subjected to nanoindentation tests. The elastic modulus was measured at 20×3 points within the gauge length of the specimen, and the average value of the measurements at the points corresponding to the necking region was taken. Then, a new specimen was used for tensile testing, and the above operation was repeated until the specimen fractured. The test parameters were set as follows: loading to 300 mN at a constant loading rate of 5 mN/s within 60 s, holding for 10 s, and then unloading within 60 s.

AD/MAX-2500/PC X-ray diffractometer was used for phase analysis of the SLM-IN718 alloy. EBSD experiments were conducted using a FEI QUANTA 650 field emission scanning electron microscope equipped with an Oxford-Instrument Nordly Nano EBSD system to generate orientation distribution maps, recrystallization distribution maps, and local misorientation angle maps. Tensile property tests were performed on three specimens in each group of SLM-IN718 alloy using an MTS LandMark universal testing machine to obtain the yield strength, tensile strength, and elongation of the specimens (in accordance with ASTM E8/E8M-21³³).

3. RESULTS AND DISCUSSION

3.1. MICROSTRUCTURAL ANALYSIS

Figure 2 shows the optical micrographs of SLM-IN718 alloy. It can be clearly observed from Figure 2(a) that the AD-0 specimen exhibits elongated scanning path traces, with an angle of approximately 67° between the scanning paths, which is consistent with the set process parameters during processing³⁴,³⁵. Fish-scale molten pool boundaries can be seen in the AD-90 specimen in Figure 2(b). During the forming process of SLM-IN718, the material grows from bottom to top, passing through multiple deposited layers, and has the characteristic of directional solidification. This is because the temperature gradient direction at the bottom of the molten pool is along the deposition direction of the material, and the laser energy presents a Gaussian distribution with low intensity at both ends and high intensity in the middle during scanning³⁵,³⁶, resulting in the layer-by-layer superposition of fish-scale molten pools. After heat treatment, it can be observed that the elongated scanning path traces in HT-0 (Figure 2(c)) disappear, and the fish-scale molten pool boundaries in HT-90 (Figure 2(d)) dissolve, but obvious banded melting tracks still remain. This is because during SLM forming, the temperature gradient direction at the bottom of the molten pool is along the material deposition direction; for the 90°-oriented specimens, due to the difference in the initial molten pool boundary direction from that of the 0°-oriented specimens, there are differences in the dissolution and recrystallization behavior of molten pool boundaries during heat treatment, leading to the retention of obvious banded melting tracks.

Figure 2
Optical microscope images of SLM-IN718 alloy: (a) AD-0; (b) AD-90; (c) HT-0; (d) HT-90.

Figure 3 shows the SEM morphologies of SLM-IN718 alloy. A large number of irregular Laves phases are distributed in the interdendritic regions and overlapping areas of adjacent scanning paths in the AD specimens. This is because IN718 alloy contains elements that are highly prone to segregation, such as Nb and Mo. Due to the extremely fast cooling rate during forming, the matrix is in a supersaturated state. When the content of Nb element is sufficiently high, eutectic reactions occur between dendrites to form Laves phases. Laves phases are important sources of cracks and have an adverse effect on the tensile properties of Inconel 718 alloy³⁵. It can be seen from Figures 3(a) and (b) that the AD-0 specimen exhibits relatively regular honeycomb-like coarse dendrites and elongated fine dendrite growth morphologies, while the AD-90 specimen shows fish-scale molten pool boundaries and relatively disordered dendrite distribution.

Figure 3
SEM image of SLM-IN718 alloy: (a) AD-0; (b) AD-90; (c) HT-0; (d) HT-90.

Figures 3(c) and (d) are the SEM images of SLM-IN718 alloy after heat treatment. It can be observed from the figures that compared with the AD specimens, the Laves phases in the SLM-IN718 alloy after HT are reduced. The HT-0 specimen has more precipitates with smaller sizes at the grain boundaries than the HT-90 specimen, and the δ phases in the HT-90 specimen are more sparsely spaced. Since the precipitation rate of δ phases is the fastest at 900°C, during the solution process at 980°C, the Laves phases dissolve and release a large amount of Nb elements, which diffuse to the grain boundaries and intragranular regions to precipitate δ phases in two forms: rod-like and needle-like. These δ phases play a role in dislocation pinning at the grain boundaries, thereby producing a dispersion strengthening effect. Due to the poor corrosion resistance of δ phases, the grain boundary contours after HT are clear³⁷,³⁸.

Figure 4 shows the XRD patterns of SLM-IN718 alloy. It can be seen from the figure that there is no significant difference in the phase composition of SLM-IN718 alloy with the two build directions. Diffraction peaks of the matrix γ phase at (111), (200), and (220) planes are detected in all specimens³⁹. δ phases are detected in HT-0 and HT-90 specimens.

Figure 4
XRD patterns of SLM-IN718 alloy.

Figure 5 shows the grain orientation distribution maps of SLM-IN718 alloy. It can be seen from the figure that the as-deposited SLM-IN718 alloy grains with both build directions grow through several stacked layers, presenting a columnar distribution. However, they are not all continuous and exhibit a certain fractured columnar morphology. This is because the 100 μm scanning spacing during SLM forming is too large, resulting in a shallow molten pool depth, which provides a basis for the directional growth of <001>-oriented columnar grains. In addition, the 67° interlayer rotation during processing makes epitaxial growth more difficult, leading to the fractured growth morphology of grains⁴⁰,⁴¹. The grains of both 0° and 90° build direction specimens are regularly arranged in fine strips, and most grains have a large aspect ratio. The aspect ratios of specimens AD-0, AD-90, HT-0, and HT-90 are 2.37, 4.66, 2.23, and 4.52, respectively. The 0°-oriented specimens have a smaller aspect ratio than the 90°-oriented specimens. This is because during horizontal forming, heat flow is mainly conducted in the direction perpendicular to the laser scanning direction, which makes the cooling rate of each layer relatively fast. This rapid cooling helps form a more uniform grain structure, thereby improving the strength of the material. In vertical forming, however, heat flow is mainly conducted along the forming direction, causing the grains in the specimen to present a longer fibrous structure due to the influence of thermal gradient, resulting in an increase in their aspect ratio¹⁰. The grains show a regular alternating distribution of coarse and fine grains. This is because the remelted regions during SLM forming are subjected to multiple heating by high-energy lasers, leading to remelting and grain refinement in these regions. These refined grains and the unremelted parts in the path form an alternating distribution of coarse and fine grain regions. After heat treatment, as can be seen from Figures 5(c) and (d), the grain orientation distribution of different printing directions still retains some original characteristics. After heat treatment, although partial recrystallization occurs, the <001> orientation advantage of columnar grains remains significant, especially in the structure with vertical build direction.

Figure 5
Orientation distribution of SLM-IN718 alloy: (a) AD-0; (b) AD-90; (c) HT-0; (d) HT-90.

Figures 6(a) and (b) are the recrystallization distribution maps of AD-0 and AD-90 SLM-IN718 alloys. The average misorientation of recrystallized grains is set to less than 1°, that of substructured grains is set to between 1° and 7.5°, and that of deformed grains is set to greater than 7.5°. In the figures, red, yellow, and blue represent deformed grains, substructured grains, and recrystallized grains, respectively. The proportions of different grains are shown in Table 3. AD specimens are mainly composed of deformed grains and recrystallized grains, with the proportion of deformed grains reaching 56.3% in AD-0 specimens and 58.8% in AD-90 specimens. This is because during the SLM forming process, the high-energy laser rapidly melts the powder, and the thermal expansion and cold contraction of the solidified phase lead to grain deformation. The proportion of recrystallized grains is relatively small, accounting for 9.4% and 7.3% respectively. Figures 6(c) and (d) are the recrystallization distribution maps of HT-0 and HT-90 alloys, respectively. It can be seen from the figures that after heat treatment, the residual stress inside the grains is released to a certain extent, the proportion of deformed grains decreases, and the structure is dominated by substructured grains. The proportion of substructured grains reaches 57.8% in HT-0 specimens and 54.4% in HT-90 specimens.

Figure 6
Recrystallization distribution of SLM-IN718 alloy: (a) AD-0; (b) AD-90; (c) HT-0; (d) HT-90.

Table 3
Percentage of different grains.

Specimen	Recrystallized grains (%)	Deformed grains (%)	Substructured grains (%)
AD-0	9.4	56.3	34.3
AD-90	7.3	58.8	33.9
HT-0	5.0	37.2	57.8
HT-90	6.5	39.1	54.4

Figure 7 shows the Kernel Averaged Misorientation (KAM) maps of SLM-IN718 alloy. It can be seen from Figures 7(a) and (b) that significant residual stress exists inside the as-deposited SLM-IN718 alloys with both build directions, and the residual stress is mainly concentrated inside the deformed grains and at the grain boundaries, which is consistent with the pattern in the recrystallization distribution maps. During the SLM forming process, the high-energy laser causes the powder to melt and solidify rapidly, and most of the deposited parts undergo repeated cycles of remelting and resolidification. As layer-by-layer deposition continues to accumulate, compressive stress keeps accumulating inside the formed material, resulting in significant residual stress within the material. It can be seen from Figures 7(c) and (d) that the residual stress inside the SLM-IN718 alloy is significantly improved after heat treatment, and the improvement degree of the 0°-oriented specimens is greater than that of the 90°-oriented specimens.

Figure 7
KAM of SLM-IN718 alloy: (a) AD-0; (b) AD-90; (c) HT-0; (d) HT-90.

3.2. MECHANICAL PROPERTY ANALYSIS

3.2.1. TENSILE PROPERTY ANALYSIS

Room-temperature tensile tests were conducted on SLM-IN718 alloy specimens, and the tensile test results are presented in Figure 8 and Table 4. Figure 8(a) shows the stress-strain curve of SLM-IN718, and its mechanical properties are displayed in Figure 8(b). It can be seen from the figures that the AD-0 specimen has higher tensile strength but slightly lower elongation than the AD-90 specimen. This is because the AD-0 specimen features a more regular grain distribution and a smaller grain aspect ratio, resulting in greater resistance to dislocation slip and thus higher tensile strength; in contrast, the AD-90 specimen has a disordered grain distribution and a larger aspect ratio, leading to lower resistance to dislocation slip, higher elongation, but lower tensile strength. After heat treatment, the yield strength and tensile strength of SLM-IN718 alloys with different build directions are both higher than those in the as-deposited state, while the elongation slightly decreases. Specifically, the tensile strength of the 0°-oriented specimen increases by 37.4% with a 10.7% decrease in elongation; the tensile strength of the 90°-oriented specimen increases by 26.4% with a 5.5% decrease in elongation. This is due to the fact that after heat treatment, the HT-0 specimen has more and smaller precipitates at the grain boundaries compared to the HT-90 specimen, and the δ phases in the HT-90 specimen are more sparsely spaced. Therefore, the increase in tensile strength of the 0°-oriented specimen after heat treatment is greater than that of the 90°-oriented specimen, and the decrease in elongation is also larger than that of the 90°-oriented specimen.

Figure 8
Mechanical properties of SLM-IN718 alloy: (a) stress-strain curves of the tensile test; (b) graph of UST, 0.2%YS, and elongation.

Table 4
Results of the Tensile Test.

Specimen	AD-0	AD-90	HT-0	HT-90
UTS(MPa)	1001±19	940±25	1375±22	1188±18
0.2%YS(MPa)	716±16	692±17	1187±20	994±22
Elongation(%)	17.7±0.2	19.9±0.3	15.8±0.3	18.8±0.3

To accurately evaluate the comprehensive performance of the specimens, the strength-ductility product is introduced as a comprehensive mechanical property index, which is a metric for measuring the combined strength and plasticity of materials. The strength-ductility products of SLM-IN718 alloys are presented in Table 5. It can be seen from the table that the HT-90 specimen exhibits the largest strength-ductility product, indicating that the alloy has relatively excellent comprehensive performance.

Table 5
Product of Ultimate Tensile Strength (UTS) and fracture elongation of SLM-IN718 alloy (GPa%).

AD-0	AD-90	HT-0	HT-90
17.7	18.7	21.8	22.3

3.2.2. FRACTOGRAPHY ANALYSIS

Figure 9 shows the tensile fracture morphologies of SLM-IN718 alloy. It can be observed from Figures 9(a) and (b) that the dimples are small and closely arranged, exhibiting characteristics of ductile fracture. Moreover, the fracture dimples of AD-90 are slightly larger and deeper than those of AD-0, thus the plasticity of AD-90 specimen is superior to that of AD-0 specimen. From the observations of Figures 9(c) and (d), after heat treatment, some small-sized dimples in the material gradually evolve into larger dimple morphologies through coalescence, and the depth of the dimples also increases significantly. Further comparison reveals that the dimples of HT-90 specimen are larger and deeper than those of HT-0 specimen. On one hand, the larger the dimple size, the greater the plastic deformation of the material during deformation, and the stronger the material's ability to resist deformation, resulting in better plasticity. On the other hand, after heat treatment, inclusions precipitate inside the alloy, which are undissolved carbides formed during material shaping. As brittle phases, these inclusions have weak bonding with the matrix and form cores of microcracks or pores during tensile deformation. These defects become local stress concentration points, promoting early nucleation of pores and leading to fracture of the material at low strain, thus reducing macroscopic plasticity⁴². Under the combined effect of these two factors, heat treatment causes a certain degree of decrease in the plasticity of SLM-IN718 alloy specimens prepared along both build directions. It is found that both before and after heat treatment, the tensile fracture of 90°-oriented specimens exhibits more obvious ductile fracture characteristics compared with 0°-oriented specimens. From the perspective of microscopic fracture mechanism, ductile fracture characteristics are closely related to plasticity, and more significant ductile fracture characteristics often imply better plasticity. Therefore, 90°-oriented specimens are more likely to absorb energy through ductile deformation during tensile process, thereby showing better plasticity than 0°-oriented specimens.

Figure 9
Tensile fracture morphology of SLM-IN718 alloy: (a) AD-0; (b) AD-90; (c) HT-0; (d) HT-90.

3.3. DAMAGE EVOLUTION ANALYSIS

To investigate the damage of SLM-IN718 alloy, the elastic modulus at different strain stages was measured using a nanoindenter. The elastic modulus of the specimens gradually decreases with increasing strain. Compared with the as-deposited specimens, the elastic modulus of the heat-treated specimens is slightly lower. Figure 10 shows the variation trend of the elastic modulus of SLM-IN718 alloy with macroscopic strain. It can be seen from the figure that the elastic modulus of the undeformed as-deposited specimens is approximately 210 GPa, while that of the undeformed heat-treated specimens increases to about 220 GPa. Overall, the variation law of the elastic modulus with macroscopic strain for the four groups of specimens is basically consistent. In the early stage of loading, the elastic modulus of SLM-IN718 alloy decreases slowly with increasing macroscopic strain, and the damage to the alloy is slight at this time. In the later stage of loading, the elastic modulus begins to decrease rapidly with increasing macroscopic strain, and the damage degree of the alloy increases sharply until failure. Among them, the elastic modulus of the AD-90 specimen decreases gently with increasing macroscopic strain in the early stage of loading and enters the rapid decrease stage later; the elastic modulus of the AD-0 specimen decreases steeply with increasing macroscopic strain in the early stage of loading and enters the rapid decrease stage earlier than the AD-90 specimen. After heat treatment, compared with the AD specimens, the elastic modulus of the HT specimens decreases more steeply with increasing macroscopic strain in the early stage of loading, and the specimens with 0° build direction are more significantly affected by heat treatment.

Figure 10
Modulus of elasticity versus macroscopic strain for SLM-IN718 alloy.

According to Lemaitre et al.'s⁴³ elasticity law, to more accurately reflect the actual stress state inside the material and thereby better predict the deformation and strength characteristics of the material, the direct state coupling derived from the concept of effective stress is:

$$\left\{\begin{array}{c}{\epsilon }_e=\frac{\sigma }{E_0},D=0\\ {\epsilon }_e=\frac{\sigma }{E_0(1-D)}\mathrm{,0}<D<1\end{array}\right.$$ (1)

where ${\epsilon }_e$ is the uniaxial elastic strain (dimensionless), $E_0$​ is the intrinsic elastic modulus of the undamaged material (GPa), $\sigma$ is the applied uniaxial stress (MPa), and $D$ is the continuum damage factor (0 ≤ $D$ < 1).

The degraded elastic modulus $E$ of the damaged material, derived from Hooke's law $\sigma =E{\epsilon }_e$, follows:

$E=E_0\left(1-D\right)$ (2)

where $E$ is the operational elastic modulus under mechanical loading (GPa).

This formulation establishes an inverse relationship between damage evolution and modulus degradation. The damage variable can be explicitly defined as:

$D=1-\frac{E}{E_0}$ (3)

This formulation allows for quantitative characterization of damage progression during mechanical loading. The progressive accumulation of internal damage manifests as a gradual reduction in elastic modulus $E$, and corresponds to incremental elevation of damage factor $D$. The modulus-damage coupling relationship facilitates quantitative prediction of strain-path-dependent fracture mechanisms in SLM-IN718 alloy under complex loading regimes.

Figure 11 shows the evolution of the damage factor $D$ for SLM-IN718 alloy. The results indicated that $D$ increased progressively with macroscopic strain until fracture, reflecting cumulative damage accumulation. This originates from the gradual degradation of elastic modulus during tensile deformation, which diminishes the material's resistance to crack initiation and propagation. Notably, the 0°-oriented specimen enter the rapid damage accumulation phase earlier than the 90°-oriented specimen. This discrepancy stems from the more homogeneous microstructure and stabilized grain orientation distribution in 90°-oriented specimens, which is attributed to laser melting characteristics during fabrication. Such microstructural uniformity enhances stress redistribution, mitigates localized stress concentration, and thereby delays rapid damage progression. Post-heat treatment, the onset of rapid damage accumulation occurs slightly earlier. This is attributed to partial elimination of microstructural segregation, which enhances tensile strength but retains brittle Laves phases due to insufficient dissolution at the lower heat treatment temperatures. These residual phases exacerbate stress concentration, accelerating crack nucleation and propagation. The 0°-oriented specimens exhibit greater sensitivity to heat treatment, as evidenced by their more pronounced advancement in the rapid damage accumulation phase compared to 90°-oriented specimens.

Figure 11
Evolution curves of damage factor D for the SLM-IN718 alloy.

To quantify the damage progression from pristine ($D$=0) to fully fractured state ($D$=1), Equation (3) is normalized to derive the elastic modulus-based damage factor $D_e$:

$D_e=\frac{E_0-E}{E_0-E_f}$ (4)

where $D_e$ is the dimensionless normalization factor, and $E_f$ is the elastic modulus at fracture (GPa).

The evolution curves of the normalization factor $D_e$ for SLM-IN718 alloy are plotted in Figure 12. Comparative analysis reveals that the normalized factor $D_e$ of all four specimen groups exhibits a similar evolution pattern to the damage factor $D$ shown in Figure 11, both demonstrating monotonic increase with macroscopic strain. The parameter $D_e$ is mathematically bounded within the dimensionless interval [0,1]. When $D_e$=0, SLM-IN718 alloy is in the initial state, at this time the alloy is not damaged; when 0<$D_e$<1, SLM-IN718 alloy is in the process of stretching, at this time the degree of damage of SLM-IN718 alloy increases with the increase of $D_e$; when $D_e$=1, it reaches the state of complete damage, and SLM-IN718 alloy breaks. Among them, 0°-oriented specimen with the increase of strain, normalized damage factor to grow to 1 before the 90°-oriented specimen.

Figure 12
Evolution curves of normalization factor $D_e$ for the SLM-IN718 alloy.

The normalized damage factor of SLM-IN718 alloy was fitted using an exponential function, yielding the following damage evolution equation:

AD-0:

$D_e=0.02195e^{\frac{\epsilon }{0.04608}}-0.02189$ (5)

AD-90:

$D_e=0.00699e^{\frac{\epsilon }{0.04011}}-0.00549$ (6)

HT-0:

$D_e=0.04106e^{\frac{\epsilon }{0.04889}}-0.04061$ (7)

HT-90:

$D_e=0.00865e^{\frac{\epsilon }{0.03954}}-0.00548$ (8)

Through parametric substitution of the fitted equations, a universal formulation is derived to describe the damage evolution:

$D_e=A{e}^{\frac{-\epsilon }{B}}+C$ (9)

where $A$, $B$, and $C$ are material-specific constants.

To distinguish between the gradual and accelerated damage accumulation phases in SLM-IN718 alloy, this study defines a critical damage factor $D_{ec}$ as the critical threshold. The damage evolution is characterized as follows:

• $0<D_e<D_{ec}$: Corresponds to the initial deformation stage with gradual damage accumulation.

• $D_{ec}<D_e<1$: Indicates the accelerated damage accumulation phase preceding fracture.

• $D_e=1$: Represents complete material failure through fracture.

The Normalized Cockcroft & Latham damage model⁴⁴ was used to calculate the critical damage factor $D_{ec}$ for SLM-IN718 alloy:

$D_{ec}={{\displaystyle \int }}_0^{{\tilde{\epsilon }}_f}\left(\frac{{\sigma }_1}{{\sigma }_u}\right)d\tilde{\epsilon }$ (10)

where $D_{ec}$ is the critical damage factor; ${\tilde{\epsilon }}_f$ is the true strain at fracture; ${\sigma }_1$ is the maximum principal stress in tension; ${\sigma }_u$ is the stress at the onset of necking in tension (tensile strength); and $\tilde{\epsilon }$ is the true strain of the material.

The $D_{ec}$ of the SLM-IN718 alloy can be obtained from Equation (10) and Table 6, and the corresponding critical strain ${\epsilon }_c$ can be obtained by substituting each group of $D_{ec}$ into the corresponding Equations (5), (6), (7), and (8), respectively. The values of $D_{ec}$ and ${\epsilon }_c$ for each group are shown in Figure 13. From the figure, it can be seen that the $D_{ec}$ and ${\epsilon }_c$ of the 90°-oriented specimen are larger than the 0°-oriented specimen, which is consistent with the results of the previous analysis of Figure 11. The $D_{ec}$ and ${\epsilon }_c$ of the alloys decreased slightly after heat treatment. The $D_{ec}$ and ${\epsilon }_c$ were the largest in the AD-90 group, which entered the stage of rapid damage accumulation the latest.

Table 6
Values used in equation (10).

Specimen	AD-0	AD-90	HT-0	HT-90
${\tilde{\epsilon }}_f$	0.16290	0.18149	0.14693	0.17227
${\sigma }_1/MPa$	1169.19576	1122.59153	1583.24780	1404.55131
${\sigma }_u/MPa$	1001.56537	939.78268	1375.37443	1188.28368

Figure 13
Critical damage factor $D_{ec}$ and critical strain ${\epsilon }_c$ for SLM-IN718 alloy.

4. CONCLUSIONS

This study investigates the effects of build direction and heat treatment on the microstructure, mechanical properties, and damage evolution behavior of SLM-IN718 alloy. The main conclusions are as follows:

1. Build direction significantly affects the microstructure and mechanical properties of as-deposited alloys. AD-0 specimens exhibit regular elongated scanning paths, with higher tensile strength but lower elongation; AD-90 specimens show fish-scale molten pool boundaries, with higher elongation but slightly lower tensile strength. Both contain Laves phases.

2. Heat treatment effectively improves the alloy properties. After 980°C/1h air cooling treatment, Laves phases dissolve and promote the precipitation of δ phases at grain boundaries. The tensile strength of HT-0 specimens increases by 37.4%, and that of HT-90 specimens increases by 26.4%. The strength-ductility product of HT-90 specimens reaches 22.3 GPa%, showing the optimal comprehensive performance.

3. The damage evolution behavior is direction-dependent. 90°-oriented specimens show slower elastic modulus degradation in the early stage of damage, with a higher critical damage factor, and enter the stage of rapid damage accumulation later. Heat treatment advances the stage of rapid damage accumulation.

Data Availability

The full dataset supporting the findings of this study is available upon request to the corresponding author – Lei Li, [e-mail]

5. ACKNOWLEDGMENTS

This work was supported by the National Natural Science Foundation [grant numbers 12262029]; the Inner Mongolia Natural Science Foundation [grant numbers 2023MS01007]; the Inner Mongolia Basic Research Operations [grant numbers JY20230010]; the Postgraduate Research Innovation Project of Inner Mongolia Autonomous Region [grant numbers S20231126Z, KC2024040S]; and the Fundamental Research Funds for the Higher Education Institutions Directly under the Inner Mongolia Autonomous Region [grant numbers ZTY2025044].

6. REFERENCES

• ¹ Va S, Na B. A study on microstructure and mechanical properties of Inconel 718 Superalloy Fabricated by Novel CMT-WAAM Process. Mater Res. 2024;27:e20230258. http://doi.org/10.1590/1980-5373-mr-2023-0258
  » http://doi.org/10.1590/1980-5373-mr-2023-0258
• ² Venkatesh G, Subramanian R, Abuthakir J, Berchmans LJ. Wear Analysis of Plasma Sprayed Calcium and Strontium Zirconates on Inconel 718. Mater Res. 2024;27:e20220282. http://doi.org/10.1590/1980-5373-mr-2022-0282
  » http://doi.org/10.1590/1980-5373-mr-2022-0282
• ³ Ferrarotti A, Giuffrida F, Sharghivand E, Mussino G, Vedani M, Baricco M, et al. Mechanical and microstructural properties of IN718 additively manufactured lattice structures. Mater Sci Eng A. 2025;919:147491. http://doi.org/10.1016/j.msea.2024.147491
  » http://doi.org/10.1016/j.msea.2024.147491
• ⁴ Sana M, Ali MA, Ehsan S, Tlija M, Khan AM. Investigation of EDM erosion behavior for Ni-based superalloy using experimental and machine learning approach. Mater Today Commun. 2024;41:110819. http://doi.org/10.1016/j.mtcomm.2024.110819
  » http://doi.org/10.1016/j.mtcomm.2024.110819
• ⁵ Sonar T, Balasubramanian V, Malarvizhi S, Venkateswaran T, Sivakumar D. An overview on welding of Inconel 718 alloy-Effect of welding processes on microstructural evolution and mechanical properties of joints. Mater Charact. 2021;174:110997. http://doi.org/10.1016/j.matchar.2021.110997
  » http://doi.org/10.1016/j.matchar.2021.110997
• ⁶ Nie Y, Xu C, Liu Z, Yang L, Li T, He Y. Investigation of support structure configurations for selective laser melting of In718. Alex Eng J. 2025;112:281-92. http://doi.org/10.1016/j.aej.2024.11.006
  » http://doi.org/10.1016/j.aej.2024.11.006
• ⁷ Sefene EM. State-of-the-art of selective laser melting process: a comprehensive review. J Manuf Syst. 2022;63:250-74. http://doi.org/10.1016/j.jmsy.2022.04.002
  » http://doi.org/10.1016/j.jmsy.2022.04.002
• ⁸ Kladovasilakis N, Charalampous P, Tsongas K, Kostavelis I, Tzovaras D, Tzetzis D. Influence of selective laser melting additive manufacturing parameters in Inconel 718 superalloy. Materials (Basel). 2022;15(4):1362. http://doi.org/10.3390/ma15041362 PMid:35207901.
  » http://doi.org/10.3390/ma15041362
• ⁹ Wang R, Chen C, Liu M, Zhao R, Xu S, Hu T, et al. Effects of laser scanning speed and building direction on the microstructure and mechanical properties of selective laser melted Inconel 718 superalloy. Mater Today Commun. 2022;30:103095. http://doi.org/10.1016/j.mtcomm.2021.103095
  » http://doi.org/10.1016/j.mtcomm.2021.103095
• ¹⁰ Heo S, Lim Y, Kwak N, Jeon C, Choi M, Jo I. Impact of heat treatment and building direction on tensile properties and fracture mechanism of Inconel 718 produced by SLM process. Metals (Basel). 2024;14(4):440. http://doi.org/10.3390/met14040440
  » http://doi.org/10.3390/met14040440
• ¹¹ Hartunian P, Eshraghi M. Effect of build orientation on the microstructure and mechanical properties of selective laser-melted Ti-6Al-4V alloy. Journal of Manufacturing and Materials Processing. 2018;2(4):69. http://doi.org/10.3390/jmmp2040069
  » http://doi.org/10.3390/jmmp2040069
• ¹² Singh PK, Kumar S, Jain PK, Dixit US. Effect of build orientation on metallurgical and mechanical properties of additively manufactured Ti-6Al-4V alloy. J Mater Eng Perform. 2024;33(7):3476-93. http://doi.org/10.1007/s11665-023-08218-4
  » http://doi.org/10.1007/s11665-023-08218-4
• ¹³ Sun W, Ma YE, Li P, Moumni Z, Zhang W. Effects of build direction and heat treatment on the defect characterization and fatigue properties of laser powder bed fusion Ti6Al4V. Aerospace (Basel). 2024;11(10):854. http://doi.org/10.3390/aerospace11100854
  » http://doi.org/10.3390/aerospace11100854
• ¹⁴ Kuntoğlu M, Salur E, Canli E, Aslan A, Gupta MK, Waqar S, et al. A state of the art on surface morphology of selective laser-melted metallic alloys. Int J Adv Manuf Technol. 2023;127(3):1103-42. http://doi.org/10.1007/s00170-023-11534-7
  » http://doi.org/10.1007/s00170-023-11534-7
• ¹⁵ Shi JJ, Zhou ZQ, Xu K, Zhou GY, Zhou ZJ, Li CP, et al. Effect of heat treatment on microstructure and small punch creep property of selective laser melted Inconel 718 alloy. Mater Sci Eng A. 2022;853:143748. http://doi.org/10.1016/j.msea.2022.143748
  » http://doi.org/10.1016/j.msea.2022.143748
• ¹⁶ Wang W, Chen Z, Lu W, Meng F, Zhao T. Heat treatment for selective laser melting of Inconel 718 alloy with simultaneously enhanced tensile strength and fatigue properties. J Alloys Compd. 2022;913:165171. http://doi.org/10.1016/j.jallcom.2022.165171
  » http://doi.org/10.1016/j.jallcom.2022.165171
• ¹⁷ Zhao R, Zhao Z, Bai P, Du W, Zhang L, Qu H. Effect of heat treatment on the microstructure and properties of Inconel 718 alloy fabricated by selective laser melting. J Mater Eng Perform. 2022;31(1):353-64. http://doi.org/10.1007/s11665-021-06212-2
  » http://doi.org/10.1007/s11665-021-06212-2
• ¹⁸ Diepold B, Vorlaufer N, Neumeier S, Gartner T, Göken M. Optimization of the heat treatment of additively manufactured Ni-base superalloy IN718. Int J Miner Metall Mater. 2020;27(5):640-8. http://doi.org/10.1007/s12613-020-1991-6
  » http://doi.org/10.1007/s12613-020-1991-6
• ¹⁹ Wang W, Wang S, Zhang X, Chen F, Xu Y, Tian Y. Process parameter optimization for selective laser melting of Inconel 718 superalloy and the effects of subsequent heat treatment on the microstructural evolution and mechanical properties. J Manuf Process. 2021;64:530-43. http://doi.org/10.1016/j.jmapro.2021.02.004
  » http://doi.org/10.1016/j.jmapro.2021.02.004
• ²⁰ Aripin MA, Sajuri Z, Syarif J, Baghdadi AH, Mohamed IF. Evaluation of microstructure and porosity for 3D printed stainless steel. Mater Today Proc. 2022;66:3082-6. http://doi.org/10.1016/j.matpr.2022.07.396
  » http://doi.org/10.1016/j.matpr.2022.07.396
• ²¹ Aripin MA, Sajuri Z, Jamadon NH, Baghdadi AH, Mohamed IF, Syarif J, et al. Microstructure and mechanical properties of selective laser melted 17–4 PH stainless steel; Build direction and heat treatment processes. Mater Today Commun. 2023;36:106479. http://doi.org/10.1016/j.mtcomm.2023.106479
  » http://doi.org/10.1016/j.mtcomm.2023.106479
• ²² Aripin MA, Sajuri Z, Jamadon NH, Baghdadi AH, Syarif J, Mohamed IF, et al. Effects of build orientations on microstructure evolution, porosity formation, and mechanical performance of selective laser melted 17-4 PH stainless steel. Metals (Basel). 2022;12(11):1968. http://doi.org/10.3390/met12111968
  » http://doi.org/10.3390/met12111968
• ²³ Wang X, Zhu T, Lu L, Zhang J, Ding H, Xiao S, et al. A damage sequence interaction model for predicting the mechanical property of in-service aluminium alloy 6005A-T6. Eng Fract Mech. 2023;291:109565. http://doi.org/10.1016/j.engfracmech.2023.109565
  » http://doi.org/10.1016/j.engfracmech.2023.109565
• ²⁴ Chen H, Yang F, Wu Z, Yang B, Huo J. A nonlinear fatigue damage accumulation model under variable amplitude loading considering the loading sequence effect. Int J Fatigue. 2023;177:107945. http://doi.org/10.1016/j.ijfatigue.2023.107945
  » http://doi.org/10.1016/j.ijfatigue.2023.107945
• ²⁵ Zhang SJ, Zhou C, Xia QX, Chen SM. Quantification and characterization of full field ductile damage evolution for sheet metals using an improved direct current potential drop method. Exp Mech. 2015;55(3):611-21. http://doi.org/10.1007/s11340-014-9982-z
  » http://doi.org/10.1007/s11340-014-9982-z
• ²⁶ Mao K, Nie S, Yang B, Tang M, Chen Z, Elchalakani M. Experimental study on residual monotonic mechanical properties of high-performance Q690 steel with pre-fatigue damage. Thin-walled Struct. 2023;192:111167. http://doi.org/10.1016/j.tws.2023.111167
  » http://doi.org/10.1016/j.tws.2023.111167
• ²⁷ Xie H, Xie H, Zhang Z, Yao Q, Cao Z, Gao H, et al. Fatigue fracture behaviors and damage evolution of coal samples treated with drying–wetting cycles investigated by acoustic emission and nuclear magnetic resonance. Int J Rock Mech Min Sci. 2025;185:105976. http://doi.org/10.1016/j.ijrmms.2024.105976
  » http://doi.org/10.1016/j.ijrmms.2024.105976
• ²⁸ Pei C, Zeng W, Yuan H. A damage evolution model based on micro-structural characteristics for an additive manufactured superalloy under monotonic and cyclic loading conditions. Int J Fatigue. 2020;131:105279. http://doi.org/10.1016/j.ijfatigue.2019.105279
  » http://doi.org/10.1016/j.ijfatigue.2019.105279
• ²⁹ Ma Z, Zhang C, Gamage RP, Zhang G. Uncovering the creep deformation mechanism of rock-forming minerals using nanoindentation. Int J Min Sci Technol. 2022;32(2):283-94. http://doi.org/10.1016/j.ijmst.2021.11.010
  » http://doi.org/10.1016/j.ijmst.2021.11.010
• ³⁰ Ying P, Xia Y. Elastic modulus identification of particles in particulate composite through nanoindentation. Int J Mech Sci. 2023;260:108660. http://doi.org/10.1016/j.ijmecsci.2023.108660
  » http://doi.org/10.1016/j.ijmecsci.2023.108660
• ³¹ ASTM: American Society for Testing and Materials. ASTM E3-11: Standard Guide for Preparation of Metallographic Specimens. West Conshohocken: ASTM; 2025.
• ³² ASTM: American Society for Testing and Materials. ASTM E2546-15: Standard Practice for Instrumented Indentation Testing. West Conshohocken: ASTM; 2023.
• ³³ ASTM: American Society for Testing and Materials. ASTM E8/E8M-21: Standard Test Methods for Tension Testing of Metallic Materials. West Conshohocken: ASTM; 2021.
• ³⁴ Zheng Z, Sun B, Mao L. Effect of scanning strategy on the manufacturing quality and performance of printed 316L stainless steel using SLM process. Materials (Basel). 2024;17(5):1189. http://doi.org/10.3390/ma17051189 PMid:38473660.
  » http://doi.org/10.3390/ma17051189
• ³⁵ Cao Y, Bai P, Liu F, Hou X. Investigation on the precipitates of IN718 alloy fabricated by selective laser melting. Metals (Basel). 2019;9(10):1128. http://doi.org/10.3390/met9101128
  » http://doi.org/10.3390/met9101128
• ³⁶ Dong P, Li Z, Li D, Yu Y, Zhang W. Track trajectory, molten pool and defect characterization in NiTi single-track selective laser melting (SLM) experiments. J Mater Res Technol. 2024;33:5210-22. http://doi.org/10.1016/j.jmrt.2024.10.112
  » http://doi.org/10.1016/j.jmrt.2024.10.112
• ³⁷ Chen M, Du Q, Shi R, Fu H, Liu Z, Xie J. Phase field simulation of microstructure evolution and process optimization during homogenization of additively manufactured Inconel 718 alloy. Front Mater. 2022;9:1043249. http://doi.org/10.3389/fmats.2022.1043249
  » http://doi.org/10.3389/fmats.2022.1043249
• ³⁸ Kakehi K, Chowdhury HT, Shinoda Y, Naidu PT, Kakuta N, Ishisako S. Effects of base plate temperature on microstructure evolution and high-temperature mechanical properties of IN718 processed by laser powder bed fusion using simulation and experiment. Int J Adv Manuf Technol. 2024;130(11):5777-93. http://doi.org/10.1007/s00170-024-13028-6
  » http://doi.org/10.1007/s00170-024-13028-6
• ³⁹ Cao Y, Bai P, Wei A, Liu F, Hou X. Effect of heat treatment on microstructure and evolution of precipitated phase of IN718 alloy for laser additive manufacturing. Heat Treatment of Metals. 2023;48(2):180-9.
• ⁴⁰ Thijs L, Verhaeghe F, Craeghs T, Van Humbeeck J, Kruth JP. A study of the microstructural evolution during selective laser melting of Ti–6Al–4V. Acta Mater. 2010;58(9):3303-12. http://doi.org/10.1016/j.actamat.2010.02.004
  » http://doi.org/10.1016/j.actamat.2010.02.004
• ⁴¹ Choi JP, Shin GH, Yang S, Yang DY, Lee JS, Brochu M, et al. Densification and microstructural investigation of Inconel 718 parts fabricated by selective laser melting. Powder Technol. 2017;310:60-6. http://doi.org/10.1016/j.powtec.2017.01.030
  » http://doi.org/10.1016/j.powtec.2017.01.030
• ⁴² Liu F, Lin X, Yang G, Huang C, Chen J, Huang W. Microstructure and mechanical properties of nickel-based superalloy Inconel 718 for laser stereolithography in different atmospheres. Chin Shu Hsueh Pao. 2010;46(9):1047-54. http://doi.org/10.3724/SP.J.1037.2010.00046
  » http://doi.org/10.3724/SP.J.1037.2010.00046
• ⁴³ Lemaitre J. Tutorial in injury mechanics. Beijing: Science Press; 1996.
• ⁴⁴ Zhang J, Wu J, Luo Y, Zhao B, Guo D, Zhao S, et al. Determination of critical damage values of Ti600 alloy based on Normalized Cockcroft & Latham toughness damage criteria. Mater Des. 2019;47(7):121-5.

Edited by

• Associate Editor:
  Aloisio Klein.

• Editor-in-Chief:
  Luiz Antonio Pessan.

Publication Dates

• Publication in this collection
  24 Oct 2025
• Date of issue
  2025

History

• Received
  13 July 2025
• Accepted
  06 Sept 2025