Improving Structural Integrity of Direct Laser-Deposited Ni–Co–Cr–Al–Y Superalloys by Alloy Modification

Ni-based superalloys find applications in aerospace and power generation systems where high temperatures (600–1100 °C; 1112–2012 °F) are encountered [1–7]. The mechanical properties of these alloys rapidly drop as the temperature increases beyond 900 °C (1652 °F) [8]. Additions of Co to these Ni superalloys can help improve the performance beyond 1000 °C (1382 °F) by increasing the solidus temperature [9–12]. A386 is an example of a Ni–Co superalloy that is commercially available and applicable for high-temperature applications up to 1050 °C (1922 °F) [13]. Many such alloy chemistries have been developed for protective thermal barrier coatings (TBC), used in gas turbine components such as blades, stators, and rotors [14–16]. They also serve as protective claddings for components that are exposed to combined thermal and mechanical loads with simultaneous reactive environments, such as those encountered in oil and gas extraction equipment and nuclear irradiation environments [17,18]. The M–Cr–Al–Y family of alloys (M = Ni and/or Co) used in cladding configurations are typically deposited using thermal and plasma spraying methods, which can yield high deposition rates and high structural integrity with minimal defects [19–23]. However, these deposition techniques have not yet demonstrated the ability to fabricate bulk, three-dimensional (3D) structures from these alloys. Therefore, it is necessary to find suitable processing routes for bulk component (thicknesses > 3 mm (0.118 in.)) fabrication of alloy chemistries that have been traditionally limited to thin coatings (thicknesses < 3 mm (0.118 in.)). Such advancements would enable fabrication of components that could withstand higher service temperatures and operational loads for longer periods of time. Gas turbine components fabricated completely from M–Cr–Al–Y alloys would exhibit greater high-temperature corrosion resistance than any Ni-based laser-fabricable alloys and mechanical properties similar to or greater than common aerospace component base materials such as CMSX, Inconel, and René alloys [24,25].

Additive manufacturing (AM) techniques have been developed to address the shortcomings of conventional manufacturing methods. Conventional manufacturing processes are capable of fabricating semi-complex structures, but specialized machinery and a series of different operations such as hot/cold working, subtractive machining, heat treatment, and assembly are required [26–29]. Each additional operation increases the lead time and manufacturing cost for a component [30,31]. Machines and tooling for conventional manufacturing processes are expensive and require a large initial investment. If a lesser number of components is desired, it may not be cost-effective to use conventional manufacturing techniques, and AM would be considered. Additive manufacturing equipment/machines also require a large initial investment, but these machines reduce the overall number of machines and operations needed to produce a complete component.

Unlike conventional techniques, AM can produce intricate and complex parts such as medical implants [32–34], aerospace components [29,35,36], and functional coatings [37–39], with a very short lead time. Conventional manufacturing methods are not able to process specialty metallic materials such as Ni, Co, or Cr alloys. Laser-based AM processes have been utilized extensively to research and develop new useful metallic alloy compositions and 3D structures. Previously unfabricable materials exhibiting improved mechanical properties and corrosion resistance allow designers to increase performance and thermal stability of mechanical systems [40–47]. The laser directed energy deposition (laser-DED) AM technique has been used as a tool to research and develop novel alloy compositions [48,49]. Laser-DED is a powder or wire-fed, freeform laser processing technique that fabricates features and components from metallic and metal–ceramic mixtures using focused laser energy [45,50–55]. Laser-DED is a versatile manufacturing technique that offers the advantages of high deposition rates [54], large build envelopes [56], and can be used to rapidly fabricate near net shaped 3D components as well as recycling worn or damaged components via maintenance, refurbishment, and overhaul (MRO) operations [26–31].

The laser-DED process, and its commercial variants such as the LENS^TM technique, have been used for processing many commercial materials based on Fe, Ni, and Ti metals, as well as specialty materials based on ceramics, metal–ceramic mixtures, and refractory alloys [57–59]. Metal alloys with high Al content, such as the aforementioned M–Cr–Al–Y alloy family, present challenges for laser-based manufacturing processes due to their poor absorptivity of laser radiation at wavelengths of 900–1100 nm (3.54e⁻⁵–4.33e⁻⁵ in.) (common wavelengths for AM processing), and affinity for forming brittle intermetallic compounds which increase an alloys susceptibility to hot cracking during laser-based manufacturing processes [41,60–62]. Therefore, there is a need to explore the effects of alloy chemistry on the material’s absorption of laser power and its effects on the relative amounts of each phase.

The primary objective of this study is to investigate the effects of tailoring the chemical composition of A386, a commercially available Ni–Co–Cr–AlY bond coat material, on the structural integrity of laser-DED-fabricated structures. Alloy modification can lead to high structural integrity by promoting the formation of desirable phases and improving laser absorptivity. Any changes in the feedstock powder material will lead to changes in the melt pool chemistry and detrimental compositions can be carefully avoided. Alloy chemistry is modified by varying the relative amounts of Ni and Cr in the alloy by adding powdered Ni–Cr to the parent alloy powder. A secondary objective of this study is to assess the feasibility of laser-DED techniques to fabricate bulk components (dimensions on the order of tens of centimeters) from A386. Laser-DED was chosen for this study because it has the ability to fabricate 3D structures and has deposition rates comparable to those of thermal spray processes. Thermodynamic phase modeling and experimental materials characterization tools are applied to provide insights into the phase and microstructural evolution of the alloy. A novel, laser-fabricable Ni–Co-based superalloy chemistry, containing a higher concentration of Al than currently used Ni-based alloys, is reported.

Laser-DED is an AM process that utilizes concentrated laser energy to melt and deposit metallic and/or ceramic powder on a substrate in a layer-wise fashion to form 3D structures. A 3D CAD model was designed and subsequently “sliced” into individual horizontal layers using a proprietary “slicing” software from Optomec Inc. Each slice represents a toolpath that the CNC controller interprets to fabricate the selected layer. Powdered feedstock material is fluidized by the carrier argon gas and blown through nozzles directly into the laser melt pool. The laser rapidly melts the incoming powder and a deposit is formed upon solidification. As the toolpath follows the deconstructed 3D model, repeated raster scanning, and layer-wise increments lead to a complete component.

Individual A386 alloy and NivCr alloy powders (Oerlicon Metco—MI, USA), with particle size in the range of 45–125 µm (0.00177–0.00492 in.), were pre-mixed via manual hand mixing and used as the feedstock material in this study. Table 1 shows the weight proportions of the constituent alloying elements for A386, Ni–Cr, and the compositions of the alloys modified with the addition of 10%, 20%, and 30% Ni–Cr. The A386 + Ni–Cr alloy coupons were fabricated using a Laser Engineered Net Shaping (LENS^™) print engine (Optomec Inc.—NM) equipped with a YLS 1000 W 1070 nm (4.21e⁻⁵ in.) laser (IPG Photonics, MA) within a glove box containing an argon environment. During fabrication, the moisture and oxygen levels were continuously monitored and maintained below 10 ppm. The powder mixtures were processed at 500 W laser power with a 900-µm (0.0354 in.) diameter laser spot size, resulting in a power density of approximately 800 MW/m². Coupons were fabricated with the four different compositions (by weight) listed in Table 1. A cuboidal geometry was selected with a square cross section, 25 mm × 25 mm (0.984 in. × 0.984 in.), that was built to a height of 6 mm (0.236 in.), for all the coupons. The deposition speed used was 20 mm/s (0.787 in./s) along with hatch spacing of 0.5 mm (0.0197 in.) and layer thickness of 0.5 mm (0.0197 in.). These coupons were deposited on a 6 mm (0.236 in.) thick stainless steel 304 (SS 304) substrate.

The phase with the highest melting point along the solidification path will solidify first. C_L will shift to a new composition as the liquid is consumed. The solidification/formation of a new phase (e.g., β or γ/γ′) is dictated by the solidus composition. Thermo-Calc facilitates the implementation of the Scheil–Gulliver model for multi-component alloys, by applying Eq. (1) for all constituent elements and phases in the alloy. All minor alloy elements, except for carbon and iron, which only occurred in trace amounts (<0.11%), were included in the Thermo-Calc simulations [68]. Exclusion of these two elements helped numerical stability. The Ni-based TCNI9 thermodynamics and MOBNI5 mobility databases were used for all simulations. Classic Scheil solidifications began above the alloy melting points and finished when 99.5% of the alloy was solidified.

To validate the non-equilibrium Scheil–Gulliver assumption, additional Thermo-Calc simulations with finite cooling rates were conducted that accounted for the back diffusion of solid phases during solidification. The results indicate that above 100 K/s the final phase fractions upon solidification converge with the Scheil–Gulliver model within 0.1%. Although cooling rates may vary between 10³ and 10⁵ K/s in actual DED processes, convergence results indicate that the Scheil model is a valid simplifying assumption for this study because of the relatively slow rates of solid-phase back diffusion. This Scheil assumption is also consistent with prior solidification modeling studies of Ni-based superalloys [68,69].

Individual laser-DED-fabricated coupons were cut from the SS 304 substrate and sectioned with an abrasive cutoff saw (Presi Mecatome T260—Eybens, FRANCE). One sample was harvested from each coupon in the build direction. Each coupon was mounted in thermosetting polymer mounting compound and wet ground from 60 to 1200 grit. The coupons were polished using a 0.06-µm (2.36e⁻⁶ in.) colloidal silica solution until all scratches were removed. After polishing, the coupons were cleaned in a 100% ethanol ultrasonic bath (SharperTEK XP PRO—MI, USA).

The laser-DED processed coupons were characterized to assess the structural integrity, microstructural features and phase evolution using optical microscopy, scanning electron microscopy (SEM), energy dispersive spectroscopy (SEM–EDS), X-ray diffraction (XRD), and microhardness. SEM imaging was performed on a field emission scanning electron microscope operating at 15 kV (JEOL 6610LV—Tokyo, JAPAN). The chemical analysis was conducted using an Energy Dispersive Spectroscopy (EDS) detector (Oxford Instruments X-Max detector—Abingdon, UK). Optical microscopy was performed using a Keyence microscope and Keyence analysis software (Keyence VHX S650E—Osaka, Japan) to achieve a composite image with high magnification (1000×) and large field-of-view (3 mm × 3 mm) (0.118 in. × 0.118 in.). Phase analysis was conducted using Rigaku Smartlab System (Tokyo, Japan) with Cu K_β filter at 40 kV for 2θ from 20 deg to 100 deg with a 0.05-deg step size. The surface roughness of the samples and all other XRD analysis parameters were kept constant in between different samples to allow for comparison of the XRD data. A Vickers microhardness tester (Phase II—NJ, USA) was used to measure the hardness of each coupon. Hardness tests were performed with a constant load of 100 g (0.220 lb) and a 15-s dwell time. Before hardness testing, the surface of each coupon was ground flat, thoroughly polished, and cleaned in a 100% ethanol ultrasonic bath.

Figure 1 shows that all bulk coupons fabricated by laser-DED are of the desired dimensions of 25 mm × 25 mm × 6 mm (0.984 in. × 0.984 in. × 0.236 in.) and deposited on the SS 304 substrate. The sample comprised of 100% A386 showed severe cracking and delamination during deposition. Even with modifications to the process parameters, such as laser power (varied between 400 and 700 W), deposition speed (varied between 17 and 25 mm/s; 0.669 in./s and 0.984 in./s), and the CAD deconstruction parameters such as hatch spacing and layer thickness (both varied between 0.3 and 0.5 mm; 0.0118 in. and 0.0197 in), it was not possible to fabricate a structurally sound coupon from the 100% A386 alloy powder. As Ni–Cr was added to A386 in increasing amounts, the cracking and delamination were reduced and the coupons showed improved structural integrity. The circle markers in Figs. 1(a) and 1(b) highlight the visible zones of defects such as hot cracking and delamination that occurred during fabrication. Even with the reduced tendency to delaminate and crack with the addition of 10% Ni–Cr, these coupons cracked and failed during sectioning. Small sections or pieces from 100% A386 and A386 + 10% Ni–Cr alloy mixtures were used for microstructural and phase analysis, but owing to their poor structural integrity, A386 and A386 + 10% Ni–Cr are inferred to be less suitable for fabrication via laser-DED processing.

The laser-DED fabrication of the mixtures with higher Ni–Cr content such as A386 + 20% Ni–Cr and A386 + 30% Ni–Cr was more successful. Upon visual inspection, these coupons (Figs. 1(c) and 1(d)) showed no external cracks or delamination. The coupons were structurally sound and did not fracture during sectioning, grinding, or polishing. A386 + 20% Ni–Cr revealed internal cracking similar to the A386 and A386 + 10% Ni–Cr samples when imaged using optical microscopy (Fig. 2(c)).

The lack of severe defects in the A386 + 20% Ni–Cr and A386 + 30% Ni–Cr coupons can be attributed to the increase in γ/γ′ as Ni–Cr additions increase. Although the predominant phase transitions from β to γ/γ′ between the A386 and A386 + 10% Ni–Cr mixtures, its stabilizing effects are not observed until the Ni–Cr addition reaches 20%. Since the A386 + 30% Ni–Cr coupon was free of structural defects, it was deemed a good choice for bulk, freeform depositions using direct laser fabrication. It may be possible that mixture compositions between 20% and 30% of Ni–Cr in A386 also lead to defect-free fabrication. However, due to the limited supply of specialty alloy powders, these intermediate compositions were not evaluated. A porosity analysis was performed using optical micrographs and Keyence analysis software, which found the planar porosity to be less than 0.05%. These results demonstrate the substantial influence of the alloy composition on the structural integrity of laser DED-manufactured coupons. Among the tested mixtures, the composition of A386 + 30% Ni–Cr addition resulted in the most successful as-fabricated coupons that were devoid of bulk cracking and interfacial delamination.

The Scheil Module in Thermo-Calc was used to explain why the structural stability of A386 improves with increasing Ni–Cr content [63]. Figure 3 illustrates the differences in the phase evolution behavior for all four A386 + Ni–Cr compositions. The 100% A386 begins solidifying around 1425 °C (2590 °F) as the β phase (BCC) shown in Fig. 3(a). As solidification progresses, the γ/γ′ phase (FCC) forms and becomes the matrix enveloping the β phase. Since yttrium has no solubility in the γ/γ′ or β phases, the last remaining liquid is yttrium rich, which results in the solidification of Ni₅Y (with Al solubility) within the grain boundaries. Solidification concludes at more than 100 °C (212 °F) below the equilibrium solidus line, 1120 °C (2048 °F) versus 1300 °C (2370 °F), showing the impact of non-equilibrium processes on alloy solidification. When 10% Ni–Cr is added, the β phase still forms first (Fig. 3(b)). For the A386 + 20% Ni–Cr and 30% Ni–Cr mixtures, the γ/γ′ phase solidifies first (Figs. 3(c) and 3(d)) at a lower liquidus temperature than for the 100% A386, which is facilitated by the increasing Ni content. Subsequently, the β-phase is formed as the matrix which holds the γ/γ′ phase, and finally, Ni₅–Y forms as solidification complete. The model predicts that as the wt% of Ni–Cr increases in the parent alloys, the weight fraction of the β phase decreases from 53.9 wt% in the 100% A386, to 15.5 wt% in the A386 + 30% Ni–Cr mixture. This trend is indicative of reduced cracking and improved structural integrity of the coupons as the weight percent of Ni–Cr increases. The model also predicts that around 2 wt% Ni₅Y for each alloy will form with the remaining composition being γ/γ′.

Hot cracking in Ni-based superalloys has become a complex challenge for manufacturers to overcome. Hot cracking refers to cracking which occurs above an alloy's solidus temperature. Hot cracking can be broken down into two different types of cracking, solidification cracking, and liquation cracking. Hot cracking typically occurs in hot manufacturing processes such as welding, casting, forging, and laser or electron beam additive manufacturing techniques. Solidification cracking occurs when the amount of liquid remaining in a melt pool is insufficient to fill gaps between dendrites as they grow during solidification [70]. These gaps between dendrites act as crack nucleation points which coalesce to form solidification cracks when the solidifying material is unable to accommodate the strain from thermal shrinkage. Liquation cracking refers to cracking caused by the partial melting of previously deposited material during hot manufacturing processes [71]. When partial melting occurs, low melting point eutectic phases (Ni₅Y for this alloy) liquefy and form an intergranular film which allows liquation cracks to form and propagate [72,73]. These films reduce microstructural stability and initiate liquation cracks when subjected to sufficient thermal stresses [74].

Hot cracking is common during DED fabrication of “unweldable” Ni-based superalloys which contain more than 6% of Al [75,76]. Varying temperature gradients incurred within a deposit during laser-DED lead to uneven thermal contraction throughout a deposit and result in hot cracking [77]. Major cracking in the A386 + 10% Ni–Cr coupon was found to be caused by liquation [74]. SEM micrographs in Fig. 4 show (a) intergranular cracking and (b) liquation features that initiated the hot crack formation.

Additional Thermo-Calc simulations were conducted to explore the impact of higher Ni–Cr content on crack susceptibility (Table 2). Up to A386 + 60%Ni–Cr, θ decreases indicating a lower likelihood of solidification cracking. At higher Ni–Cr content, the crack susceptibility slope increases again before decreasing for pure Ni–Cr. This indicates there may be a point beyond which adding Ni–Cr may be detrimental to crack susceptibility. Adding Ni–Cr content also significantly reduces the amount of aluminum in the alloy, so the alloy mixtures with increased Ni–Cr content likely have inferior corrosion resistance to the base A386. Further experimental investigation would be needed to confirm these trends.

In the Ni–Co–Cr–Al alloy system, γ/γ′ and β are the two prevailing phases [79,80]. The reported microstructures of the A386 + Ni–Cr mixtures are similar to other laser-processed Ni–Co–Cr–Al–Y microstructures [81,82]. The SEM backscatter images in Fig. 5 reveal the microstructure of the A386 + Ni–Cr mixtures consisting of the light contrast γ/γ′ phase and the dark contrast β phase. From these four micrographs, it is clear that the relative proportion of β is reduced as more Ni–Cr is added to the Amdry 386 + Ni–Cr mixtures. The microstructure of laser-processed alloys with lesser Ni–Cr alloy has a matrix of β phase with sparse, interconnected γ/γ′ phases Figs. 5(a)–5(c). As the Ni–Cr content increases, the γ/γ′ phases are stabilized and form coarse dendritic features Fig. 5(d).

XRD was used to understand the nature of these contrasting phases and relate them with the changes in the structural integrity of the material. An XRD comparison of A386 powder, Ni–Cr powder, and all four DED-processed mixtures of the parent alloys is shown in Fig. 6. The Ni–Cr alloy powder contains only γ/γ′ phases, whereas the feedstock A386 powder is rich in β phases with lesser amounts of γ/γ′. After laser DED fabrication, these β and γ/γ′ phases are retained in the A386 + Ni–Cr mixtures (Figs. 6(c)–6(f)). The results show strong γ/γ′ phase peaks and weak β phase peaks in the A386 + 30% Ni–Cr sample (Fig. 6(f)).

CALPHAD models predict the formation of a third phase, Ni₅Y, which comprises roughly 2 wt% of the laser-processed microstructure. The relatively small amount of Ni₅Y is anticipated to produce low-intensity XRD peaks. However, all Ni₅Y XRD peaks coincide with high-intensity peaks from either of the two dominant phases, which makes the peaks indistinguishable. Therefore, the presence of Ni₅Y even when theoretically established cannot be directly verified using XRD measurements alone. While the Ni₅Y phase is predicted to be formed by the Thermo-Calc model (Fig. 8), additional analysis including transmission electron microscopy (TEM) is needed to confirm the presence of the Ni₅Y phase.

The model predictions for the increase in γ/γ′ phases with increasing Ni-content due to the addition of Ni–Cr to the base A386 alloy are further verified by comparing the predictions with experimental measurements for the relative percentage of β phase in Fig. 7. Error bars represent the standard deviation in the measured data due to image processing and analyses. The A386 alloy contained a relative β phase area of 58.48 ± 2.25%, and its proportion reduced to 33.8 ± 3.02% in the A386 + 30% Ni–Cr due to the concurrent rise in γ/γ′. Figure 7 depicts that the model predictions for the effect of the composition of the alloy on the dependence of the relative β phase area are consistent with experimental measurements. Compared to the base A386 alloy, which is chemically 45.8% Ni, 22.8% Co, 17.3% Cr, and 12.4% Al, the addition of 30% Ni–Cr alloy to A386 causes the wt% of Ni and Cr to increase to 55.8% and 17.9%, respectively, whereas the Al and Co were reduced to 8.69% and 16.0%, respectively (Table 1). Therefore, as Ni–Cr is added to A386, the fraction of γ/γ′ begins to dominate the microstructure and the structural integrity is increased.

While the trends predicted by the models and the image analyses converge as the Ni–Cr content decreases, the model consistently underestimates the amount of the β (BCC) phase (Fig. 7). A reason for the model underestimation may be due to the Scheil model neglecting solid-state phase transformation. During laser-DED processes, each time an additional layer is deposited and melted there is a temperature increase in the layers below. This results in either remelting or reheating deposited layers, which may result in significant solid-state phase transformation [68,83–85].

While some additional diffusion may occur at lower temperatures, most phase transformation will occur at elevated temperatures as solid-state diffusion rates tend to increase exponentially with temperature [86]. To show the expected phases below the solidus, an equilibrium phase diagram of A386 + 30% Ni–Cr is plotted in Fig. 8. While an equilibrium phase diagram by itself cannot predict solid-phase diffusion of the non-equilibrium laser-DED process, the phase diagram can still shed insight into the thermodynamic driving forces if solid-phase transformation were to occur.

The results reveal that as the temperature decreases below the solidus, the β phase fraction increases while the γ/γ′ (FCC) fraction decreases with the β fraction approaching the experimental results in Fig. 7 at 1000 °C. This result suggests that if any solid-phase diffusion occurs, the β fraction would likely increase, which would give better agreement between the model and experiments for high Ni–Cr content alloys. Phase transformations of up to several percent for laser-based processes are possible although it depends on the thermal profile and partition coefficients of the alloy [69,86–89]. To do a more comprehensive phase analysis, the thermal history of the solidification process is needed as well as a more complex thermal model that accounts for dendritic structures, such as a phase field model coupled with the DICTRA diffusion model [63,86]. This will be considered in future research work.

The phase transformation and subsequent improvement in the structural integrity of A386 + Ni–Cr alloys are also supported by SEM–EDS analysis. In 100% A386, major elements such as Ni, Co, and Cr are evenly distributed whereas Al is segregated at the microscopic scale, with stronger signal of Al corresponding to the darker β phase (Fig. 9). Similarly, in the A386 + 30% Ni–Cr coupon, Al is richer in the β phase regions (Fig. 10). However, in the A386 + 30% Ni–Cr alloy, these darker β phase regions are significantly less than that of 100% A386. The EDS maps of Figs. 9 and 10 show segregation of yttrium, which coincides with the brightest spots in the microstructure. This result has previously been identified with the Ni₅Y phase experimentally [88,89].

Microhardness measurements obtained from the center of each coupon are shown in Table 3. The hardness of the 100% A386 coupon is the highest at 606 ± 13.6 HV. After the additions of 10% and 20% Ni–Cr alloy, the hardness reduces to 547 ± 13.6 HV and 475 ± 15.8, respectively, continuing to decrease to 412 ± 19.1 HV with 30% Ni–Cr addition. This gradual reduction in the hardness can be attributed to the reduction of β phase.

The feasibility of the novel A386 + 30% Ni–Cr alloy mixture for making practical components without any cracking or delamination failure is further evaluated by depositing two types of geometries. The first geometry is a plate with an area of 125 mm × 75 mm (4.92 in. × 4.95 in.) and a thickness of 6 mm (0.236 in.). The second geometry is a bar with a build height of 120 mm (4.72 in.) and a planar area of 19 mm × 12.5 mm (0.748 in. × 0.498 in.; Fig. 11). These geometries are deposited without any interruptions in the DED process and are free of any visible cracks or delamination defects.

A386 has a composition similar to other common Ni-based superalloys listed in Table 4, but the alloy’s high Al content gives it superior corrosion resistance. Ni–Cr additions to the A386 parent alloy do reduce the relative amount of Al in the mixture, but even the A386 + 30% Ni–Cr has 1.4 times more Al than the commercially available laser-fabricable alumina forming Ni-based superalloys such as Inconel 713 and René N5 (6–6.2%). Therefore, the alloy composition reported here could be a strong candidate for applications where high temperatures and corrosive environments are present. The ability to fabricate near-net-shaped structures using laser-based AM processes such as laser-DED further enables the deployment of this alloy in novel structures.

The present work investigates the ability to create bulk samples of the Ni–Co-based, dual phase superalloy A386 using laser-DED. The base alloy chemistries were modified by controlled additions of a Ni–Cr alloy to create mixtures with varying relative amounts of Ni, Co, Cr and Al. The effects of alloy compositions on the fabricability, structural integrity, and hardness were evaluated through a combination of thermodynamic phase modeling and experimental materials characterization and testing. The as-received A386 and Ni–Cr feedstock powders were comprised of mostly β with smaller amounts of γ/γ′ phases, and almost completely γ/γ′, respectively. The constituent phases of A386 powder were retained in relatively the same proportions after laser-DED processing. The dominating phase switches from β to γ/γ′ between the 0% Ni–Cr and 10% Ni–Cr mixtures. This microstructural switch from β to γ/γ′ increases the structural integrity as the wt% of Ni–Cr approaches 30%. Additionally, Thermo-Calc modeling results based on Scheil solidification and thermodynamic phase calculations match the experimentally observed trends for the extent of β-phase reduction with increasing Ni–Cr additions to the base alloy. The modified alloy chemistries allowed for the successful fabrication of larger-scale components. Specifically, A386 + 30% Ni–Cr is reported to be a viable mixture for direct laser fabrication of near-net shaped components. A plate and a bar geometry were successfully fabricated using laser-DED, with the largest dimensions on the order of tens of centimeters. Modeling predictions demonstrate that Ni–Cr addition beyond 60% will lead to diminishing improvements to the crack susceptibility of the base alloy. Future work will explore high-temperature mechanical characterization of A386 + 30% Ni–Cr alloy chemistry in the as-processed state as well as after post-processing heat treatment.

Kinzer acknowledges software support for Thermo-Calc offered by CAEN at the University of Michigan. Authors are also grateful to Dr. Alexandra Zevalkink for access to XRD analysis facilities. Authors also acknowledge the equipment support from Center for Advanced Microscopy (CAM) at Michigan State University.

There are no conflicts of interest.

The datasets generated and supporting the findings of this article are obtainable from the corresponding author upon reasonable request.

This work was supported by the Department of Energy Advanced Research Projects Agency-Energy (ARPA-E) under co-operative agreement DE-AR0001123 with Michigan State University and the University of Michigan. Authors Sahasrabudhe and O’Neil also acknowledge startup support from Michigan State University. Bala Chandran and Kinzer acknowledge startup funding from the College of Engineering and the Department of Mechanical Engineering at the University of Michigan.