# A New Family of Ternary Intermetallic Compounds with Dualistic Atomic Ordering – The ZIP Phases

Matheus A Tunes, Sean M Drewry, Franziska Schmidt, James A Valdez, Matthew M Schneider, Caitlin A Kohnert, Tarik A Saleh, Saryu Fensin, Stuart A Maloy, Cláudio G Schön, et al.

► To cite this version:

Matheus A Tunes, Sean M Drewry, Franziska Schmidt, James A Valdez, Matthew M Schneider, et al. A New Family of Ternary Intermetallic Compounds with Dualistic Atomic Ordering – The ZIP Phases. Advanced Materials, In press, 10.1002/adma.202308168. [https://doi.org/10.1002/adma.202308660](https://doi.org/10.1002/adma.202308660)

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# A New Family of Ternary Intermetallic Compounds with Dualistic Atomic Ordering – The ZIP Phases

Matheus A. Tunes,* Sean M. Drewry, Franziska Schmidt, James A. Valdez, Matthew M. Schneider, Caitlin A. Kohnert, Tarik A. Saleh, Saryu Fensin, Stuart A. Maloy, Cláudio G. Schön, Sylvain Dubois, Omri Tabo, Anna Eyal, Amit Keren, Asaf Pesach, Ganesh K. Nayak, Stavros-Richard G. Christopoulos, Marco Molinari, Marcus Hans, Nick Goossens, Shuigen Huang, Jochen M. Schneider, Per O. Å. Persson, Jozef Vleugels, and Konstantina Lambrinou*

First, this work is dedicated to H. Kudielka and H. Rohde of the Max Planck Institute, Germany, and W. Jeitschko, H. Nowotny, and F. Benesovsky of the University of Vienna, Austria, who made important early strides in the synthesis of ternary early transition metal carbides and nitrides of the $M_2AX$ type in the 1960s. A second dedication goes to M.W. Barsoum of Drexel University, USA, whose group introduced the MAX phases, associating their unique properties with their nanolaminated hexagonal crystal structure in the 1990s/2000s; the Drexel group also discovered the 2D derivatives of the MAX phases, known as MXenes, in 2011. A third dedication goes to E.I. Gladyshevskii, Yu. B. Kuz’ma and P.I. Kripyakevich of the Ivan Franko National University of Lviv, Ukraine, for their pioneering work on ternary intermetallic compounds of the $Nb_3Ni_2Si$ type in the 1960s.

A new family of nanostructured ternary intermetallic compounds — named the ZIP phases — is introduced in this work. The ZIP phases exhibit dualistic atomic ordering, i.e., they form two structural variants: one with the fcc diamond cubic structure (space group *Fd\bar{3}m*) and one with the hexagonal structure (space group *P6\bar{3}/mmc*). They are also characterized by metallic behavior, ionic bonding, and atomic zigzagging. Powder metallurgical routes involving pressure-assisted densification are adopted to demonstrate ZIP phase synthesis in the Nb-Si-Ni, Nb-Si-Co, Ta-Si-Ni, V-Si-Ni, and Nb-Si-Fe ternary systems. Crucially, reactive hot pressing is capable of producing high-purity ZIP phase materials after the judicious, elemental system-specific optimization of the processing route. Synthesis of phase-pure materials — demonstrated in the Nb-Si-Ni ternary system by the synthesis of quasi phase-pure $Nb_3SiNi_2$ and $Ni_3SiNb_2$ ZIP phase-based materials — is a steppingstone to the prospective exploitation of the ZIP phases. Characterization of $Nb_3SiNi_2$ and $Ni_3SiNb_2$ involves crystal structure determination, spatially resolved chemical analysis, and determination of select thermal, electrical, magnetic, mechanical, and physical properties. Density functional theory is used to assess the stability of $Nb_3SiNi_2$ & $Ni_3SiNb_2$ and derivative binary compounds at different temperatures, also exploring the exfoliation of these two ZIP phases along specific surfaces to produce 2D derivatives.

M. A. Tunes
Department Metallurgy
Chair of Nonferrous Metallurgy
Montanuniversität Leoben
Franz-Josef-Strasse 18, Leoben 8700, Austria
E-mail: matheus.tunes@unileoben.ac.at

The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/adma.202308168

© 2025 The Author(s). Advanced Materials published by Wiley-VCH GmbH. This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original work is properly cited.

DOI: 10.1002/adma.202308168

K. Lambrinou
Department of Engineering
School of Computing and Engineering
University of Huddersfield
Queensgate, Huddersfield HD1 3DH, UK
E-mail: k.lambrinou@hud.ac.uk

S.-R. G. Christopoulos
Department of Computer Science
School of Computing and Engineering
University of Huddersfield
Queensgate, Huddersfield HD13DH, UK
# 1. Introduction

A revolution in materials science was undoubtedly achieved through the microstructural manipulation of materials on the nanoscale.[1,2] When the dimensions of the basic building blocks of an engineering material are confined in the 1–100 nm range, the material can be defined as nanostructured.[3] The nanoscale manipulation of matter via the (application-driven) tailoring of structural “modules”, such as ordered atomic clusters, precipitates and grains, engineered grain boundaries, complex molecules, and atomically thin layers is an effective approach

to achieve a set of exceptional properties in advanced materials, well beyond what is possible by conventional processing routes that aim at modifications on the microscale, mesoscale, and/or macroscale.[1–4] Transferring the superior properties of nanostructured materials to the macroscale relies on judicious multiscale materials engineering and remains one of the key challenges of contemporary materials science.[4]

Gleiter, the scientist who coined the term “nanostructured materials” in 1992, predicted that this novel research field would be greatly impacted by the emergence of two classes of nanostructured materials, i.e., ceramics and intermetallic phases.[2] Gleiter’s prediction came partly true in 1996, when Barsoum and his group made significant strides in the development and characterization of a new class of nanolaminated ceramics with the $M_{n+1}AX_n$ general stoichiometry, where M is an early transition metal, A is an element mainly from groups 13–15 in the periodic table, X is C or N, and n is an integer (commonly, 1, 2, or 3); these ceramics became known as the MAX phases.[5,6] The 1996 breakthrough by Barsoum and El-Raghy[5] consisted in the key synthesis of quasi phase-pure $Ti_3SiC_2$ via reactive hot pressing; notably, the existence of the $Ti_3SiC_2$ compound, which is regarded as the H-phase of the ternary Ti-Si-C system,[7] had first been reported in 1967 by Jeitschko and Nowotny.[8] Starting by the successful experimental synthesis and characterization of $Ti_3SiC_2$, i.e., the 312 MAX phase of the Ti-Si-C system,[7] Barsoum realized in the early 2000s that $Ti_3SiC_2$ and other (similar in nature) nanolaminated ternary carbides/nitrides form an entire class of hexagonal solids (space group $P6_3/mmc$) obeying the $M_{n+1}AX_n$ stoichiometric rule and in possession of exceptional physicochemical and mechanical properties.[6] To date, about 155 ternary carbides and nitrides of early transition metals, as well as numerous higher order (chemically complex) solid solutions with the $M_{n+1}AX_n$ overarching stoichiometry and precisely tailored properties have been synthesized.[9,10] The exceptional properties of the MAX phases are attributed to their nanolaminated hexagonal crystal structure, where n $M_6X$ (ceramic-like) octahedra are interleaved with a single A (metallic-like) atomic layer, and the M-atoms form a trigonal prism with the A-element in its center.[10] Therefore, the MAX phase structure can be perceived as a combination of two building blocks, i.e., $M_6X$ octahedra with the rock-salt (NaCl-type) structure and $M_6A$-trigonal prisms with the hexagonal closed packed (hcp) structure.[10] Since their inception, the MAX phases[11,12] and their 2D derivatives known as MXenes[13–18] have been proposed for promising niche applications in diverse fields, including energy storage, fuel cells, biomedical applications, photonics/phononics,[19–31] as well as advanced fusion and fission nuclear systems.[32–43]

This work introduces a new family of nanostructured ternary intermetallic compounds (IMCs), herein aptly named “zigzag

M. Molinari
Department of Physical and Life Sciences
School of Applied Sciences
University of Huddersfield
Queensgate, Huddersfield HD1 3DH, UK
M. A. Tunes, S. M. Drewry, F. Schmidt, J. A. Valdez, M. M. Schneider, C. A. Kohnert, T. A. Saleh, S. Fensin
Materials Science and Technology Division
Los Alamos National Laboratory (LANL)
Bikini Atoll Blvd, Los Alamos, NM 87544, USA

S. A. Maloy
Pacific Northwest National Laboratory (PNNL)
902 Battelle Blvd, Richland, WA 99352, USA

C. G. Schön
Department of Metallurgical and Materials Engineering
Universidade de São Paulo Escola Politécnica
Av. Prof. Mello Moraes 2463, São Paulo CEP 05508-030, Brazil

S. Dubois
Institut PPRIME
Département de Physique et Mécanique des Matériaux
CNRS

Université de Poitiers
ENSMA UPR 3346, SP2MI, Téléport 2, Boulevard Marie et Pierre Curie,
Futuroscope, Chasseneuil Cedex, 30179, Poitiers, Nouvelle-Aquitaine
86962, France

O. Tabo, A. Eyal, A. Keren
Faculty of Physics
Technion – Israel Institute of Technology
Haifa 32000, Israel

A. Pesach
Department of Physics
Nuclear Research Centre – Negev
P.O. Box 9001, Beer Sheva 84190, Israel

G. K. Nayak, M. Hans, J. M. Schneider
Materials Chemistry
RWTH Aachen University
Kopernikusstr. 10, D-52074 Aachen, Germany

P. O. Å. Persson
Thin Film Physics
Department of Physics
Chemistry and Biology
Linköping University
Linköping SE-581 83, Sweden

N. Goossens, S. Huang, J. Vleugels
Department of Materials Engineering

KU Leuven
Kasteelpark Arenberg 44, Leuven 3001, Belgium
K. Lambrinou
Italian Institute of Technology (IIT)
Via Morego 30, Genova 16163, Italy

S. M. Drewry
Department of Materials Science and Engineering
University of Tennessee
414 Ferris Hall, 1508 Middle Drive, Knoxville, TN 37996-2100, USA

F. Schmidt
Department of Nuclear Engineering
University of California Berkeley
4153 Etcheverry Hall, 1730, Berkeley, CA 94720-1730, USA

N. Goossens
Laboratory of High-Performance Ceramics
Swiss Federal Laboratories for Materials Science Technology
EMPA
Überlandstrasse 129, Dübendorf CH-8600, Switzerland
intermetallic compounds", or simply the ZIP phases, due to their
distinctive "zigzag" atomic arrangement. Each member of the
ZIP phase family has two structural variants, i.e., one with the fcc
diamond cubic structure (space group Fd3m) and one with the
hexagonal structure (space group P63/mmc). Transitioning be-
tween ZIP phase variants requires elemental redistribution in the
"312" stoichiometry describing these materials. For example, the
two ZIP phases in the Nb-Si-Ni system are the fcc Nb3SiNi2 IMC,
also regarded as the H-phase of the ternary Nb-Si-Ni system, and
the hexagonal Ni3SiNb2 IMC.[44-47] The hexagonal variants of the
ZIP phases exhibit striking similarities with the MAX phases, in-
cluding crystal symmetry (i.e., they share space group P63/mmc),
crystal structure nanolamination, and the 312 MAX phase stoi-
chiometry; however, the extent of this resemblance, especially in
terms of potential applications, cannot be properly assessed un-
til materials with high phase purity have been synthesized. This
work reports, for the first time, the experimental synthesis of
quasi phase-pure Nb3SiNi2 and Ni3SiNb2 via reactive hot press-
ing, enabling the detailed characterization of ZIP phases in the
Nb-Si-Ni ternary system. ZIP phases have also been synthesized
in the Nb-Si-Co, Ta-Si-Ni, V-Si-Ni, and Nb-Si-Fe ternary systems
by means of reactive hot pressing and spark plasma sintering.
Moreover, basic aspects of the potential formation of 2D deriva-
tives of the ZIP phases – similar to the 2D derivatives of the MAX
phases broadly known as MXenes – have been explored using
density functional theory in conjunction with the temperature-
dependent effective potential method. The introduction of the
ZIP phase family of ternary IMCs in this work validates Gleiter’s
1992 prediction regarding the emergence and diverse potential
benefits of “nanostructured intermetallics”.[1,2]

**2. Results and Discussion**

**2.1. The Gateway to the Discovery of the ZIP Phases – The Nb-Si-Ni System**

2.1.1. Synthesis of Nb₃SiNi₂ by Arc Melting

Arc melting was initially selected to experimentally synthesize
bulk materials with the Nb₃SiNi₂ stoichiometry. The microstruc-
tural inspection of the as-cast sample (Figure 1A) by means
of scanning electron microscopy (SEM) revealed the presence
of four distinct phases, a finding supported by both energy-
dispersive X-ray spectroscopy (EDS) and X-ray diffraction (XRD).
Phase identification by XRD (Figure 1C) revealed the co-existence
of the Nb₃SiNi₂ H-phase (first candidate ZIP phase; fcc variant),
the Nb₇Ni₆ µ-phase, the Ni₃SiNb₂ ternary Laves (L) phase (first
candidate ZIP phase; hexagonal variant), and the Nb₄NiSi T-
phase. These four phases were peak-indexed in accordance with
data found in the ICSD database.[48] whilst they were also iden-
tified as candidate phases during the experimental investigation
of phase equilibria in the ternary Nb-Si-Ni system of refractory
alloys.[47] The nominal stoichiometries and measured EDS ele-
mental compositions of the four phases in the as-cast sample
are provided in Table S1 (Supporting Information), which shows
that all phases formed as off-stoichiometric solid solutions dur-
ing melt solidification rather than as strictly stoichiometric com-
pounds. The formation of solid solutions in the Nb-Si-Ni system
has been previously reported by dos Santos et al.,[47] who provided

strong evidence of solid solution formation associated with the
dissolution of significant amounts of silicon (Si) in the Nb₇Ni₆
µ-phase,[47] as also indicated in Table S1 (Supporting Informa-
tion). Even though arc melting aimed at the production of sam-
ples primarily made of the Nb₃SiNi₂ IMC, the as-cast alloy was
a mixture of four phases, i.e., the maximum number of phases
in equilibrium for a ternary alloy at constant pressure according
to Gibbs' phase rule.[49] In particular, the formation of the Nb-
rich Nb₄NiSi T-phase suggests the loss of substantial amounts
of nickel (Ni) during arc melting, as Ni is volatile in vacuum
above its melting point (1728 K). Considering the CALPHAD
phase equilibria calculated in this work (Figure 2C) in conjunc-
tion with the microstructure of the as-cast alloy (Figure 1A), it
is reasonable to assume that (rapid) melt solidification formed
solid solutions of three principal phases, i.e., the Nb₃SiNi₂ H-
phase, the Nb₇Ni₆ µ-phase, and the Ni₃SiNb₂ L-phase, whereas
the Ni-impoverished residual melt yielded the Nb₄NiSi T-phase
(see the 1423 K isotherm in Figure 2C). This hypothesis is sup-
ported by the small fraction of the Nb₄NiSi T-phase in the as-cast
alloy (Figure 1A).

A homogenization annealing of the as-cast sample was per-
formed at 1421 K to investigate whether the four-phase assem-
bly could change into an alloy with a higher Nb₃SiNi₂ IMC con-
tent. Annealing was also aimed at assessing the thermal stability
of the Nb₃SiNi₂ phase of interest. The SEM/EDS/XRD analysis
of the annealed alloy (Figure 1B,C) showed that it consisted of
the same four phases, with a slightly increased fraction of the
Nb₃SiNi₂ H-phase that appears stable at these conditions; how-
ever, the marginal increase in Nb₃SiNi₂ fraction showed that the
particular annealing treatment could not transform the as-cast
alloy into a material primarily made of Nb₃SiNi₂. This suggests
that solid-state elemental diffusion is not adequately fast at 1421
K to achieve alloy homogenization within 336 h, indicating the
need for prolonged annealing treatments at higher temperatures.
Moreover, the loss of Ni during arc melting – as confirmed by the
presence of the Nb₄NiSi T-phase – is bound to prevent the for-
mation of phase-pure Nb₃SiNi₂, irrespective of the temperature
and/or duration of the annealing treatment. Figure 1B shows
only slight modifications in grain morphology due to anneal-
ing; for example, L-Ni₃SiNb₂ and µ-Nb₇Ni₆ grains appear more
equiaxed than in the as-cast sample. The elemental compositions
of the four phases in the annealed sample were also determined
by EDS (see Table S1, Supporting Information). Additional SEM
micrographs of the arc-melted (as-cast & annealed) Nb-Si-Ni al-
loys are provided in Figure S1 (Supporting Information). The
unsuccessful attempt to synthesize phase-pure Nb₃SiNi₂ via arc
melting showed that this processing route is difficult to mas-
ter, mainly due to the poorly controlled losses of constituting el-
ements above their melting points, T<sub>m</sub> (i.e., T<sub>m</sub>(Ni) = 1728 K,
T<sub>m</sub>(Nb) = 2750 K, T<sub>m</sub>(Si) = 1683 K). Moreover, calculations of
the equilibrium phase mixtures in the Nb-Si-Ni ternary system
at various temperatures are needed to better understand melt so-
lidification, as prior studies examined only the 1073, 1323, and
1473 K isotherms.[44-47]

The results of this work contradict the early work of Glady-
shevskii et al.,[44] who reported that melting a feedstock of ele-
mental powders (Nb, Si, Ni) under an inert gas atmosphere (pre-
sumably, argon) produced quasi phase-pure Nb₃SiNi₂ with only
a small amount of a ternary Laves phase (possibly, the Ni₃SiNb₂
<figure><img src="image_3.png" /><figcaption>**Figure 1.** Characterization of arc-melted Nb-Si-Ni alloy samples. SEM/EDS analysis of the A) as-cast and B) annealed (1421 K, 336 h) samples shows four phases, i.e., H-Nb₃SiNi₂, L-Ni₃SiNb₂, µ-Nb₇Ni₆, and T-Nb₄SiNi. C) XRD analysis of the as-cast and annealed samples confirms the presence of the same four phases. The phase of interest, i.e., the Nb₃SiNi₂ ternary IMC (H-phase), is highlighted in both SEM and XRD data using distinctive colors.</figcaption></figure>

IMC also detected in this work). These authors used light optical microscopy to confirm the phase purity of their arc-melted alloys, and powder XRD to propose the existence of a group of ternary IMCs with the H-phase structure (space group *Fd\*3m*) and the R₃SiNi₂ general stoichiometry, with R being the early transition metals Mn, Cr, V, Nb, and Ta.[44] Apart from demonstrating a much better command of the arc-melting processing route, the 1964 pioneering work by Gladyshevskii et al.[44] paved the way for the study of the ZIP phases introduced in this work. Similarly, the 1967 work by Jeitschko et al.[7] opened the way for the discovery of the MAX phases family by Barsoum et al. almost 40 years later.[5,6]

A thorough thermodynamic assessment of the ternary Nb-Si-Ni system was only recently conducted – 50 years after Gladyshevskii's pioneering work – in a set of works published by dos Santos et al.[45–47] In an attempt to reproduce the results of Gladyshevskii et al.,[44] dos Santos et al.[46] synthesized twelve different Nb-Si-Ni alloy compositions via arc melting and subsequently homogenized the as-cast samples at 1073 K for up to 1000 h. Curiously, the Nb₃SiNi₂ phase was not detected in any of the twelve

homogenized samples characterized by means of SEM/EDS and XRD. The reason for this discrepancy between the results of Gladyshevskii et al.[44] and dos Santos et al.[46] is presently not understood. Based on their data, dos Santos et al. suggested a homogenization isothermal annealing at higher temperatures to explore the thermodynamic equilibria in this system.[46] In a more extensive follow-up study, dos Santos et al.[47] investigated 48 alloy compositions in the Nb-Si-Ni system, implementing two different heat treatment approaches: (a) 1323 K for 336 h, and (b) 1473 K for 504 h. Using SEM/EDS and XRD after homogenization annealing, Nb₃SiNi₂ was only observed in seven samples, but none of these were phase-pure Nb₃SiNi₂. All seven samples exhibited triphasic microstructures, each comprising two or three of the four principal phases also detected in this work, i.e., the ternary Laves, µ-, T-, and H-phases; small fractions of the NbNi₃ IMC or the unknown Z phase were also reported in three of the seven samples. The Z phase is one of three phases – tentatively designated X, Y, and Z – with unresolved crystal structures found by dos Santos et al.[47] after annealing certain Nb-Si-Ni alloy compositions at 1473 K for 504 h.
<figure><img src="image_4.png" /><figcaption>**Figure 2.** Characterization of the RHP Nb-Si-Ni alloy samples by means of A) XRD and B) EPMA/WDS shows the presence of mainly the two ZIP phases, i.e., H-Nb₃SiNi₂ and L-Ni₃SiNb₂, along with small amounts of oxidic inclusions (NbO₂, NiO). Microstructural homogeneity and phase purity (in terms of the total ZIP phase content) increased with the sintering temperature. C) CALPHAD calculations of phase equilibria in the ternary Nb-Si-Ni system on the 1423, 1523, 1623, and 1723 K isotherms.</figcaption></figure>

### 2.1.2. Synthesis of High Phase Purity Nb₃SiNi₂ by RHP

Careful consideration of the prior attempts by dos Santos et al.[46,47] and the findings of this work demonstrated that the synthesis of phase-pure Nb₃SiNi₂ via arc melting is challenging. Even though Gladyshevskii et al.[44] reported that the production of quasi phase-pure Nb₃SiNi₂ is possible via the melting of the corresponding mixture of elemental powders in an inert gas atmosphere, high-vacuum plasma arc melting methods have been empirically known to face challenges during the

synthesis of compounds with strict stoichiometries,[50] such as the ternary IMCs designated as the ZIP phases in this work, or the early transition metal ternary carbides/nitrides known as the MAX phases. As already mentioned, one important challenge is the evaporation of alloying elements in the vacuum chamber.[51] Given the strict stoichiometry of the ZIP phases and the 312 MAX phases, the processing route and the starting powder feedstock must always be carefully adjusted to produce materials with high phase purity. For example, Goossens et al.[52] explained only recently why using early transition metal hydride powders
produces MAX phase-based ceramics with higher phase puri-
ties than using elemental powders. These authors proved that
the use of metal hydride powders lowers the formation temper-
ature of MAX phase precursor phases (e.g., binary IMCs), in-
creases the formation temperature of competing phases (e.g.,
binary carbides), and suppresses the oxidation of easily oxidiz-
able powders (e.g., zirconium, Zr) by providing a reducing (H₂-
rich) atmosphere created by powder dehydrogenation, thus al-
lowing the use of finer powder feedstocks. Based on the above, it
was decided in this work to employ reactive hot pressing (RHP)
of carefully adjusted powder feedstocks consisting of hydride
(NbH₀.₈₉) and elemental powders (Ni and Si) to produce phase-
pure Nb₃SiNi₂. The synthesis efforts were successful at three dif-
ferent sintering temperatures (i.e., 1523, 1623, and 1723 K), pro-
ducing quasi phase-pure Nb₃SiNi₂ with a small fraction of the
ternary Ni₃SiNb₂ Laves phase, and traces of oxidic inclusions
(NbO₂ & NiO) that presumably stemmed from the limited oxida-
tion of the raw powders. Electron probe microanalyzer (EPMA)
images and wavelength-dispersive X-ray spectroscopy (WDS) el-
emental maps of all Nb₃SiNi₂-based materials are shown in
Figure 2B, together with their XRD patterns (Figure 2A). Addi-
tional EPMA/WDS data from all RHP quasi phase-pure Nb₃SiNi₂
materials are shown in Figure S2 (Supporting Information). The
XRD pattern of the pulverized RHP sample sintered at 1723
K was Rietveld-refined (Figure S3 and Table S2, Supporting
Information).

Figure 2C shows CALPHAD calculations of phase equilibria in the ternary Nb-Si-Ni system at 1423 K (i.e., close to the temperature of isothermal annealing, 1421 K, of the arc-melted alloy), 1523, 1623, and 1723 K. These calculations suggest that melt solidification during vacuum arc melting yielded a four-phase mixture comprising the L-Ni₃SiNb₂ line compound, the H-Nb₃SiNi₂ point compound (i.e., the targeted IMC), and the µ-Nb₇Ni₆ phase-field compound with enhanced Si solid solubility. Due to the anticipated loss of Ni during vacuum arc melting, a small fraction of the melt was Ni-poor and solidified as the T-Nb₄SiNi point compound, which appears only in the 1423 K isotherm (Figure 2C, top). On the other hand, all three alloys synthesized via RHP at 1523/1623/1723 K resulted in quasi binary phase mixtures made of the H-Nb₃SiNi₂ and L-Ni₃SiNb₂ IMCs. The NbO₂ and NiO oxide impurities present in the RHP alloys were not encountered in the arc-melted alloys due to the much higher processing temperatures associated with the complete melting of all raw elemental powders and their oxide films during arc melting synthesis.

## 2.2. Characterization of ZIP Phases in the Nb-Si-Ni System

### 2.2.1. *Crystal Structure Determination by TEM – The ZIP Phase Variants*

Even though the arc-melted sample had a small Nb₃SiNi₂ frac-
tion, focused ion beam (FIB) was used to lift out thin foils from
a large Nb₃SiNi₂ grain in the annealed sample (Figure S4A, Sup-
porting Information) to study its crystal structure by means of
scanning transmission electron microscopy (STEM) and selected

area diffraction (SAED). STEM/EDS analysis of this grain re-
vealed a composition close to the Nb₃SiNi₂ stoichiometry, with
limited deviations from the nominal Ni (5 at.% excess) and Si
(5 at.% deficiency) contents (Figure S4B, Supporting Informa-
tion). Atomically resolved BF-STEM images of the same grain
revealed a “zigzag” atomic arrangement, observed in low-index
crystallographic orientations (Figure S5A,B,C, Supporting Infor-
mation). Using published crystallographic data (CIF file ICSD-
2044323)[⁴⁴] and the SingleCrystal software package,[⁵³] a kine-
matic simulation of a selected area diffraction pattern (SAEDP)
in the low-index [112] zone axis (Z.A.) was compared with
the recorded SAEDP (Figure S5B,D, Supporting Information).
The excellent agreement between experimental and simulated
SAEDPs confirmed that the crystal structure of Nb₃SiNi₂ is di-
mond cubic (Fdbar, space group, SG, 227), as reported by Glady-
shevskii et al.[⁴⁴]

The successful RHP synthesis of quasi phase-pure Nb-Si-Ni
alloy samples allowed the study of FIB foils lifted out from
large grains of Nb₃SiNi₂ and Ni₃SiNb₂, i.e., the two main
phases (Figure 2A,B), by means of high-angle annular dark-field
(HAADF) STEM, using a double Cs-corrected FEI Titan³ 60–300.
Atomically resolved HAADF STEM images of both phases are
shown in **Figure 3**. The crystal structure of Nb₃SiNi₂ (Figure 3A)
was again identified as fcc (Fdbar, SG 227), whereas that of
Ni₃SiNb₂ (Figure 3B) was found to be hexagonal (P6₃/mmc, SG
194). Figure 3C shows the schematic representations of the same
Z.A. of both phases depicted in Figure 3A,B. The two variants of
the ternary Nb-Si-Ni IMCs, i.e., the fcc Nb₃SiNi₂ and the hexago-
nal Ni₃SiNb₂, are considered from this point onward the two vari-
ants of the Nb-Si-Ni ZIP phases. It will be further demonstrated
that the ZIP phase concept of ternary IMCs extends beyond the
Nb-Si-Ni system, thus constituting a new family of “nanostruc-
tured materials”, in alignment with the prediction of Gleiter.[2]

Version 3.5.8 of the VESTA software[54] was utilized to depict
different Z.A. in the Nb₃SiNi₂ and Ni₃SiNb₂ ZIP phase variants,
thus illustrating the complexity and size of their unit cells, as
well as the atomic “zigzagging” inferred by the herein introduced
term “ZIP phases”. **Figure 4** shows the [001], [110], and [111] Z.A.
of Nb₃SiNi₂ and the [0001] and [1120] Z.A. of Ni₃SiNb₂. It also de-
picts the [0001] and [1120] Z.A. of the well-known Ti₃SiC₂ MAX
phase to compare its crystal structure complexity with that of the
hexagonal Ni₃SiNb₂ ZIP phase variant; direct size comparison is
possible as both Ti₃SiC₂ and Ni₃SiNb₂ contain Si.

Figure 4 illustrates the remarkable differences in crystal struc-
ture complexity and size of unit cell for the fcc Nb₃SiNi₂ ZIP
phase (96 atoms per unit cell; Figure 4A), the hexagonal Ni₃SiNb₂
ZIP phase (12 atoms per unit cell; Figure 4B), and the hexagonal
Ti₃SiC₂ MAX phase (6 atoms per unit cell; Figure 4C). Unit cell
volume, theoretical density, and Wyckoff parameters for Ti₃SiC₂,
Nb₃SiNi₂, and Ni₃SiNb₂ are compared in Table S3 (Supporting
Information). All three nanostructured solids are characterized
by centrosymmetric Bravais lattices, i.e., the Nb₃SiNi₂ ZIP phase
is fcc (Fdbar, SG 227), whereas the Ni₃SiNb₂ ZIP phase and the
Ti₃SiC₂ MAX phase are both hexagonal (P6₃/mmc, SG 194); it
appears that these particular Bravais lattices favor the formation
of nanostructured solids (i.e., ternary ceramics & ternary IMCs)
with characteristic atomic zigzagging.
<figure><img src="image_5.png" /><figcaption>**Figure 3.** Atomically resolved HAADF STEM images of A) Nb<sub>3</sub>SiNi<sub>2</sub> (fcc, SG Fd&#773;3m) and B) Ni<sub>3</sub>SiNb<sub>2</sub> (hexagonal, SG P6<sub>3</sub>/mmc). A) Z.A. [110] (left) and [001] (right) of Nb<sub>3</sub>SiNi<sub>2</sub>. B) Z.A. [1&#773;20] of Ni<sub>3</sub>SiNb<sub>2</sub>. C) Schematics depicting the [110] and [001] Z.A. of Nb<sub>3</sub>SiNi<sub>2</sub> and the [1&#773;20] Z.A. of Ni<sub>3</sub>SiNb<sub>2</sub>; these schematics were made with the VESTA software.[54]</figcaption></figure>

### 2.2.2. Compositional Reconstruction of the Nb-Si-Ni ZIP Phase Variants by APT

The compositions of the Nb<sub>3</sub>SiNi<sub>2</sub> and Ni<sub>3</sub>SiNb<sub>2</sub> ZIP phase variants were spatially resolved by means of atom probe tomography (APT). Details of the site-specific FIB lift-outs can be found in Figure S6 (Supporting Information). Reconstructions of the Nb, Ni, Si, and O atomic positions, as well as the composition profiles of Nb<sub>3</sub>SiNi<sub>2</sub> and Ni<sub>3</sub>SiNb<sub>2</sub> are presented in Figure 5A,B. Apart from the principal Nb, Ni, and Si elements, O, C, Ta, and W impurities were also identified, and the mass spectra are shown in Figure S7 (Supporting Information) for Nb<sub>3</sub>SiNi<sub>2</sub> and Figure S8 (Supporting Information) for Ni<sub>3</sub>SiNb<sub>2</sub>. The sum of C, Ta, and W impurities was <1 at.% in both phases. The average composition of Nb<sub>3</sub>SiNi<sub>2</sub> was 45 ± 1 at.% Nb, 16 ± 1 at.% Si, 32 ± 1 at.% Ni and 7 ± 1 at.% O (Figure 5A). The O impurities were tentatively associated with Nb, as a significant fraction of molecular NbO ions is evident in the mass spectrum of Figure S7 (Supporting Information). Neglecting the O impurities, a Nb<sub>2.8</sub>SiNi<sub>2.0</sub> stoichiometry was obtained, in excellent agreement with the nominal composition of Nb<sub>3</sub>SiNi<sub>2</sub>. The average composition of Ni<sub>3</sub>SiNb<sub>2</sub> was 35±1 at.% Nb, 21 ± 1 at.% Si and 44 ± 1 at.% Ni, whilst the O content was <1 at.% (Figure 5B), corresponding to a stoichiometry of Ni<sub>2.1</sub>SiNb<sub>1.7</sub>.

Magnifications of the atomic positions in the Nb<sub>3</sub>SiNi<sub>2</sub> and Ni<sub>3</sub>SiNi<sub>2</sub> ZIP phases in a rectangular volume of 20 × 20 × 15 nm³ are shown in Figure 5C. These data provide evidence of atomic “layering” in Nb<sub>3</sub>SiNi<sub>2</sub>, whereas in the case of Ni<sub>3</sub>SiNb<sub>2</sub>, the experimental proof of nanolamination (see Figure 4B) is probably at the resolution limit of APT. The O content data shown in Figure 5A,B are supported by data collected from a third sample that contained both Ni<sub>3</sub>SiNb<sub>2</sub> and Nb<sub>3</sub>SiNi<sub>2</sub> grains separated by a slanted grain boundary, GB (Figure S9, Supporting Information). The APT data in Figure S9 (Supporting Information) show the presence of O in the Nb<sub>3</sub>SiNi<sub>2</sub> grain and its absence from the neighboring Ni<sub>3</sub>SiNb<sub>2</sub> grain. The better accommodation of O in Nb<sub>3</sub>SiNi<sub>2</sub> suggests that the O solid solubility in the fcc ZIP phase variant is higher than that in the Ni<sub>3</sub>SiNb<sub>2</sub> hexagonal ZIP phase variant. The O impurities measured by APT cannot, thus, be only attributed to the oxidation of Nb, as Nb is strongly present in both ZIP phase variants. The reason for the difference in O content might be sought in the size and “open spaces” available in the unit cells of the two ZIP phase variants (Figure 4). The hypothesis regarding O solubility has been tested with ZIP phases in other ternary systems explored in this work, i.e., Nb-Si-Co, Ta-Si-Ni, V-Si-Ni, and Nb-Si-Fe.
<figure><img src="image_6.png" /><figcaption>**Figure 4.** A) Schematics of three Z.A. of the fcc Nb₃SiNi₂ ZIP phase: [001], [110], and [111]. The confines of one unit cell (96 atoms) are outlined using thicker black lines (top row). Representation of the Nb₃SiNi₂ crystal structure resulting from the stacking of 2 × 2 × 2 unit cells along the [001], [110], and [111] Z.A. (bottom row). B) Schematics of two Z.A. of the hexagonal Ni₃SiNb₂ ZIP phase: [0001] and [1120] (left half of the image); the right half of the image shows perspectives of the same Z.A. after a minor counterclockwise rotation. The confines of the unit cell (12 atoms) are outlined by thicker black lines. C) Schematics of two Z.A. of the hexagonal Ti₃SiC₂ MAX phase: [0001] and [1120] (left half of the image) and 3D perspectives thereof after a minor counterclockwise rotation (right half of the image). The confines of one unit cell (6 atoms) are outlined using thicker black lines. In the 3D perspectives of the [1120] Z.A. of the Ni₃SiNb₂ ZIP phase and the Ti₃SiC₂ MAX phase, black arrows pinpoint the atomic layers, the removal of which is likely to yield 2D derivatives. Dashed orange lines and arrows are combined to highlight atomic zigzagging in both fcc Nb₃SiNi₂ (Z.A. [110]) and hexagonal Ni₃SiNb₂ (Z.A. [1120]) ZIP phases.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>**Figure 5.** Spatially resolved composition of the two variants of the Nb-Si-Ni ZIP phases, as determined by APT. Reconstructions of the Nb, Ni, Si, and O atomic positions and elemental profiles for Nb₃SiNi₂ A) and Ni₃SiNb₂ B). The elemental profiles stem from cylinders (Ø 10 nm) as shown in the maps of the O atomic positions. C) Atomic positions within a rectangle (20 × 20 × 15 nm³) for Nb₃SiNi₂ and Ni₃SiNb₂. The mass spectra of Nb₃SiNi₂ and Ni₃SiNb₂ are found in Figures S7 and S8 (Supporting Information), respectively.</figcaption></figure>

### 2.2.3. Thermal, Electrical, and Magnetic Properties

Thermal, electrical, and magnetic properties of the Nb₃SiNi₂ and Ni₃SiNb₂ ZIP phases have been determined according to the experimental methodologies described in Section “Thermal, Electrical, and Magnetic Properties”. Figure 6A,B shows the variation of the thermal conductivity, $\kappa$, of Nb₃SiNi₂ (Figure 6A) and Ni₃SiNb₂ (Figure 6B) as a function of temperature. The electronic contribution to the total thermal conductivity was estimated based on the Wiedemann-Franz law to be approximately

50% of the total thermal conductivity of Nb₃SiNi₂ and Ni₃SiNb₂ at room temperature (RT); this is much higher than the electronic contribution to the total thermal conductivities of the MAX phases, which has been calculated to vary in the 10–20% range.

Figure 6C shows the variation of the electrical resistivity, $\rho$, of Nb₃SiNi₂ and Ni₃SiNb₂ as a function of temperature. Both Nb₃SiNi₂ and Ni₃SiNb₂ exhibit metallic behavior, showing a reduction in electrical resistivity with decreasing temperature. The residual resistivity (≈1 $\mu\Omega$ cm) is reached at around 35 K for Nb₃SiNi₂; this resistivity is associated with the scattering of the
<figure><img src="image_8.png" /><figcaption>**Figure 6.** Temperature dependence of the thermal conductivity, $κ$, for Nb$_3$SiNi$_2$ A) and Ni$_3$SiNb$_2$ B). The electronic contribution to the thermal conductivity for both ZIP phases has been calculated using the Wiedemann-Franz law. C) Temperature dependence of the electrical resistivity, $ρ$, for Nb$_3$SiNi$_2$ and Ni$_3$SiNb$_2$. D) Temperature dependence of the Seebeck coefficient, S, for Nb$_3$SiNi$_2$ and Ni$_3$SiNb$_2$. Temperature dependence of the molar heat capacity, $c_V$, for Nb$_3$SiNi$_2$ E) and Ni$_3$SiNb$_2$ F). A Debye-Einstein model (Equation 1 in Section "Thermal, Electrical, and Magnetic Properties") shown as solid blue lines, was fitted to the experimental $c_V$ data (squares). The dashed red lines correspond to the upper $c_V$ limit, dictated by the Dulong-Petit law.</figcaption></figure>

charge carrier by defects and/or impurities in this sample. On the other hand, the residual resistivity of Nb$_3$SiNb$_2$ does not seem to be reached at 10 K, which suggests that less defects were present in this sample.

Figure 6D shows the variation of the Seebeck coefficient, S, of Nb$_3$SiNi$_2$ and Ni$_3$SiNb$_2$ as a function of temperature. The Seebeck coefficients are radically different for these two ZIP phases. First, their signs are opposite, i.e., positive for Nb$_3$SiNi$_2$ and negative for Ni$_3$SiNb$_2$ at high temperatures. Second, the sign of the Seebeck coefficient changes at ≈340 K and ≈190 K for Nb$_3$SiNi$_2$ and Ni$_3$SiNb$_2$, respectively. A sign change in the Seebeck coefficient implies that the band structures are probably not parabolic, and that a two-band model is likely necessary to account for the transport properties of the Nb$_3$SiNi$_2$ and Ni$_3$SiNb$_2$ ZIP phases. Magnetoresistance initially measured at 10 K and 9 T was less than 0.15%.

Molar heat capacity, $c_V$, measured as a function of temperature (2-325 K range) showed that the recorded curves for both Nb$_3$SiNi$_2$ and Ni$_3$SiNb$_2$ exhibited an asymptotic behavior, i.e., the $c_V$ values approached asymptotically the value of ~146 ± 3 J/(mol K) (Figure 6E,F). This agrees with the Dulong-Petit law, which states that at T → ∞, $c_V$ = 3NR ≈ 149.7 J/(mol K), where N is the number of atoms per formula unit (N = 6) and R = 8.314 J/(mol K). Fitting a Debye-Einstein model[55,56] (Equation 1 in Section "Thermal, Electrical, and Magnetic Properties") to these profiles yields the Debye, $θ_D$, and Einstein, $θ_E$, temperatures, as well as the Sommerfeld coefficient, $γ$, for each sample (Table 1). The measured molar heat capacities and fitted model for the Ni$_3$SiNb$_2$ and Nb$_3$SiNi$_2$ samples are shown in Figure 6E,F.

Temperature-dependent DC magnetization susceptibility, $χ(T)$, curves were recorded as a function of the applied
<table><caption>**Table 1.** Fitted Debye ($\theta_D$) and Einstein ($\theta_E$) temperatures for Nb$_3$SiNi$_2$ and Ni$_3$SiNb$_2$, using the Debye-Einstein model (Equation 1 in Section "Thermal, Electrical, and Magnetic Properties"). The Sommerfeld coefficient, $\gamma$, and the number of Einstein nodes, $\alpha$, are also given. HC is the molar heat capacity in J/(mol K).</caption><thead><tr><th>ZIP phase</th><th>HC at RT<br/>[J/mol K]</th><th>γ [J/mol K<sup>2</sup>]</th><th>θ<sub>E</sub> [K]</th><th>θ<sub>D</sub> [K]</th><th>α</th></tr></thead><tbody><tr><td>Nb<sub>3</sub>SiNi<sub>2</sub></td><td>143</td><td>15 ± 1</td><td>192 ± 2</td><td>426 ± 2</td><td>0.19 ± 0.02</td></tr><tr><td>Ni<sub>3</sub>SiNb<sub>2</sub></td><td>148</td><td>35 ± 1</td><td>183 ± 3</td><td>377 ± 5</td><td>0.14 ± 0.01</td></tr></tbody></table>

magnetic field for Ni$_3$SiNb$_2$ and Nb$_3$SiNi$_2$ (Figure 7). As may be seen, Ni$_3$SiNb$_2$ was tested in the low magnetic fields of 10 Oe (Figure 7A) and 100 Oe (Figure 7B) as well as in the high magnetic field of 7 T (Figure 7C), whereas Nb$_3$SiNi$_2$ was tested in the low magnetic field of 100 Oe (Figure 7D) and the high magnetic field of 7 T (Figure 7E). Figure 7A,B shows a sudden change in the measured $\chi(T)$ of Ni$_3$SiNb$_2$ at low magnetic fields, indicating a transition from ferromagnetic (FM) to paramagnetic (PM) behavior with a Curie temperature, $T_C$, of ≈14 K. At the weak magnetic field of 10 Oe, $\chi(T)$ decreases sharply below ≈5 K, whilst at the strong magnetic field of 7 T (Figure 7C), Ni$_3$SiNb$_2$ remains paramagnetic within the explored temperature range. Similarly, the measured $\chi(T)$ of Nb$_3$SiNi$_2$ at the weak magnetic field of 100 Oe (Figure 7D) indicates that the FM-PM transition occurs at $T_C$ ≈ 22 K, whereas at the strong magnetic field of 7 T, the sample remains paramagnetic within the measured temperature range.

As expected, magnetic field-dependent magnetization curves, $m(H)$, measured below the Curie temperature, $T_C$, for both Ni$_3$SiNb$_2$ (Figure 8A) and Nb$_3$SiNi$_2$ (Figure 8B) showed clear ferromagnetic behavior. As shown in Figure 8D, the Ni$_3$SiNb$_2$ ZIP phase exhibits a clear hysteresis loop at 1.8 K, whereas the coercive field in the $m(H)$ curve of the Nb$_3$SiNi$_2$ ZIP phase at 1.8 K is negligible (Figure 8C).

Interestingly, magnetization measurements performed on the Nb$_3$SiNi$_2$ ZIP phase using a He3 refrigerator under both FC (field cooling) and ZFC (zero-field cooling) conditions (Figure 7F) captured clear evidence of superconductivity with a critical temperature, $T_c$, of 1.4 K. Even though it is difficult to estimate the volume fraction of superconductivity in powder samples similar to the ones used in this study, the near overlap of the FC and ZFC curves in Figure 7F suggests that it is close to 100%.

### 2.2.4. Other Properties

The Archimedes density of the Nb$_3$SiNi$_2$-based material was measured to be 7.630 g cm$^{-3}$, whilst that of the Ni$_3$SiNb$_2$-based material was measured to be 7.787 g cm$^{-3}$.

The Vickers hardness, $H_v$, of the quasi phase-pure Nb$_3$SiNi$_2$ material was measured by means of indentation to be $H_v = 12.6 \pm 0.1$ GPa.

The flexural strength, $\sigma_f$, of the quasi phase-pure Nb$_3$SiNi$_2$ material was measured by means of four-point bending (4 PB) tests to be $\sigma_f = 251 \pm 10$ MPa.

The dynamic Young's modulus, $E$, of the quasi phase-pure Nb$_3$SiNi$_2$ material was determined by the impulse excitation technique (IET) to be $E = 253.3 \pm 0.4$ GPa. Moreover, the damp-

ing or internal friction, $Q^{-1}$, of the same material was determined by means of IET to be $Q^{-1} = 0.0033 \pm 0.0002\%$, which corresponds to a characteristic resonance frequency of 7224 ± 6 Hz.

### 2.2.5. Stability of Nb-Si-Ni ZIP Phases and their Binary/2D Derivatives by DFT

The phonon band structures of Nb$_3$SiNi$_2$ (Figure 9A) and Ni$_3$SiNb$_2$ (Figure 9B) at 0 K, as calculated from density functional theory (DFT) simulations, show no imaginary modes at the Gamma (Γ) point or at any high-symmetry point, indicating thermodynamic and dynamic stability. Moreover, the lack of imaginary modes at any high-symmetry point in the phonon band structures of Nb$_3$SiNi$_2$ (Figure 9C) and Ni$_3$SiNb$_2$ (Figure 9D) at 1723 K suggests thermodynamic and dynamic stability up to this rather elevated temperature. Additional computational results on the phase stability of Nb$_3$SiNi$_2$ & Ni$_3$SiNb$_2$ and of derivative binary compounds produced without altering the elemental ratios in their stoichiometries are presented in Section S3.1 (Supporting Information). Assessing the stability of candidate binary compounds derived from “parent” ternary IMCs was a first step in the tentative exploration of the possible production of ZIP phase 2D derivatives. A second step explored the thermodynamic stability of surfaces formed via the etching of Nb$_3$SiNi$_2$ and Ni$_3$SiNb$_2$ along specific planes (Section S3.3, Supporting Information). Despite the scientific interest of a protracted discussion on the formation and stability of 2D derivatives, this topic is only addressed in the Supporting Information, as this work focuses on the ZIP phase ternary IMCs and not their 2D derivatives.

The electronic density of states (eDOS) in the hexagonal Ni$_3$SiNb$_2$ ZIP phase (Figure 9E) suggests metallic behavior, as calculated using hybrid DFT (HSE06 functional). Figure 9F depicts the reduced charge density in the hexagonal Ni$_3$SiNb$_2$ ZIP phase variant, offering a visualization of the interactions between atoms. As there is no charge density between atoms, the bonding network is considered ionic. Additional information on bond analysis is provided in Section S3.2 (Supporting Information).

## 2.3. Expanding the ZIP Phase Concept to Other Ternary Systems

### 2.3.1. Synthesis of ZIP Phases in the Nb-Si-Co, Ta-Si-Ni, and V-Si-Ni Systems

As this work proposes the existence of a whole new family of nanostructured ternary IMCs, each member of which exhibits “dualistic atomic ordering”, thus forming one fcc and one hexagonal variant, it was deemed imperative to synthesize ZIP phases in other ternary systems. Due to the sheer magnitude of this quest, especially in the absence of equilibrium phase diagrams, it has been decided to use spark plasma sintering (SPS) to make selected ternary IMCs previously reported in literature. For example, apart from the fcc Nb$_3$SiNi$_2$ ternary IMC considered as the “ternary H-phase” of the Nb-Si-Ni system,[44, 46, 47] other ternary IMCs with similar “312” stoichiometry and the same crystal structure (Fd$\bar{3}$m, SG 227) have been reported in the Mn-Si-Ni (Mn$_3$SiNi$_2$),[44, 57, 58] V-Si-Ni (V$_3$SiNi$_2$),[44] Cr-Si-Ni (Cr$_3$SiNi$_2$),[44] Ta-Si-Ni (Ta$_3$SiNi$_2$),[44,59] Nb-Si-Co (Nb$_3$SiCo$_2$),[60]
<figure><img src="image_9.png" /><figcaption>**Figure 7.** Temperature-dependent DC magnetization susceptibility, χ(T), at applied magnetic fields of A) 10 Oe, B) 100 Oe, and C) 7 T for Ni₃SiNb₂; and D) 100 Oe and E) 7 T for Nb₃SiNi₂. F) At an applied magnetic field of 20 Oe, the negative magnetization of Nb₃SiNi₂ below 1.4 K is a proof of superconductivity; (Z) FC = (zero)-field cooling.</figcaption></figure>

Na-Au-In (Na₃AuIn₂),[¹²] Na-Ag-In (Na₃AgIn₂),[¹²] and Mg-Ga-Ni (Mg₃GaNi₂)[⁶¹] systems. Few ternary IMCs with the “312” stoichiometry and the hexagonal crystal structure (P6₃/mmc, SG 194) have also been sporadically reported in prior studies. Examples include hexagonal ternary IMCs in the Co-Si-Nb (Co₃SiNb₂),[⁶⁰] Ni-Si-Nb (Ni₃SiNb₂),[⁴⁵⁻⁴⁷] and Cu-Si-Mg (Cu₃SiMg₂)[⁶⁰] systems.

In this work, SPS was employed to synthesize intermetallic alloys in the Nb-Si-Co (Figure 10A), Ta-Si-Ni (Figure 10B), and V-Si-Ni (Figure 10C) systems; the main aim of these trial synthesis efforts was to demonstrate the formation of ZIP phases in these three ternary systems. Characterization of all SPS intermetallic alloy samples by means of EPMA/WDS showed that de-

spite the low phase purity (expected, in view of the limited synthesis trials), the two ZIP phase variants formed in each ternary system, usually appearing intimately linked with each other (also expected, as they comprise the same elements in different ratios). HAADF STEM/EDS analysis of FIB foils from all intermetallic alloy samples, combined with Rietveld-refined XRD analysis of the SPS samples (Figures S16, S17, and S18, Supporting Information), confirmed the formation of the two ZIP phase variants in these ternary systems; moreover, the fcc ZIP phase variants showed invariably much higher O contents than their hexagonal counterparts, in agreement with the ZIP phases in the Nb-Si-Ni system (see Figures 2B and 5A,B; Figure S9, Supporting
<figure><img src="image_10.png" /><figcaption>**Figure 8.** Magnetic field-dependent magnetization of Ni₃SiNb₂ A) and Nb₃SiNi₂ B) at 1.8, 50, 100, 200, and 300 K. The *m*(*H*) curves of both Ni₃SiNb₂ and Nb₃SiNi₂ at 1.8 K show ferromagnetic behavior C) with a pronounced hysteresis loop for Ni₃SiNb₂ D).</figcaption></figure>

Information). Producing phase-pure materials in the Nb-Si-Co, Ta-Si-Ni, V-Si-Ni or other ternary systems falls outside the scope of this work; however, high phase purity is vital for the accurate determination of material properties. The systematic calculation of equilibrium phase diagrams in ZIP phase-forming ternary systems will surely assist and guide the synthesis of phase-pure ZIP phase materials.

### 2.3.2. Synthesis of ZIP Phases in the Nb-Si-Fe System

It is worthwhile noting that in all hitherto reported *fcc* ternary IMCs now considered as candidate ZIP phases, the electronegativity, $\chi$, of the Z-element is invariably smaller than the electronegativities of both I- and P-elements (see Table S12, Supporting Information). Moreover, the atomic radius of the Z-element is larger than the atomic radii of both I- and P-elements, with the exception of Cr₃SiNi₂, where the atomic radius (125 pm) of the Z-element, Cr, is equal to that of the P-element, Ni (Table S12, Supporting Information). One interesting way to look at the ZIP phases is by considering that the hexagonal variant is the “compositionally inverted” *fcc* variant, in the sense that the two ZIP phase variants are produced by simply interchanging the Z-element with the P-element whilst continuing to respect the “312” stoichiometry that is also shared by the 312 MAX phases (i.e., M<sub>n+1</sub>AX<sub>n</sub> where n = 2). Based on the above considerations, the electronegativity and atomic radii trends that govern the formation of the *fcc* ZIP phase variants will also be sort of “inverted” in the hexagonal ZIP phase counterparts. This means that the electronegativity of the P-element will be smaller than the electronegativities of both Z- and I-elements, whereas the atomic radius of the P-element will be larger than the atomic radii of the Z- and I-elements in the hexagonal ZIP phases. For example, for

the Nb₃SiNi₂ and Ni₃SiNb₂ ZIP phases, $\chi$(Nb) = 1.60 < $\chi$(Si) = 1.90, $\chi$(Ni) = 1.91 and *r*(Nb) = 143 pm > *r*(Si) = 118 pm, *r*(Ni) = 125 pm (Table S12, Supporting Information), irrespective of whether Nb occupies the Z-site in the *fcc* variant or the P-site in the hexagonal one. Comparing the hexagonal Ni₃SiNb₂ ZIP phase variant (Figure 4B) with the hexagonal Ti₃SiC₂ MAX phase (Figure 4C) – they both crystallize in SG *P*6₃/*mmc* – shows that these nanolaminated solids follow opposite electronegativity and atomic radii trends, i.e., $\chi$(Nb) < $\chi$(Si), $\chi$(Ni) and *r*(Nb) > *r*(Si), *r*(Ni) for Ni₃SiNb₂ versus $\chi$(C) = 2.55 > $\chi$(Si) = 1.90, $\chi$(Ti) = 1.54 and *r*(C) = 77 pm < *r*(Si) = 118 pm, *r*(Ti) = 145 pm for Ti₃SiC₂ (Table S12, Supporting Information).

Based on the herein identified trends in electronegativity and atomic radii that govern ZIP phase formation, it was decided to attempt the synthesis of ZIP phases in the Nb-Si-Fe ternary system, as no ZIP phases have been previously synthesized and/or reported in that system. Succeeding in this effort would reinforce our claim for the existence of a whole new family of ternary IMCs with dualistic atomic ordering, named the ZIP phases.

Characterization of the RHP Nb-Si-Fe intermetallic alloy samples sintered at 1723 K by means of EPMA/WDS (Figure 11) and Rietveld-refined XRD (Figure S19, Supporting Information) confirmed the formation of the two ZIP phase variants in the Nb-Si-Fe system, as predicted based on the herein identified electronegativity and atomic radii trends vis-à-vis ZIP phase formation. XRD analysis suggested that the formation of the hexagonal Fe₃SiNb₂ ZIP phase variant might have been enabled by the formation of a Fe-rich solid solution of the pseudo-binary Nb(Fe,Si)₂ IMC. The catalytic role of “precursor” solid solution IMCs in the formation of ZIP phases – also observed by Goossens et al.[52] during MAX phase formation – is likely to become a topic of future research. Former studies in support of this prediction include the work of dos Santos et al.,[47] who mentioned appreciable Si solid
<figure><img src="image_11.png" /><figcaption>**Figure 9.** Predicted phonon dispersion and atom-projected phonon density of states (PDOS) without the inclusion of temperature effects (i.e., at 0 K) A,B) and at 1723 K C,D) for the stoichiometric *fcc* Nb₃SiNi₂ A,C) and hexagonal Ni₃SiNb₂ B,D) ZIP phase variants. E) The electronic density of states (eDOS) in hexagonal Ni₃SiNb₂, using HSE06 functional (the Fermi energy is at 0 eV). F) Reduced density gradient map of hexagonal Ni₃SiNb₂.</figcaption></figure>

solubility in the Nb₇Ni₆ binary IMC; the work of Gladyshevskii et al.,[44] who reported the formation of Si solid solutions based on the Nb₃Fe₂ binary IMC; and the work of Schön and Tenório,[62] who observed that the dissolution of small amounts of phosphorus (P) and sulfur (S) in Fe-Nb alloys stabilized a pseudo-binary Nb₃Fe₂ IMC with the *fcc* diamond cubic structure (SG *Fd*3*m*). Based on their electronegativity and atomic radius values, P (χ = 2.19, *r* = 110 pm) and S (χ = 2.58, *r* = 103 pm) are “allowed” to occupy the I-site in ZIP phases containing Nb and Fe (possibly even forming solid solutions on this site). Demonstrating ZIP phase

synthesis in the Nb-(P,S)-Fe system – a candidate topic of future research – would support the idea that “precursor” binary IMCs (e.g., Nb₃Fe₂) promote ZIP phase formation.

## 2.4. Comparing the ZIP Phases with the MAX Phases

In the search for nanostructured materials with exceptional properties, both the herein proposed ZIP phases and the MAX phases share the same historical roots, dating back to the advent of space
<figure><img src="image_12.png" /><figcaption>**Figure 10.** Microstructural characterization of SPS intermetallic alloys in the Nb-Si-Co A), Ta-Si-Ni B), and V-Si-Ni C) systems. Despite the low phase purity of these three alloys, EPMA/WDS (top rows) and HAADF STEM/EDS (bottom rows) data confirm the formation of the Nb₃SiCo₂ and Co₃SiNb₂ ZIP phases in the Nb-Si-Co alloy A); the Ta₃SiNi₂ and Ni₃SiTa₂ ZIP phases in the Ta-Si-Ni alloy B); and the V₃SiNi₂ and Ni₃SiV₂ ZIP phases in the V-Si-Ni alloy C). The HAADF STEM insets in the bottom right-hand corner of the SEM micrographs of the three alloys correspond to the areas analyzed by HAADF STEM/EDS.</figcaption></figure>

exploration in the 1960s.[7, 8, 44, 46, 47, 63] Despite the fact that they have been both discovered in the same period of time, there is very little, if any, knowledge of the properties and possible applications of the ternary IMCs named “the ZIP phases” and their potential 2D derivatives, as opposed to the already extensively studied ternary early transition metal carbides and nitrides named “the MAX phases”, their higher order solid solutions, and their 2D derivatives known as MXenes. Therefore, this section will attempt to discuss the rather mature MAX phases in association with the emerging ZIP phases, identifying potential new opportunities arising from the advent of the ZIP phases based on the experience gained through the MAX-phase research over the past 30 years.

The herein introduced family of ternary IMCs, named the ZIP phases, and the family of ternary early transition metal carbides/nitrides, known as the MAX phases, constitute a set of nanostructured solids with the same overarching stoichiometry that spans almost the entire periodic table of elements, comprising jointly constituent elements from groups 1–16. **Figure 12** is a color-coded periodic table of elements showing that the MAX phases are made of M-elements from groups 3–7 (early transition metals); A-elements that belong mainly to groups 13–15

(post-transition metals, metalloids & non-metals) with limited expansion to groups 8–12 (late transition metals) and group 16 (non-metals); and X-elements that are either C or N (non-metals in groups 14 and 15, respectively). On the other hand, the so far synthesized fcc ZIP phase variants utilize Z-elements from groups 1–2 (alkali metals & alkaline earth metals) and groups 5–7 (early transition metals); I-elements from group 11 (late transition metals) and groups 13–14 (post-transition metals & metalloids); and P-elements from groups 8–10 (late transition metals) and group 13 (post-transition metals). As discussed earlier, both ZIP and MAX phases share the “312” MAX phase stoichiometry, i.e., $M_{n+1}AX_n$ where n = 2. A review by Goossens et al.[10] pointed out that the MAX phases typically belong either to one of the three well-known orders 211, 312 and 413, or exhibit one of the following exceptional stoichiometries: i) $M_{n+1}AX_n$-like structures with n > 3, such as Ta₆AlC₅ and Ti₇SnC₆, or ii) “hybrid” orders $M₅A₂C₃$ (e.g., Hf₅AlC₃,[64] or Ti₅A₂C₃ with A = Si, Ge, and Al) and $M₇A₂C₅$ (e.g., Ti₇Si₂C₅). The “hybrid” orders $M₅A₂C₃$ and $M₇A₂C₅$ have been regarded as combinations of a 312 atomic stacking with a 211 and a 413 order, respectively, as they do not obey the $M_{n+1}AX_n$ general MAX phase stoichiometry. However, considering carefully the two “312” ZIP phase variants and all MAX phase orders,
<figure><img src="image_13.png" /><figcaption>**Figure 11.** Characterization of an RHP Nb-Si-Fe intermetallic alloy sample by means of A) SEM and B) EPMA/WDS. WDS identified the Nb₃SiFe₂ and Fe₃SiNb₂ ZIP phases in the phase assembly. The O solubility in Nb₃SiFe₂ appears higher than in Fe₃SiNb₂. The inset SEM micrograph shows the formation of a eutectic structure in this system.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>**Figure 12.** Distribution of the M, A, and X elements (found in reported MAX phases) and the Z, I, and P elements (as encountered in synthesized fcc ZIP phases) in the periodic table of elements. Together, the two classes of solids span almost the entire periodic table, comprising jointly elements from groups 1-16. The M-elements of the MAX phases appear in two shades of red: the darker indicates full M-site occupancy, the lighter partial occupancy (solid solutions). The A-elements also appear in two shades of blue, depending on the degree of A-site occupancy.</figcaption></figure>
including the “hybrid” ones and the only so far identified member of the 514 order, i.e., the Mo<sub>x</sub>ValC<sub>4</sub> compound,<sup>[65]</sup> suggests that it is possible to describe both classes of nanostructured solids by the overarching stoichiometric rule P<sub>x+y</sub>A<sub>x</sub>N<sub>y</sub> (PAN), where x = 1 or 2, and y = 1, 2, 3, 4, 5, or 6, according to all “PAN phases” so far experimentally synthesized. The herein proposed overarching PAN stoichiometric rule reconciles the “hybrid” orders M<sub>2</sub>A<sub>2</sub>C<sub>3</sub> and M<sub>7</sub>A<sub>2</sub>C<sub>5</sub> with the common MAX phase orders 211, 312 & 413, answering affirmatively to the question whether it is conceptually correct to include the “hybrid” orders in the MAX phase family.

Another important aspect is that the two families of nanostruc-
tured solids have specific crystal structure similarities, i.e., both
the MAX phases and one of the two ZIP variants are hexagonal
nanolaminated solids with identical crystal structure (P63/mmc,
SG 194). In this sense, the hexagonal ZIP phase variants ex-
pand the MAX phases concept of crystal structure nanolamina-
tion beyond the early transition metals and into the late tran-
sition metals (e.g., Fe, Co, Ni) and post-transition metals (e.g.,
In). This heralds the potential fabrication of higher-order solid
solutions and 2D derivatives of the hexagonal ZIP phases, sim-
ilar to the complex MAX phase solid solutions and their 2D
derivatives (MXenes). As MAX phase solid solution synthesis
has proven a very potent tool, allowing the fabrication of ceram-
ics tailored to the property requirements of the targeted applica-
tion(s), it is reasonable to expect similar benefits from the fur-
ther development of the ZIP phases. This expectation is based
on the fact that the ZIP phases comprise elements that do not
favor MAX phase formation (e.g., alkali & alkaline earth met-
als, late transition metals), thus paving the way for the syn-
thesis of novel nanostructured solids with properties atypical
for the MAX phases (e.g., magnetic properties; superconduc-
tivity; tailorable ductility; etc.). This hypothesis finds first sup-
port in the work of Kolenda et al.,[58] who reported magnetic
properties in the Mn<sub>3</sub>SiNi<sub>2</sub> member of the ZIP phase family.
Magnetic properties were also determined for the Nb<sub>3</sub>SiNi<sub>2</sub> and
Ni<sub>3</sub>SiNb<sub>2</sub> ZIP phases produced in this work (see Figures 7 and 8),
whereas superconductivity was discovered in the Nb<sub>3</sub>SiNi<sub>2</sub> ZIP
phase (Figure 7F). By contrast, magnetic properties and super-
conductivity have been particularly difficult to impart and/or
measure in MAX phase compounds. In the plethora of studies
on the properties of the MAX phases, only a handful of stud-
ies have measured magnetic properties in experimentally synthe-
sized MAX phases, i.e., the (Cr<sub>0.75</sub>, Mn<sub>0.25</sub>)<sub>2</sub>GeC MAX phase solid
solution,[66] rare-earth-containing i-MAX phase solid solutions
(i.e., MAX phase compounds with in-plane ordered M-elements)
of the (Mo<sub>2/3</sub>, RE<sub>1/3</sub>)<sub>2</sub>AlC stoichiometry, where RE = Ce, Pr, Nd,
Sm, Gd, Tb, Dy, Ho, Er, Tm, and Lu,[24,67] and (Ti<sub>1-x</sub>, Fe<sub>x</sub>)<sub>3</sub>AlC<sub>2</sub>
MAX phase solid solutions, where x = 0, 0.025, 0.05, and 0.075.[68]
Superconductivity has only been reported for the Lu<sub>2</sub>SnC ternary
MAX phase compound.[69] It is reasonable to assume that the
new range of properties achieved by the ZIP phases (e.g., mag-
netic properties, superconductivity) might be transferable to their
2D derivatives (assuming these can be synthesized), giving rise to
the next generation of nano-tailored advanced materials for use
in flexible and wearable electronics, batteries, quantum comput-
ing, high-sensitivity magnetic sensors, functionalized templates
for biomedical applications, etc.

A unique feature of the ZIP phases is that the I- and P-
elements can belong to the same group of the periodic table. For

example, the P-element (In) in Na₃AuIn₂ and Na₃AgIn₂, as well
as the I-element (Ga) in Mg₃GaNi₂, belong to group 13. Such free-
dom in choosing the elements making up the ZIP phases does
not exist in the MAX phases, where only two X-elements are al-
lowed (i.e., C or N), and these are not known to be replaced by
any of the candidate M- or A-elements. The only known element
with flexibility to occupy either the M- or A-site in specific MAX
phases is iron (Fe), which belongs to group 8 of late transition
metals. More specifically, Fe has been either reported to form
the (Ti₁₋ₓ, Feₓ)₃AlC₂ solid solutions on the M-site,[68] or to oc-
cupy fully the A-site of the ternary MAX phases Ta₂FeC, Ti₂FeN,
and Nb₂FeC.[70] The allowable combinations of elemental con-
stituents in the ZIP phases could tentatively be associated with
relative differences in the electronegativity and atomic radius val-
ues of the Z/I/P-elements; however, before strict rules governing
ZIP phase formation are proposed, further systematic research
is needed to better map, in terms of attainable chemical compo-
sitions, the ZIP phases' "landscape".

A cornerstone in the prospective exploitation of the ZIP phases is the synthesis of phase-pure materials, as high phase purity is a prerequisite for the determination of material properties, such as elastoplastic/fracture behavior, corrosion resistance, electromagnetic properties, radiation tolerance, etc., thereby exploring the possible applications of the ZIP phases and/or their 2D derivatives. Challenges in producing phase-pure (bulk) materials have also been reported for the Zr/Hf-based MAX phases.[20, 21, 38, 40, 71] Overcoming these challenges involved either the synthesis of complex double solid solutions based on steric stability criteria[72,73] or the use of early transition metal hydride starting powders rather than elemental ones.[52] Several processing methods, other than RHP, can be explored to synthesize phase-pure ZIP phases in different forms (bulk, coating) to make progress with property determination and exploration of the application space for both ternary IMCs and their 2D derivatives. For example, physical vapor deposition (PVD) methods, such as magnetron sputtering, are most likely suitable for the deposition of ZIP phases (i.e., IMCs with very precise stoichiometries and appreciable structural complexity) on different substrates, suppressing the formation of competing phases, and producing strongly textured, phase-pure coatings that might be desirable for various high-tech applications.

**3. Conclusion and Outlook**

This work introduced a new family of nanostructured ternary in-
termetallic compounds (IMCs), aptly named the ZIP phases due
to their distinctive “zigzag” atomic arrangements. Each member
of the ZIP phase family manifests itself in two variants, one with
the fcc diamond cubic structure (Fd3m, SG 227) and one with the
hexagonal structure (P63/mmc, SG 194); the latter is identical to
the structure of the hexagonal MAX phases. All ZIP phases obey
the 312 MAX phase stoichiometric rule, i.e., Mn<sub>n+1</sub>AXn with n =
2, or, alternatively, the herein proposed overarching ZIP/MAX
stoichiometric rule P<sub>x+y</sub>A<sub>x</sub>N<sub>y</sub> (PAN) with x = 1 and y = 2. The
hitherto synthesized ZIP phases (sixteen ternary IMCs, in to-
tal) comprise Z-elements from groups 1-2 (alkali metals & al-
kaline earth metals) and groups 5-7 (early transition metals); I-
elements from group 11 (late transition metals) and groups 13-14
(post-transition metals & metalloids); and P-elements from groups 8-10 (late transition metals) and group 13 (post-transition metals). The MAX phases (ternary early transition metal carbides & nitrides) and the ZIP phases (ternary IMCs) constitute a set of nanostructured solids that spans almost the entire periodic table of elements, comprising jointly elements from groups 1-16. This implies that the hexagonal ZIP phase variants expand the MAX phases concept of crystal structure nanolamination beyond the early transition metals and well into the late transition and post-transition metals, thus enabling the synthesis of new nanostructured solids and potential 2D derivatives with properties (e.g., magnetism) that continue to elude the MAX phases and their 2D derivatives known as MXenes.

In this work, ZIP phase-based materials in the Nb-Si-Ni, Nb-Si-Co, Ta-Si-Ni, V-Si-Ni, and Nb-Si-Fe ternary systems were synthesized by means of reactive hot pressing (RHP) and spark plasma sintering (SPS), proving the formation of the two ZIP variants in all five ternary systems. Quite importantly, RHP was successfully and reproducibly able to synthesize quasi phase-pure $Nb_3SiNi_2$ and $Ni_3SiNb_2$ materials, thus demonstrating (scalable) synthesis feasibility and paving the way for the further development and exploitation of the ZIP phases. The experimentally synthesized $Nb_3SiNi_2$ and $Ni_3SiNb_2$ ZIP phase variants were characterized in terms of crystal structure; spatially resolved composition and reconstructed atomic positions; thermal, electrical, and magnetic properties; and selected mechanical and physical properties. Moreover, density functional theory (DFT) and the temperature-dependent effective potential (TDEP) method were employed to assess the stability of the ZIP ternary IMCs and candidate derivative binary compounds as a function temperature, also exploring the possible exfoliation of the ZIP phase variants along specific planes in order to produce 2D derivatives; the latter analysis showed that the hexagonal ZIP phase variant is more prone to exfoliation than its *fcc* counterpart. The ZIP phases are characterized by metallic overall behavior and ionic bonding.

The advent of the ZIP phases is expected to create a new ecosystem of nanostructured ternary IMCs with dualistic atomic ordering and potential 2D derivatives with properties (e.g., magnetism, tailorable ductility, superconductivity, etc.) that cannot be found in other nanostructured or nanolaminated solids, such as the MAX phases and their 2D derivatives (MXenes). This paves the way for the exploration of a host of deep tech applications in diverse technology sectors, including flexible/wearable electronics; batteries; quantum computing; highly sensitive magnetic sensors; biomedical applications aiming at smart drug delivery; etc. The structural complexity of the ZIP phases (e.g., the unit cell of $Nb_3SiNi_2$ has 96 atoms) renders them, after a fashion, comparable to high-entropy alloys (HEAs), the chemical complexity of which imparts them with remarkable radiation tolerance. At this stage, it is impossible to make accurate predictions as to the full range of applications of the ZIP phases. The first step toward the further exploitation of this remarkable new family of materials is the synthesis and full characterization of phase-pure materials in various ZIP phase-forming ternary systems. A second step of great importance would undoubtedly be the experimental synthesis of proof-of-concept 2D derivatives of the ZIP phases.

## 4. Experimental Section

**Synthesis of $Nb_3SiNi_2$ via Arc Melting:** Conventional vacuum arc melting of high-purity (>99.99%) Nb, Ni, and Si elemental raw powders was initially employed to synthesize the Nb-Si-Ni intermetallic alloy with the targeted H-phase $Nb_3SiNi_2$ stoichiometry, i.e., 50.0Nb-16.7Si-33.3Ni in at.%. Arc melting yielded a 200-g disc of low phase purity, even though the ratio of elemental powders in the powder feedstock had been adjusted to match the $Nb_3SiNi_2$ stoichiometry. Additional information on the arc melting synthesis of alloys in the Nb-Si-Ni ternary system may be found elsewhere.[45-47]

The thermal and microstructural stability of the as-cast Nb-Si-Ni intermetallic alloy was assessed via an isothermal annealing treatment at 1421 K for 2 weeks (336 h) in a vacuum furnace (ThermoLyne furnace, model 46 100). The as-cast alloy samples were pre-encapsulated in a quartz ($SiO_2$) tube, which was pressurized with argon (Ar) gas to 1 atm; both heating and cooling rates were fixed at ≈14.2 K/min.

**Synthesis of $Nb_3SiNi_2$ via Reactive Hot Pressing (RHP):** The failure to produce a phase-pure Nb-Si-Ni intermetallic alloy by arc melting redirected the synthesis efforts in this work toward a powder metallurgical processing route. Discs of the Nb-Si-Ni intermetallic alloy (~26 g each) with high phase purity were successfully produced via reactive hot pressing (RHP), using hydride and elemental starting powders, i.e., $NbH_{0.89}$ (<40 µm, CBMM, Brazil), Ni (Vale, T123, 3-7 µm, UK), and Si (A10, 2.1 µm, HC Starck, Germany). To optimize the mixing of the starting powders and improve powder particle contact, the as-received $NbH_{0.89}$ powder was first refined by low-energy ball milling in isopropanol for 24 h (Turbula T2), using WC-6 wt.% Co (WC-6Co) milling balls with diameters of 5 and 10 mm. The refined $NbH_{0.89}$ powder was dried in a rotating evaporator (Heidolph 4010) and then mixed with the Ni and Si elemental powders in isopropanol for another 24 h, using WC-6Co milling balls. The ratio of $NbH_{0.89}$, Ni, and Si in the powder feedstock was adjusted to the targeted 3Nb:1Si:2Ni. After mixing, the feedstock powder suspension was dried in a rotating evaporator.

The powder feedstock was cold-pressed into "green" pellets (Ø 30 mm) under a load of 30 MPa in a graphite die. Exploratory RHP runs were conducted in a hot press (W100/150-2200-50 LAX, FCT Systeme, Frankenblick, Germany). The powder compacts were heated in vacuum (0.4 mbar) at 10 K/min to different sintering temperatures (i.e., 1523, 1623, and 1723 K), where densification was achieved under a uniaxial pressure of 30 MPa for 60 min. The sintered discs (Ø 30 mm, 4.5 mm in thickness) were allowed to cool naturally by switching off the power supply of the hot press. RHP was also used to produce a larger disc of $Nb_3SiNi_2$ (Ø 56 mm, 4.5 mm in thickness) to assess mechanical properties and hardness; this disc was sintered at 1723 K for 60 min under a uniaxial pressure of 15 MPa.

**Synthesis of ZIP Phases in Other Ternary Systems—Synthesis of Previously Reported Ternary IMCs of the ZIP Phase Type:** Spark plasma sintering (SPS) and RHP were employed to produce discs with ternary IMCs in the Nb-Si-Co, Ta-Si-Ni, V-Si-Ni, and Fe-Si-Nb systems, thus proving that the ZIP phases are indeed a family of ternary IMCs with two variants, i.e., an *fcc* (SG *Fd\overline{3}m*) variant and a hexagonal (SG *P6\overline{3}/mmc*) variant. The starting powders used to synthesize the Nb-Si-Co intermetallic alloy with the $Nb_3SiCo_2$ stoichiometry were $NbH_{0.89}$ (<40 µm, CBMM, Brazil), Co (Co-HMP, FSSS 0.55 µm, Umicore, Belgium), and Si (A10, 2.1 µm, HC Starck, Germany). The $NbH_{0.89}$ powder was refined by low-energy ball milling in isopropanol for 24 h (Turbula T2), using WC-6Co milling balls (diameters of 5 mm and 10 mm). The $NbH_{0.89}$, Co, and Si powders were mixed in isopropanol for another 24 h, using WC-6Co milling balls. After mixing, the feedstock powders were dried in a rotating evaporator. The dried powders were densified by means of SPS (HP D 25, FCT Systeme, Frankenblick, Germany), in vacuum, at 1723 K, under a uniaxial pressure of 30 MPa for 2 min; the powder compact was heated to 1723 K at 100 K/min and cooled by switching off the power supply. The sintered disc was 30 mm in diameter and 4.5 mm in thickness.

The feedstock powders used to synthesize the Ta-Si-Ni intermetallic alloy of the $Ta_3SiNi_2$ stoichiometry were Ta (FTa-1, 2 µm, Ningxia Orient Tantalum Industry Co., Ltd., China), Ni (Vale, T123, 3-7 µm, UK), and Si
(A10, 2.1 µm, HC Starck, Germany). The Ta, Ni, and Si powders were dry-mixed for 3 min at 1800 rpm, using a high-speed mixer (DAC 150.1 FWZ-K, Flacktek, Germany). The powders were densified by SPS (HP D 25, FCT Systeme, Frankenblick, Germany), in vacuum, at 1623 K, under a uniaxial pressure of 30 MPa for 10 min; the powder compact was heated to 1623 K at 100 K/min and cooled by switching off the power supply. The sintered disc was 30 mm in diameter and 4.5 mm-thick.

The starting powders used to synthesize the V-Si-Ni intermetallic alloy with the $V_3SiNi_2$ stoichiometry were $V_2H$ (<20 µm, lab grade), Ni (Vale, T123, 3-7 µm, UK), and Si (A10, 2.1 µm, HC Starck, Germany). The $V_2H$ powder was synthesized via the hydrogenation of V powder (<44 µm, Alfa Aesar, Germany) in a horizontal tube furnace (Carbolite type CTF), under continuous flow of hydrogen ($H_2$) gas (Air Liquide N40). Powder hydrogenation was achieved by slowly heating (5 K/min) the V powder to 1093 K, holding for 5 h, and slowly cooling (5 K/min) to room temperature (RT). Switching the gas flow to technical argon (Ar; 3-5 ppm $O_2$, Nippon Gases Ltd.) after cooling allowed the removal of any remaining $H_2$ without risking the spontaneous oxidation of $V_2H$ due to its sudden exposure to oxygen ($O_2$). Vibratory ring milling (Retsch RS200) for 30 s at 1100 rpm of the brittle $V_2H$ powder resulted in powder refinement; the powder was sieved and only its finer fraction (<20 µm) was used for further processing. The feedstock powders were mixed in isopropanol for 24 h using WC-6Co milling balls and were subsequently dried in a rotating evaporator. The dried powders were densified by SPS (HP D 25, FCT Systeme, Frankenblick, Germany), in vacuum, at 1523 K, under a uniaxial pressure of 30 MPa for 10 min; the powder compact was heated to 1523 K at 100 K/min and cooled by switching off the power supply. The sintered disc diameter was 30 mm, and its thickness was 4.5 mm.

**Synthesis of ZIP Phases in Other Tertiary Systems—Synthesis of Predicted ZIP Phases in the Nb-Si-Fe System:** The starting powders used to synthesize the Nb-Si-Fe intermetallic alloy with the $Nb_3SiFe_2$ stoichiometry were $NbH_{0.89}$ (<40 µm, CBMM, Brazil), Fe (Carbonyl Iron, 3.6 µm, BASF, Germany), and Si (A10, 2.1 µm, HC Starck, Germany). The $NbH_{0.89}$, Fe, and Si powders were dry-mixed for 24 h without milling balls (Turbula T2). The feedstock powders were densified in a hot press (W100/150-2200-50 LAX, FCT Systeme, Frankenblick, Germany), in vacuum, at 1723 K, under a uniaxial pressure of 30 MPa for 120 min; the powder compact was heated to 1723 K at 10 K/min and cooled by switching off the power supply. The sintered disc had a diameter of 30 mm and a thickness of 4.5 mm.

**Materials Characterization—X-Ray Diffraction (XRD):** X-ray diffraction (XRD) was used to identify the phases present in all intermetallic alloy samples at RT. The as-cast and annealed Nb-Si-Ni alloy samples produced by arc melting were analyzed on a Bruker AXS D8 Advance X-ray diffractometer, using Cu-Kα radiation (40 kV, 40 mA) in a Bragg-Brentano geometry; the XRD spectra were recorded at steps of 0.02° in the 20-90° 2θ range. All RHP and SPS Nb-Si-Ni, Nb-Si-Co, Ta-Si-Ni, V-Si-Ni, and Nb-Si-Fe alloy samples were first analyzed in bulk form on a Rigaku SmartLab SE XRD system with 1D/teX Ultra 250 HE detector, using Cu-Kα radiation (40 kV, 40 mA) in a Bragg-Brentano geometry; the XRD spectra were recorded at steps of 0.05° in the 5-90° 2θ range, at a scan speed of 0.1° per min. One of the quasi phase-pure RHP Nb-Si-Ni alloy discs (Ø 30 mm) sintered at 1723 K was pulverized for XRD analysis, followed by full Rietveld refinement of the XRD pattern. The powder was obtained by ring milling (Retsch RS200; 30 s, 1100 rpm; stainless steel container), whilst the Rietveld refinement (weighted-profile R-factor, $R_{wp}$ < 2.84%) was performed using the Profix version 5.4.1 software package. The Rietveld refinement results, including the sample constituent phases and their lattice parameters (also deduced for the “impurity phases” $NbO_2$, $SiO_2$, and Si), the atomic site positions, occupancies & displacements, and the temperature factors, are summarized in Table S2 (Supporting Information). Due to the important insights gained by Rietveld-refined XRD analysis of the quasi phase-pure Nb-Si-Ni alloy, it was decided to do Rietveld refinement on powder XRD samples produced in a similar manner from the lower phase purity SPS Nb-Si-Co, Ta-Si-Ni, V-Si-Ni, and RHP Nb-Si-Fe alloy samples. Rietveld refinement was again performed using the Profix version 5.4.1 software package, and the results (constituent phases and lattice parameters, atomic site positions, and occupancies) are summarized in Table S8 (Supporting Information) for the Nb-Si-Co alloy, Table S9 (Supporting Information) for the Ta-Si-Ni alloy, Table S10 (Supporting Information) for the V-Si-Ni alloy, and Table S11 (Supporting Information) for the Nb-Si-Fe alloy.

**Materials Characterization—Scanning Electron Microscopy (SEM) & Electron Probe Microanalysis (EPMA):** Metallographic cross-sections were prepared from all intermetallic alloy samples that were produced via arc melting, RHP, and SPS; all cross-sections were polished to a mirror surface finish with a colloidal silica ($SiO_2$) suspension (known as OPS) in the final step. The microstructure of as-cast and annealed Nb-Si-Ni alloy samples made by arc melting was studied by means of scanning electron microscopy (SEM; Thermo Fisher Apreo SEM) equipped with energy-dispersive X-ray spectroscopy (EDS; EDAX Octane Elite Super detector). Elemental mapping of all RHP and SPS alloy samples was performed by means of electron probe microanalysis (EPMA; JXA-8530F, JEOL Ltd, Japan) equipped with wavelength-dispersive X-ray spectroscopy (WDS; 1-5 eV, 8 eV at Fe-Kα, full scanner type spectrometer). EDS elemental analysis (1 SDD/UTW, 129 eV at Mn-Kα) of the identified phases was also performed on the same device at fixed conditions (15 kV accelerating voltage, 100 nA beam current).

**Materials Characterization—Focused Ion Beam (FIB):** Focused ion beam (FIB; FEI Helios 600 dual-beam FIB/SEM) was used to lift out thin foils (thickness < 100 nm) from as-cast and annealed Nb-Si-Ni alloy samples for inspection by means of transmission electron microscopy (TEM), utilizing the classic lift-out method.[74] Combined platinum/carbon (Pt/C) protective capping layers were deposited on the sites of interest to protect the thin foils from ion beam damage during lift-out. The thin foils were ion-milled to electron transparency at 30 keV, with the final thinning/ion beam damage removal step at 2 keV.

Additional FIB foils were prepared from the Nb-Si-Co, Ta-Si-Ni, and V-Si-Ni alloy samples, using a FEI Helios Nanolab 660 dual-beam FIB/SEM. The regions of interest were covered first with a thin (~100 nm) Pt layer deposited by the electron beam, followed by a ~2 µm Pt protective capping layer deposited using the ion beam. Final ion milling was carried out at 5 kV acceleration voltage, and the produced thin foils were thinner than 50 nm.

**Materials Characterization—Transmission Electron Microscopy (TEM):** TEM analysis of FIB foils lifted out of the arc-melted and annealed Nb-Si-Ni alloy samples was performed with a FEI Tecnai TF30 and a FEI Titan 80/300, both operated at 300 kV. The latter was equipped with an EDAX Octane Elite T EDS detector, whilst quantitative EDS analysis was done based on the Cliff-Lorimer method.[75,76] Selected area electron diffraction patterns (SAEDPs) of the $Nb_3SiNi_2$ ZIP phase were indexed using data from the Inorganic Crystal Structure Database (ICSD).[48] The CrystalMaker, SingleCrystal, and CrystalDiffraction software packages were used to simulate electron diffraction patterns and aid the indexing of experimentally acquired SAEDPs.[53]

Atomically resolved TEM images of the $Nb_3SiNi_2$ and $Ni_3SiNb_2$ ZIP phases were acquired by means of high-angle annular dark-field scanning transmission electron microscopy (HAADF STEM) on a double Cs-corrected FEI Titan³ 60-300 operated at 300 kV and equipped with a Super-X EDS system.

A Thermo Fisher Scientific Talos F200X S/TEM operating at 200 keV was used for HAADF STEM imaging of FIB foils from the RHP Nb-Si-Co, Ta-Si-Ni, and V-Si-Ni alloy samples. STEM/EDS elemental maps were acquired using the Super-X EDS system in the same TEM.

**Materials Characterization—Atom Probe Tomography (APT):** FIB (FEI Helios Nanolab 660 dual-beam FIB/SEM) was employed to make needle-shaped samples from the Nb-Si-Ni intermetallic alloy for further inspection by means of atom probe tomography (APT). The APT samples were made following a standard protocol[77] using gallium (Ga) ions at 30 kV on the FIB. The site-specific sample lift-out was guided by EDS (EDAX Octane Elect detector; 5 kV acceleration voltage, 1.6 nA current). Final annular milling of the needle-shaped specimens was done at 5 kV and 40 pA for 9 min to achieve a radius at the apex of ~40 nm and a shank angle of ~30°. Individual steps of APT specimen preparation are presented in Figure S6 (Supporting Information).

The composition of the Nb-Si-Ni intermetallic alloy was spatially-resolved with APT, using voltage pulsing in a CAMECA local electrode atom
probe (LEAP) 4000X HR. The pulse fraction, pulse rate, base temperature, and detection rate were 20%, 125 kHz, 60 K, and 0.5%, respectively. At least 15 × 10⁶ ions were collected, data analysis was carried out with the AP Suite 6.1 software (CAMECA, Madison, WI, USA), and the shank angle protocol was used for reconstructions.

**Materials Characterization—Thermal, Electrical, and Magnetic Properties:** The Nb-Si-Ni intermetallic alloy samples were cut into rectangular parallelepipeds of 2 × 2 × 8 mm³. Electrical resistivity, ρ, measurements were performed with a four-probe technique, using the BRT (Bridge, Rotator, and Tq-Mag) option of the physical properties measurement system (PPMS; Quantum Design, San Diego, CA, USA). A current of 8 mA was applied to reach a resistance of ≈1 mΩ at RT. The temperature was reduced slowly from RT to 5 K, thus measuring continuously the electrical resistivity. The thermal conductivity, κ, and the Seebeck coefficient, S, were measured using the TTO (Thermal Transport Option) system of the PPMS. The TTO system measures thermal conductivity by applying heat from the heater shoe, creating a user-specified temperature differential between the two thermometer shoes. The TTO system models dynamically the thermal response of the sample to a low-frequency, square-wave heat pulse, thus expediting data acquisition. The TTO system can then calculate thermal conductivity directly from the applied heater power (~50 mW at RT), the resulting ΔT (~10% of the temperature of the sample), and the sample geometry. With respect to the Seebeck coefficient, the voltage drop created between the two thermometer shoes was also monitored.

Molar heat capacity, c_V, measurements were conducted on Nb-Si-Ni intermetallic alloy samples as a function of temperature in the 2-325 K range, using a Quantum Design PPMS Dyna Cool measurement system. Fitting a Debye-Einstein model (Equation 1) to the molar heat capacity curves of Nb₃SiNi₂ and Ni₃SiNb₂ yielded the Debye (θ_D) and Einstein (θ_E) temperatures, as well as the Sommerfeld coefficient, γ, for each ZIP phase:

$$c_V = \gamma T + 9NR(1-\alpha) \left(\frac{T}{\theta_D}\right)^3 \left(\frac{\theta_D}{T}\right)^{\frac{\alpha}{\gamma}} \frac{\chi^4 e^\alpha}{(e^\alpha - 1)^2} dx + 3NR\alpha \left(\frac{\frac{\theta_E}{T}}{e^\alpha - 1}\right)^2 \quad (1)$$

where T is the temperature of the sample, α is the number of Einstein modes, N = 6 is the number of atoms per formula unit (i.e., Nb₃SiNi₂ and Ni₃SiNb₂), and R = 8.314 J/(mol K).

Temperature- and applied-magnetic-field-dependent magnetization measurements were conducted on Nb-Si-Ni intermetallic alloy samples, using the Quantum Design MPMS3 measurement system equipped with a He3 refrigerator. Temperature-dependent DC magnetic susceptibilities, χ(T), were obtained for both Nb₃SiNi₂ and Ni₃SiNb₂ ZIP phases from zero-field-cooled (ZFC) magnetization measurements at low magnetic fields, H, of 10 Oe and 100 Oe, as well as at a high magnetic field of 70 kOe (equal to 7 T). Field-dependent magnetization measurements, m(H), were carried out at 1.8 K, 50 K, 100 K, 200 K, and 300 K. Moreover, the temperature-dependent magnetic moment, m(T), was measured between 0.4 K and 2 K for the Nb₃SiNi₂ ZIP phase at an applied magnetic field of 20 Oe under ZFC and field-cooled (FC) conditions.

**Materials Characterization—Other Properties:** Density was determined on the smaller RHP discs (Ø 30 mm) of Nb₃SiNi₂ and Ni₃SiNb₂ based on Archimedes' principle. The larger RHP disc (Ø 56 mm) of Nb₃SiNi₂ was cut into rectangular bars (43 × 3 × 2.3 mm³) that were polished prior to testing for mechanical and physical property determination. Vickers hardness (H_V) was determined using an indentation load of 3 kg for 15 s (Future Tech model FV700 indenter). Four-point bending (4 PB) testing was employed to determine flexural strength (σ_f); five rectangular bars (43 × 3 × 2.3 mm³) were tested on an Instron 4467 test machine with fixed crosshead displacement speed (0.5 mm/min) and fixed span length (20 mm inner span, 40 mm outer span). Dynamic elastic properties (Young's modulus, damping, and the help of the impulse excitation technique (IET; GrindoSonic MK).

**Thermodynamic and Crystal Structure Simulations:** Thermodynamic simulations were performed to predict the equilibrium phases at the sin-

tering temperatures of the RHP Nb-Si-Ni alloys (i.e., 1523 K, 1623 K, and 1723 K), as well as at 1423 K (i.e., close to the homogenization annealing temperature, 1421 K, of the arc-melted alloy). These simulations were performed with version V2020a of the Thermo-Calc software (Thermo-Calc Software Inc., Solna, Sweden).[78] based on the database published by dos Santos et al.[45] The unit cells of the ZIP phases in the Nb-Si-Ni system and the unit cell of the Ti₃SiC₂ MAX phase were drawn using version 3.5.8 of the VESTA software[54] and utilizing the Wyckoff parameters provided in Table S3 (Supporting Information) for precise crystal structure representation. The two ZIP phases (i.e., Nb₃SiNi₂ and Ni₃SiNb₂) in the Nb-Si-Ni intermetallic alloy were compared in terms of unit cell sizes and structural complexity, whilst the crystal structure of the hexagonal Ni₃SiNb₂ ZIP phase variant was further compared to the (hexagonal) Ti₃SiC₂ MAX phase.

**Computational Methodology—Energy Minimization Calculations of Bulk Phases:** Non-spin polarized calculations were run using the Vienna Ab initio Simulation Package (VASP 5 code)[79,80] It employs the frozen-core projector-augmented wave (PAW) method[81,82] and a plane-wave basis set with a 650-eV cutoff, which was checked for convergence. The electron configurations were Si (2s²2p²), Ni (4d⁸4s²), and Nb (4s²4p⁶4d³5s²). 3D boundary conditions were used. The exchange correlation functional was the Perdew-Burke-Ernzerhof (PBE) generalized gradient approximation (GGA).[83] All models were generated using the METADISE code.[84] The full bulk 12-atom unit cell for the hexagonal Ni₃SiNb₂, and the 96-atom unit cell for the fcc Nb₃SiNi₂ were simulated with electronic and ionic convergence criteria of 1 × 10⁻⁹ eV and 1 × 10⁻³ eV Å⁻¹, with the Brillouin zone sampled using Γ-centered 5 × 5 × 3 and 4 × 4 × 4 k-point meshes, respectively. The k-point meshes were checked for convergence. To check the stability of any binary compounds that could arise from the removal of one of the elements in the ternary IMCs, the hexagonal Ni₃Si, Ni₃Nb₂, SiNb₂, and fcc Nb₃Si, Nb₃Ni₂, SiNi₂ were simulated using the same settings as for the ternary IMCs. The Phonopy code[85] was used to calculate second-order harmonic interatomic force constants using the finite-displacement method, with the VASP 5 code used as the force calculator. The phonon dispersions of all the stoichiometric structures were determined using a transformation matrix to the primitive cell, as implemented in the Phonopy code. The phonon dispersions of all bulk models and the atom-projected phonon density of states (PDOS) curves for the stoichiometric cells were calculated on 16 × 16 × 16 and 30 × 30 × 30 Γ-centered q-point grids for the fcc and hexagonal structures, respectively.

**Computational Methodology—Energy Minimization Calculations of Surfaces:** All surface models were generated using the METADISE code. They were modelled using the slab method, according to which two identical surfaces were created. A vacuum gap of 15 Å perpendicular to the surface was introduced to minimize the interaction between images. The surfaces were simulated with electronic and ionic convergence criteria of 1 × 10⁻⁵ eV and 1 × 10⁻² eV Å⁻¹, with the Brillouin zone sampled using a Γ-centered k-point mesh of 3 × 3 × 1 and 4 × 6 × 1 for the hexagonal (001) and (110) Ni₃SiNb₂ surfaces, and 2 × 2 × 1 for the fcc Nb₃SiNi₂ surfaces, with the third vector perpendicular to the surface plane. For the hexagonal Ni₃SiNb₂, the (001) and (110) surfaces (planes) were chosen as representative models of possible etching of the material. The (001) models were a (2 × 2) in-plane expansion of the full surface unit cell. For the (001), the thickness of the slab was 2, 4, and 6 surface repeat units with 48 (8 Ni₃SiNb₂ units), 96 (16 Ni₃SiNb₂ units), and 144 (24 Ni₃SiNb₂ units) atoms, respectively. A surface with 1 repeat surface unit could not be stabilized for the (001) surface. To remove the dipole of the (001) slab, half of the surface Si atoms were moved from one side to the other of the slab, generating highly symmetric surface models. The (110) models were a (1 × 1) in-plane expansion of the full surface unit cell. For the (110), the thickness of the slab was 1, 2, and 3 surface repeat units with 12 (2 Ni₃SiNb₂ units), 24 (4 Ni₃SiNb₂ units), and 36 (6 Ni₃SiNb₂ units) atoms, respectively. For the fcc Nb₃SiNi₂, the (010) and (101) surfaces (planes) were chosen as representative models of possible etching of the material. Both models were a (1 × 1) in-plane expansion of the primitive surface unit cell. For both surfaces, the thickness of the slab was 1, 2, and 3 surface repeat units with 24 (4 Ni₃SiNi₂ units), 48 (8 Ni₃SiNi₂ units), and 72 (12 Ni₃SiNi₂ units) atoms, respectively. Two different terminations were
used for the (010), and to remove the dipole, half of the surface Nb and Si atoms, respectively, were moved from one side to the other of the slab, generating highly symmetric surface models. To remove the dipole of the (101) slab, half of the surface Si atoms were moved from one side to the other of the slab, generating highly symmetric surface models. The surface energy was calculated as $\gamma_{surf} = (E_{surf} - E_{bulk})/2A$, where $E_{surf}$ and $E_{bulk}$ are the energies of the surface model and the bulk structure with the same number of units as the surface model, respectively; A is the surface area; and 2 accounts for the usage of the slab method.

**Computational Methodology—Bond Analysis:** Hybrid density func-
tional theory (DFT) was also performed, using the Heyd-Scuseria-
Ernzerhof (HSE06) functional[86,87] for the geometry optimization of
Ni₃SiNb₂, computed with a mesh of 5 × 5 × 3 k-points, and the electronic
density of states (eDOS) was computed with a mesh of 8 × 8 × 4 k-points.
The structure was simulated with electronic and ionic convergence criteria
of 1 × 10⁻⁵ eV and 1 × 10⁻² eV Å⁻¹. The Bader analysis was performed
by partitioning the hybrid DFT core and valence charge density grids.[88,89]
The critical points (CP) and the non-covalent interactions (NCI) of the hy-
brid DFT charge densities were computed using the Critic2 code.[90,91]

**Computational Methodology—Temperature-Dependent Phonon Dispersion:** The phonon dispersion relations were determined at finite temperatures from the ab initio molecular dynamics (AIMD) simulations using the temperature-dependent effective potential (TDEP) method. The TDEP method was explained in detail elsewhere.[92,93] The basis of the TDEP method was to obtain the interatomic force constant matrices at finite-temperatures from the AIMD together with Hooke's law, provided the structure has a valid symmetry. It was necessary to use a large supercell to achieve the force convergence for TDEP. Therefore, a supercell with 324 atoms for hexagonal Ni₃SiNb₂ (a 3 × 3 × 3 expansion of the full unit cell) and a supercell with 768 atoms for fcc Nb₃SiNi₂ (a 2 × 2 × 2 expansion of the full unit cell) were used to run the AIMD simulations. Similarly, we constructed supercells for the binary compounds with atoms 270 for hexagonal Ni₃Nb₂, 216 for hexagonal Ni₃Si, 162 for hexagonal SiNb₂, 640 for fcc Nb₃Ni₂, 512 for fcc Nb₃Si, and 384 for fcc SiNi₂, respectively. The lattice parameters for the supercells were adopted from the 0 K optimized structure of the corresponding compounds. The AIMD simulations were carried out using the VASP code with identical pseudopotentials (electron configurations) and exchange-correlation functionals, as described earlier. However, the Brillouin zone integration was performed using the Γ point of the supercell. The temperature of both ternary compounds was set to 1723 K during the AIMD simulations, as they were found to be stable at this elevated temperature and in agreement with the experiments. However, all the binary compounds were simulated at 300 K except hexagonal SiNb₂, which was simulated at 600 K. A Nosé-Hoover thermostat was used to control the temperature. Since a very large simulation box was chosen, the simulations were run about 2000 steps for fcc structures and 5000 for hexagonal structures with a time step of 2 fs to obtain the converged forces.

Supporting Information

Supporting Information is available from the Wiley Online Library or from
the author.

Acknowledgements

M.A.T. acknowledges support by the Laboratory Directed Research and Development (LDRD) program of Los Alamos National Laboratory (LANL), primarily under the Directors' Fellowship project number 20200689PRD2 entitled "Highly Ordered Refractory Intermetallics: The ZIA-Phases Project". LANL is an affirmative action equal opportunity employer, managed by Triad National Security, LLC, for the U.S. Department of Energy's National Nuclear Security Administration, under contract No 89233218CNA000001. M.A.T. would also like to thank Dr. Kurt E. Sickafus, LANL, for valuable discussions on the crystallography of complex solid-state phases. Moreover, M.A.T. is grateful for funding by the Aus-

trian Research Promotion Agency (FFG) within the project 3DnanoAna-
lytics (FFG-No 858040). C.G.S. acknowledges financial support provided
by the Conselho Nacional de Desenvolvimento Científico e Tecnológico
(CNPq, Brasília-DF, Project 307627/2021-7). K.L. and J.V. acknowledge
knowledge exchange on chemically complex MAX phases with the H2020
IL TROVATORE project (Euratom research and training program 2014-
2018; Grant Agreement No. 740415). N.G. acknowledges the Fund for
Scientific Research Flanders (FWO-Vlaanderen) for funding his PhD un-
der the Fundamental Research Fellowship No 1118120N. All authors fur-
ther acknowledge the Swedish Research Council and Swedish Foundation
for Strategic Research for access to ARTEMI, the Swedish National Infras-
tructure in Advanced Electron Microscopy (2021-00171 and RIF21-0026).
M.H. and J.M.S. acknowledge funding from Deutsche Forschungsgemein-
schaft (DFG) within the Collaborative Research Centre SFB 1394, "Struc-
tural and Chemical Atomic Complexity-From Defect Phase Diagrams to
Materials Properties" (project ID 409476157). M.M. and S.R.C. performed
calculations using the ARCHER2 UK National Supercomputing Services
via M.M.'s membership of the UK HEC Materials Chemistry Consor-
tium (MCC; EPSRC EP/X035859/1). G.K.N. performed calculations on the
CLAIX-2018 supercomputer at RWTH Aachen, under an NHR-N project
with No. p0020883.

**Conflict of Interest**

The authors declare no conflict of interest.

**Data Availability Statement**

The data that support the findings of this study are available from the
corresponding author upon reasonable request.

Keywords

Intermetallic compounds (IMCs), MAX phases, MXenes, nanolaminated hexagonal solids & 2D derivatives, nanostructured solids, ZIP phases

Received: August 12, 2023

Revised: August 7, 2025

Published online:

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