# Transcending the MAX phases concept of nanolaminated early transition metal
carbides/nitrides – the ZIA phases

M.A. Tunes<EMAIL_ADDRESS>S.M. Drewry  F. Schmidt  J.A. Valdez  M.M.
Schneider  C.A. Kohnert  T.A. Saleh  C.G. Schön  S. Fensin  O. El-Atwani  N.
Goossens  S. Huang  J. Vleugels  S.A. Maloy  K. Lambrinou
<EMAIL_ADDRESS>Materials Science and Technology Division, Los Alamos
National Laboratory, United States of America  Department of Materials Science
and Engineering, University of Tennessee, United States of America  Department
of Nuclear Engineering, University of California Berkeley, United States of
America  Materials Physics and Applications Division, Los Alamos National
Laboratory, United States of America  Department of Metallurgical and
Materials Engineering, Universidade de São Paulo, Brazil  Reactor Materials
and Mechanical Design, Pacific Northwest National Laboratory, United States of
America  Department of Materials Engineering, KU Leuven, Belgium  School of
Computing and Engineering, University of Huddersfield, United Kingdom

###### Abstract

A new potential class of nanolaminated and structurally complex materials,
herein conceived as the Zigzag IntermetAllic (ZIA) phases, is proposed. A
study of the constituent phases of a specific Nb–Si–Ni intermetallic alloy
revealed that its ternary H-phase, i.e., the Nb3SiNi2 intermetallic compound
(IMC), is a crystalline solid with the close-packed fcc Bravais lattice, the
312 MAX phase stoichiometry and a layered atomic arrangement that may define
an entire class of nanolaminated IMCs analogous to the nanolaminated ceramic
compounds known today as the MAX phases. The electron microscopy investigation
of the Nb3SiNi2 compound – the first candidate ZIA phase – revealed a
remarkable structural complexity, as its ordered unit cell is made of 96
atoms. The ZIA phases extend the concept of nanolaminated crystalline solids
well beyond the MAX phases family of early transition metal carbides/nitrides,
most likely broadening the spectrum of achievable material properties into
domains typically not covered by the MAX phases. Furthermore, this work
uncovers that both families of nanolaminated crystalline solids, i.e., the
herein introduced fcc ZIA phases and all known variants of the hcp MAX phases,
obey the same overarching stoichiometric rule $P_{x+y}A_{x}N_{y}$, where $x$
and $y$ are integers ranging from 1 to 6.

###### keywords:

MAX Phases, Electron Microscopy, Nanolaminated Crystalline Solids, ZIA Phases

††journal: arXiv

a4paper,margin=20mm

## Special Dedication

This manuscript is dedicated to V. Nowotny, E. Reiffenstein, F. Benesovsky and
W. Jeitschko of the University of Vienna, Austria, who first synthesized
complex ternary carbides in the 1960/1970s, and to M.W. Barsoum of Drexel
University, USA, who in the 1990/2000s revolutionized this field with his
research on the MAX phases and the discovery of their 2D derivatives known as
MXenes in 2011. Another dedication goes to E.I. Gladyshevskii, Yu.B. Kuz’ma
and P.I. Kripyakevich of the Ivan Franko L’vovsk State University, former
USSR, for their pioneering work on complex ternary intermetallics of the
Nb3Ni2Si type in the 1960/1970s. The herein proposed “ZIA phases” are named
after the ZIA indigenous nation of New Mexico, USA, where the Los Alamos
National Laboratory is located.

##  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 building blocks of a material are confined in the 1-100 nm
range, the material can be defined as nanostructured [3]. The nanoscale
manipulation of matter through the engineering of structural modules, such as
atomic clusters, precipitates, crystallites/grains, grain boundaries,
molecules, and/or atomic layers has been an effective approach to achieve a
set of extraordinary and unique properties in nanostructured materials, well
beyond what is possible by conventional processing routes that target changes
on the micro- and macroscale [1, 2, 3, 4]. Transferring the superior
properties of nanostructured materials to the bulk scale, a process demanding
multiscale materials engineering, is one of the key challenges of contemporary
materials science [4].

H. Gleiter – the scientist who introduced the term “nanostructured materials”
[2] – predicted in 1992 [2] that this new research field would be
significantly impacted by the emergence of two classes of nanostructured
materials: ceramics and intermetallic phases. Gleiter’s prediction was partly
realized in 1996 by M.W. Barsoum and his group, who made significant strides
in the development and characterization of a class of nanolaminated ceramics
with the Mn+1AXn 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 = 1, 2, or 3; these ceramic materials are known today as the MAX phases
[5, 6]. The 1996 breakthrough by Barsoum and El-Raghy [5] consisted in the
pioneering synthesis of quasi-phase-pure Ti3SiC2 via reactive hot pressing;
notably, the existence of the Ti3SiC2 ternary carbide, which is 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 synthesis and characterization of
Ti3SiC2, i.e., the 312 ternary carbide of the Ti–Si–C system [7], Barsoum
realized early in the 2000s that the Ti3SiC2 compound and other nanolaminated
ternary carbides/nitrides form an entire new class of hexagonal close-packed
(hcp) ceramics described by the Mn+1AXn strict stoichiometric rule and
characterized by exceptional physicochemical and mechanical properties [6]. To
date, around 155 ternary carbides/nitrides of early transition metals as well
as numerous chemically complex solid solutions with the Mn+1AXn strict
stoichiometry and precisely tailored properties have been experimentally
synthesized [9, 10]. The exceptional properties of the MAX phases are
attributed to their ordered hcp atomic arrangement of M6X ceramic-like
octahedra and intercalated A metallic-like layers, which results in
nanolaminated unit cells with long-range translational symmetry. Since their
inception, the MAX phases [11, 12] and their 2D derivatives known as MXenes
[13, 14, 15, 16, 17] have been proposed for various applications in diverse
industrial/technological fields, including energy storage, fuel cells,
biomedical, photonics/phononics [18, 19, 20, 21, 22, 23, 24, 25, 26, 27, 28,
29, 30], as well as in advanced nuclear fusion and fission reactor systems
[31, 32, 33, 34, 35, 36, 37, 38, 39, 40, 41, 42].

This work focuses on a single member of a potential class of nanolaminated
intermetallic compounds (IMCs), herein named the “zigzag intermetallic” phases
or simply the ZIA phases. The first candidate ZIA phase introduced in this
work is the Nb3SiNi2 ternary IMC, which is considered the H-phase of the
Nb–-Si–-Ni system [43, 44, 45, 46]. This work reports the experimental
synthesis of quasi-phase-pure Nb3SiNi2 via reactive hot pressing and also
addresses, for the first time, the aspects of crystal structure nanolamination
and ‘zigzag’ atomic arrangement in the Nb3SiNi2 IMC. At first glance, the ZIA
phases share many similarities with the MAX phases, such as crystal structure
nanolamination and obeying the Mn+1AXn stoichiometric rule with n = 2;
however, their true potential cannot yet be reliably assessed due to the
scarcity of available data. We believe that the nanolaminated ZIA phases and
their possible 2D derivatives might have a potential commensurate with that of
the nanolaminated MAX phases and their 2D derivatives known as MXenes. In case
future research validates this hypothesis, then a whole new research field
will have been inaugurated and the prediction of Gleiter [1, 2] on the
existence of ‘nanostructured IMCs’ will have come true, at the same time.

##  2 Results and discussion

###  2.1 Synthesis of phase-pure Nb3SiNi2

Initially, arc melting was employed to produce bulk materials made of the
Nb3SiNi2 IMC ZIA phase. The microstructure of the as-cast Nb–Si–Ni
intermetallic alloy was first studied by SEM with the aid of a backscattered
electron (BSE) detector. The compositional contrast of the BSE image in Fig.
1A indicates the presence of four distinct phases in the as-cast alloy
microstructure, a finding supported by both EDX elemental mapping and XRD.
Phase identification by XRD (Fig. 1C) revealed the (metastable) co-existence
of the Nb3SiNi2 H-phase (candidate ZIA phase), the Nb7Ni6 $\mu$-phase, the
Ni3SiNb2 ternary Laves (L) phase and the Nb4NiSi T-phase. These four phases
were peak-indexed in accordance with data provided by the ICSD database [47]
and were also identified as possible phases during the experimental
investigation of phase equilibria in the Nb–Si–Ni refractory alloy system
[46]. The nominal stoichiometries and measured EDX elemental compositions of
the four phases in the as-cast alloy are provided in Table 1, showing that all
phases formed as imperfect (off-stoichiometric) solid solutions during melt
solidification rather than as stoichiometric compounds. The formation of solid
solutions in the Nb–Si–Ni system has been previously reported by dos Santos et
al. [46]; a strong indication of solid solution formation is encountered in
the dissolution of significant amounts of Si in the Nb7Ni6 $\mu$-phase, as
previously reported [46] and also seen in Table 1. Even though arc-melting
targeted the formation of bulk materials (discs) made primarily of the
Nb3SiNi2 IMC ZIA phase, the presence of four distinct phases in the as-cast
alloy microstructure is a clear violation of the Gibbs’ phase rule for a
ternary alloy [48], indicating that thermodynamic equilibrium has not been
reached via arc melting synthesis. This is most likely caused by the loss of
substantial amounts of Ni, which is volatile in vacuum above its melting point
(1728 K). By looking at the CALPHAD phase equilibria calculated in this work
(Fig. 2C) in conjunction with the microstructure of the as-cast alloy (Fig.
1A), it is reasonable to assume that the (rapid) solidification of the melt
produced solid solutions of three main phases, i.e., the Nb3SiNi2 H-phase, the
Nb7Ni6 $\mu$-phase and the Ni3SiNb2 L-phase, while the Ni-impoverished
residual melt formed the Nb4NiSi T-phase (see the 1423 K isotherm in Fig. 2C).
This hypothesis is supported by the small fraction of the Nb4NiSi T-phase in
the as-cast alloy (Fig. 1A).

Table 1: Nominal stoichiometries vs. actual compositions of the phases
identified by SEM/EDS in the arc-melted (as-cast & annealed) and RHP Nb–Si–Ni
intermetallic alloy samples.

Arc-Melted Nb–Si–Ni Alloy Samples
---
Phases Identified | Stoichiometry | Nb [at.%] | Si [at.%] | Ni [at.%]
H-phase (as-cast) |  Nb3SiNi2 (ZIA phase) | 46.5$\pm$3.2 | 18.4$\pm$4.0 | 35.1$\pm$1.9
H-phase (annealed) | 46.4$\pm$3.0 | 18.6$\pm$3.7 | 35.0$\pm$1.5
H-phase (expected) | 50.0 | 17.0 | 33.0
Laves phase (as-cast) | Ni3SiNb2 | 31.9$\pm$2.7 | 21.8$\pm$3.5 | 46.3$\pm$1.8
Laves phase (annealed) | 31.3$\pm$2.4 | 20.1$\pm$3.1 | 48.6$\pm$1.2
Laves phase (expected) | 33.0 | 17.0 | 50.0
$\mu$-phase (as-cast) | Nb7Ni6* | 45.2$\pm$3.1 | 13.5$\pm$3.8 | 41.3$\pm$1.9
$\mu$-phase (annealed) | 43.8$\pm$3.0 | 13.5$\pm$3.7 | 42.7$\pm$1.5
$\mu$-phase (expected) | 53.8 | - | 46.2
T-phase (as-cast) | Nb4NiSi | 62.8$\pm$3.8 | 25.8$\pm$4.7 | 11.4$\pm$2.5
T-phase (annealed) | 62.6$\pm$3.8 | 25.3$\pm$4.7 | 12.0$\pm$1.9
T-phase (expected) | 66.0 | 17.0 | 17.0
RHP Nb–Si–Ni Alloy Samples
Sintering T [K] | Phases Identified | Nb [at.%] | Si [at.%] | Ni [at.%] | O [at.%]
1523 | H-Nb3SiNi2 | 53.86$\pm$0.16 | 12.71$\pm$0.03 | 33.43$\pm$0.12 | -
L-Ni3SiNb2 | 38.60$\pm$0.12 | 17.63$\pm$0.03 | 43.77$\pm$0.12 | -
$\mu$-Nb7Ni6 | 71.99$\pm$0.13 | 17.17$\pm$0.03 | 10.85$\pm$0.07 | -
1623 | H-Nb3SiNi2 | 53.57$\pm$0.15 | 13.26$\pm$0.03 | 33.17$\pm$0.12 | -
L-Ni3SiNb2 | 36.25$\pm$0.15 | 17.49$\pm$0.04 | 46.27$\pm$0.16 | -
$\mu$-Nb7Ni6 | - | - | - | -
1723 | H-Nb3SiNi2 | 52.80$\pm$0.17 | 13.12$\pm$0.04 | 34.08$\pm$0.14 | -
L-Ni3SiNb2 | 36.32$\pm$0.11 | 17.47$\pm$0.03 | 46.21$\pm$0.11 | -
$\mu$-Nb7Ni6 | 65.96$\pm$0.21 | 8.43$\pm$0.04 | 4.65$\pm$0.09 | 20.96$\pm$0.05

* Si has been reported to dissolve in the binary Nb7Ni6 IMC, forming $\mu$-phase solid solutions [46].

The homogenization annealing at 1421 K was performed, first and foremost, to
investigate whether the four-phase assembly of the as-cast Nb–Si–Ni
intermetallic alloy would converge to an alloy with a higher fraction of the
Nb3SiNi2 IMC ZIA phase. The annealing aimed also at a first assessment of the
thermal stability of the Nb3SiNi2 ZIA phase of interest in this work. The
SEM/EDS/XRD analysis of the annealed alloy (Figs. 1B & 1C) showed that it
still consisted of the same four phases with a slightly increased fraction of
the Nb3SiNi2 H-phase that appears able to sustain these conditions; the
increase of the latter was not sufficient to suggest that the selected
annealing conditions were capable of transforming the phase assembly of the
as-cast alloy into an alloy primarily composed of the Nb3SiNi2 IMC. This
suggests that solid-state diffusion at 1421 K is relative slow, and that
annealing treatments much longer than 336 h would be required to achieve alloy
homogenization at 1421 K; moreover, the fact that the alloy composition after
solidification of the melt appears to be off the Nb3SiNi2 stoichiometry
(primarily due to the loss of Ni) will prevent the formation of a phase-pure
Nb3SiNi2 material, irrespective of the annealing duration and temperature.
Fig. 1B reveals that the microstructure of the annealed alloy was slightly
modified, suggesting limited coarsening and change of grain morphology (e.g.,
the Nb7Ni6 $\mu$-phase grains appear more angular and equiaxed in the annealed
alloy than in the as-cast alloy). The limited microstructural changes indicate
that the thermodynamically metastable as-cast alloy strives to achieve
equilibrium during isothermal annealing, forming more Nb3SiNi2 and
transitioning towards a slightly coarser microstructure, however, the imposed
annealing conditions are not conducive to achieving equilibrium. The elemental
compositions of the four phases in the annealed alloy, as determined by EDX
analysis, are provided in Table 1. Additional BSE images of the
microstructures of both as-cast and annealed arc-melted samples are given in
Fig. S1 (Supplementary Information). The unsuccessful attempt to synthesize
phase-pure Nb3SiNi2 via arc melting showed that this processing route is
difficult to master, mainly due to the poorly controlled losses of the
constituting elements above their melting points (e.g., Tm(Ni) = 1728 K,
Tm(Nb) = 2750 K, Tm(Si) = 1683 K). Moreover, it became clear that it is
necessary to calculate the equilibrium phase mixtures at the annealing
temperature (1421 K) and at higher temperatures (in order to understand melt
solidification), as prior studies have only examined the 1073, 1323 and 1473 K
isotherms of the Nb–Si–Ni system [43, 44, 45, 46].

The results of this work contradict the early work of Gladyshevskii et al.
[43], who reported that by melting a feedstock of elemental powders under an
inert gas atmosphere (presumably, argon), they were able to produce quasi-
phase-pure Nb3SiNi2 with only a small amount of a ternary Laves phase
(presumably, the Ni3SiNb2 ternary 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
(atomically ordered) ternary IMCs with the H-phase structure (space group
$Fd\bar{3}m$) and the R3SiNi2 general stoichiometry, with R being the early
transition metals Mn, Cr, V, Nb and Ta [43]. Apart from demonstrating a much
better command of the arc-melting processing route, the 1964 pioneering work
by Gladyshevskii et al. [43] paved the way for the study of the ZIA phases
proposed in this work. In a similar manner, the 1967 work by Jeitschko et al.
[7] opened the way for the further development of the class of early
transition metal ternary carbides/nitrides known as the MAX phases almost 40
years later by Barsoum et al. [5, 6].

A thorough thermodynamic assessment of the 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. [44, 45, 46]. In an attempt to reproduce the
results of Gladyshevskii et al. [43], dos Santos et al. [45] synthesized 12
different alloy compositions in the Nb–Si–Ni system via arc melting and
subsequently homogenized the as-cast samples at 1073 K for longer times (1000
h). Curiously, the Nb3SiNi2 phase was not found in any of the 12 homogenized
samples characterized by means of SEM/EDX and XRD. It is presently not
understood whether the difference in atmosphere during melting – reducing (H2)
in Gladyshevskii et al. [43] and possibly slightly oxidizing (argon, Ar, of
unknown purity) in dos Santos et al. [44, 45, 46] – is accountable for the
difference in the produced arc-melted alloys. Based on their data, dos Santos
et al. suggested a homogenization isothermal annealing at higher temperatures
to explore the thermodynamic equilibrium in this system [45]. In a more
extensive follow-up study, dos Santos et al. 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 [46]. Using SEM/EDX
and XRD after homogenization annealing, the Nb3SiNi2 phase was only observed
in seven samples, but none of these were phase-pure Nb3SiNi2. All 7 samples
exhibited triphasic microstructures, each comprising two or three of the four
principal phases, i.e., the ternary Laves, $\mu$, T- and H-phases; small
fractions of the NbNi3 IMC or the unknown Z phase were also reported in 3 of
the 7 samples. The Z phase is one of three phases with unresolved crystal
structure – tentatively designated X, Y and Z – found by dos Santos et al.
after annealing select Nb-Si-Ni alloy compositions at 1473 K for 504 h [46].

Careful consideration of both the prior attempts by dos Santos et al. [45, 46]
and the outcome of this work demonstrated that the synthesis of phase-pure
Nb3SiNi2 via arc melting is challenging. Even though Gladyshevskii et al. [43]
reported that the production of quasi-phase-pure Nb3SiNi2 is possible via the
melting of an elemental powder feedstock with identical composition 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, such as the ternary IMCs designated as the ZIA phases
in this work, or the ternary carbides/nitrides known as the MAX phases [49].
As already mentioned, one important challenge is the evaporation of alloying
elements in the vacuum chamber [50]. Given the strict stoichiometry of the ZIA
phases and the 312 MAX phases, the processing route as well as the starting
powder feedstock must be always carefully adjusted to produce materials with
high phase purity. For example, Goossens et al. [51] explained rather recently
why the use of early transition metal hydride powders produces MAX phase-based
ceramics with higher phase purities than the use of elemental starting
powders. These authors proved that the use of metal hydride powders lowers the
formation temperature of MAX phase precursor phases (e.g., binary IMCs),
increases the formation temperature of competing phases (e.g., binary
carbides), and suppresses the oxidation of oxidation-prone powders (e.g., Zr)
by providing a reducing (H2-rich) atmosphere caused by powder dehydrogenation,
thus allowing 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 & Si) to produce phase-pure Nb3SiNi2. The synthesis efforts proved
to be successful at three different sintering temperatures (i.e., 1523 K, 1623
K, and 1723 K), producing quasi-phase-pure Nb3SiNi2 with a small fraction of
the ternary Ni3SiNb2 Laves phase, and tiny amounts of oxidic inclusions (NbO &
NiO), presumably stemming from the limited oxidation of the raw powders. BSE
images of the microstructures and WDXS elemental maps of all quasi-phase-pure
Nb3SiNi2 materials are shown in Fig. 2B, next to their XRD patterns (Fig. 2A).
Additional BSE micrographs and EDS elemental maps of all RHP alloy samples are
shown in Fig. S2 (Supplementary Information).

The results of CALPHAD calculations of the 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 K, 1623 K and 1723 K are
shown in Fig. 2C. These calculations suggest that melt solidification during
vacuum arc melting produced mainly a ternary phase mixture comprising the
L-Ni3SiNb2 line compound, the H-Nb3SiNi2 point compound (targeted ZIA phase),
and the $\mu$-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 (Fig. 2C). 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 (targeted ZIA phase) and L-Ni3SiNb2 compounds.
The NbO and NiO oxide impurities present in the RHP alloys are not encountered
in the arc-melted alloys due to the much higher processing temperatures
associated with the complete melting of all raw elemental powders during arc
melting synthesis.

Figure 1: SEM/EDX and XRD characterization of the arc-melted Nb–Si–Ni
intermetallic alloy in the as-cast state and after homogenization annealing at
1421 K for 336 h. BSE imaging and EDX elemental mapping of both as-cast (A)
and annealed (B) samples show that the (thermodynamically metastable) arc-
melted alloy comprises four distinct phases, i.e., H-Nb3SiNi2, L-Ni3SiNb2,
$\mu$-Nb7Ni6 and T-Nb4SiNi. (C) XRD patterns of the as-cast and annealed alloy
samples confirm the presence of these four phases. The phase-of-interest,
i.e., the ternary H-phase Nb3SiNi2 IMC, is highlighted in both BSE micrographs
and XRD spectra. Figure 2: XRD (A) and SEM/WDXS (B) characterization of the
RHP Nb–Si–Ni intermetallic alloy, showing that all alloy samples comprise
mainly H-Nb3SiNi2 and L-Ni3SiNb2, with small fractions of oxidic inclusions
(NbO, NiO) presumably stemming from the limited powder particle oxidation.
Microstructural homogeneity and phase purity (in terms of the desired Nb3SiNi2
ZIA phase) seem to increase with sintering temperature. (C) CALPHAD
calculations of phase equilibria in the Nb–Si–Ni system on the 1423 K, 1523 K,
1623 K and 1723 K isotherms.

###  2.2 TEM/STEM characterization of the Nb3SiNi2

Even though arc melting produced a bulk sample with a low Nb3SiNi2 content,
FIB was used to lift out thin foils from sufficiently large Nb3SiNi2 grains in
the annealed sample (Fig. 3A) for characterization of the Nb3SiNi2 ZIA phase
by means of (S)TEM/SAED. This characterization was decided before quasi-phase-
pure bulk samples were produced via RHP, so as to gain a first in-depth
understanding of the crystal structure of the Nb3SiNi2 ZIA phase. A bright-
field STEM (BF-STEM) cross-sectional micrograph of such a Nb3SiNi2 grain is
shown in Fig. 3A. The defects observed within the first 500 nm beneath the
sample surface (protected by a combined Pt/C strip during FIB lift-out) are
attributed to mechanical damage induced during grinding/polishing of the bulk
annealed sample. The presence of sub-surface mechanical damage did not affect
the chemistry of the analyzed grain, which was found by STEM/EDS to be
compositionally homogeneous and close to the Nb3SiNi2 stoichiometry with
limited deviations from the expected Ni (5 at.% excess) and Si (5 at.%
deficiency) contents (Fig. 3B).

Figure 3: STEM/EDX analysis of a FIB foil lifted-out from a large Nb3SiNi2 IMC
grain in the annealed arc-melted alloy. (A) BF-STEM micrograph of the Nb3SiNi2
IMC grain showing sub-surface mechanical damage caused by sample
grinding/polishing. The inset is a BSE image of the Nb3SiNi2 grain selected
for this FIB foil lift-out. (B) STEM/EDS analysis of the area inside the green
rectangle of Fig. 3A shows the homogeneous distribution of Nb, Si and Ni
(Nb/Si/Ni elemental maps, bottom), confirming that the average chemical
composition of this grain within the analyzed area deviates slightly from the
Nb3SiNi2 stoichiometry (Nb/Si/Ni profiles, top). Figure 4: STEM/SAED
characterization of the single Nb3SiNi2 IMC grain of Figure 3A. (A) High-
resolution BF-STEM micrograph of the Nb3SiNi2 IMC grain viewed in the low-
index $[\bar{1}12]$ zone axis. (B) Experimental SAEDP and (C) Fast Fourier
Transform (FFT) pattern of the $[\bar{1}12]$ zone axis of the Nb3SiNi2 IMC.
(D) Simulated SAEDP of the $[\bar{1}12]$ zone axis of the Nb3SiNi2 IMC; the
inset presents the unit cell of the Nb3SiNi2 in the $[\bar{1}12]$ zone axis.
The simulated SAEDP was produced using SingleCrystal [52] and experimental
data that are available in the ICSD database [43]. The experimental and
simulated SAEDPs of Figures 4B and 4D, respectively, confirm that the Nb3SiNi2
candidate ZIA phase has the diamond cubic crystal structure (fcc Bravais
lattice; $Fd\bar{3}m$; space group 227).

Given the rather satisfactory chemistry match with the targeted Nb3SiNi2
stoichiometry, further characterization of the same Nb3SiNi2 grain by STEM &
SAED (selected area diffraction) was deemed necessary to elucidate the crystal
structure. Figure 4A is an atomically resolved BF-STEM image of the Nb3SiNi2
ZIA phase, showing the nanolayered atomic arrangement in a “zigzag” pattern,
which was observed in specific low-index crystallographic orientations. Figure
4B shows a SAEDP recorded after orienting the sample onto the low-index
[$\bar{1}12$] zone axis. Using published crystallographic data (CIF file code
ICSD-2044323 [43]) and the SingleCrystal software package [52], a kinematic
simulation of the SAEDP (Fig. 4D) was made and compared with the recorded
SAEDP (Fig. 4B). Experimental and simulated SAEDPs were in excellent
agreement, whilst all collected data suggested that the crystal structure of
the experimentally synthesized Nb3SiNi2 is diamond cubic ($Fd\bar{3}m$, space
group 227) [43]. This crystal structure characterizes not only the ternary
H-phase (Nb3SiNi2) of the Nb–Si–Ni system [43, 45, 46], but also the ternary
IMCs with identical stoichiometry in the Mn–Si–Ni (Mn3SiNi2) [43, 53, 54],
V–Si–Ni (V3SiNi2) [43], Cr–Si–Ni (Cr3SiNi2) [43], Ta–Si–Ni (Ta3SiNi2) [43,
55], Nb–Si–Co (Nb3SiCo2) [56], Na–Au–In (Na3AuIn2) [12], Na–Ag–In (Na3AgIn2)
[12], and Mg–Ga–Ni (Mg3GaNi2) [57] systems, apparently being the preferred
crystal structure for the entire class of the (312) ZIA phases, just like the
hexagonal close-packed (hcp; $P6_{3}/mmc$, space group 194) crystal structure
encompasses all MAX phase compounds [6, 10]. Having said that, it is also
worthwhile adding that few IMCs with the 312 stoichiometry and the hcp
($P6_{3}/mmc$, space group 194) crystal structure have been sporadically
identified, possibly suggesting the existence of a not yet fully identified
and/or explored overarching trend in the formation of nanolaminated solids,
such as the known MAX phases and the herein proposed ZIA phases. Examples of
ternary 312 IMCs with the hcp crystal structure have been previously reported
in the Co–Si–Nb (Co3SiNb2) [56], Ni–Si–Nb (Nb3SiNi2) [43] and Cu–Si–Mg
(Cu3SiMg2) [56] systems. In view of the scarcity of available data, it would
be premature to speculate on the relation between the fcc ternary IMCs
identified in this work as the ZIA phases and the largely unexplored hcp
ternary IMCs, or to assume that the latter are also characterized by crystal
structure nanolamination, even though such a hypothesis seems quite plausible.

###  2.3 Opportunities arising from the advent of the ZIA phases

In the search for new refractory materials with exceptional properties, both
ZIA and MAX phases are nanolaminated crystalline solids that share the same
historical roots, dating back to the advent of space exploration in the 1960s
[7, 8, 58, 43, 45, 46]. 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 IMC ZIA phases and their
(not yet synthesized) 2D derivatives, as opposed to the extensively studied
MAX phase ternary carbides/nitrides, their higher order solid solutions and 2D
derivatives known as MXenes. Therefore, this section will attempt to describe
the basic differences and similarities of the rather mature MAX phases vs. the
emerging ZIA phases, discussing potential applications of the latter based on
the experience gained through the MAX-phase research over the past 30 years.

Using data compiled by the Materials Genome Project [59], the first candidate
ZIA phase, i.e., the Nb3SiNi2 IMC, is herein discussed with respect to the
extensively studied Ti3SiC2 MAX phase. First, they are both ternary compounds
that belong to different elemental systems, i.e., Nb3SiNi2 forms in the
Nb–Si–Ni system, while Ti3SiC2 in the Ti–Si–C system. Even though both
compounds obey the 312 MAX phase stoichiometric rule, i.e., Mn+1AXn with n =
2, Nb3SiNi2 is clearly an IMC, while Ti3SiC2 is classified as a ceramic.
Moreover, there is a remarkable difference in crystal structure complexity,
with the unit cell of the Nb3SiNi2 ZIA phase (Fig. 5A) being significantly
larger than the unit cell of the Ti3SiC2 MAX phase (Fig. 5B). More
specifically, the Ti3SiC2 unit cell is made of 6 atoms, while the Nb3SiNi2 one
is made of 96 atoms, resulting in unit cell volumes that differ in size by one
order of magnitude (Table S1 in the Supplementary Information). Both ternary
compounds are characterized by centrosymmetric close-packed Bravais lattices,
i.e., the Ti3SiC2 MAX phase is hcp ($P6_{3}/mmc$, space group 194) whilst the
Nb3SiNi2 ZIA phase is fcc (face-centered cubic; $Fd\bar{3}m$, space group
227); apparently, such close-packed atomic configurations seem to favor
crystal structure nanolamination.

Figure 5: Comparison of the unit cells of (A) the herein synthesized and
characterized fcc Nb3SiNi2 ZIA phase with (B) the well-known hcp Ti3SiC2 MAX
phase [8, 5]. Fig. 5A (bottom) shows 1 unit cell (96 atoms/unit cell) of the
Nb3SiNi2 ZIA phase, as well as the [110] (upper left) and [111] (upper right)
zone axes using 4 unit cells. Red arrows pinpoint atomic zigzagging in the
[110] zone axis, yellow arrows indicate crystal structure nanolamination. Fig.
5B (bottom) shows 8 unit cells of the Ti3SiC2 MAX phase (6 atoms/unit cell),
as well as the $[11\bar{2}0]$ (upper left) and [0001] (upper right) zone axes.
(C) The ZIA phases extend the concept of crystal structure nanolamination
beyond the early transition metal carbides/nitrides known as the MAX phases.
Together, the two classes of nanolaminated solids span almost the whole period
table, comprising jointly elements from groups 1-16. The M-elements of the MAX
phases are shown using 2 shades of red: the darker indicates full occupation
of the M-site, the lighter indicates partial occupation (solid solutions). The
same holds for the 2 shades of blue dedicated to the A-elements of the MAX
phases. Note: the periodic table in Fig. 5C is a modified version of the
original American Chemical Society’s periodic table.

A groundbreaking impact of the herein proposed ZIA phases is that they
actually extend the concept of crystal structure nanolamination much beyond
the early transition metal carbides/nitrides known as the MAX phases. As may
be seen in the color-coded periodic table of elements of Fig. 5C, the
M-elements used to form ternary MAX phases are early transition metals (groups
3-7), the A-elements belong mainly to groups 13-15 (post-transition metals &
metalloids) with limited expansion to groups 8-12 (late transition metals) and
group 16 (reactive non-metals), while the X-element is either C or N (reactive
non-metals in groups 14 and 15, respectively). On the other hand, the ZIA IMCs
that have been hitherto identified utilize Z-elements from groups 1 (alkali
metals) and 2 (alkaline earth metals) and groups 5-7 (early transition
metals), I-elements from group 11 (late transition metals) and groups 13 and
14 (post-transition metals & metalloids), while the A-elements belong either
to groups 9 and 10 (late transition metals) or group 13 (the post-transition
metal indium, In). In short, the synthesis of nanolaminated crystalline solids
(MAX and ZIA phases) requires different combinations of constituent elements
from groups 1-16 in the periodic table. This heralds the potential fabrication
of ternary compounds and higher-order solid solutions with unique properties
and very broad applicability, not only in the form of the 3D ZIA and MAX
phases, but also in the form of their 2D derivatives, which have so far only
been explored for the MAX phases in the form of the rapidly developing MXenes.
Especially the discovery of ZIA phases comprising elements that have hitherto
not been introduced in the MAX phases (e.g., alkali and alkaline earth metals,
several late transition metals) paves the way for the fabrication of new
nanolaminated crystalline solids with certain properties (e.g., ferromagnetic
properties, enhanced/tailorable ductility, superconductivity, etc.) that are
typically not associated with the MAX phases. This hypothesis finds support in
the work of Kolenda et al. [54], who reported on the magnetic properties of
the Mn3SiNi2 candidate ZIA phase. The new range of properties achieved by the
ZIA phases is also expected to be reflected in their 2D derivatives (assuming
these can first be produced experimentally), probably giving rise to the next
generation of 2D materials for use in flexible and portable microelectronics,
batteries, sensors, etc.

It is worthwhile noting that in all hitherto reported fcc ternary IMCs now
considered as candidate ZIA phases, the electronegativity, $\chi$, of the
Z-element is invariably smaller than the electronegativities of both I- and
A-elements (Table S2 in the Supplementary Information). Moreover, the atomic
radius of the Z-element is larger than the atomic radii of both I- and
A-elements with the exception of Cr3SiNi2, where the atomic radius (125 pm) of
the Z-element, Cr, is equal to that of the A-element, Ni (Table S2 in the
Supplementary Information). Interestingly, the structurally unexplored 312 hcp
ternary IMCs Co3SiNb2 and Ni3SiNb2 could be considered as ‘compositionally
inverted’ Nb3SiCo2 and Nb3SiNi2 ZIA phases, respectively, produced by
interchanging the Z- and A-elements while still respecting the strict
stoichiometric rule of the 312 MAX phases. For these two hcp IMCs as well as
for the Cu3SiMg2 one, the electronegativity of the third element (Nb or Mg) is
smaller than the electronegativities of the first two elements (Co, Ni or Cu).
Also, the atomic radius of the third element (Nb or Mg) is larger than the
atomic radii of the first two elements (Co, Ni or Cu). Based on these
structural considerations and the fact that the hcp ternary IMCs Co3SiNb2,
Ni3SiNb2 and Cu3SiMg2 are characterized by the same hexagonal close-packed
lattice ($P6_{3}/mmc$, space group 194) as the MAX phases, one could indeed
expect crystal structure nanolamination in these ternary IMCs as well. This
hypothesis, however, merits further dedicated investigation.

A unique feature of the ZIA phases is that the I- and A-elements could be
chosen from the same group in the periodic table. For example, both the
A-element (In) in Na3AuIn2, and Na3AgIn2 as well as the I-element (Ga) in
Mg3GaNi2 belong to group 13. This freedom in the choice of the elements making
up the ZIA phases does not characterize the MAX phases, where only two
X-elements are allowed (i.e., C or N), and these are not known to be replaced
by any of the candidate M- and A-elements. The only known exception of an
element that has been rather flexible in occupying the M- or A-site of
specific MAX phase compounds is the late transition metal iron (Fe), which
belongs to group 8 of the periodic table. More specifically, Fe has been
fairly recently reported to form the (Ti1-x,Fex)3AlC2 solid solutions on the
M-site [60], or to occupy fully the A-site of the ternary MAX phases Ta2FeC,
Ti2FeN and Nb2FeC [61]. With respect to the ZIA phases, the allowable
combinations of elemental constituents could probably be associated with
differences in the electronegativity and atomic radius values of the Z-, I-
and A-elements; however, before definitive rules governing the ZIA phase
formation can be proposed, further systematic research is needed to better map
– in terms of attainable chemical compositions – the ZIA phases ‘landscape’.

Even though the main objective of this work was to introduce the class of
nanolaminated ZIA phases via the experimental synthesis and characterization
of its Nb3SiNi2 first candidate member as well as to demonstrate the
feasibility of producing phase-pure materials based on the Nb3SiNi2 ternary
IMC, an overarching stoichiometric rule governing both the MAX and the ZIA
phases was discovered and is proposed here. As pointed out by Goossens et al.
in a rather recent review [10], the nanolaminated MAX phases are mainly
encountered either as the three well-known Mn+1AXn phase orders 211, 312 and
413, or as the following exceptional structures: (i) the Mn+1AXn-like
structures with n $>$ 3, such as Ta6AlC5 and Ti7SnC6, and (ii) the “hybrid”
orders M5A2C3 and M7A2C5, such as Ti5A2C3 with A = Si, Ge and Al. To date, the
“hybrid” orders M5A2C3 and M7A2C5 have been considered as combinations of a
312 atomic stacking with a 211 and a 413 order, respectively, as they did not
obey the generic Mn+1AXn phase stoichiometry. However, the careful
consideration of both the 312 ZIA phases and all the MAX phase variants,
including the “hybrid” orders and the only identified member of the 514 order,
i.e., i.e., the Mo4VAlC4 compound [62], revealed that both classes of
nanolaminated crystalline solids with fcc and hcp close-packed lattices can be
described by the general stoichiometric rule Px+yAxNy (PAN), where x = 1 or 2,
and y = 1, 2, 3, 4, 5, or 6, according to all the ‘PAN phases’ that were
experimentally synthesized to date. The herein proposed global stoichiometric
rule governs the constitution of all nanolaminated crystalline solids with
atomically close-packed lattices, such as the fcc ZIA phases, the hcp MAX
phases and the ternary hcp IMCs, such as Co3SiNb2, Ni3SiNb2 and Cu3SiMg2, even
though crystal structural nanolamination has still not been confirmed for the
latter. Moreover, this global stoichiometric rule reconciles the “hybrid”
orders M5A2C3 and M7A2C5 with the common MAX (Mn+1AXn) phase orders 211, 312
and 413, providing an affirmative answer to the question whether it was
conceptually correct to include all identified MAX phase variants in the same
class of materials.

A major challenge addressed in this work was to demonstrate that the synthesis
of phase-pure bulk ZIA phase-based materials is feasible. Producing phase-pure
bulk materials is a prerequisite for the systematic investigation of material
properties (e.g., elastoplastic and fracture behavior; oxidation and corrosion
resistance; electrical and magnetic properties; radiation tolerance; etc.) or
the possible applications of the ZIA phases and their 2D derivatives, provided
that these can be extracted. Even though the production of phase-pure Ni3SiNb2
via arc melting proved extremely challenging, presumably due to the
volatilization of Ni and/or its compounds above a certain temperature, the
synthesis of quasi-phase-pure bulk Ni3SiNb2 via RHP was successful. Similar
challenges in producing phase-pure bulk materials have been previously
reported in association with the synthesis of specific MAX phases, such as the
Zr- and Hf-based MAX phases [63, 19, 20, 37, 64, 39]; typically, overcoming
these challenges involved the synthesis of complex double solid solutions
based on steric stability criteria [65, 39], or the use of early transition
metal hydride raw powders rather than elemental ones [51]. Apart from RHP,
several other processing methods will be explored to synthesize phase-pure ZIA
phases in various forms (bulk, coatings) so as to make progress with
structural characterization, property determination and possible applications.
Physical vapor deposition (PVD) coating techniques, such as magnetron
sputtering, are expected to be able to deposit coatings of very precise
compositions, such as the herein studied Ni3SiNb2 ZIA phase onto substrates of
choice, suppressing the formation of competing phases and producing phase-pure
coatings with strong textures that might be desired by certain high-tech
applications.

##  3 Conclusions

Nanostructured materials have revolutionized modern materials science, even
though their potential is far from being fully exploited for the benefit of
the Society. Nanolaminated ceramics, such as the MAX phases and their 2D
derivatives known as MXenes, already changed the way we perceive materials
design and functionalization on the nanoscale, thereby producing innovative
materials that are characterized by a wide variety of unique properties,
precisely tailored to the requirements of the targeted applications.
Identifying market niches that could use nanostructured materials, including
applications in extreme service environments (e.g., high temperatures and
pressures, corrosive media, irradiation, etc.), is an important driver towards
the commercial deployment of innovative materials with sophisticated
nanodesigns and carefully engineered functionalities. In this work, we propose
a novel class of nanolaminated crystalline solids, the ‘zigzag intermetallic’
phases or, simply, the ZIA phases. The herein introduced ZIA phases are
characterized by an fcc ($Fd\bar{3}m$, space group 227) close-packed lattice
of appreciable structural complexity, while obeying the stoichiometric rule of
the 312 MAX phases, i.e., Mn+1AXn with n = 2. The ZIA phases and all known MAX
phase orders, including the “hybrid” M5A2C3 and M7A2C5 ones mentioned above,
appear to respect the herein proposed global stoichiometric rule Px+yAxNy,
where x and y are integers in the 1-6 range.

The first candidate ZIA phase, i.e., the Nb3SiNi2 IMC, is identified as the
H-phase of the Nb–Si–Ni system and has been found in bulk samples fabricated
by both arc melting (low Nb3SiNi2 content) and RHP (quasi-phase-pure
Nb3SiNi2). The STEM/SAED characterization of Nb3SiNi2 confirmed the expected
diamond cubic crystal structure, revealing a remarkable structural complexity,
as reflected in an exceptionally large unit cell made of 96 atoms with
nanolayered arrangement. Considering that, apart from the Nb3SiNi2 IMC, other
ternary 312 IMCs (i.e., Mn3SiNi2, V3SiNi2, Cr3SiNi2, Ta3SiNi2, Nb3SiCo2,
Na3AuIn2, Na3AgIn2, and Mg3GaNi2) with the same diamond cubic fcc lattice have
already been experimentally synthesized, it is reasonable to assume that these
IMCs constitute an entire class of nanolaminated crystalline solids, herein
designated as the ZIA phases. Confirming the nanolaminated nature of the other
ternary 312 IMCs is the aim of ongoing investigations.

The ZIA phases extend the concept of crystal structure nanolamination beyond
the early transition metal carbides/nitrides known as the MAX phases. Despite
sharing similarities and the same historical roots with the MAX phases, ZIA
phases such as the Nb3SiNi2 IMC are characterized by appreciably more complex
crystal structures and, by extension, configurational entropies than ternary
MAX phases such as the Ti3SiC2 ceramic. Therefore, the nanolaminated ZIA
phases are conceptually positioned, in terms of characteristics and potential
applications, between the nanolaminated MAX phases (with unique hybrid
metallic/ceramic properties, and proven capability of producing 2D
derivatives, known as MXenes, with tailored functionalities) and other dense
and complex multicomponent alloys, such as high-entropy alloys (HEAs)
characterized by large chemical complexity and remarkable radiation tolerance.
This work succeeded in synthesizing phase-pure bulk Nb3SiNi2 via reactive hot
pressing (RHP), which is considered an important achievement as the production
of phase-pure ZIA phase-based materials is not only a prerequisite for the
(mechanical, electrical, magnetic, etc.) property determination of these
largely unknown compounds, but also to explore the possibility of producing
and characterizing their (potential) 2D derivatives, as well as assessing ways
to tailor their properties (e.g., ductility) in an application-driven manner.
It is likely that the new class of ZIA phases will create a new ecosystem of
nanolaminated/nanostructured (3D/2D) materials with amazing new properties and
fascinating applications in diverse technological areas. The first step
towards the exploration of this hitherto unexplored ecosystem is the
successful synthesis of many materials with high phase purity, starting from
the already identified candidate ZIA phases (e.g., Nb3SiNi2, Mn3SiNi2,
V3SiNi2, Cr3SiNi2, Na3AuIn2, Mg3GaNi2, etc.) and proceeding with theoretically
predicted (e.g., by means of thermodynamic modeling, density functional
theory, etc.) but not yet experimentally made IMCs. The second step in this
exploratory work should include attempts to produce 2D derivatives of the ZIA
phases, as well as ternary hcp IMCs (e.g., Co3SiNb2, Ni3SiNb2, Cu3SiMg2,
etc.), in order to find out whether they are also characterized by crystal
structure nanolamination, as expected, or not.

##  4 Experimental Section

###  4.1 Synthesis 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 Nb3SiNi2 stoichiometry, i.e.,
50.0Nb-16.7Si-33.3Ni111The term “Nb–Si–Ni intermetallic alloy” is used for
this alloy composition throughout the manuscript., in at.%. A disc (200 g) of
low phase purity was produced via arc melting, despite the fact that the ratio
of elemental powders in the powder feedstock had been adjusted to the Nb3SiNi2
stoichiometry. Additional information on the arc melting synthesis of alloys
in the Nb–Si–Ni system may be found elsewhere [44, 45, 46].

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 46100).
The as-cast alloy samples were pre-encapsulated in a quartz (SiO2) tube, which
was pressurized with argon (Ar) gas to 1 atm; both heating and cooling rates
were fixed at about 14.2 K/min.

###  4.2 Synthesis via Reactive Hot Pressing

The inability to produce a phase-pure Nb–Si–Ni intermetallic alloy by means of
arc melting redirected the synthesis efforts in this work towards a powder
metallurgical processing route. Quasi phase-pure discs of the Nb–Si–Ni
intermetallic alloy were successfully produced via reactive hot pressing
(RHP), using hydride and elemental starting powders, i.e., NbH0.89 ($<$40
$\mu$m, CBMM, Brazil), Ni (Vale, T123™, 3-7 $\mu$m, UK) and Si (A10, 2.1
$\mu$m, HC Starck, Germany). To optimize the mixing of the starting powders
and improve powder particle contact, the as-received NbH0.89 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 (diameters of 5 and 10 mm). The refined
NbH0.89 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 balls. The ratio of the NbH0.89, Ni and Si powders in the powder
feedstock was adjusted to the targeted 3Nb:1Si:2Ni. After mixing, the powder
suspension was dried in a rotating evaporator.

The powder feedstock was cold-pressed into ‘green’ pellets (30 mm in diameter,
4-5 mm thickness) under a load of 30 MPa in a graphite die. Exploratory RHP
runs were carried out 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 consolidation occurred under an uniaxial pressure of 30 MPa for a dwell
period of 60 min. The sintered discs (30 mm in diameter) were allowed to cool
naturally by switching off the power supply of the hot-press.

###  4.3 Materials Characterization

Metallographic cross-sections were prepared from all Nb–Si–Ni intermetallic
alloy samples produced via arc melting and RHP; all cross-sections were
polished to a mirror surface finish with a colloidal silica (OPS) suspension
in the final step. The microstructure of as-cast and isothermally annealed
samples produced via arc melting was studied by means of a Thermo Fisher Apreo
SEM (scanning electron microscope) equipped with an EDX (energy-dispersive
X-ray spectroscopy) detector (EDAX Octane Elite Super). Elemental mapping in
the RHP samples was done by means of an electron probe microanalyzer (EPMA;
JXA-8530F, JEOL Ltd, Japan) equipped with a WDXS (wavelength-dispersive X-ray
spectroscopy) detector (1-5 eV, 8 eV at Fe-Kα, full scanner type
spectrometer). EDX elemental analysis (1 SDD/UTW, 129 eV at Mn-Kα) of the
identified phases was performed on the same device at fixed conditions (15 kV
accelerating voltage, 100 nA beam current).

X-ray diffraction (XRD) was employed to identify the phases present in all
Nb–Si–Ni intermetallic alloy samples at room temperature. The as-cast and
annealed samples produced via 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∘ 2$\theta$ range. The RHP samples were analyzed 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∘ 2$\theta$ range, at a scan speed of 0.1∘ per min.

An FEI Helios 600 dual-beam FIB/SEM was used to lift out TEM (transmission
electron microscopy) thin foils from sufficiently large Nb3SiNi2 grains in the
arc-melted alloy samples, utilizing the classic lift-out method [66]. Combined
Pt/C protective capping layers were deposited on the sites of interest to
protect the foils from ion beam damage during lift-out. The FIB (focused ion
beam) foils were ion-milled to electron transparency at 30 keV, with final
thinning/ion beam damage removal at 2 keV.

TEM analysis of the FIB foils lifted out from the arc-melted alloy samples was
performed by means of 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 EDX detector,
and all EDX quantification work was conducted based on the Cliff-Lorimer
method [67, 68]. Selected area electron diffraction patterns (SAEDPs) of the
Nb3SiNi2 IMC ZIA phase were indexed using data from the Inorganic Crystal
Structure Database (ICSD) [43]. The CrystalMaker, SingleCrystal and
CrystalDiffraction software packages were used to simulate electron
diffraction patterns and aid the indexing of experimentally acquired SAEDPs
[52].

###  4.4 Thermodynamic & Crystal Structure Simulations

Thermodynamic simulations were performed to predict the equilibrium phases at
the sintering temperatures of the RHP 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) [69], based on the database published by dos
Santos et al. [44]. The unit cells of the fcc Nb3SiNi2 ZIA phase and the hcp
Ti3SiC2 MAX phase were drawn using version 3.5.8 of the VESTA software [70],
utilizing the Wyckoff parameters provided in Table S1 (Supplementary
Information) for precise crystal structure representation. These two phases
were compared in terms of relative structural complexity and unit cell sizes,
whilst certain zone axes of the herein addressed Nb3SiNi2 ZIA phase were
selected to depict ‘zigzag’ atomic arrangement and crystal structure
nanolamination.

## Acknowledgments

The Los Alamos National Laboratory (LANL), an affirmative action equal
opportunity employer, is managed by Triad National Security, LLC, for the U.S.
Department of Energy’s National Nuclear Security Administration, under
contract number 89233218CNA000001. LANL provided research support to MAT via
the Laboratory Directed Research and Development (LDRD) program, under project
number 20200689PRD2. OEA acknowledges funding from his Early Career Program
supported by LANL’s LDRD, under contract number 20210626ECR. CGS thanks the
financial support provided by the Conselho Nacional de Desenvolvimento
Científico e Tecnológico (CNPq, Brasília-DF, Project 307627/2021-7). MAT would
like to thank Dr. Kurt E. Sickafus, LANL, for valuable discussions on the
crystallography of complex solid-state phases. KL and JV acknowledge funding
from the Euratom research and training program 2014-2018 under Grant Agreement
No. 740415 (H2020 IL TROVATORE). NG acknowledges the Fund for Scientific
Research Flanders (FWO-Vlaanderen) for funding his PhD Fundamental Research
Fellowship No 1118120N.

## Author Contributions

The initially devised experimental plan that employed arc melting to explore
the precursor idea of the ZIA phases was formulated by MAT as a project for a
Director’s fellowship at LANL. MAT and KL share the conceptualization of the
major scientific ideas associated with the class of nanolaminated ternary IMCs
herein identified as the ZIA phases as well as their correlation with the
class of nanolaminated early transition metal carbides/nitrides known as the
MAX phases. SH, NG, and JV succeeded in producing quasi-phase-pure Nb3SiNi2
materials by means of RHP. All other authors have smaller contributions with
experiments, data analysis, research supervision, and funding. All authors
have read and approved the manuscript. MAT started the first draft and KL
extensively revised it, contributing with key ideas and RHP data, the latter
in collaboration with KU Leuven.

## References

## References

  * [1] R. W. Cahn, Nanostructured materials, Nature 348 (6300) (1990) 389–390.
  * [2] H. Gleiter, Nanostructured materials, Advanced Materials 4 (7-8) (1992) 474–481.
  * [3] P. Moriarty, Nanostructured materials, Reports on Progress in Physics 64 (3) (2001) 297.
  * [4] F. J. Martin-Martinez, K. Jin, D. López Barreiro, M. J. Buehler, The rise of hierarchical nanostructured materials from renewable sources: learning from nature, ACS Nano 12 (8) (2018) 7425–7433.
  * [5] M. W. Barsoum, T. El-Raghy, Synthesis and characterization of a remarkable ceramic: Ti3SiC2, Journal of the American Ceramic Society 79 (7) (1996) 1953–1956.
  * [6] M. W. Barsoum, T. El-Raghy, The MAX phases: Unique new carbide and nitride materials: Ternary ceramics turn out to be surprisingly soft and machinable, yet also heat-tolerant, strong and lightweight, American Scientist 89 (4) (2001) 334–343.
  * [7] W. Jeitschko, H. Nowotny, F. Benesovsky, Kohlenstoffhaltige ternare Verbindungen (H-phase), Monatshefte fur Chemie und verwandte Teile anderer Wissenschaften 94 (4) (1963) 672–676.
  * [8] W. Jeitschko, H. Nowotny, Die kristallstruktur von Ti3SiC2—ein neuer komplexcarbid-typ, Monatshefte fur Chemie-Chemical Monthly 98 (2) (1967) 329–337.
  * [9] M. Sokol, V. Natu, S. Kota, M. W. Barsoum, On the chemical diversity of the MAX phases, Trends in Chemistry 1 (2) (2019) 210–223.
  * [10] N. Goossens, B. Tunca, T. Lapauw, K. Lambrinou, J. Vleugels, MAX Phases, Structure, Processing, and Properties, in: M. Pomeroy (Ed.), Encyclopedia of Materials: Technical Ceramics and Glasses, Elsevier, Oxford, 2021, pp. 182–199.
  * [11] M. Dahlqvist, A. Petruhins, J. Lu, L. Hultman, J. Rosén, Origin of chemically ordered atomic laminates (i-MAX): expanding the elemental space by a theoretical/experimental approach, ACS Nano 12 (8) (2018) 7761–7770.
  * [12] B. Li, J. D. Corbett, Participation of sodium in the bonding of anionic networks: Synthesis, structure, and bonding of Na3MIn2 (M = Au, Ag), Inorganic Chemistry 44 (19) (2005) 6515–6517.
  * [13] M. Naguib, M. Kurtoglu, V. Presser, J. Lu, J. Niu, M. Heon, L. Hultman, Y. Gogotsi, M. W. Barsoum, Two-dimensional nanocrystals produced by exfoliation of Ti3AlC2, Advanced Materials 23 (37) (2011) 4248–4253.
  * [14] M. Naguib, O. Mashtalir, J. Carle, V. Presser, J. Lu, L. Hultman, Y. Gogotsi, M. W. Barsoum, Two-dimensional transition metal carbides, ACS Nano 6 (2) (2012) 1322–1331.
  * [15] Y. Gogotsi, B. Anasori, The rise of MXenes, ACS Nano 13 (8) (2019) 8491–8494.
  * [16] Z. Wu, T. Shang, Y. Deng, Y. Tao, Q.-H. Yang, The assembly of MXenes from 2D to 3D, Advanced Science 7 (7) (2020) 1903077.
  * [17] Y. Long, Y. Tao, T. Shang, H. Yang, Z. Sun, W. Chen, Q.-H. Yang, Roles of metal ions in MXene synthesis, processing and applications: a perspective, Advanced Science 9 (12) (2022) 2200296.
  * [18] M. Ghebouli, B. Ghebouli, M. Fatmi, A. Bouhemadou, Theoretical prediction of the structural, elastic, electronic and thermal properties of the MAX phases X2SiC (X = Ti and Cr), Intermetallics 19 (12) (2011) 1936–1942.
  * [19] T. Lapauw, J. Halim, J. Lu, T. Cabioch, L. Hultman, M. Barsoum, K. Lambrinou, J. Vleugels, Synthesis of the novel Zr3AlC2 MAX phase, Journal of the European Ceramic Society 36 (3) (2016) 943–947.
  * [20] T. Lapauw, K. Lambrinou, T. Cabioch, J. Halim, J. Lu, A. Pesach, O. Rivin, O. Ozeri, E. N. Caspi, L. Hultman, et al., Synthesis of the new MAX phase Zr2AlC, Journal of the European Ceramic Society 36 (8) (2016) 1847–1853.
  * [21] J. Xu, M.-Q. Zhao, Y. Wang, W. Yao, C. Chen, B. Anasori, A. Sarycheva, C. E. Ren, T. Mathis, L. Gomes, L. Zhenghua, Y. Gogotsi, Demonstration of Li-ion capacity of MAX phases, ACS Energy Letters 1 (6) (2016) 1094–1099.
  * [22] H. Lin, Y. Chen, J. Shi, Insights into 2D MXenes for versatile biomedical applications: current advances and challenges ahead, Advanced Science 5 (10) (2018) 1800518.
  * [23] H. Tong, S. Lin, Y. Huang, P. Tong, W. Song, Y. Sun, Difference in physical properties of MAX-phase compounds Cr2GaC and Cr2GaN induced by an anomalous structure change in Cr2GaN, Intermetallics 105 (2019) 39–43.
  * [24] Q. Tao, J. Lu, M. Dahlqvist, A. Mockute, S. Calder, A. Petruhins, R. Meshkian, O. Rivin, D. Potashnikov, E. N. Caspi, H. Shaked, A. Hoser, C. Opagiste, R.-M. Galera, R. Salikhov, U. Wiedwald, C. Ritter, A. R. Wildes, B. Johansson, L. Hultman, M. Farle, M. W. Barsoum, J. Rosén, Atomically layered and ordered rare-earth i-MAX phases: a new class of magnetic quaternary compounds, Chemistry of Materials 31 (7) (2019) 2476–2485.
  * [25] L. Naslund, P. O. Persson, J. Rosén, X-ray photoelectron spectroscopy of Ti3AlC2, Ti3C2Tz, and TiC provides evidence for the electrostatic interaction between laminated layers in MAX-Phase materials, The Journal of Physical Chemistry C 124 (50) (2020) 27732–27742.
  * [26] M. Sanna, S. Ng, J. V. Vaghasiya, M. Pumera, Fluorinated MAX Phases for Photoelectrochemical Hydrogen Evolution, ACS Sustainable Chemistry & Engineering 10 (8) (2022) 2793–2801.
  * [27] H. Shao, S. Luo, A. Descamps-Mandine, K. Ge, Z. Lin, P.-L. Taberna, Y. Gogotsi, P. Simon, Synthesis of MAX Phase Nanofibers and Nanoflakes and the Resulting MXenes, Advanced Science 10 (1) (2023) 2205509.
  * [28] X.-H. Zha, X. Ma, S. Du, R.-Q. Zhang, R. Tao, J.-T. Luo, C. Fu, Role of the A-element in the structural, mechanical, and electronic properties of Ti3AC2 MAX phases, Inorganic Chemistry 61 (4) (2021) 2129–2140.
  * [29] M. A. Saeed, A. Shahzad, K. Rasool, F. Mateen, J.-M. Oh, J. W. Shim, 2D MXene: a potential candidate for photovoltaic cells? A critical review, Advanced Science 9 (10) (2022) 2104743.
  * [30] L. Chen, Y. Li, B. Zhao, S. Liu, H. Zhang, K. Chen, M. Li, S. Du, F. Xiu, R. Che, et al., Multiprincipal Element M2FeC (M = Ti, V, Nb, Ta, Zr) MAX Phases with Synergistic Effect of Dielectric and Magnetic Loss, Advanced Science (2023) 2206877.
  * [31] D. J. Tallman, E. N. Hoffman, E. N. Caspi, B. L. Garcia-Diaz, G. Kohse, R. L. Sindelar, M. W. Barsoum, Effect of neutron irradiation on select MAX phases, Acta Materialia 85 (2015) 132–143.
  * [32] C. Ang, C. Silva, C. Shih, T. Koyanagi, Y. Katoh, S. J. Zinkle, Anisotropic swelling and microcracking of neutron irradiated Ti3AlC2–Ti5Al2C3 materials, Scripta Materialia 114 (2016) 74–78.
  * [33] D. J. Tallman, L. He, B. L. Garcia-Diaz, E. N. Hoffman, G. Kohse, R. L. Sindelar, M. W. Barsoum, Effect of neutron irradiation on defect evolution in Ti3SiC2 and Ti2AlC, Journal of Nuclear Materials 468 (2016) 194–206.
  * [34] C. Ang, C. M. Parish, C. Shih, C. Silva, Y. Katoh, Microstructure and mechanical properties of titanium aluminum carbides neutron irradiated at 400–700∘C, Journal of the European Ceramic Society 37 (6) (2017) 2353–2363.
  * [35] D. J. Tallman, L. He, J. Gan, E. N. Caspi, E. N. Hoffman, M. W. Barsoum, Effects of neutron irradiation of Ti3SiC2 and Ti3AlC2 in the 121–1085∘C temperature range, Journal of Nuclear Materials 484 (2017) 120–134.
  * [36] M. A. Tunes, R. W. Harrison, S. E. Donnelly, P. D. Edmondson, A transmission electron microscopy study of the neutron-irradiation response of Ti-based MAX phases at high temperatures, Acta Materialia 169 (2019) 237–247.
  * [37] T. Lapauw, B. Tunca, J. Joris, A. Jianu, R. Fetzer, A. Weisenburger, J. Vleugels, K. Lambrinou, Interaction of Mn+1AXn phases with oxygen-poor, static and fast-flowing liquid lead-bismuth eutectic, Journal of Nuclear Materials 520 (2019) 258–272.
  * [38] D. Bowden, J. Ward, S. Middleburgh, S. de Moraes Shubeita, E. Zapata-Solvas, T. Lapauw, J. Vleugels, K. Lambrinou, W. Lee, M. Preuss, et al., The stability of irradiation-induced defects in Zr3AlC2, Nb4AlC3 and (Zr0.5,Ti0.5)3AlC2 MAX phase-based ceramics, Acta Materialia 183 (2020) 24–35.
  * [39] B. Tunca, T. Lapauw, C. Callaert, J. Hadermann, R. Delville, E. N. Caspi, M. Dahlqvist, J. Rosén, A. Marshal, K. G. Pradeep, et al., Compatibility of Zr2AlC MAX phase-based ceramics with oxygen-poor, static liquid lead–bismuth eutectic, Corrosion Science 171 (2020) 108704.
  * [40] M. A. Tunes, M. Imtyazuddin, C. Kainz, S. Pogatscher, V. M. Vishnyakov, Deviating from the pure MAX phase concept: Radiation-tolerant nanostructured dual-phase Cr2AlC, Science advances 7 (13) (2021) eabf6771.
  * [41] M. A. Tunes, S. M. Drewry, J. D. Arregui-Mena, S. Picak, G. Greaves, L. B. Cattini, S. Pogatscher, J. A. Valdez, S. Fensin, O. El-Atwani, S. E. Donnelly, T. A. Saleh, P. D. Edmondson, Accelerated radiation tolerance testing of Ti-based MAX phases, Materials Today Energy (2022) 101186.
  * [42] B. Tunca, G. Greaves, J. Hinks, P. O. Persson, J. Vleugels, K. Lambrinou, In situ He+ irradiation of the double solid solution (Ti0.5,Zr0.5)2(Al0.5,Sn0.5)C MAX phase: Defect evolution in the 350–800∘C temperature range, Acta Materialia 206 (2021) 116606.
  * [43] E. Gladyshevskii, Y. B. Kuz’ma, P. Kripyakevich, The crystal structures of Mn3Ni2Si, V3Ni2Si, Nb3Ni2Si and related Cr and Ta compounds, Journal of Structural Chemistry 4 (3) (1964) 343–355.
  * [44] V. O. dos Santos, H. M. Petrilli, C. G. Schön, L. T. Eleno, Thermodynamic modelling of the Nb–Ni–Si phase diagram based on the 1073 K isothermal section using ab initio calculations, Calphad 51 (2015) 57–66.
  * [45] V. O. dos Santos, M. A. Tunes, L. T. Eleno, C. G. Schön, K. W. Richter, Experimental investigation of phase equilibria in the Nb–Ni–Si refractory alloy system at 1073 K, Scripta Materialia 164 (2019) 96–100.
  * [46] V. O. dos Santos, L. T. Eleno, C. G. Schön, K. W. Richter, Experimental investigation of phase equilibria in the Nb–Ni–Si refractory alloy system at 1323 K, Journal of Alloys and Compounds 842 (2020) 155373.
  * [47] M. Hellenbrandt, The inorganic crystal structure database (ICSD)—present and future, Crystallography Reviews 10 (1) (2004) 17–22.
  * [48] D. A. Porter, K. E. Easterling, Phase transformations in metals and alloys (revised reprint), CRC press, 2009.
  * [49] S. Coulombe, J.-L. Meunier, Theoretical prediction of non-thermionic arc cathode erosion rate including both vaporization and melting of the surface, Plasma Sources Science and Technology 9 (3) (2000) 239.
  * [50] R. Radhakrishnan, J. Williams, M. Akinc, Synthesis and high-temperature stability of Ti3SiC2, Journal of Alloys and Compounds 285 (1-2) (1999) 85–88.
  * [51] N. Goossens, T. Lapauw, K. Lambrinou, J. Vleugels, Synthesis of MAX phase-based ceramics from early transition metal hydride powders, Journal of the European Ceramic Society 42 (16) (2022) 7389–7402.
  * [52] D. C. Palmer, Visualization and analysis of crystal structures using CrystalMaker software, Zeitschrift fur Kristallographie-Crystalline Materials 230 (9-10) (2015) 559–572.
  * [53] K. Kuo, C. Dong, D. Zhou, Y. Guo, Z. Hei, D. Li, A Friauf-Laves (Frank-Kasper) phase related quasicrystal in a rapidly solidified Mn3Ni2Si alloy, Scripta Metallurgica 20 (12) (1986) 1695–1698.
  * [54] M. Kolenda, A. Szytula, J. Leciejewicz, A. Pawlukojc, H. Ptasiewicz-Bak, Magnetic structure of Mn3Ni2Si, Journal of Magnetism and Magnetic Materials 89 (1-2) (1990) 26–28.
  * [55] C. Wang, L. Huang, M. Yang, X. Huang, J. Zhang, S. Pan, S. Yang, X. Liu, Experimental investigation of phase equilibria in the Ni–Ta–Si refractory alloy system, Journal of Alloys and Compounds 888 (2021) 161467.
  * [56] Y. B. Kuz’ma, E. Gladyshevskii, D. Byk, Crystal structure of some ternary compounds in the Nb–Co–Si system, Journal of Structural Chemistry 5 (4) (1964) 518–522.
  * [57] N. Pavlyuk, G. Dmytriv, V. Pavlyuk, B. Rozdzynska-Kielbik, G. Cichowicz, M. K. Cyranski, I. Chumak, H. Ehrenberg, New cubic cluster phases in the Mg–Ni–Ga system, Acta Crystallographica Section B: Structural Science, Crystal Engineering and Materials 76 (4) (2020) 534–542.
  * [58] M. Radovic, M. W. Barsoum, MAX phases: bridging the gap between metals and ceramics, American Ceramics Society Bulletin 92 (3) (2013) 20–27.
  * [59] A. Jain, S. P. Ong, G. Hautier, W. Chen, W. D. Richards, S. Dacek, S. Cholia, D. Gunter, D. Skinner, G. Ceder, Commentary: The Materials Project: A materials genome approach to accelerating materials innovation, APL Materials 1 (1) (2013) 011002.
  * [60] S. Sun, Y. Yu, S. Sun, Q. Wang, T. Chen, J. Chen, Y. Zhang, W. Cui, Magnetic properties and microstructures of Fe-doped (Ti1-xFex)3AlC2 MAX phase and their mxene derivatives, Journal of Superconductivity and Novel Magnetism 34 (2021) 1477–1483.
  * [61] Y. Li, J. Liang, H. Ding, J. Lu, X. Mu, P. Yan, X. Zhang, K. Chen, M. Li, P. O. Persson, et al., Near-room temperature ferromagnetic behavior of single-atom-thick 2D iron in nanolaminated ternary MAX phases, Applied Physics Reviews B 8 (3) (2021) 031418.
  * [62] G. Deysher, C. E. Shuck, K. Hantanasirisakul, N. C. Frey, A. C. Foucher, K. Maleski, A. Sarycheva, V. B. Shenoy, E. A. Stach, B. Anasori, et al., Synthesis of Mo4VAlC4 MAX phase and two-dimensional Mo4VC4 MXene with five atomic layers of transition metals, ACS Nano 14 (1) (2019) 204–217.
  * [63] T. Lapauw, B. Tunca, T. Cabioc’h, J. Lu, P. O. Persson, K. Lambrinou, J. Vleugels, Synthesis of MAX phases in the Hf–Al–C system, Inorganic Chemistry 55 (21) (2016) 10922–10927.
  * [64] B. Tunca, T. Lapauw, R. Delville, D. R. Neuville, L. Hennet, D. Thiaudière, T. Ouisse, J. Hadermann, J. Vleugels, K. Lambrinou, Synthesis and characterization of double solid solution (Zr,Ti)2(Al,Sn)C MAX phase ceramics, Inorganic Chemistry 58 (10) (2019) 6669–6683.
  * [65] T. Lapauw, B. Tunca, D. Potashnikov, A. Pesach, O. Ozeri, J. Vleugels, K. Lambrinou, The double solid solution (Zr,Nb)2(Al,Sn)C MAX phase: a steric stability approach, Scientific reports 8 (1) (2018) 12801.
  * [66] L. A. Giannuzzi, F. A. Stevie, A review of focused ion beam milling techniques for TEM specimen preparation, Micron 30 (3) (1999) 197–204.
  * [67] G. Cliff, G. Lorimer, Quantitative analysis of thin metal foils using EMMA-4, the ratio technique, in: Proceedings of the 5th European Congress on Electron Microscopy, The Institute of Physics, London, 1972, pp. 140–141.
  * [68] G. Cliff, G. W. Lorimer, The quantitative analysis of thin specimens, Journal of Microscopy 103 (2) (1975) 203–207.
  * [69] B. Sundman, B. Jansson, J.-O. Andersson, The thermo-calc databank system, Calphad 9 (2) (1985) 153–190.
  * [70] K. Momma, F. Izumi, Vesta 3 for three-dimensional visualization of crystal, volumetric and morphology data, Journal of applied crystallography 44 (6) (2011) 1272–1276.