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Article

Elevated-Temperature Tensile Behavior and Properties of Inconel 718 Fabricated by In-Envelope Additive–Subtractive Hybrid Manufacturing and Post-Process Precipitation Hardening

by
Sheida Sarafan
1,*,
Priti Wanjara
1,
Roger Pelletier
1,
Sila Ece Atabay
1,
Javad Gholipour
1,
Josh Soost
2,
Robert Amos
3 and
Prakash Patnaik
1
1
Transportation and Manufacturing Division, National Research Council Canada, Ottawa, ON K1A 0R6, Canada
2
Matsuura Machinery USA Inc., St. Paul, MN 55102, USA
3
Department of National Defence, Directorate of Technical Airworthiness and Engineering Support (DTAES), Ottawa, ON K1A 0K2, Canada
*
Author to whom correspondence should be addressed.
J. Manuf. Mater. Process. 2024, 8(6), 297; https://doi.org/10.3390/jmmp8060297 (registering DOI)
Submission received: 19 November 2024 / Revised: 16 December 2024 / Accepted: 19 December 2024 / Published: 21 December 2024
(This article belongs to the Special Issue Industry 4.0: Manufacturing and Materials Processing)
Browse Figures
Figure 1
<p>(<b>a</b>,<b>b</b>) Morphology and (<b>c</b>) cohesive index of the starting IN718 powder.</p> "> Figure 2
<p>Process flow detailing the different stages in the experimental methodology.</p> "> Figure 3
<p>(<b>a</b>) CAD layout of the build plate with 24 vertically built tensile specimens. (<b>b</b>) The 24 vertically built tensile specimens after the build. (<b>c</b>) A sleeve-shaped support structure designed with a small gap to ease removal of the tensile specimens. (<b>d</b>) Easy support removal after EDM from the build plate. (<b>e</b>) Tensile specimen geometry based on ASTM E8M-22 [<a href="#B49-jmmp-08-00297" class="html-bibr">49</a>].</p> "> Figure 4
<p>Vertically built tensile specimens fabricated to have three surface finish conditions in the gauge section: AB (<b>left</b>), hybrid (<b>middle</b>) and PM (<b>right</b>).</p> "> Figure 5
<p>Different precipitation-hardening heat treatment (PHT) cycles used in this study.</p> "> Figure 6
<p>Map of the surface topography of vertically built IN718 specimens with (<b>a</b>) AB, (<b>b</b>) hybrid and (<b>c</b>) PM surfaces.</p> "> Figure 7
<p>Porosity inspections of vertically built IN718 specimens with (<b>a</b>,<b>b</b>) AB, (<b>c</b>,<b>d</b>) hybrid and (<b>e</b>,<b>f</b>) PM conditions.</p> "> Figure 7 Cont.
<p>Porosity inspections of vertically built IN718 specimens with (<b>a</b>,<b>b</b>) AB, (<b>c</b>,<b>d</b>) hybrid and (<b>e</b>,<b>f</b>) PM conditions.</p> "> Figure 8
<p>Differential distribution of the pore volume fraction and number fraction as a function of the distance R from the specimen outer surface.</p> "> Figure 9
<p>Representative (<b>a</b>) engineering stress–strain and (<b>b</b>) true stress–strain curves of vertically built IN718 with the different surface conditions.</p> "> Figure 10
<p>DIC analysis of the local strain distribution maps of the gauge section of the vertically built IN718 tensile specimens just before fracture: (<b>a</b>) AB, (<b>b</b>) hybrid and (<b>c</b>) PM surface conditions.</p> "> Figure 11
<p>Tensile properties at 650 °C for the vertically built IN718 with the different precipitation-hardening conditions: (<b>a</b>) PHT1, (<b>b</b>) PHT2 and (<b>c</b>) PHT3; and different surface conditions: (<b>d</b>) AB, (<b>e</b>) hybrid and (<b>f</b>) PM.</p> "> Figure 12
<p>Representative (<b>a</b>–<b>c</b>) engineering stress–strain and (<b>d</b>–<b>f</b>) true stress–strain curves at 650 °C for the vertically built IN718 with the different surface conditions and PHTs.</p> "> Figure 13
<p>µXCT cross-section of vertically built IN718 specimens tested at 650 °C with (<b>a</b>) AB, (<b>b</b>) hybrid and (<b>c</b>) PM surface conditions.</p> "> Figure 14
<p>Fractographs after room-temperature tensile testing of the vertically built IN718 specimens with (<b>a</b>) AB (<b>b</b>) hybrid and (<b>c</b>) PM surface finish conditions.</p> "> Figure 15
<p>High-magnification fractographs after room-temperature tensile testing of the vertically built IN718 specimens with (<b>a</b>) AB (<b>b</b>) hybrid and (<b>c</b>) PM surface finish conditions.</p> "> Figure 16
<p>Fractographs after high-temperature (650 °C) tensile testing of vertically built IN718 specimens with an AB surface finish and under (<b>a</b>) PHT1, (<b>b</b>) PHT2 and (<b>c</b>) PHT3 conditions; with a hybrid surface finish and under (<b>d</b>) PHT1, (<b>e</b>) PHT2 and (<b>f</b>) PHT3 conditions; as well as a with a PM surface finish and under (<b>g</b>) PHT1, (<b>h</b>) PHT2 and (<b>i</b>) PHT3 conditions.</p> "> Figure 17
<p>High magnification of fractographs after high-temperature (650 °C) tensile testing of vertically built IN718 specimens with an AB surface finish and under (<b>a</b>) PHT1, (<b>b</b>) PHT2 and (<b>c</b>) PHT3 conditions; with a hybrid surface finish and under (<b>d</b>) PHT1, (<b>e</b>) PHT2 and (<b>f</b>) PHT3 conditions; as well as a with a PM surface finish and under = (<b>g</b>) PHT1, (<b>h</b>) PHT2, (<b>i</b>) PHT3 conditions.</p> ">
Versions Notes

Abstract

:
The present study focuses on advancing one of the most popular AM techniques, namely, laser powder bed fusion (LPBF) technology, which has the ability to produce complex geometry parts with minimum material waste but continues to face challenges in minimizing the surface roughness. For this purpose, a novel hybrid manufacturing technology, which applies in a single setup (in-envelope) both LPBF technology and high-speed machining, was examined in this research for the fabrication of tensile specimens with three different surface finish conditions: as-built, hybrid (in-envelope machining) and post-machining (out-of-envelope) on Inconel® alloy 718, hereafter referred to as IN718. As the application of the IN718 alloy in service is typically specified in the precipitation-hardened condition, three different heat treatments were applied to the tensile specimens based on the most promising thermal cycles identified previously for room-temperature tensile properties by the authors. The as-built (AB) specimens had the highest average surface roughness (Ra) of 5.1 μm ± 1.6 μm, which was a significant improvement (five-fold) on the hybrid (1.0 μm ± 0.2 μm) and post-machined (0.8 μm ± 0.5 μm) surfaces. The influence of this surface roughness on the mechanical properties was studied both at ambient temperature and at 650 °C, which is close to the maximum service temperature of this alloy. Regardless of the surface conditions, the room-temperature mechanical properties of the as-fabricated IN718 specimens were within the range of properties reported for standard wrought IN718 in the annealed condition. Nonetheless, detailed examination of the strain localization behavior during tensile testing using digital image correlation showed that the IN718 specimens with AB surfaces exhibited lower ductility (global and local) relative to the hybrid and post-machined ones, most likely due to the higher surface roughness and near-surface porosity in the former. At 650 °C, even though the mechanical properties of all the heat-treated IN718 specimens surpassed the minimum specifications for the wrought precipitation-hardened IN718, the AB surface condition showed up to 4% lower strength and 33–50% lower ductility compared with the hybrid and PM surface conditions. Microfocus X-ray computed tomography (µXCT) of the fractured specimens revealed the presence of numerous open cracks on the AB surface and a predisposition for the near-surface pores to accelerate rupture, leading to premature failure at lower strains.

1. Introduction

Currently, manufacturing is in a renaissance era, as engineers and professionals transition from conventional production operations based on a subtractive methodology to learning and applying new skills in digital technologies such as additive manufacturing (AM), which is poised to be a key enabler for building creative solutions layer by layer with advanced materials/alloys. Among the current state-of-the-art in metal AM technologies [1,2,3,4,5,6], a key process is laser powder bed fusion (LPBF), where laser energy is used to selectively melt and solidify a thin layer of metal powder along a pre-defined laser scan path [7,8,9]. Layer-upon-layer building with LPBF processing offers design flexibility and the high resolution needed for fabricating complex net and/or near-net-shape geometry parts with reduced raw material usage and wastage [10,11].
These advantages of LPBF AM are especially significant for the shaping of hard or difficult-to-machine materials [12], which experience manufacturing challenges regardless of the approach used to remove material and to produce a finished product with the desired dimensional requirements. One such group of materials are nickel-based superalloys that exhibit poor machining performance due to their high strength at elevated temperatures, poor thermal conductivity, high toughness and high degree of strain hardening [13]. However, the same properties that limit the machinability of these materials provide superior resistance to mechanical and chemical degradation, even when operating at ultra-high homologous temperature conditions [14]. Although nickel-based superalloys were originally developed for hot gas components in aerospace and industrial gas turbines, the demand for them has expanded to many other critical structures in defense, oil and gas mining, petrochemical processing, metalworking and power generation, to name a few. Conventionally, superalloy parts are manufactured using three general routes—cast, wrought and powder metallurgy processing [15]—followed by a sequence of secondary operations that include heat treatment and machining to the final geometry. Additive processing presents a promising alternate manufacturing path for these materials, especially for those with chemistries that provide reasonably good weldability, such as Inconel® 718 (IN718), the workhorse alloy of the industry [16,17,18,19].
IN718 is a precipitation-strengthened polycrystalline nickel-based superalloy that consists of a complex multiphase microstructure [20,21]. The primary phase is the austenitic gamma (γ) matrix, which is a face-centered cubic (FCC) Ni–Fe–Cr solid solution. After appropriate heat treatment, two main coherent strengthening phases form within the γ matrix: principal metastable γ″ (Ni3Nb) precipitates that are disc-shaped with a body centered tetragonal (BCT) structure and auxiliary spherical or cuboidal γ′ (Ni3(Al,Ti)) precipitates that have an ordered L12 structure. Depending on the applied heat treatment, the microstructure can also contain plate- or needle-like δ phases (Ni3Nb), the irregularly shaped Laves phase (Ni,Fe,Cr)2(Nb,Mo,Ti) and fine carbides (MC, M23C6 and M6C) [22,23].
Research on LPBF processing of IN718 over the past decade [24] has reported progress on process optimization for the fabrication of high-density crack-free parts [25,26,27,28,29,30], as well as on suitable post-process thermal treatments for improving the mechanical properties [31,32,33,34,35,36,37]. However, less attention has been paid to studying the thermal performance of LPBF-processed IN718 [17,38,39,40,41,42]. As well, the scientific research is still far from developing hybrid additive–subtractive processes that integrate LPBF and machining operations into a single processing setup (i.e., in-envelope) to realize a suitable machined surface finish [43] on exterior surfaces, internal microchannels and thin walls. For conventionally manufactured IN718, the surface finish has been key to realizing the necessary surface characteristics for achieving fatigue-critical performance [44]. As well, the surface finish is known to impact the laminar-to-turbulent transition behavior [45], thermal efficiency and oxidation/corrosion resistance [46] of wrought IN718. Unsurprisingly, LPBF-processed IN718 usually requires post-process machining to meet the surface roughness requirements for long-term or high-temperature service, and the time and cost associated with these subsequently applied secondary operations can obliterate the benefits from additive processing [47]. This is contributing to a growing interest in developing additive–subtractive hybrid manufacturing technologies [7,8,12], but the only study available in the open literature for LPBF-fabricated IN718 [48] is related to the research work by the current authors on hybrid process development and on establishing suitable heat treatments to produce optimal tensile properties at room temperature.
The present study is a continuation of our previous endeavors to further the development of additive–subtractive hybrid manufacturing by examining the influence of surface roughness and post-process heat treatment on the tensile response of IN718 under service conditions, namely, at a high temperature of 650 °C. IN718 tensile specimens were fabricated with three different surface finishes: as-built (AB), hybrid and post-machined (PM). These specimens were subjected to three different heat treatments: two based on standard practices for wrought IN718 and one non-standard process specially adapted for IN718 fabricated by LPBF processing. The surface finish of the specimens in the AB, hybrid and PM conditions was examined first. Finally, the tensile properties at room temperature and 650 °C, as well as the fracture behavior, were studied for the AB, hybrid and PM IN718 with and without the post-process heat treatments.

2. Materials and Methods

The starting material used in this study was commercially available argon gas-atomized IN718 powder from Matsuura (St. Paul, MN, USA), with a nominal particle size of −45/+10 µm and an elemental composition as given in Table 1. The morphology of the IN718 powder particles was analyzed using a JCM-7000 NeoScope™ Benchtop (Fukuoka, Japan) scanning electron microscope (SEM). As shown in Figure 1a,b, the majority of the as-received powder particles were spherical in shape, with a few irregularities and satellites attached to their surfaces. The particle size distribution (PSD) of the powder was measured using an LA-920 Horiba (Kyoto, Japan) laser-scattering particle size analyzer. The PSD analysis showed a normal distribution, with D10, D50 and D90 values as given in Table 2. Flowability and the apparent density (AD) of the powder were assessed using Hall and Carney funnels according to the specifications in ASTM B213 [23] and ASTM B964 [24], and the results (Table 2) indicate that the powder has suitable flowability to be used in the LPBF process. Also, dynamic flowability analysis was conducted using a GranuDrum (Avans, Belgium) rotating-drum instrument. Figure 1c shows the cohesive index (CI) of the powder as a function of the drum rotational speed from 2 rpm to 30 rpm. The measured CI values, which ranged between 18 and 22, were statistically similar for the rotational speeds tested in this study. It is suggested that metal powders having CI values lower than 24 show good flowability and spreadability characteristics, resulting in a uniform powder layer for LPBF processing [25,26].
Figure 2 presents the process flow of the experimental methodology at the different stages (e.g., 3D computer-aided design (CAD) model, manufacturing, surface roughness and µXCT characterization, and tensile testing) in this study. The CAD layout of the tensile specimens on the build plate is given in Figure 3a. Each build, as shown in Figure 3b, comprised twenty-four tensile specimens manufactured in a vertical orientation with sleeve-shaped support structures added for the overhanging edges, as illustrated in Figure 3c,d. Leaving a small gap between the support sleeves and the tensile specimens eased support removal without leaving any surface residue or scrapes when sliding the specimens from the sleeves, as shown in Figure 3d. The tensile specimens conformed to a standard sub-size geometry according to ASTM E8M-22 [49], with a gauge length of 25 mm and diameter of 4 mm, as shown in Figure 3e.
Three different builds, each with twenty-four vertically built tensile specimens, were manufactured in the present study to investigate the three surface conditions of interest: AB, hybrid and PM. The first build comprised tensile specimens fabricated using LPBF processing only, which produced an AB surface finish. For the second build, a hybrid surface finish along the gauge length of the tensile specimens was produced by additive–subtractive hybrid manufacturing, which combined LPBF processing with (in-envelope) micro-milling of the surface after every ten deposition layers. The third build involved fabricating the tensile specimens using LPBF processing followed by out-of-envelope (or sequential, after completing the LPBF processing) machining of the gauge section to render a PM surface finish. After completion of the three builds, the tensile specimens were separated from the build plate using electro-discharge machining (EDM) (FANUC Robocut C400iB, Oshino-mura, Yamanashi, Japan) with a brass wire having a diameter of 0.2 mm. Figure 4 shows representative tensile specimens with AB, hybrid and PM surface conditions in the gauge section.
The three builds were manufactured with a LUMEX Avance-25 system (supplied by Matsuura, MN, USA) that combines LPBF processing with a high-speed micro-milling system in-envelope (i.e., in a single manufacturing cell). This system is equipped with a single Yb fiber laser having a maximum power output of 400 W and a beam diameter at a focus of 200 μm. Each build consisted of twenty-four vertically built specimens that were manufactured under a protective nitrogen gas atmosphere (with less than 1% oxygen) on a precision-ground demagnetized 4140 steel build plate maintained at 50 °C. A laser power of 240 W, a hatch distance of 110 µm, a laser scanning speed of 700 mm/s and a layer thickness of 50 µm were used for the LPBF process in all three builds (AB, hybrid, PM) for both the infill and contour scans. In the second build (hybrid additive–subtractive processing), the entire reduced cross-section of the tensile specimens underwent in-envelope milling using an iteration of ten layers of LPBF deposition followed by dry milling at a feed rate of 225 mm/min of the surface. The third build (post-machined) was manufactured by LPBF processing followed by EDM removal of the tensile specimens from the build plate; then, the entire reduced cross-section of each tensile specimen underwent machining using out-of-envelope turn-milling with coolant at 1000 rpm with a four-flute endmill, cutting fluid and an axial depth of cut of ~0.4 mm.
The surface quality of the tensile specimens with the three different surface conditions was assessed by measuring the linear (Ra and Rz) and areal (Sa and Sz) roughness parameters. These parameters were evaluated using a three-dimensional laser scanning confocal microscope (Keyence VK-X250, Osaka, Japan) as per ISO 25178-2 [50,51]. The arithmetic mean height (Ra) and maximum height (Rz) values were derived from a primary line profile with a length of 6 mm by suppressing the long-wave component using a high-pass filter with a cut-off of λc = 0.8 mm. In regard to the areal roughness parameters, Sa and Sz are the counterparts of Ra and Rz. Measurement of Sa and Sz was performed over a definition area of 1 mm by 0.1 mm on the surfaces in the reduced section of the different tensile specimens (AB, hybrid and PM). Also, the bulk density was measured using the Archimedes method with a theoretical density value of 8.19 g/cm3 [29] to calculate the relative density of the fabricated IN718 tensile specimens.
Microfocus X-ray computed tomography (µXCT) was used to characterize the porosity in the specimens with the three surface conditions (AB, hybrid and PM). For each of these, one specimen that was not tensile tested and one tensile-tested specimen at 650 °C were CT-scanned. The field of view during these scans was limited to a 5 mm long region close to the middle of the specimens, in the gauge length. The resulting voxel size was 2.8 µm. The system used for the scans was a NIKON HMXST 225 (Tokyo, Japan) configured to operate with a reflection source and a tungsten target. The acquisition conditions were as follows: potential, 165 kV; current, 50 µA; integration time, 1 s; panel gain, 30 dB; projections, 3142; image averaging, 4:1, ring artefact removal function, activated and with the X-ray beam filtered through 0.25 mm of silver. Image reconstruction was performed using the CT Pro 3D version 3.1.12 software, and the image improvement capabilities were set to level 2 for beam-hardening correction, level 3 for noise reduction and 4% for scatter reduction.
Image analysis was performed using the Dragonfly version 2022.2.0.1399 software. Porosity in the µXCT images was segmented using a global gray-value thresholding method. The gray values in the six models were between 0 and approximately 1 on the pseudo-density scale generated by the reconstruction software. Initial attempts to use Otsu’s method [52] to select the threshold value to separate the pores from the matrix did not yield adequate results, as the vast majority of the smaller pores, which appeared lighter gray, were not captured. To select the threshold gray value, the six models were segmented at several gray values above the one obtained by Otsu’s method, and a selection was made with the objective of maximizing the number of true pores and minimizing the number of false pores resulting from image noise and non-uniformities. Since the focus of this work is more on the spatial distribution and the relative size of the individual pores rather than on the exact determination of the absolute porosity, the final choice of the threshold gray value was less critical. Also, to reduce the number of false pores in the analysis, detected pores with a volume smaller than 8 voxels (Deq = 7.7 µm) were discarded. Finally, key attributes of individual pores (number of voxels, volume, maximum Feret diameter and the x–y–z position of the center of mass) were exported for further analysis.
To examine the influence of post-process heat treatments on the mechanical properties of the IN718 specimens with AB, hybrid and PM surface conditions, three thermal cycles were evaluated, as shown schematically in Figure 5. These precipitation (age)-hardening cycles—as detailed in Table 3—were selected based on two main types of heat treatments that are applied to wrought IN718 products, namely, AMS 5663 [53] and AMS 5664 [54], as well as a non-standard process specially adapted for LPBF-fabricated IN718 [55]. Four representative tensile specimens per surface condition (AB, hybrid and PM) were subjected to the three precipitation-hardening heat treatments (PHT), namely, PHT1, PHT2 and PHT3, followed by microstructural analysis and tensile testing. Detailed microstructural analysis was conducted for each of the PHTs examined in this study and can be found in a previous publication by the current authors [48].
The room-temperature tensile properties of IN718 with different surface conditions (AB, hybrid and post-machined) were determined using a 250 kN MTS load frame integrated with a laser extensometer and a non-contact optical 3D deformation measurement system (often referred to as digital image correlation (DIC)), Aramis® GOM-Trilion Quality Systems, (King of Prussia, PA, USA). The surface of the tensile specimen was first painted with a white background, and then a high-contrast random pattern of black speckles was applied. As the functionality of the Aramis® system is sensitive to the quality of this speckle pattern, verification of pattern recognition was performed before tensile testing to ensure proper strain recording along the entire gauge length, as described in [35,36]. Also, before tensile testing, two pieces of retro-reflective tape were attached to the tensile specimen to distinguish the gauge section for the laser extensometer during testing.
Tensile tests were conducted at a strain rate of 0.005 mm/mm/min up to yield and 0.05 mm/mm/min until fracture. To obtain the stress–strain curves and the related mechanical properties—yield strength (YS), ultimate tensile strength (UTS), and fracture strain—of the specimens, the load data, collected by the tensile testing machine, were used to calculate the engineering stress, while data collected by the laser extensometer were used to calculate the global strain or elongation (EL), based on the assumption that the entire gauge length is under uniform or homogeneous deformation. For each condition (AB, hybrid, PM), three tensile specimens were tested at room temperature to measure the average properties. The strain maps captured by the Aramis® system were used to assess the local deformation behavior of the specimens during tensile testing and to measure the local strain just before fracture. As well, the material toughness (TM)—also referred to as the modulus of toughness—was calculated from the area under the engineering stress–strain curve for each specimen condition to evaluate the strain energy density (strain on a unit volume of material) that the LPBF-processed IN718 can absorb before fracture.
Elevated-temperature tensile testing was conducted under the guiding principles of ASTM E21 [49], and four specimens were tested for each condition in this study. The tests were performed using a Tinius Olsen Super L testing system equipped with an infrared radiant-heating furnace. For each test, individual tensile specimens were heated at a rate of 18 °C/min to 650 °C and then held at this temperature for 30 min to reach thermal uniformity along the total length of the specimen. For this purpose, two K-type thermocouples were placed in direct contact with the specimen—close to the upper and lower ends of the reduced section—to monitor the thermal distribution along the gauge length. Test specimens equilibrated at 650 °C were then pre-loaded with a constant force of 222 N (50 lb) to prevent buckling, which is possible due to the thermal expansion of the IN718 during heating to the test temperature. The specimens were then tensile tested at a fixed strain rate of 0.005 mm/mm/min before yielding and subsequently at 0.05 mm/mm/min until fracture. The tensile stress–strain data were used to calculate the YS, UTS and EL properties. After fracture, the tensile specimens were air-cooled, and then the minimum cross-sectional area was measured and compared to the original cross-sectional area to calculate the reduction in area (RA). Also, fractographic analysis was conducted on select fractured tensile specimens using a JCM-7000 NeoScope™ Benchtop SEM.

3. Results and Discussion

3.1. Effect of Different Heat Treatments on the Microstructure

Detailed microstructural analysis was conducted for each of the PHTs examined in this study, details of which can be found in the authors’ previous publication [48]. Table 4 summarizes the size of the observed microstructural constituent for each condition to emphasize the similarities and differences between them.
As depicted in Table 4, all specimens had similar grain morphology and size regardless of the PHTs they were subjected to. It is also worth noting that these values are similar to the values observed in the as-fabricated IN718, showing that the grain morphology was stable during the applied heat treatments. Even though the grain structure remained unchanged during the heat treatments, there were significant microstructural differences observed between the different heat treatments.
The solidification conditions in LPBF resulted in the formation of a cellular sub-grain structure. The boundaries of the cells are known to have high dislocation density and provide additional strengthening. Precipitation of the strengthening phases was suppressed due to the high cooling rates experienced during LPBF processing. However, formation of the Laves phase and the presence of globular carbide particles with an average size of 59.2 ± 24.3 nm at the cell boundaries were observed.
The observed cellular structure was dissolved during all the PHTs except for the PHT3 specimen condition. The shorter solutionizing time applied in this heat-treatment cycle was not sufficient for dissolution; hence, the cellular structure from LPBF processing was maintained. Additionally, dissolution of the Laves phase was observed for all three PHTs, releasing large amounts of Nb into the surrounding area. As shown in Table 4, δ phase formation is observed only in PHT1, as the solutionizing temperature of PHT1 was below the solvus temperature of the δ phase. For the other two PHTs, growth of the carbides, present in the as-fabricated state, was observed. The differences in the size of these particles can be attributed to differences in the solutionizing times and temperatures applied during these two PHT conditions. In particular, PHT2 was subjected to a solutionizing treatment at a higher temperature for a holding time that was four times longer than that of PHT3, which resulted in coarser carbides in PHT2 specimens.
Finally, formation of strengthening γ′ and γ″ precipitates was reported for the PHTs. The size of these precipitates was also measured and compared, even though it was not possible to differentiate between the two types of precipitates. The size of the precipitates also showed variations between the different heat-treatment cycles. Among the three heat-treatment cycles, PHT3 had the finest precipitate size distribution, while PHT2 had the coarsest. This variation was associated with the higher aging temperatures of PHT2 promoting a higher diffusion rate and precipitate coarsening.

3.2. Inspection of Surface Finish and Density

The surface roughness and relative densities of vertically built IN718 with the different surface conditions (AB, hybrid, PM) are given in Table 5. The average linear (Ra) and areal (Sa) roughness measurements for the IN718 specimens with an AB surface condition were 5.1 μm ± 1.6 μm and 4.9 μm ± 0.9 μm, while the Rz and Sz were 34.4 μm ± 11.1 μm and 38.4 μm ± 11.4 μm, respectively. Specimens processed to have hybrid and PM surface conditions displayed significantly lower surface roughness values due to the integration of machining steps, either in-envelope (iteratively with LPBF processing) or out-of-envelope (sequentially after LPBF processing). For the hybrid surface condition, the Ra and Sa values were ~80% lower, while the Rz and Sz values were ~75% lower than specimens with AB surface conditions. The roughness parameters measured for the PM specimens showed moderately better surface finish conditions relative to the hybrid specimens. Here, it is worth mentioning that the machining conditions selected in this study for the hybrid specimens were based on the manufacturer’s recommended parameters, while the machining parameters for the PM specimens were according to the best (optimal) in-house practices for IN718. However, with optimization, the surface quality of the hybrid specimens can be further improved, as reported recently by Sommer et al. [56] in their study, which used the Taguchi method to reach Ra and Rz values of 0.47 μm and 3.73 μm, respectively, for IN718.
The surface topography of representative vertically built IN718 specimens in the AB, hybrid and PM conditions is compared in Figure 6 using a laser scanning confocal microscope to generate 3D images of the surface asperities. The AB surface consists of perturbations along the build direction (BD) that occur as a result of the layer-by-layer processing involved in LPBF AM. Additionally, unmelted powder particles attached to the AB surface are clearly visible in Figure 6a. Relative to the AB specimens, the hybrid and PM surfaces, as shown in Figure 6b,c, have similar surface topography, with considerably lower asperities and height fluctuations due to the machining finish applied to the additive surfaces.
Regardless of the surface conditions, all the specimens exhibited similar relative densities, with average values above 99.94% and standard deviations below ±0.02%, as listed in Table 5. Even so, a small effect of the surface condition could be discerned, with the AB specimens having the lowest average value for the relative density, while the hybrid specimens had the highest. A possible reason for the slightly higher average relative density of the hybrid specimens compared with the AB and PM specimens can be attributed to the nature of the hybrid process, which reduces the surface asperities regularly during the building process. Specifically, in hybrid processing, the intermittent machining steps, conducted in-envelope and after every ten layers of deposition, systematically achieve more even surfaces for the subsequent deposition layers; this enables higher consistency in the spreading of the powder layers by the recoater and results in more uniform conditions for densification (powder melting and solidification).

3.3. Inspection of Porosity

Figure 7 presents a visualization of the porosity in vertically built IN718 specimens in the AB, hybrid and PM conditions. In these figures, the color scale corresponds to the maximum Feret diameter, which is the length between the two most opposite points of a pore. This attribute is believed to be more representative of the severity of a pore. The first observation is that the amount of porosity in the AB specimen is several times higher than in the hybrid and PM specimens. as the calculated porosities based on the segmented images are 0.03%, 0.001% and 0.002%, respectively. More importantly, it is the spatial distribution of the pores that is drastically different between the AB condition and the two other conditions. The porosity in the AB condition is mainly located close to and under the surface of the specimen. This is better illustrated in Figure 8, where the radial distribution of the porosity is shown. In the AB specimen, around 98% of the total porosity is located within 0.2 mm of the surface as opposed to being almost uniformly distributed in the two other specimens. In addition, pores in the AB condition are significantly bigger than in the hybrid and PM conditions. For illustration, the maxF50, the value of the max Feret at 50% cumulative passing porosity, is 79 µm for the AB condition compared with 24 µm and 17 µm for, respectively, the hybrid and PM conditions. The maximum values measured for the max Feret are 306 µm, 47 µm and 57 µm, respectively, for the AB, hybrid and PM conditions. Such porosity beneath the surface is typically linked to the contour parameters, and it has also been attributed to causes such as the laser spot slowing before turning around, which increases the power density leading to keyhole formation, and a large offset between the hatch and contour tracks leading to a lack of fusion [57]. Given their size and spatial distribution, these near-surface pores in the specimens having an AB surface condition may influence the tensile deformation behavior.

3.4. Effect of Different Surface Conditions on Room-Temperature Tensile Properties

The average room-temperature tensile properties of vertically built IN718 with the different surface conditions are given in Table 6. The average YS, UTS and EL of the IN718 specimens with an AB surface condition were 575.4 ± 53.3 MPa, 966.1 ± 1.4 MPa and 27.5% ± 2.7%, respectively. Specimens processed to have hybrid and PM surface conditions displayed slightly higher tensile strength and ductility at room temperature. These tensile properties fall within the range of properties (minimum to typical) reported for standard wrought IN718 in the annealed condition, namely, a YS of 314–534 MPa, a UTS of 683–934 MPa and an EL of 26–62% [58,59]. Overall, the properties measured in the present study for vertically built IN718 agree well with previously reported observations of a YS of 572 ± 44 MPa, UTS of 904 ± 22 MPa and EL of 19% ± 4% [22,60,61] for LPBF-processed IN718 oriented in the BD. Nonetheless, when compared with the tensile properties reported previously by the current authors for horizontally built IN718 (YS of 790.8 ± 6.5 MPa, UTS of 1007.7 ± 5.8 MPa and EL of 34.0% ± 2.5%) with a hybrid/PM surface condition, the YS, UTS and EL in vertically built specimens were lower by ~25%, 5% and 20%, respectively. This influence of build orientation on the properties has been reported previously, and, typically, the tensile strength for IN718 along the BD (i.e., vertically built) is lower than that perpendicular to the BD (i.e., horizontally built) [62]. However, the trend for the EL is less clear, with some reporting an inverse relationship [62] while others have obtained relatively unchanged values that range from 20% to 35% with the incline angle [61,63,64]. In general, these previous findings support the tensile properties observed in the present study for vertically built IN718.
The room-temperature mechanical response of the vertically built IN718 with the different surface conditions is shown in Figure 9. Overall, the three surface conditions examined in the present study showed similar elastic and plastic behaviors, although there is a trend toward lower ductility in the vertically built IN718 having an AB surface condition. Thus, detailed examination was undertaken using the DIC technique to map the local strain distribution within the gauge section just before tensile fracture of representative vertically built IN718 specimens having AB, hybrid and PM surface conditions, as shown in Figure 10. In this figure, the tensile-loading direction is parallel to the vertical direction of the images, and the color-coded strain map on the entire gauge section of the specimens corresponds to the value of normal strain in the tensile direction. The region of strain localization, indicated as the area of highest intensity in the tensile specimen gauge section, corresponds to the location of the fracture that occurred immediately afterward. From these maps, it can be seen that the average strain in the majority of each specimen (AB, hybrid and PM) reaches around 25–30% and rises to about twice this value at the eventual fracture location for the hybrid and PM specimens. By contrast, the local strain just before fracture was about 15% lower for the AB specimen compared with the hybrid and PM ones. Also, in the specimens with an AB surface, the plastic zone size in the gauge section of the tensile specimens, over which the strain concentrates, is more confined compared with the hybrid and PM specimen conditions, which showed a more uniform development of plastic instabilities. These findings complement the higher ductility (by 10%) obtained for the hybrid and PM specimens relative to the AB ones. This is also evident from the TM values that are provided in Table 6, which shows the lowest value for the AB specimens (243 MPa) compared with the hybrid (259 MPa) and PM (262 MPa) ones. Specifically, these TM values were calculated from the area under the engineering stress–strain curves for each specimen condition and point to the higher strain energy (strain on a unit volume of material) that the IN718 specimens with machined surfaces (hybrid and PM) absorb before fracture relative to the AB ones.
The authors of the present study was able to find one recent study [65] on the influence of surface machining (hybrid) on the tensile properties of LPBF-processed IN718. Specifically, Sommer et al. [65] reported an increase in the UTS from 849 MPa for an AB surface condition (Ra = 13.37 μm) to 947 MPa for IN718 with a hybrid surface (Ra = 0.73 μm); in addition, the EL was roughly 5% higher for the latter surface condition. However, as no details were provided on the specimen geometry, tensile test conditions and the number of test specimens, it is difficult to assess the reproducibility of their data and to gauge the measurement errors without any reported standard deviation values.
Notwithstanding this, Zhao et al. [66] also recently studied the role of machining on the tensile properties of IN718 produced by electron-beam melting (EBM). Their study showed that even with a contouring strategy applied to the surface during EBM, their AB IN718 specimens exhibited a rough surface (Rz = 270 ± 80 μm), with partially sintered powder particles fused on the surface contributing to a high level of surface porosity (~15%). Such surface defects can cause local stress concentrations and act as stress raisers that lead to early crack initiation during the tensile loading of full-density parts fabricated by EBM [2]. Zhao et al. observed that shallow machining to reduce the fraction of surface voids (to 2.1%) and the surface roughness (to Rz = 5.7 ± 0.6 μm) was not as effective in improving the tensile properties of EBM IN718—most especially the ductility—relative to deep machining of the surface, which also removed the fraction of near-surface voids (to 0.4%) but without significant improvement in the surface roughness (Rz = 5 ± 0.8 μm). In the present study, the inherently smoother surface finish possible during LPBF processing with finer IN718 powder feedstock (i.e., D50 of 26 μm) compared with EBM (i.e., D50 of 90 μm) [66]—combined with the optimized LPBF processing conditions that were effective in limiting surface and near-surface defects (<0.03% as seen through µXCT)—brought about the high tensile performance of all the vertically built IN718 specimens, with even slightly better performances for the hybrid and PM ones relative to the AB. In this regard, the application of in-envelope hybrid surface machining appears to be an effective strategy for improving the AB surface finish and for rendering LPBF-processed IN718 to have comparable room-temperature tensile properties to out-of-envelope PM parts. As well, there may be additional potential for improving the tensile mechanical properties by optimizing the machining parameters in the hybrid process, which can reduce the surface roughness further.

3.5. Effect of Surface Conditions and Heat-Treatment Cycles on High-Temperature Tensile Properties

To minimize mechanically and thermally induced wear of tooling, IN718 is ordinarily shaped (machined, formed, etc.) without any heat treatment [67], but for use, this alloy is typically specified in the precipitation-hardened state. As such, the elevated-temperature tensile properties were studied only after applying PHTs, for which three were selected in the present study based on our previous investigations [48]. Table 7 lists the average tensile properties at 650 °C for vertically built IN718 with the different surface conditions and PHTs examined in the present study. For comparison, the tensile properties at 650 °C of precipitation-hardened wrought [59] and LPBF-processed IN718, as reported by Trosch et al. [17], Zhao et al. [38], and Gao et al. [68], respectively, are also provided in Table 7.
Overall, both the strength (YS and UTS) and ductility (EL and RA) of the vertically built IN718 in the present study surpassed the minimum specification for wrought IN718 as stated in the AMS 5662 [69] and 5663 [53] standards. This is a significant outcome, as previous studies that examined the high-temperature tensile properties of IN718 fabricated by LPBF processing reported lower strength and/or ductility compared to the wrought-equivalent alloy. For instance, Trosch et al. [17] studied LPBF processing of IN718 and reported on the tensile properties at 650 °C for round specimens built in three directions (vertical, horizontal and at 45°) that were precipitation-hardened under conditions similar to PHT1 in the present study. Their vertically built IN718 exhibited YS (860 MPa) and UTS (992 MPa) values below the minimum specifications, and although the UTS was slightly higher (and just above the specified value) in both the horizontal and 45° inclinations, the EL values (3.6% and 5.8%, respectively) were quite low [17]. When averaging the tensile properties of their IN718, the YS (862 MPa) and UTS (1026 MPa) met the minimum specifications, but the average EL (7.9%) remained considerably below the wrought specification, as compared in Table 7. A similar study by Zhao et al. [38] on horizontally built flat specimens also showed considerable challenges for reaching the specified minimum EL at 650 °C; for the three precipitation-hardening treatments examined in their study, the highest EL was ~5.5% under conditions similar to PHT1 in the present study. By contrast, Gao et al. [68] reported lower YS (773 MPa) and UTS (992 MPa) properties after PHT1, but their EL (18%) was well above the minimum value of 12% specified for wrought IN718 at 650 °C. Their work reported that precipitation hardening, with the addition of a homogenization step according to AMS 5383, achieved the highest strength (YS = 965 MPa and UTS = 1126 MPa) and ductility (EL = 21%) at 650 °C for their horizontally built IN718 specimens. Strößner et al. [40] examined both horizontally and vertically built round tensile specimens at 650 °C; they reported tensile strengths (UTS) in the range of 1100 MPa but with a much wider EL range of 8–18% that depended on the specimen orientation and heat treatment (i.e., PHT1 or PHT2). Recently, Fayed et al. [41] investigated the influence of both homogenization and solution treatment conditions on the tensile properties at 650 °C using vertically built IN718 flat specimens. Overall, all five of the applied heat treatments led to YS and UTS values above the wrought minimum, but this was at the expense of a significant ductility loss, with reported average EL values of 7.3 to 11.8%. Further work by Fayed et al. [70] showed that despite considerable efforts to optimize the post-process heat treatment, the ductility loss at 650 °C in their vertically built IN718 specimens could not be recovered. It is noteworthy that the surfaces of the tensile specimens in these reported studies [17,40,66,70] were prepared by machining after LPBF processing, and their observations on the high-temperature embrittlement of IN718 was related to different microstructural aspects, such the coarsening of carbides, excessive formation of δ phase, overaging of the strengthening precipitates, etc. By contrast, the present study aimed to examine the role of the surface condition on the high-temperature properties, which presents an important gap for maturing the application of LPBF technology for IN718 in high-temperature service environments, especially as not all features on near-net-shape parts can be post-machined.
Figure 11 graphically presents the tensile properties at 650 °C for the vertically built IN718 in the different precipitation-hardened conditions. In general, for all the PHTs examined in the present study, the tensile specimens with an AB surface exhibited lower YS, UTS, EL and RA values compared with those with the hybrid and PM surface conditions, as illustrated in Figure 11a–c for PHT1, PHT2 and PHT3, respectively. Considering the standard deviation values detailed in Table 7, the differences are relatively minor (up to 4%) for the strength properties (YS and UTS), but they are definitely evident for the ductility, with the EL increasing by 33–68% and the RA nearly doubling for the machined (hybrid and PM) surfaces. By contrast, the role of the PHTs on the tensile properties at 650 °C is less obvious, as illustrated in Figure 11d–f. For instance, the PHT2 specimens having an AB surface show the best combination of tensile strength and ductility properties at 650 °C (Figure 11d), while PHT1 specimens with a PM surface are marginally better than the PHT2 (Figure 11f) ones. Considering that both the AB and PM specimens were fabricated using equivalent LPBF process parameters, the observed difference in the tensile mechanical response to the applied PHTs points to an effect from the variation of specimen surface conditions (rather than microstructural), which can occur even within each specimen type (i.e., AB, hybrid and PM). This is further supported by the hybrid specimens (Figure 11e), which exhibited practically equivalent tensile strengths (YS and UTS) as well as ductility (EL and RA) at 650 °C for all the PHTs examined in the present study.
The high temperature (650 °C) engineering and true stress-strain curves are given in Figure 12 for representative IN718 specimens with AB, hybrid and PM surfaces and in the PHT1, PHT2 and PHT3 conditions. In contrast to the room-temperature mechanical response (as shown in Figure 9), the tensile behavior at 650 °C shows some differences, predominately in the lower stresses and strains of the vertically built IN718 specimens with AB surfaces relative to the hybrid and PM ones. Also, this trend in the tensile behavior was consistent for all the PHTs applied in the present study, which reinforces the important role of the surface condition on the performance of IN718 at or close to the maximum temperatures experienced by this alloy in service. Differences in the mechanical response between the specimens with hybrid or PM surfaces are less obvious, with the most notable dissimilarities occurring after the onset of necking.
Hence, to examine the role of the surface condition on the strain localization behavior at elevated temperatures, future research should examine the concomitant application of DIC techniques during the tensile mechanical property assessment of IN718, which is an innovative but challenging methodology to develop and implement at the relevant service temperatures for this alloy. An alternate traditional approach to examine the plasticity and fracture path is to section the fractured tensile specimens parallel to the loading direction and to metallographically prepare them for microscopic examination. In the present study, it was conceived that the role of the surface condition on the fracture behavior could be more thoroughly investigated using µXCT inspection, which allows for direct examination, i.e., without any sectioning or preparation of the specimen. Thus, specimens of all three surface conditions were inspected by µXCT after being tensile-tested at 650 °C to visualize the influence of the near-surface porosity on the rupture behavior. Cross-sections of these specimens are shown in Figure 13. Cracks emerging from the surface are visible in the specimens with all three surface conditions, but their number and size are much more significant in the AB specimen. The near-surface pores present in the AB condition clearly act as stress raisers that visibly initiate rupture at numerous sites and predispose the specimen to premature/early fracture. Also, the formation of such open cracks at an earlier than normal stage of deformation explains the observed lower EL and RA values at rupture for the specimens with AB surfaces.

3.6. Fractography

Representative images of the fracture surface after room-temperature tensile testing are shown in Figure 14a–c for the vertically built IN718 specimens with AB, hybrid and PM surface conditions, respectively. The fracture surface morphologies in these low-magnification images provide a signature of the failure mechanism for the IN718 specimens with these different surface conditions. Specifically, the specimens with hybrid and PM surface conditions showed the characteristic appearance of cup-and-cone-shaped failure. As displayed in the low-magnification images given in Figure 14b,c, both fracture surfaces exhibited a relatively smooth shear lip along the edge of the fracture surface and a relatively rough central rupture zone in the cup, where final fracture occurred from tensile overload. By contrast, specimens having an AB surface condition sometimes showed a cup-and-cone-shaped failure, but more often, fracture occurred through tensile crack propagation from an edge location (Figure 14a), probably due to stress concentration at any of the open voids on the specimen surface, near-surface pores and/or the valleys of the rough surface that were identified in the µXCT (Figure 7, Figure 8 and Figure 13) and surface roughness (Figure 6 and Table 5) results. This premise is supported by the research of Choo et al. [71] on tensile crack propagation in 316L stainless steel specimens with surface and internal voids that related a much higher local stress concentration factor to flat voids oriented perpendicular to the tensile-loading direction compared with elongated voids oriented parallel to the loading direction. Thus, the rough AB surface on the vertically built IN718 specimens that is predisposed to early stress concentration at surface irregularities—as also explained by the strain localization behavior through DIC (Figure 10a)—brings about variation in the tensile fracture characteristics, depending on the magnitude of and local variations in the surface roughness. On the whole, this manifests in a lower average ductility and higher standard deviation for the specimens with AB surfaces relative to the hybrid and PM ones and explains the observed differences in the fracture morphologies.
At high magnification, Figure 15 shows that, regardless of the surface conditions, the fracture surface consists of a large number of homogeneously distributed dimple colonies along with a small number of flat cleavage marks and a few remnant small gas pores. Also visible on these fracture surfaces is the cellular sub-grain structure, which points to void formation starting at the cell boundaries, which are inherently present in the as-fabricated LPBF microstructure of IN718 (i.e., without any PHT) as reported in [72]. These characteristic ductile features on all the fracture surfaces correspond well with the high elongation at fracture observed for all the specimens.
Representative macroscopic images of the fractured surfaces after tensile testing at 650 °C are shown in Figure 16 and Figure 17 for the different IN718 specimens conditions examined in this study. From the low-magnification images, it can be observed that the fracture surface morphologies are closely related to the ductility of the IN718 specimens. For instance, vertically built IN718 specimens with an AB surface under PHT1, PHT2 and PHT3 conditions exhibit a lower reduction in the area of the gauge section when compared with the corresponding hybrid and PM ones.
Also, the lower ductility of the IN718 specimens with an AB surface is evidenced by the relatively smooth fracture surface for the PHT1, PHT2 and PHT3 conditions (Figure 16a–c), which is in contrast to the higher amount of ductile deformation around the crack tip of the cup-and-cone fractured surfaces observed in the corresponding hybrid (Figure 16d–f) and PM (Figure 16g–i) specimens. In addition, a portion (~120 μm) of the edge periphery on the fractured surface of the specimens with an AB surface showed brittle features (arising from numerous near-surface pores) for all three PHTs, which were not seen on any of the hybrid and PM specimens.
Even so, at high magnification (Figure 17), the fracture surfaces for all three surface conditions and all three PHTs displayed the presence of dimples formed through void nucleation, growth and coalescence, indicating a ductile tensile fracture mode at 650 °C. Usually tensile fracture at high temperature (>0.3TM, where TM is the melting temperature) occurs through three main modes—transgranular creep fracture, intergranular creep fracture and rupture with dynamic recrystallization—and each of these is based on the growth and coalescence of voids. Typically, the rupture failure mode is primarily observed at large reductions in area, while a transition from transgranular to intergranular creep fracture is seen when the level of stress is reduced and maintained for a sufficiently long period [73]. Considering the test conditions in the present study, ductile tensile failure of the IN718 at 650 °C progressed by transgranular creep fracture, which is supported by the observation of similar ductility levels (EL ~20–25%) and fracture characteristics for the hybrid and PM specimens tested at 650 °C relative to those reported previously by the authors [48] for the room-temperature tensile testing of machined and precipitation-hardened specimens. In the case of the AB surface condition, the near-surface pores were noticed to have an accelerating effect on fracture progression that resulted in early failure (lower EL and RA) during tensile testing, which was less noticeable at room temperature but was clearly evident at 650 °C.
Overall, these observations on the ductility of the specimens from the fractured surfaces corroborate well with the specific values detailed in Table 7 for the average EL and RA for different surface and PHT conditions. In any case, all of the vertically built specimens in the present study had relatively good tensile ductility at 650 °C that exceeded the specified minimum value (of 12%) for wrought IN718. This is almost certainly attributable to the absence of cracks and lack-of-fusion defects in the specimens, as well as the visibly low defect content, consisting of a few small dispersed spherical gas pores on the fractured surfaces, as evidenced in the high-magnification images of the fracture surfaces for the different IN718 specimen conditions tested in the present study (Figure 17) as well as the µXCT results (in Figure 6 and Figure 7). Nonetheless, the lower ductility properties identified in the present study for the AB surface of IN718 produced by LPBF processing strongly advocates for advancing research into reducing surface/sub-surface defects of additively built materials, as well as to investigate the properties with and without machining to understand the resulting gains/debits in performance under service-relevant conditions.

4. Conclusions

This research study investigated the application of a hybrid additive–subtractive method to fabricate vertically built IN718 tensile specimens with three different surface conditions—as-built (AB), hybrid (in-envelope machining) and post-machining (PM) or out-of-envelope. After fabrication, select specimens with each surface condition were tensile tested at room temperature. The remaining specimens were subjected to three different post-process heat treatments and the tensile behavior and properties were studied at a high temperature of 650 °C. The following conclusions can be drawn from the current study:
  • The vertically built IN718 in the PM and AB conditions exhibited the lowest and highest surface roughness values, respectively. For the hybrid surface condition, the Ra and Sa values were ~80% lower, while the Rz and Sz values were ~75% lower than specimens with AB surface conditions.
  • Regardless of the surface conditions, all the specimens exhibited similar relative densities, with average values above 99.94% and standard deviations below ±0.02%. A small effect of the surface condition on the density was observed, with the AB specimens having the lowest average value for the relative density while the hybrid specimens had the highest.
  • As revealed by µXCT inspection, regardless of the surface condition, all the specimens contained a very small amount of uniformly distributed porosity. In addition to this general porosity, compared with the specimens with hybrid and PM surface finishes, the specimens in the AB condition contained a significantly higher (10 times) amount of porosity that was concentrated in a region 100 µm to 200 µm below the surface.
  • The room-temperature tensile properties of all the vertically built IN718 specimens were within the range of properties reported for standard wrought IN718 in the annealed condition. The three surface conditions examined in the present study showed similar elastic and plastic behaviors, but the IN718 specimens with AB surfaces exhibited lower ductility relative to the hybrid and PM ones, which was explained through differences in strain localization behavior through digital image correlation, the near-surface porosity and differences in the fracture morphology.
  • The strength and ductility of the vertically built IN718 at a high temperature of 650 °C surpassed the minimum specifications for wrought IN718. Even so, the AB surface exhibited lower yield strength (YS), ultimate tensile strength (UTS), elongation (EL) and a reduction in area (RA) values compared with the hybrid and PM surface conditions for all the post-process heat treatments examined in the present study. Compared with the AB surface, machined (hybrid and PM) surfaces exhibited slightly higher strength (YS and UTS) properties by about 4%. But the difference was more significant for the ductility, with the EL increasing 33–68% and the RA nearly doubling for the machined (hybrid and PM) surfaces relative to AB ones.
  • After tensile testing at 650 °C, µXCT images showed that the specimens with an AB surface condition contained a lot more open cracks at their surface and that these cracks initiated at the near-surface pores. This strongly points to the significant role of the near-surface pores in having an accelerating effect on the transgranular creep fracture process and a negative impact on the ductility properties.
  • The application of in-envelope hybrid surface machining appears to be an effective strategy for improving the AB surface finish and for rendering LPBF-processed IN718 to have a comparable tensile mechanical performance (at room and high temperatures) to out-of-envelope PM parts.

Author Contributions

Conceptualization, S.S., P.W., S.E.A., J.G., J.S., R.A. and P.P.; methodology, S.S., P.W., S.E.A., R.P., J.G. and J.S.; software, R.P. and J.S.; validation, R.P. and R.A.; formal analysis, S.S., S.E.A., P.W. and R.P.; investigation, S.S., P.W., R.P. and S.E.A.; resources, J.G., J.S. and P.P.; data curation, S.S., P.W., S.E.A. and R.P.; writing—original draft, S.S.; writing—review and editing, P.W., S.S., S.E.A., J.G., J.S., R.A., P.P. and R.P.; visualization, S.S.; supervision, P.W. and J.G.; project administration, S.S., P.W. and J.G.; funding acquisition, P.W., J.G., R.A. and P.P. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the Defence Technology Sustainment Program of the National Research Council Canada and the Department of National Defence (DND) under project A1-018703.

Data Availability Statement

The authors confirm that the data supporting the findings of this study are available within the article.

Acknowledgments

The authors wish to thank X. Pelletier (NRC) for supporting the EDM removal of tensile specimens from the build plate, D. O’Keefe for the post-machining of the select tensile specimen surfaces, C. MacDowell (McGill) for imaging the fractured surfaces, and M. Guerin (NRC) for supporting the tensile testing.

Conflicts of Interest

Josh Soost was employed by the company Matsuura Machinery USA Inc. The remaining authors declare no conflicts of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.

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Figure 1. (a,b) Morphology and (c) cohesive index of the starting IN718 powder.
Figure 1. (a,b) Morphology and (c) cohesive index of the starting IN718 powder.
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Figure 2. Process flow detailing the different stages in the experimental methodology.
Figure 2. Process flow detailing the different stages in the experimental methodology.
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Figure 3. (a) CAD layout of the build plate with 24 vertically built tensile specimens. (b) The 24 vertically built tensile specimens after the build. (c) A sleeve-shaped support structure designed with a small gap to ease removal of the tensile specimens. (d) Easy support removal after EDM from the build plate. (e) Tensile specimen geometry based on ASTM E8M-22 [49].
Figure 3. (a) CAD layout of the build plate with 24 vertically built tensile specimens. (b) The 24 vertically built tensile specimens after the build. (c) A sleeve-shaped support structure designed with a small gap to ease removal of the tensile specimens. (d) Easy support removal after EDM from the build plate. (e) Tensile specimen geometry based on ASTM E8M-22 [49].
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Figure 4. Vertically built tensile specimens fabricated to have three surface finish conditions in the gauge section: AB (left), hybrid (middle) and PM (right).
Figure 4. Vertically built tensile specimens fabricated to have three surface finish conditions in the gauge section: AB (left), hybrid (middle) and PM (right).
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Figure 5. Different precipitation-hardening heat treatment (PHT) cycles used in this study.
Figure 5. Different precipitation-hardening heat treatment (PHT) cycles used in this study.
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Figure 6. Map of the surface topography of vertically built IN718 specimens with (a) AB, (b) hybrid and (c) PM surfaces.
Figure 6. Map of the surface topography of vertically built IN718 specimens with (a) AB, (b) hybrid and (c) PM surfaces.
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Figure 7. Porosity inspections of vertically built IN718 specimens with (a,b) AB, (c,d) hybrid and (e,f) PM conditions.
Figure 7. Porosity inspections of vertically built IN718 specimens with (a,b) AB, (c,d) hybrid and (e,f) PM conditions.
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Figure 8. Differential distribution of the pore volume fraction and number fraction as a function of the distance R from the specimen outer surface.
Figure 8. Differential distribution of the pore volume fraction and number fraction as a function of the distance R from the specimen outer surface.
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Figure 9. Representative (a) engineering stress–strain and (b) true stress–strain curves of vertically built IN718 with the different surface conditions.
Figure 9. Representative (a) engineering stress–strain and (b) true stress–strain curves of vertically built IN718 with the different surface conditions.
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Figure 10. DIC analysis of the local strain distribution maps of the gauge section of the vertically built IN718 tensile specimens just before fracture: (a) AB, (b) hybrid and (c) PM surface conditions.
Figure 10. DIC analysis of the local strain distribution maps of the gauge section of the vertically built IN718 tensile specimens just before fracture: (a) AB, (b) hybrid and (c) PM surface conditions.
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Figure 11. Tensile properties at 650 °C for the vertically built IN718 with the different precipitation-hardening conditions: (a) PHT1, (b) PHT2 and (c) PHT3; and different surface conditions: (d) AB, (e) hybrid and (f) PM.
Figure 11. Tensile properties at 650 °C for the vertically built IN718 with the different precipitation-hardening conditions: (a) PHT1, (b) PHT2 and (c) PHT3; and different surface conditions: (d) AB, (e) hybrid and (f) PM.
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Figure 12. Representative (a–c) engineering stress–strain and (d–f) true stress–strain curves at 650 °C for the vertically built IN718 with the different surface conditions and PHTs.
Figure 12. Representative (a–c) engineering stress–strain and (d–f) true stress–strain curves at 650 °C for the vertically built IN718 with the different surface conditions and PHTs.
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Figure 13. µXCT cross-section of vertically built IN718 specimens tested at 650 °C with (a) AB, (b) hybrid and (c) PM surface conditions.
Figure 13. µXCT cross-section of vertically built IN718 specimens tested at 650 °C with (a) AB, (b) hybrid and (c) PM surface conditions.
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Figure 14. Fractographs after room-temperature tensile testing of the vertically built IN718 specimens with (a) AB (b) hybrid and (c) PM surface finish conditions.
Figure 14. Fractographs after room-temperature tensile testing of the vertically built IN718 specimens with (a) AB (b) hybrid and (c) PM surface finish conditions.
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Figure 15. High-magnification fractographs after room-temperature tensile testing of the vertically built IN718 specimens with (a) AB (b) hybrid and (c) PM surface finish conditions.
Figure 15. High-magnification fractographs after room-temperature tensile testing of the vertically built IN718 specimens with (a) AB (b) hybrid and (c) PM surface finish conditions.
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Figure 16. Fractographs after high-temperature (650 °C) tensile testing of vertically built IN718 specimens with an AB surface finish and under (a) PHT1, (b) PHT2 and (c) PHT3 conditions; with a hybrid surface finish and under (d) PHT1, (e) PHT2 and (f) PHT3 conditions; as well as a with a PM surface finish and under (g) PHT1, (h) PHT2 and (i) PHT3 conditions.
Figure 16. Fractographs after high-temperature (650 °C) tensile testing of vertically built IN718 specimens with an AB surface finish and under (a) PHT1, (b) PHT2 and (c) PHT3 conditions; with a hybrid surface finish and under (d) PHT1, (e) PHT2 and (f) PHT3 conditions; as well as a with a PM surface finish and under (g) PHT1, (h) PHT2 and (i) PHT3 conditions.
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Figure 17. High magnification of fractographs after high-temperature (650 °C) tensile testing of vertically built IN718 specimens with an AB surface finish and under (a) PHT1, (b) PHT2 and (c) PHT3 conditions; with a hybrid surface finish and under (d) PHT1, (e) PHT2 and (f) PHT3 conditions; as well as a with a PM surface finish and under = (g) PHT1, (h) PHT2, (i) PHT3 conditions.
Figure 17. High magnification of fractographs after high-temperature (650 °C) tensile testing of vertically built IN718 specimens with an AB surface finish and under (a) PHT1, (b) PHT2 and (c) PHT3 conditions; with a hybrid surface finish and under (d) PHT1, (e) PHT2 and (f) PHT3 conditions; as well as a with a PM surface finish and under = (g) PHT1, (h) PHT2, (i) PHT3 conditions.
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Table 1. Chemical composition of the starting IN718 powder feedstock.
Table 1. Chemical composition of the starting IN718 powder feedstock.
ElementCrMoNbAlCMnNiFe
Wt.%19.193.075.250.640.040.1452.37Balance
ElementSPCoTiSi-ON
Wt.%0.0010.0030.040.930.15ppm14050
Table 2. Characteristics of the starting IN718 powder feedstock.
Table 2. Characteristics of the starting IN718 powder feedstock.
PSD (μm)Flowability per 50 g (s)AD (g/cm3)
D10D50D90HallCarneyHallCarney
2637552543.923.93
Table 3. Details of the specimen conditions and the applied precipitation-hardening heat treatment (PHT) cycles.
Table 3. Details of the specimen conditions and the applied precipitation-hardening heat treatment (PHT) cycles.
DesignationABPHT1PHT2PHT3
Standard-AMS 5663 [53]AMS 5664 [54]Non-standard [55]
Solution Treatment (ST)
Temperature, °C-98010651020
Time, h-110.25
Cooling-ACACWQ
Aging (A) Step 1
Temperature, °C-720760720
Time, h-81024
Cooling-Controlled FC to 620 °C at 55 °C/hFC to 650 °CAC
Aging (A) Step 2
Temperature, °C-620650-
Time, h-88-
Cooling-ACAC-
AC = Air-cooled; FC = Furnace-cooled; WQ = Water-quenched.
Table 4. Microstructural characteristics of vertically built IN718 with the applied precipitation-hardening treatments (PHTs).
Table 4. Microstructural characteristics of vertically built IN718 with the applied precipitation-hardening treatments (PHTs).
PHT1PHT2PHT3
Average grain diameter (µm)33.2 ± 32.532.2 ± 30.931.7 ± 26.8
Grain aspect ratio2.7 ± 1.53.5 ±2.23.3 ± 1.9
Carbides (nm)-307.8 ± 172.584.4 ± 20.3
δ phase—major axis (nm)900.6 ± 314.0--
δ phase—minor axis (nm)195.4 ± 93.5--
Strengthening precipitates (nm)19.0 ± 4.033.0 ± 9.013.0 ± 3.0
Table 5. Measured relative density (Archimedes) as well as linear and areal roughness values from the vertically built IN718 specimens.
Table 5. Measured relative density (Archimedes) as well as linear and areal roughness values from the vertically built IN718 specimens.
Specimen
Surface		Linear Roughness (μm)Areal Roughness (μm)Relative Density (%)
RaRzSaSz
AB5.1 ± 1.634.4 ± 11.14.9 ± 0.938.4 ± 11.499.94 ± 0.01
Hybrid1.0 ± 0.26.4 ± 2.01.2 ± 0.29.5 ± 3.899.98 ± 0.02
PM0.8 ± 0.54.8 ± 2.60.9 ± 0.59.2 ± 4.099.96 ± 0.02
Table 6. Average room-temperature tensile mechanical properties of vertically built IN718.
Table 6. Average room-temperature tensile mechanical properties of vertically built IN718.
DesignationSurface ConditionYS
(MPa)		STD
(MPa)		UTS
(MPa)		STD
(MPa)		EL
(%)		STD
(%)		TM
(MPa)
ABAs-built LPBF only575.453.3966.11.427.52.7242.6
HybridIn-envelope additive–subtractive582.15.7957.31.829.61.5258.5
PMOut-of-envelope machined595.521.7955.16.230.00.7262.0
Wrought [31,58,59]
annealed		Out-of-envelope machined314–534-683–934-26–62--
LPBF [22,60,61,64]Out-of-envelope machined5724490422194-
Table 7. Average tensile mechanical properties at 650 °C for vertically built IN718 in precipitation-hardened conditions.
Table 7. Average tensile mechanical properties at 650 °C for vertically built IN718 in precipitation-hardened conditions.
Surface ConditionHeat TreatmentYS
(MPa)		STD
(MPa)		UTS
(MPa)		STD
(MPa)		EL
(%)		STD
(%)		RA
(%)		STD
(%)
AB
As-built LPBF only		PHT1947.36.51076.55.814.70.917.20.6
PHT2968.77.01092.32.718.21.321.72.8
PHT3914.418.31026.523.115.72.925.92.4
Hybrid
In-envelope additive–subtractive		PHT1980.416.01099.95.224.71.237.52.7
PHT2987.07.51110.43.824.21.337.52.7
PHT3977.77.61105.65.623.31.843.18.9
PM
Out-of-envelope machined		PHT1953.231.41104.56.724.61.142.64.9
PHT2925.010.31111.42.025.91.246.57.7
PHT3935.415.81109.24.921.51.241.07.7
Wrought [53,69]
Out-of-envelope PM		AMS 5662 [69]
5663 [53]		862-1000-12-15-
LPBF [17]
Out-of-envelope PM round specimens		PHT1862-1026-7.9---
LPBF [38]
Flat specimens		PHT1915-1025 5.5---
LPBF [68]
Out-of-envelope PM round specimens		PHT1773 992 18
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Sarafan, S.; Wanjara, P.; Pelletier, R.; Atabay, S.E.; Gholipour, J.; Soost, J.; Amos, R.; Patnaik, P. Elevated-Temperature Tensile Behavior and Properties of Inconel 718 Fabricated by In-Envelope Additive–Subtractive Hybrid Manufacturing and Post-Process Precipitation Hardening. J. Manuf. Mater. Process. 2024, 8, 297. https://doi.org/10.3390/jmmp8060297

AMA Style

Sarafan S, Wanjara P, Pelletier R, Atabay SE, Gholipour J, Soost J, Amos R, Patnaik P. Elevated-Temperature Tensile Behavior and Properties of Inconel 718 Fabricated by In-Envelope Additive–Subtractive Hybrid Manufacturing and Post-Process Precipitation Hardening. Journal of Manufacturing and Materials Processing. 2024; 8(6):297. https://doi.org/10.3390/jmmp8060297

Chicago/Turabian Style

Sarafan, Sheida, Priti Wanjara, Roger Pelletier, Sila Ece Atabay, Javad Gholipour, Josh Soost, Robert Amos, and Prakash Patnaik. 2024. "Elevated-Temperature Tensile Behavior and Properties of Inconel 718 Fabricated by In-Envelope Additive–Subtractive Hybrid Manufacturing and Post-Process Precipitation Hardening" Journal of Manufacturing and Materials Processing 8, no. 6: 297. https://doi.org/10.3390/jmmp8060297

APA Style

Sarafan, S., Wanjara, P., Pelletier, R., Atabay, S. E., Gholipour, J., Soost, J., Amos, R., & Patnaik, P. (2024). Elevated-Temperature Tensile Behavior and Properties of Inconel 718 Fabricated by In-Envelope Additive–Subtractive Hybrid Manufacturing and Post-Process Precipitation Hardening. Journal of Manufacturing and Materials Processing, 8(6), 297. https://doi.org/10.3390/jmmp8060297

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