Abstract

Ni- and Co-based M–Cr–Al–Y (M = Ni and/or Co), dual phase β and γ/γ′, β—Ni–Al, γ—Ni(Co, Cr), and γ′—Ni3–Al superalloys exhibit several beneficial high-temperature (>1000 °C) (2012 °F) strength and corrosion resistance properties. Our study investigates the feasibility of fabricating a commercially available Ni–Co–Cr–AlY alloy (tradename “Amdry 386”) using laser directed energy deposition (laser-DED). A significant challenge for laser fabrication of bulk Amdry 386 (A386) structures stems from the presence of relatively higher amounts of the β phase than γ/γ′ phases. While prior studies report laser fabrication of these materials in coating and cladding configurations, bulk structures have not been realized. To investigate these challenges, this study was developed to systematically evaluate the effects of modifying the A386 alloy by adding 10, 20, and 30% by weight of a Ni–20%Cr (Ni–Cr) alloy (mainly γ/γ′). Laser-DED-fabricated A386 coupons exhibited a metastable, rapidly solidified β-rich microstructure typical to laser processes. The A386 + Ni–Cr mixtures were processed using laser-DED, and small coupons from each composition were evaluated using SEM, XRD, and microhardness to characterize the as-processed microstructure. Thermodynamic modeling was performed to determine the phase evolution as a function of the alloy composition. The dominating phase switches from β to γ/γ’ between the A386 and A386 + 10% Ni–Cr mixtures, but the increase in structural integrity is not observed until the Ni–Cr additions reach a minimum of 20%. Our results show that the alloy chemistry can be successfully modified to improve the structural integrity of laser-processed structures.

1 Introduction

Ni-based superalloys find applications in aerospace and power generation systems where high temperatures (600–1100 °C; 1112–2012 °F) are encountered [17]. 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 [912]. A386 is an example of a NiCo 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 [1416]. 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 MCrAlY 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 [1923]. 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 MCrAlY 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 [2629]. 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 [3234], aerospace components [29,35,36], and functional coatings [3739], 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 [4047]. 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 [45,5055]. 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 [2631].

The laser-DED process, and its commercial variants such as the LENSTM 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 [5759]. 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 [41,6062]. 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 NiCoCrAlY 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 NiCr 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 NiCo-based superalloy chemistry, containing a higher concentration of Al than currently used Ni-based alloys, is reported.

2 Materials and Methods

2.1 Laser-Directed Energy Deposition Fabrication.

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, NiCr, and the compositions of the alloys modified with the addition of 10%, 20%, and 30% NiCr. The A386 + NiCr 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.

Table 1

Compositions of parent alloys—A386 and NiCr as obtained from the powder supplier and the mixture composition of the coupons that were fabricated with varying amounts of NiCr added to A386

Coupons (wt%)Weight percentage of constituent elements
NiCoCrFeSiMnAlCHfY
100% A38645.822.817.30.4612.40.020.280.61
Ni20%Cr79.319.40.110.760.39
A386 + 10% NiCr49.120.617.50.010.490.0411.20.020.250.55
A386 + 20% NiCr52.518.317.70.020.520.089.940.020.220.49
A386 + 30% NiCr55.816.017.90.030.550.128.700.010.200.43
Coupons (wt%)Weight percentage of constituent elements
NiCoCrFeSiMnAlCHfY
100% A38645.822.817.30.4612.40.020.280.61
Ni20%Cr79.319.40.110.760.39
A386 + 10% NiCr49.120.617.50.010.490.0411.20.020.250.55
A386 + 20% NiCr52.518.317.70.020.520.089.940.020.220.49
A386 + 30% NiCr55.816.017.90.030.550.128.700.010.200.43

2.2 Thermodynamic Modeling of Phase Evolution.

CALculation of PHAse Diagrams (CALPHAD) modeling was performed to predict the thermodynamics of the influence of NiCr addition on the phases formed in the alloy during solidification using Thermo-Calc (version 2021b) [63]. Due to the rapid cooling rates of laser-DED processing (103–105 K/s) (>103–105 °F/s) [64], solidification will be a non-equilibrium process. These non-equilibrium effects can be accounted for using the Scheil–Gulliver model, which assumes infinitely fast species diffusion (perfect mixing) in the liquid phase, no diffusion in the solidified phases, and equilibrium of the solid–liquid interface [65,66]. Applying conservation of mass to the Scheil–Gulliver model yields Eq. (1) [67]
Cs=kC0(1fs)k1
(1)
where Co is the initial alloy chemical composition (weight percent), Cs is the solid composition (weight percent), CL is the liquid composition (weight percent), fs is the weight fraction of the individual solid phases, and k is the partition coefficient k = Cs/CL.

The phase with the highest melting point along the solidification path will solidify first. CL 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].

2.3 Characterization and Structural Integrity Assessment.

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 Kβ 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.

3 Results and Discussion

3.1 Structural Integrity of Laser-Directed Energy Deposition Fabricated Coupons.

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 NiCr 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% NiCr, these coupons cracked and failed during sectioning. Small sections or pieces from 100% A386 and A386 + 10% NiCr alloy mixtures were used for microstructural and phase analysis, but owing to their poor structural integrity, A386 and A386 + 10% NiCr are inferred to be less suitable for fabrication via laser-DED processing.

Fig. 1
Four different A386 + Ni–Cr mixtures fabricated using laser DED technique for (a) A386, (b) A386 + 10% Ni–Cr, (c) A386 + 20% Ni–Cr, and (d) A386 + 30% Ni–Cr. Circled areas show the presence of defects in (a) and (b), and the absence of visible defects in (c) and (d). All coupons measure 25 mm × 25 mm × 6 mm (0.9843 in × 0.9843 in × 0.2362 in), approximately.
Fig. 1
Four different A386 + Ni–Cr mixtures fabricated using laser DED technique for (a) A386, (b) A386 + 10% Ni–Cr, (c) A386 + 20% Ni–Cr, and (d) A386 + 30% Ni–Cr. Circled areas show the presence of defects in (a) and (b), and the absence of visible defects in (c) and (d). All coupons measure 25 mm × 25 mm × 6 mm (0.9843 in × 0.9843 in × 0.2362 in), approximately.
Close modal

The laser-DED fabrication of the mixtures with higher NiCr content such as A386 + 20% NiCr and A386 + 30% NiCr 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% NiCr revealed internal cracking similar to the A386 and A386 + 10% NiCr samples when imaged using optical microscopy (Fig. 2(c)).

Fig. 2
Optical imaging for (a) A386, (b) A386 + 10% Ni–Cr, (c) A386 + 20% Ni–Cr, and (d) A386 + 30% NiCr
Fig. 2
Optical imaging for (a) A386, (b) A386 + 10% Ni–Cr, (c) A386 + 20% Ni–Cr, and (d) A386 + 30% NiCr
Close modal

The lack of severe defects in the A386 + 20% NiCr and A386 + 30% NiCr coupons can be attributed to the increase in γ/γ′ as NiCr additions increase. Although the predominant phase transitions from β to γ/γ′ between the A386 and A386 + 10% NiCr mixtures, its stabilizing effects are not observed until the NiCr addition reaches 20%. Since the A386 + 30% NiCr 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 NiCr 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% NiCr addition resulted in the most successful as-fabricated coupons that were devoid of bulk cracking and interfacial delamination.

3.2 Phase Analyses: Thermodynamic Models and Experimental Characterization.

The Scheil Module in Thermo-Calc was used to explain why the structural stability of A386 improves with increasing NiCr content [63]. Figure 3 illustrates the differences in the phase evolution behavior for all four A386 + NiCr 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% NiCr is added, the β phase still forms first (Fig. 3(b)). For the A386 + 20% NiCr and 30% NiCr 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, Ni5Y forms as solidification complete. The model predicts that as the wt% of NiCr 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% NiCr mixture. This trend is indicative of reduced cracking and improved structural integrity of the coupons as the weight percent of NiCr increases. The model also predicts that around 2 wt% Ni5Y for each alloy will form with the remaining composition being γ/γ′.

Fig. 3
Scheil solidification profiles of A386 and NiCr for temperature versus the fraction solidified computed using Thermo-Calc for (a) 100% A386, (b) A386 + 10% NiCr, (c) A386 + 20% NiCr, and (d) A386 + 30% NiCr. The slope of the blue dashed line represents solidification crack susceptibility with θ=dT/d(fs1/2) at (fs)1/2 = 0.99.
Fig. 3
Scheil solidification profiles of A386 and NiCr for temperature versus the fraction solidified computed using Thermo-Calc for (a) 100% A386, (b) A386 + 10% NiCr, (c) A386 + 20% NiCr, and (d) A386 + 30% NiCr. The slope of the blue dashed line represents solidification crack susceptibility with θ=dT/d(fs1/2) at (fs)1/2 = 0.99.
Close modal

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% NiCr 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.

Fig. 4
Fracture surface of hot crack in A386 + 10% Ni–Cr: (a) intergranular fracture at 250X (b) magnification of (a) showing liquation features at 2000X
Fig. 4
Fracture surface of hot crack in A386 + 10% Ni–Cr: (a) intergranular fracture at 250X (b) magnification of (a) showing liquation features at 2000X
Close modal
Kou demonstrated that an alloy’s susceptibility to solidification cracking can be quantified by
θ=dT/d(fs)1/2
(2)
where θ is the slope, t is the temperature and fs is the fraction solidified [78]. The index to crack susceptibility is given by the value of θ (°C) when (fs)1/2 is 0.99 [78]. θ for each of the four alloys can be seen in Table 2 and are represented by the blue dotted lines in Fig 3. Steeper slopes correspond to greater susceptibility to solidification cracking. The crack susceptibility of the A386 + NiCr mixtures decreases with greater additions of NiCr. This result is consistent with interpretations gained from visual inspection, microhardness, and the relative amounts of each phase.
Table 2

Slope of Scheil solidification profiles at 0.99(fs)1/2

Materialθ Slope at 0.99(fs)1/2 (°C)Materialθ Slope at 0.99(fs)1/2 (°C)Materialθ Slope at 0.99(fs)1/2 (°C)
A3866,511A386 + 40% NiCr3,271A386 + 80% NiCr2,309
A386 + 10% NiCr4,933A386 + 50% NiCr2,838A386 + 90% NiCr3,127
A386 + 20% NiCr4,531A386 + 60% NiCr2,241Pure NiCr875
A386 + 30% NiCr3,984A386 + 70% NiCr4,356
Materialθ Slope at 0.99(fs)1/2 (°C)Materialθ Slope at 0.99(fs)1/2 (°C)Materialθ Slope at 0.99(fs)1/2 (°C)
A3866,511A386 + 40% NiCr3,271A386 + 80% NiCr2,309
A386 + 10% NiCr4,933A386 + 50% NiCr2,838A386 + 90% NiCr3,127
A386 + 20% NiCr4,531A386 + 60% NiCr2,241Pure NiCr875
A386 + 30% NiCr3,984A386 + 70% NiCr4,356

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

In the NiCoCrAl alloy system, γ/γ′ and β are the two prevailing phases [79,80]. The reported microstructures of the A386 + NiCr mixtures are similar to other laser-processed NiCoCrAlY microstructures [81,82]. The SEM backscatter images in Fig. 5 reveal the microstructure of the A386 + NiCr 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 NiCr is added to the Amdry 386 + NiCr mixtures. The microstructure of laser-processed alloys with lesser NiCr alloy has a matrix of β phase with sparse, interconnected γ/γ′ phases Figs. 5(a)5(c). As the NiCr content increases, the γ/γ′ phases are stabilized and form coarse dendritic features Fig. 5(d).

Fig. 5
Microstructures of DED fabricated A386 + Ni–Cr mixtures: (a) 100% A386, (b) A386 + 10% Ni–Cr, (c) A386 + 20% Ni–Cr, and (d) A386 + 30% Ni–Cr
Fig. 5
Microstructures of DED fabricated A386 + Ni–Cr mixtures: (a) 100% A386, (b) A386 + 10% Ni–Cr, (c) A386 + 20% Ni–Cr, and (d) A386 + 30% Ni–Cr
Close modal

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, NiCr powder, and all four DED-processed mixtures of the parent alloys is shown in Fig. 6. The NiCr 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 + NiCr mixtures (Figs. 6(c)6(f)). The results show strong γ/γ′ phase peaks and weak β phase peaks in the A386 + 30% NiCr sample (Fig. 6(f)).

Fig. 6
XRD analysis of feedstock powder and DED processed coupons: (a) A386 powder, (b) Ni–Cr powder, (c) A386, (d) A386 + 10% Ni–Cr, (e) A386 + 20% Ni–Cr and (f) A386 + 30% Ni–Cr
Fig. 6
XRD analysis of feedstock powder and DED processed coupons: (a) A386 powder, (b) Ni–Cr powder, (c) A386, (d) A386 + 10% Ni–Cr, (e) A386 + 20% Ni–Cr and (f) A386 + 30% Ni–Cr
Close modal

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 NiCr 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% NiCr 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% NiCr 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 NiCr is added to A386, the fraction of γ/γ′ begins to dominate the microstructure and the structural integrity is increased.

Fig. 7
Relative proportions of β phase in the A386 + Ni–Cr mixtures gathered through theoretical and experimental methods
Fig. 7
Relative proportions of β phase in the A386 + Ni–Cr mixtures gathered through theoretical and experimental methods
Close modal

While the trends predicted by the models and the image analyses converge as the NiCr 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 [68,8385].

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% NiCr 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.

Fig. 8
Equilibrium phase diagram of A386 + 30%Ni–Cr from Thermo-Calc
Fig. 8
Equilibrium phase diagram of A386 + 30%Ni–Cr from Thermo-Calc
Close modal

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 NiCr 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 [69,8689]. 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 + NiCr 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% NiCr coupon, Al is richer in the β phase regions (Fig. 10). However, in the A386 + 30% NiCr 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].

Fig. 9
SEM-EDS elemental mapping of laser DED fabricated 100% A386
Fig. 9
SEM-EDS elemental mapping of laser DED fabricated 100% A386
Close modal
Fig. 10
SEM-EDS elemental mapping of laser DED fabricated A386 + 30% Ni–Cr
Fig. 10
SEM-EDS elemental mapping of laser DED fabricated A386 + 30% Ni–Cr
Close modal

3.3 Mechanical Behavior, Scalability, and Material Comparison.

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% NiCr alloy, the hardness reduces to 547 ± 13.6 HV and 475 ± 15.8, respectively, continuing to decrease to 412 ± 19.1 HV with 30% NiCr addition. This gradual reduction in the hardness can be attributed to the reduction of β phase.

Table 3

Vicker’s micro-hardness test results of A386 + NiCr mixture coupons

MaterialA386A386 + 10% NiCrA386 + 20% NiCrA38 + 30% NiCr
Average Hardness (Vicker’s)606 ± 13.6546.8 ± 13.6475 ± 15.8411.5 ± 19.1
MaterialA386A386 + 10% NiCrA386 + 20% NiCrA38 + 30% NiCr
Average Hardness (Vicker’s)606 ± 13.6546.8 ± 13.6475 ± 15.8411.5 ± 19.1

The feasibility of the novel A386 + 30% NiCr 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.

Fig. 11
Large near-net shaped components fabricated from A386 + 30% Ni–Cr alloy mixture using laser-DED: (a) a large plate-shaped coupon and (b) a tall bar-shaped coupon
Fig. 11
Large near-net shaped components fabricated from A386 + 30% Ni–Cr alloy mixture using laser-DED: (a) a large plate-shaped coupon and (b) a tall bar-shaped coupon
Close modal

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. NiCr additions to the A386 parent alloy do reduce the relative amount of Al in the mixture, but even the A386 + 30% NiCr 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.

Table 4

Compositions of common laser processable Ni-based superalloys

MaterialWeight percentage of constituent elements
NiCoCrMoFeAlTiWTaNbHfY
A38645.822.817.312.40.280.61
NiCr79.319.40.11
A386 + 10% NiCr49.220.517.50.0111.10.250.55
A386 + 20% NiCr52.518.217.70.029.920.220.49
A386 + 30% NiCr55.916.017.90.038.680.200.43
Haynes 21475.016.03.004.500.01
Inconel 10061.415.09.53.000.505.504.75
Inconel 71375.512.03.900.185.990.631.80
Inconel 73861.58.5012.03.900.185.990.631.80
MAR-M 24757.59.489.050.785.521.1211.23.780.121.18
René N560.68.07.002.006.201.005.007.000.100.10
René 10061.015.09.503.000.505.504.20
René 14260.812.06.801.506.154.906.351.50
MaterialWeight percentage of constituent elements
NiCoCrMoFeAlTiWTaNbHfY
A38645.822.817.312.40.280.61
NiCr79.319.40.11
A386 + 10% NiCr49.220.517.50.0111.10.250.55
A386 + 20% NiCr52.518.217.70.029.920.220.49
A386 + 30% NiCr55.916.017.90.038.680.200.43
Haynes 21475.016.03.004.500.01
Inconel 10061.415.09.53.000.505.504.75
Inconel 71375.512.03.900.185.990.631.80
Inconel 73861.58.5012.03.900.185.990.631.80
MAR-M 24757.59.489.050.785.521.1211.23.780.121.18
René N560.68.07.002.006.201.005.007.000.100.10
René 10061.015.09.503.000.505.504.20
René 14260.812.06.801.506.154.906.351.50

4 Conclusions

The present work investigates the ability to create bulk samples of the NiCo-based, dual phase superalloy A386 using laser-DED. The base alloy chemistries were modified by controlled additions of a NiCr 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 NiCr 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% NiCr and 10% NiCr mixtures. This microstructural switch from β to γ/γ′ increases the structural integrity as the wt% of NiCr 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 NiCr additions to the base alloy. The modified alloy chemistries allowed for the successful fabrication of larger-scale components. Specifically, A386 + 30% NiCr 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 NiCr 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% NiCr alloy chemistry in the as-processed state as well as after post-processing heat treatment.

Acknowledgment

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.

Conflict of Interest

There are no conflicts of interest.

Data Availability Statement

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

Funding Data

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.

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