
==== Front
Nat Commun
Nat Commun
Nature Communications
2041-1723
Nature Publishing Group UK London

52261
10.1038/s41467-024-52261-1
Article
Multicomponent alloys designed to sinter
Naunheim Yannick 1
http://orcid.org/0000-0001-9856-2682
Schuh Christopher A. schuh@northwestern.edu

12
1 https://ror.org/042nb2s44 grid.116068.8 0000 0001 2341 2786 Department of Materials Science and Engineering, Massachusetts Institute of Technology, 77 Massachusetts Avenue, Cambridge, MA 02139 USA
2 https://ror.org/000e0be47 grid.16753.36 0000 0001 2299 3507 Department of Materials Science and Engineering, Northwestern University, 2145 Sheridan Road, Evanston, IL 60208 USA
13 9 2024
13 9 2024
2024
15 80281 4 2024
2 9 2024
© The Author(s) 2024
2024
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Powder sintering is a low-energy, net-shape processing route for many new products in the additive manufacturing space. We advance the viewpoint that for future manufacturing, alloys should be designed from materials science principles to sinter quickly at lower temperatures and with controlled final microstructures. Specifically, we illustrate the computational design of multinary Ni-base alloys, whose chemistries permit a low-temperature solid-state sintering scheme without any pressure- or field-assistance, as well as heat-treatability after sintering. The strategy is based on sequential phase evolutions designed to occur during sintering. The reactions involve rapid reorganization of matter to full density in cycles up to just 1200 °C, while conventional Ni alloys sintered in the solid-state require about ten times longer, or more than 250 °C degrees higher temperature. Our approach yields an alloy that benefits from precipitation hardening, has an increased strength ~50% higher than solid-state processed commercial Ni alloys, and yet exhibits extensive plasticity beyond 35% uniaxial strain. The results point to a generalizable design scheme for many other alloys designed for solid-state powder processing that can enable greater value from additive manufacturing.

In this work, the authors advance the viewpoint that powder alloys yield better processability and final properties, when they are designed from materials science principles to sinter quickly and with controlled microstructure evolution.

Subject terms

Metals and alloys
Design, synthesis and processing
https://doi.org/10.13039/100000001 National Science Foundation (NSF) DMR-1419807 DMR-1419807 Naunheim Yannick Schuh Christopher A. https://doi.org/10.13039/100006197 NASA | Marshall Space Flight Center 80MSFC19C0050 80MSFC19C0050 Naunheim Yannick Schuh Christopher A. issue-copyright-statement© Springer Nature Limited 2024
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pmcIntroduction

Sintering is an increasingly attractive alternative to traditional metal processing routes such as melt casting and working, since it allows for more geometrical flexibility and avoids melting. In an additive manufacturing context, this leads to lowered processing temperatures, less internal stress development, and potentially broader microstructural control1–3. However, sintering can also often require long time-at-temperature cycles to effectively remove enough porosity for mechanical viability, which can in turn promote undue grain growth or phase evolution4–6. Green body printing also produces different (generally lower) green densities and particle configurations than other powder routes, so the expectations of the sintering process are higher for complex parts production by 3D printing. Therefore, the current rapid expansion of sintering-based additive manufacturing is in need of new materials science-based approaches to address these challenges7,8.

Considering the trajectory of powder metallurgy specifically for additive manufacturing, it is typical to see extant alloys thrust into new processing modalities for which they are not optimized9–12. Some examples for structural alloys include the binder jetting and sintering of Inconel 625, which only reached a relative density of ~75% despite sintering in the vicinity of its solidus temperature13, or a similar process with steel 316 L that attained only ~85% density, even when sintered up to its melting point14. In both these cases, full density could only be obtained by sintering up into the range of liquification (super-solidus sintering), which creates new challenges for shape preservation15 due to slumping and distortion of the semi-solid structures.

Beyond these examples, it is generally the case that considerable efforts are needed to deal with the complexities of introducing standard alloys into an additive paradigm16–21. In the case of melting-solidification-reheating cycles (as in laser-based manufacturing), some authors have begun to specifically design alloys tuned to the new processing path22–24, but such an approach has not yet been applied to sintered (unmelted) alloys. Alloy design approaches have led to the development of new materials that overcome many classical challenges25–29 for a variety of applications30. Taking inspiration from those works, we propose the use of modern computational tools to deliberately design alloys processible by low-temperature solid-state free-sintering in the present paper. Our approach uses materials design principles to manage the phase evolution of powdered alloys, which in turn depend on the alloy thermodynamics and can be directly controlled by the alloy chemistry. More specifically, the chemical compositions of the alloys are designed to pass through a sequence of phase evolutions that accelerate mass transport, develop an interconnected monolithic structure at low temperatures, and facilitate rapid sintering at higher temperatures. After consolidation, the alloys are also designed to be homogenizable (fully soluble at higher temperatures) and heat treatable (dual-phase at lower temperatures) to effect a variety of microstructures. We elaborate this approach specifically in the design of Ni-based alloys. Not only do the newly designed alloys have advantages for the economical production of large-scale bulk components with improved mechanical properties, but the methodology is also generalizable to other families of alloys designed specifically for solid-state part production routes.

Results

Alloy design for rapid sintering

Our approach is based on the concept of nano-phase separation sintering, which is an all-solid-state phenomenon that was initially discovered to accelerate powder consolidation in some binary (refractory) alloys31–33. Briefly, the phenomenon happens when: (1) the powder is in a non-equilibrium solid solution state prior to sintering, and thus (2) thermodynamically prefers to pass through a sequence of phase evolutions during the sintering cycle. The chemical driving forces for phase evolution promote substantial mass transport, which in turn is biased by the surface area in a sintering compact, accelerating densification.

Requirement (1) speaks to a grain refinement and an out-of-equilibrium chemical state of the powder before heating. High-energy ball milling is ideal for this requirement, because it leads to refined grains34,35 and mechanically alloyed constituents that are thermodynamically immiscible in equilibrium (Fig. 1a)36. Both effects facilitate the preparation of powders that are poised to undergo rapid sintering.Fig. 1 Schematic of nano-phase separation sintering.

a Powders are brought into a non-equilibrium state via high-energy ball-milling as the grain size is refined and an immiscible solute is mechanically alloyed far beyond the equilibrium solubility. The gray color denotes the solvent-rich phase with a nanocrystalline grain size and dissolved solute atoms (orange, inset) (b) Upon heating, the excess solute is ejected from the supersaturated powder, forming a secondary phase at interparticle necks. Such structure (dashed rectangle) is shown in the micrographs in Fig. 4a. c At a late stage of the sintering cycle, chemical homogenization between the phases leads to rapid consolidation of the powder compact mitigating residual porosity, as indicated by the arrows.

For requirement (2), the alloy thermodynamics must be conducive to sequential phase decomposition and remixing upon heating. At low sintering temperatures, the powder evolves into a dual-phase microstructure from the initially supersaturated state (Fig. 1b). If the alloy is designed such that the second phase has a lower surface energy than the first, the release of the excess solute from the powders leads to preferential surface decoration and formation of interparticle necks, and the mass transport is accelerated by the shortened (grain) diffusion length scale. Once it achieves this dual-phase structure, consolidation of the powder is accelerated by chemical homogenization at higher temperature (where sintering completes, Fig. 1c). This stage typically makes up the major contribution to consolidation, with the redissolution of the second phase providing rapid chemical diffusion that accelerates densification.

We propose that the direct link between the alloy chemistry and powder consolidation in this nano-phase separation sintering process can become a powerful tool for alloy design. While there are few binary alloys in which the sequence of events from Fig. 1 happens fortuitously31–33, by contrast, we hypothesize that in a multicomponent space this sintering-accelerating phase evolution sequence can be designed for explicitly. That is the approach of the present paper: beginning with Ni as base, we proceed to design multicomponent alloys that combine phase thermodynamics and microstructure evolution to undergo nano-phase separation sintering. Specifically, our task is to design an alloy with two immiscible phases, which phase separate and then redissolve in sequence upon heating, at temperatures conducive to accelerating a sintering cycle.

There are relatively few immiscible Ni-based systems, and some like Ni-Au are not economically practical. However, one prototype that provides an excellent basis for design is Ni-Cu. Figure 2a illustrates the miscibility gap of this binary system; it meets most of the requirements to favor phase-separation sintering and is therefore attractive on the surface: it exhibits a miscibility gap at low temperature and solid-solution at high temperature. What is more, the second phase is copper-rich, has a low surface energy that would favor neck formation between particles and is a faster diffuser than Ni, which could accelerate kinetics generally. Unfortunately, the transus temperature is too low to kinetically allow for phase decomposition (Fig. 1a); effectively no diffusion happens at the temperatures in the miscibility gap (as confirmed by control experiments shown later).Fig. 2 Design of alloys for rapid solid-state sintering at low temperatures, exercised on the Ni-Cu binary couple.

a While the miscibility gap in Ni-Cu is too low to facilitate diffusion, adding a third component can elevate the transus temperature, as shown for various fixed Ni:Cu ratios while alloying with Co. b To identify the elements for this alloy addition, potential candidates preferentially have a low interaction with Ni (and thus some solubility; red shaded area, left phase diagram), and exhibit preferentially a positive heat of mixing with the second phase forming Cu (red shaded area, right phase diagram). Potential alloying candidates that exhibit these two requirements (blue arrows) are fully marked (orange arrow) in green in (c). Those candidates that pass the screening are then subjected to computational predictions (d, e) that assess the multi-dimensional thermodynamic and kinetic space of the alloy design process. The choice of solute species and its concentration affects the effective interdiffusion (d) and the surface wettability (e) of the second phase. Preferred alloys attain the green region at the top of the diagrams in part (d and e). The baseline system Ni-20Cu is denoted by a red star. f All three selected elements (yellow bars) carbothermally reduce within the densification window and therefore require a remedy to avoid a competition between densification and organic removal: oxygen scavengers that reduce at temperatures beyond the processing window are shown, among which Mn (green bar) is the only candidate that comports with the design criteria in part c (red bar otherwise)39,40.

The goal of our design exercise is therefore to identify alloying components for Ni-Cu that purposely engineer the miscibility gap temperature. Figure 2a compiles a set of miscibility gaps which we assess using computational thermodynamics with the commercial code ThermoCalc. When augmented by the third component such as Co in the current example, the Ni-Cu system has a highly tunable phase structure and is designable to optimize kinetic mass transport as highlighted for various fixed Ni-Cu ratios. We begin such an optimization process by repeating these kinds of thermodynamic calculations for many possible alloying additions.

Our screening for candidates uses tabulated thermodynamic interaction data of each element with Ni and Cu individually. As sketched in Fig. 2b, desired elements favor low chemical interactions with Ni (low heats of mixing generally comparable to the nearest competing phase formation energy), and have a rather positive heat of mixing with Cu to encourage the development of the miscibility gap. Figure 2c presents a visual summary of such screening of the binary interactions space from the ASM alloy database37; the color choice inside the box of each element X for the upper and lower part indicates general suitability based on these two interactions, Ni-X and Cu-X, respectively, with green denoting a positive screen, and red, negative.

This approach swiftly and effectively narrows the class of alloying elements to the primary candidates that will enable nano-phase separation in the Ni-Cu couple: Co, Fe, Mn, V, Cr, Nb, Ta, Mo, W, Ru, Rh, Ir. Among these, the last three are not likely economically viable, and therefore not further considered. Since the remaining set spans a full range of various crystallographic structures and melting temperatures, the next phase in our design process is to calculate quantities that speak directly to the sintering accelerating mechanism.

Densification via nano-phase separation is mainly dominated by the high-temperature phase of the thermal cycle, where interdiffusion accelerates densification. Specifically, the process is dominated by the net rate of bulk diffusion of the solvent into the solute-rich interparticle necks31–33,38. This property is quantifiable based on the relative importance of interdiffusion D~ vis-a-vis the diffusivity in the primary particle phase DNi. An alloy with a high value for D~/DNi is desirable, and we seek ratios on the scale of at least a factor of two for kinetic phenomena. We compute this quantity at the solvus temperature where dissolution would take place in each alloy (which is where rapid sintering begins)38, as described in the “Methods” section. The results for various ternary alloying additions to Ni-Cu are shown in Fig. 2d.

In order to facilitate neck formation, the lower-temperature phase of the sintering cycle is critical to nano-phase separation sintering. For this process, phase separation should lead to the preferential deposition of the secondary phase on the surface of an initially supersaturated powder. This requires the second phase to have a lower surface energy than the primary phase, so we aim to design the chemistry of both phases with the goal of maximizing the surface energy difference between them. Results are shown for prospective alloys in Fig. 2e.

Combining the trends in Fig. 2d, e (see also Supplementary Fig. 1), we see that thermodynamic considerations from the screening phase are clearly not sufficient to the design; many additives would worsen diffusion, and just a few alloys are clearly preferred in the full multidimensional design space. Interestingly, some elements that historically play substantial roles for Ni-base alloys are quite detrimental for our specific goals in powder metallurgy processing. For example, Al and W are typical ingredients for the thermodynamic and kinetic stabilization of the second phase in many superalloys. As revealed through this alloy design exercise, however, Al does not contribute to engineering of the miscibility gap, whereas W would delay the neck growth during interdiffusion and require more thermal input for powder consolidation. The top candidate based on the above analysis is Co, which will be our main focus for the remainder of the paper. In fact, Fig. 2a, d, and e also directly give guidance on the best compositions to optimize these design parameters for Co specifically. We identify a composition of Ni-40Co-20Cu as preferred option for further targeted work.

There is one final design consideration for rapid sintering, related to the burnout of organic processing agents or binders. Most powder routes involve such organics, and for alloys with accelerated consolidation, the sintering onset can compete with the critical temperature for organic removal39. In our process the use of a process control agent (ethanol), provides a source of organics that can burn out, and also carbothermally reduce native metal oxides on the powder particles, creating gas. If such gas is created at temperatures above the sintering onset, it can become entrapped and prevent consolidation. Designing the alloy to avoid gas entrapment therefore requires that we take a separate approach, for which we turn to an oxide reduction analysis based on the approach in ref. 40. Figure 2f shows the target sintering range for our process, lying between the onset of densification for nanocrystalline Ni at ~320 °C and the maximum temperature allowable for all-solid-state sintering, i.e., the nominal melting temperature of Ni. Elements (x-axis) that carbothermally reduce within this range, such as the primary constituents of our alloy, Ni, Co, and Cu, will offgas CO or CO2 in a temperature range where it can become trapped and prevent full sintering41,42.

The solution to this problem is to complement the designed alloy with a fourth component of high oxygen affinity to preferentially scavenge oxygen from those species, and prevent it from reacting with C during processing39,40. This strategy points towards the use of elements that form oxides stable up to high temperatures in Fig. 2f according to the classical Richardson-Ellingham diagram43,44; the candidate elements are Mg, Al, Mn, Y, Zr, Ti, and Hf. The complicating feature of adding these quaternary additions to the NiCoCu alloy is that they will interact with each of those species, potentially disturbing their balance for their intended function of evolving the phase evolution sequence we are targeting for sintering acceleration. The simplest way to address these new interactions and downselect many candidates is to refer to the thermodynamic screening we have performed in the other part of Fig. 2; all of these elements and their interactions with the Ni-rich and Cu-phases in our design are already articulated there. The element amongst those that is most compatible with the design criteria from Fig. 2a–e is Mn. Therefore, Mn is a preferred option for the present purposes; it will scavenge oxygen preferentially from the other elements in the alloy, and should do so at appropriate temperatures to prevent oxygen-based offgassing events according to Fig. 2f. The collection of oxygen into MnO particles is a process that occurs fully internally (internal oxidation) during sintering by diffusional precipitation; it produces a fine dispersion of particles which are non-interacting and should therefore not have any tendency to clump or aggregate.

Typical oxygen contents of powder alloys can be as high as a few thousand ppm, requiring a comparable addition of Mn to form MnO; to generously cover this range we add 5 at.% Mn; excess Mn incorporates into the phase evolution design per Fig. 2a-e. We thus arrive at a final quaternary designed composition of 37.5:37.5:20:5 (Ni:Co:Cu:Mn) under consideration of all design criteria for rapid sintering that should not entrap process gases.

The above design approach leads to a preferred alloy candidate, but it is important to emphasize that other candidates are certainly possible, and entire families of new alloys could be designed in this manner. As an illustration of the full generality of the approach, we also consider a second alloy that is compositionally independent of the first. Noting in Fig. 2 that Fe also emerges alongside Co as a viable phase-separation-facilitating addition, we identify an alloy of Ni:Fe:Cu:Mn (37.5:37.5:20:5) as a second option from our design approach. A full sintering characterization for that alloy is also provided here (more details in the Supplementary Information). Although we focus on the Co-bearing alloy in what follows as the top candidate, the Fe-bearing alloy provides a clear parallel throughout.

Experimental validation of the accelerated sintering design

Pressureless sintering was conducted on cold-compacted ball milled powders (see Supplementary Figs. 2-5 and Supplementary Note 1), typically by heating at constant rate up to 1200 °C. As shown in Fig. 3 (black curve), noticeable acceleration of sintering takes place above 600 °C in the designed alloy, reaching full density (>97%) over the ~2.5 h full cycle (confirmed by Supplementary Figs. 6–9). In contrast, solid-state sintering of microcrystalline-sized elemental Ni (red) and Inconel 718 (blue) powder showed significantly slower (and more typical) densification behavior, with final densities well below 80%.Fig. 3 Comparison of sintering behavior between designed alloy and other Ni powders.

Superior consolidation kinetics of the designed Ni alloy (black curve) as compared to conventional microcrystalline Ni (red) and commercial Inconel 718 powders (blue). The alloy reaches full density by only heating up to 1200 °C, including immediate cooling (density after the full cycle shown by the circular data marker), measured by dilatometry.

The key principle of the design approach is the phase evolution sequence laid out in Fig. 1, which in turn hinges upon the low-temperature phase separation, which assembles interparticle necks of the second phase. Validation for the formation of this decomposed structure is provided by the micrographs in Fig. 4a on a sample that was held at 600 °C for 4 h, typifying the low-temperature portion of the sintering cycle where sintering is setting on in Fig. 1 (see Supplementary Figs. 10–13 as well as Supplementary Note 2 for an analysis of the low temperature segment). The micrographs include a high-angle annular dark-field image of the microstructure region depicting two neighboring powders and several chemical maps based on energy dispersive spectroscopy (EDS). The images here contain a single interparticle region that is vertically oriented, as schematized by the frame in Fig. 1b. It is apparent in Fig. 4b that a secondary Cu-rich phase bridges a large area between the powders, as hoped for from the alloy design. In addition, Mn and O are found together, showing successful scavenging of oxygen; the nanoscale oxides seen in Fig. 4 are seen more broadly around the microstructure, fine and well-distributed.Fig. 4 Multicomponent alloy designed for sintering shows unique rapid consolidation behavior.

Nano-phase separation sintering is based on a sequence of phase evolutions in the powder compact. Phase decomposition must occur at low temperature into a dual-phase microstructure that facilitates enhanced sintering rates at the later stage of the sintering cycle, as confirmed for the present alloy by transmission electron microscopy (TEM) imaging in (a). The role of each alloy component agrees with the predicted behavior; most importantly, Cu forming a secondary phase between powder particles; vertically oriented neck of Cu in (a) is schematized by the dashed box in Fig. 1(b). The alloy chemistry that achieves both this interplay of phase evolution and oxygen management is the only combination that reaches full density (after cooling, circled point), as confirmed by the various control experiments using dilatometry in (b). Comparison of the thermal budget required to achieve a given relative density between the designed alloys and commercial Ni-based powders (c) shows that this phase evolution leads to substantially accelerated kinetics. Full markers denote structures which relative density was arrived at by solid-state sintering solely, whereas circled markers indicate the formation of a liquid phase during consolidation13,41,42,45–61.

The emergence of the dual-phase structure at an early stage of sintering (Fig. 4b) in combination with the enhanced densification rate upon chemical remixing at a later stage of the cycle (Fig. 2), establishes that the phase evolution kinetics are responsible for the rapid powder consolidation of the designer alloy. This observation is further buttressed by a set of control experiments in Fig. 4b that explore the effect of chemical variations on the sintering behavior. All samples in those control experiments were prepared from the same initial elemental powders and under the same processing conditions, which rule out any effects concerning the driving force for sintering or length scales for mass transport; the grain sizes, particle sizes, green densities, etc., are all similar amongst these experiments.

The sintering behavior of all other powders in Fig. 4b demonstrates that rapid sintering is only achieved when all components are present, and the powders follow a sequence of phase evolutions. More specifically, several of these control experiments reflect the entrapment of volatilizing gas (nanocrystalline Ni, red curve, Ni-20Cu alloy powder, orange, and Ni-Co-Cu, brown), as apparent from the loss in relative density as these sample experience ‘foaming’ and an increase in volume40. On the other hand, merely addressing gas entrapment is not sufficient for densification: nanocrystalline Ni powder with Mn as getter mitigates gas entrapment and foaming, but exhibits a lack of densification (blue), and even the addition of 20 Cu (violet) fails to accelerate sintering because the miscibility gap is at a too low temperature to allow for phase decomposition. A simple blend of powders of the same composition as the final alloy fails to accelerate sintering because it does not promote the sequence of phase evolutions that underlie the accelerated sintering process (teal). Sintering to full density is only achieved when all the alloy design principles are at play, suppressing gas entrapment and evolving the phases to accelerate neck growth and densification sequentially (black). Further confirmation is provided by Supplementary Figs. 10-16, including a full set of sintering data for the second, Fe-bearing alloy composition that validates the generality of our approach.

Alloys with chemistry conforming to these design concepts exhibit a superior densification behavior that can be quantified by the thermal budget λ, which is a measure of time-at-temperature given by1 λ=∫1Texp−QRTdt

where Q is the apparent activation energy for densification, R is the gas constant, and T is a function of time t over the course of the sintering cycle. The thermal budget incorporates all information of the heating history (temperature and time) that a powder passes through (see Supplementary Note 3). This metric permits quantitative comparisons of disparate materials and thermal cycles, as shown in Fig. 4c, where we see that the thermal budget for the designed Ni alloys (black and maroon) is reduced by at least two orders of magnitude as compared to their commercial Ni-based powder alloy counterparts, as well as elemental Ni and Cu powders13,41,42,45–61. What is more, for most of these commercial Ni-base alloys, densification is too slow in the solid state (solid data points), so some degree of melting is required for densification; the open data points show that high density typically requires melting in those alloys. In practical terms, the thermal budget advantage in Fig. 4c corresponds to a reduction of the production time by factor of 16 (at a constant sintering temperature of 1200 °C), or a reduction of the required process temperature by 276 °C (at constant process duration). Importantly, these process advantages have all been established with a very lean design approach: computation augmented by small-scale (few-gram) targeted experiments allow the design of new alloy families without the cost and duration needed to achieve large scales; such scale-up can thus be addressed separately in future work, with reduced scientific risk.

Properties of the designer alloy

Engineering of the dual-phase field is not only beneficial to the sintering kinetics of the compact, it also allows for heat treatments of the dense samples to control the phase structure and mechanical properties post-sintering; the alloy is precipitation hardenable. Figure 5 summarizes the effects of an isothermal heat treatment (starting from a solutionized condition) on the sintered Ni alloy. The Vickers hardness of the as-sintered and heat-treated samples are shown for different aging times in Fig. 5a. Whereas the as-sintered, solutionized sample exhibits a Vickers hardness of 240 Hv, the hardness value evolves in a pattern typical for precipitation treatments during isothermal annealing (at 670 °C) achieving a maximum of ~370 Hv after 10 h. Compression tests performed on bulk specimens aged to peak hardness show excellent strength and malleability, as depicted in Fig. 5b. The full true stress-strain curve is shown in Supplementary Fig. 17.Fig. 5 Heat treatment of sintered microstructure and mechanical properties.

The presented designed Ni alloy is also heat treatable to affect the microstructure and mechanical properties post-sintering. The Vickers hardness of the alloy reaches a maximum after aging at 670 °C for 10 h, tested by at least 7 indents for each condition (a), and a uniaxial yield strength of 650 MPa with considerable strain hardening and malleability during bulk mechanical compression tests (b). c The aged microstructure (high-angle annular dark-field, HAADF, micrograph) exhibits a dual-phase microstructure, with secondary Cu-rich lamellae at a thickness of about 13 nm embedded in a Ni-Co rich matrix. d Comparison of microhardness as function of the thermal budget for the designed Ni alloy and commercial Ni-base powders (full markers), focusing only on structures that were completely sintered in the solid-state. The empty markers additionally project to an equivalent strength13,57,60. The strength for Inconel 625 should be considered as upper limit for the solid-state, as it involves the formation of a liquid phase during processing. The error bars in (a) denote standard deviation.

These excellent mechanical properties are a product of the as-heat-treated microstructure, which is characterized via transmission electron microscopy in Fig. 5c. The microstructure involves a crystallographically aligned, ~13 ( ± 2) nm lamellar structure of Cu-rich precipitates that are embedded in a Ni- and Co-rich matrix (see Supplementary Fig. 18). The structure is reminiscent of spinodal decomposition, and indeed, according to computational thermodynamic analysis this composition does lie within the spinode. This structure is nearly ideal for classical precipitation strengthening, being in the vicinity of the expected dislocation cutting-to-bowing transition size for nickel-based alloys62,63. And, as seen in Fig. 5d, the resulting hardness and strength values (black) stand out amongst other commercial Ni-based alloys that were processed via a solid-state powder route, even as the alloy processing is substantially more rapid and efficient13,57,60. The present approach therefore differentiates from other more conventional strategies to accelerate powder consolidation through the use of sintering aids. These approaches, e.g., liquid phase64,65 or activated solid-state sintering66,67 focus exclusively on a faster consolidation rate at the expense of post-consolidation properties. As a result, they often involve a loss of control over the microstructure evolution during processing, such as coarsening of the grain size by several orders of magnitude32, and generally cannot sustain post-consolidation treatments or usage at elevated temperatures. The current approach, by contrast, foreshadows a future of far more sustainable, designer alloys that are tuned to both engineering property targets as well as advanced processability in modern paradigms such as solid-state additive manufacturing.

Methods

Computational thermodynamics

Thermodynamic calculations were performed in ThermoCalc version 2022a, using the following commercial databases for Fig. 2 (and Supplementary Fig. 1): TCNI11, TCCU3, TCNOBL1, and TCHEA3. For the main database TCNI11, the following phases were considered in the energy minimization: BCC_B2, BCC_B2#2, BCT_DO22, C14_LAVES, CBCC_A12, CUB_A13, DIS_MU, DIS_SIG, FCC_L10, FCC_L12, FCC_L12#2, HCP_A3, LIQUID, MU_PHASE, NI3TA_DoA, NI3TI_DO24 and SIGMA.

The design for alloying additions to Ni-Cu was performed on two main parameters, namely effective interdiffusion and surface wettability. The interdiffusion D~=xCuDNi+xNiDCu at the transus temperature was calculated by assuming (constant) phase fractions in the miscibility gap determined at 320 °C (to match with the onset of phase separation in experiments generally, Supplementary Fig. 12). The solidus temperature was calculated for both phases based on their instantaneous composition at the onset of phase decomposition. In some cases, a third phase was calculated to be present at 320 °C, which was assumed to not participate in the sintering kinetics and neglected. Therefore, phase fraction and chemical composition of both Ni- and Cu-rich phase were accordingly rescaled and recalculated, respectively. The activation energy for diffusion in the complex two-phase structure was approximated using D=D0exp(−18TS/T) with D0=10−5m2/s, and Ts the solidus temperature of each phase68. The diffusivity of each phase was quantified specifically at the transus temperature (which is where sintering acceleration occurs in nanophase separation sintering)38. The surface energy of each phase was calculated based on the alloy melting temperature at each (phase) composition69.

Powder processing

Powder particles with sizes of 3-7 µm for Ni (Thomas Scientific, purity 99.9%), 800 µm for Co (Fisher Scientific, purity 99.998%), 149 µm for Cu (Fisher Scientific, purity 99.999 %) and 44 µm for Mn (Fisher Scientific, purity 99.95%) were used in this study. All elemental and alloy powders were produced by mechanical alloying in a SPEX 8000 high-energy mill using a hardened steel vial and grinding media with a ball-to-powder ratio of 10:1 and 0.3 ml/g ethanol (C2OH6) as a process control agent. All powders (nanocrystalline Ni: nc-Ni (in Fig. 4b); Ni-5 at.%Mn: nc-Ni + Mn; Ni-40 Co-20 Cu: NiCo-20 at.% Cu; Ni-20 Cu-5 Mn: Ni-20 at.% Cu + Mn; Ni-20 Cu: Ni-20 at.% Cu) were milled for 3.5 h, if not reported otherwise (Supplementary Figs. 4 and 5). To mitigate impurity contamination from interactions of the powder with the atmosphere, high-energy ball-milling was performed in a glovebox under high-purity Ar. After milling, all powders had a typical particle size of 3-4 µm as assessed from powder spread on tape and imaged via scanning electron microscopy (SEM). The chemical composition of each powder was confirmed using EDS to also verify minimal pick-up of Fe contamination (typically ~ 0.5 at.%) from wear with the grinding media and vial during the milling process. The Ni-37.5 Co-20 Cu-5 Mn alloy was also subjected to shorter milling times (1.5, 2, and 3 h) to track the progress of mechanical alloying and the evolution of the grain size (Supplementary Figs. 2–5). More details on powder processing and characteristics are provided in Supplementary Note 1.

Specimens for the sintering experiments were cold-compacted into cylinders (6 × 1.5 mm for height and diameter, respectively) via uniaxial compression on a hydraulic press (model YLJ-15L from MTI Corporation) at a pressure of ~400–450 MPa. The initial relative density was calculated from the initial density (based on sample mass and cylinder volume) relative to the theoretical density (nominal weighted density of the elemental constituents), typically around ~60%. The sintering behavior and the change in relative density were characterized using dilatometry70. The densification of the powder was calculated by converting the raw change in length of the sample as a function of the temperature into a relative density assuming an isotropic shape change, as confirmed by caliper measurements after sintering. The change in raw length was analyzed using thermomechanical analysis (TMA, Hyperion F3 from Netzsch). Each green compact was heated at a constant rate (3, 5, 10, and 15 °C/min) from room temperature to 1200 °C followed by cooling at a rate of 40 °C/min in a forming gas of Ar containing 4% H2, under negligible contact force of 100 mN. The sintering curves at all heating rates are given in Supplementary Figs. 6 and 7. A more detailed discussion on the analysis of the sintering kinetics (see also Supplementary Figs. 10-15) is provided in Supplementary Note 2.

The sample for the microstructure in Fig. 4a was cold-compacted under the same conditions as described before, heated at a constant rate of 10 °C/min up to 600 °C, followed by isothermal annealing at 600 °C for 4 h, and then cooled at 10 °C/min back to room temperature. The sample for the microstructure in Fig. 5c was cold-compacted under the same conditions as described before, sintered up to 1200 °C, isothermally held for 2 h and then aged at 670 °C for 10 h, before cooling to room temperature.

Microstructure characterization

As-sintered samples (Supplementary Figs. 8 and 9) were prepared by standard mechanical grinding and polishing routine using a final suspension of 50 nm sized diamonds. The microstructure was imaged on a SEM in secondary electron mode. Specimens for scanning transmission electron microscopy (STEM) analyzes (Figs. 4 and 5) were prepared by a Raith VELION focused ion beam (FIB)-SEM system with an Au+ source. The initial FIB milling of the lamella was done with a 35 kV Au+ beam and the final polishing process was accomplished using 5 kV Au + . STEM and EDS were performed in a probe-corrected Thermo Fisher Themis Z, operated at 200 kV and equipped with Super-X EDS detectors. A convergence angle of 19 mrad and probe currents of 100 pA and 300 pA were used for imaging and EDS mapping, respectively. Annular dark field STEM images were collected using a collection angle of 36–200 mrad.

Mechanical testing

As-milled powders for the bulk specimen were compacted in a rectangular die set (MSE Supplies LLC) into samples with dimensions of 6 × 3 × 3 mm3 (for height, width and depth, respectively) and similar initial relative densities as before. The powder compacts were sintered passing through the same cycle as before, subsequently solid-solution homogenized at 1200 °C for ~2 h followed by an isothermal hold at 670 °C for various times between 1 and 20 h (to adjust the precipitated second phase structure) before cooling to room temperature. The aging temperature was chosen in the light of the need for reasonably short (10 h) bulk phase evolution kinetics (favors higher temperature), while maintaining a high phase fraction of precipitates (favors lower temperature). This optimization is typical in precipitation treatments to maximize the mechanical properties. The miscibility gap of the system was engineered by the alloy chemistry to be above 800 °C, and the aging temperature of 670 °C was chosen as a compromise to both aspects described above after empirical testing at a variety of heat treatment temperatures. Samples for mechanical testing were subjected to a standard mechanical grinding and polishing routine down to a suspension with 50 nm sized diamonds. Vickers hardness testing was performed on a Hardness testing system by LECO model AMH55 at a load of 200 gf and a time of 15 s. Reported hardness data is based on a minimum of 7 indentations for each condition with a spacing of more than 10 as compared to the indentation size to rule out any interference. Bulk compression testing was performed on an Instron 5984 Mechanical testing machine at a constant strain rate of 10−4/s with direct strain measurement by a linear voltage-displacement transducer. Hardness can be compared with uniaxial strength by assuming that indentation causes a characteristic 15% strain, and a Tabor factor of 3.

Supplementary information

Supplementary Information

Peer Review File

Source data

Source Data

Supplementary information

The online version contains supplementary material available at 10.1038/s41467-024-52261-1.

Acknowledgements

This work was supported by NASA Marshall Space Flight Center under grant 80MSFC19C0050 (Y.N and C.A.S) and made use of the MRSEC Shared Experimental Facilities at MIT, supported by the National Science Foundation under award number DMR-1419807 (Y.N. and C.A.S.). The authors would like to thank Dr. Yang Yu and Dr. Aubrey Penn for help with the FIB sample preparation and STEM measurements. This work was performed in part on the Raith VELION FIB-SEM in the MIT.nano Characterization Facilities.

Author contributions

Y.N. and C.A.S. proposed the research idea. Y.N. performed all experiments and conducted the computational design. C.A.S. provided overall guidance. Y.N. and C.A.S. discussed the results and contributed to writing and reviewing the manuscript.

Peer review

Peer review information

Nature Communications thanks the anonymous reviewer(s) for their contribution to the peer review of this work. A peer review file is available.

Data availability

The datasets used for the figures are available at 10.6084/m9.figshare.26548279. The datasets generated during and/or analyzed during the current study are available from the corresponding author on request. Source data are provided in this paper.

Competing interests

Multiple patents have been filed by MIT, including methods and data presented here, including PCT/US2023/062662 (Inventors: Yannick Naunheim and Christopher A. Schuh; publication date: Sept 28, 2023; state of application: in review) and 63/310,444 (Inventors: Yannick Naunheim and Christopher A. Schuh; state of application: pending). Those patents include the methodology of Figs. 1, 2, and 4, and are licensed by Foundation Alloy Technology Explorations, Inc., of which Christopher A. Schuh is a founder.

Publisher’s note Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.
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