2H-Au
[27]
Dealloying of A3-Au–Ga
[42]
Ag, Rh, Os, Ru,
Cu,
Pt−Cu,
Pt−Co, Ag−Pd,
Pt−Ag,
and
Pt−Pd−Ag
Epitaxial growth of metals
on 4H-Au nanoribbons
[18, 32, 43,
44]
14H
Co
Epitaxial growth of metals
on 4H-Au nanoribbons
[28]
hcp-type phase
with short-range
ordered
point
defects
Rh
Atomic diffusion rate
suppressed by unique
morphology
[29]
Z3-type phase
Fe(Pd,Ind)3 and
Fe(Pd,Pbd)3
Additional third element
based on inter-element
miscibility
[34]
20
(A)
2H (hcp)
3C (fcc)
4H
(B)
6H
(D)
100 nm
2 nm
Growth
(E)
(C)
2 nm
100 nm
(G)
(F)
2 nm
Intensity (a.u.)
(H)
0 0.5 1.0 1.5 2.0 2.5
Lattice Spacing (nm)
Figure 1. Mono-metal nanomaterials with ordered stacking faults. (A) Variation of
ordered stacking faults; where C is cubic, H is hexagonal, and the numerical characters
are repeating numbers of a close-packed plane. (B–E) Transmission electron microscopy
(TEM) and high resolution TEM (HRTEM) images of Au nanowires (B and D) and
nanoribbons (C and E). (F–H) HRTEM images of 4H-type Au and 14H-type Co phases
(F and G) as well as the integrated pixel intensities in G (H). Reproduced from (B–E) ref.
21
27 and with permission from (F–H) ref. 28. Copyright 2021, John Wiley and Sons.
22
(A)
(B)
(G)
[010]h/[110]h
(H)
(C)
(E)
A side view
C1
(I)
[001]h
C2
[210]h
/[110]h
[010]h/[110]h
(F)
A tilted side view
(J)
(D)
C4
(K)
1h
C3
3h
6h
14 h
(L)
Rh atoms
Slow diffusion
Quick
diffusion
Figure 2. Formation of hcp-type Rh with short-range ordered point defects. (A) Model of
the unit cells for two-type vacated Barlow packing (VBP-1 and VBP-2) of hcp-type Rh;
where Barlow packing is a general term for e.g. 2H, 3C, 4H, and 6H. (B) Model of hcp-
Rh structure without ordered point defects from [010]h or [110]h and its fast Fourier-
transform (FFT) image. (C–F) High-angle annular dark-field scanning transmission
electron microscopy (HAADF–STEM) images (C and D) observed from two angles (E
and F, respectively). (G–J) FFT images (left images) obtained from C1–C4 in the
HAADF–STEM images (C and D), respectively, and the VBP (middle) and FFT (right)
10
images corresponding to the experimental FFT images (left images). (K) Transmission
11
electron microscopy images of Rh NMs synthesized by time-dependent experiments.
12
(Scale bar: 50 nm) (L) Schematic of atomic diffusion rate limited by atomic-level
23
intercalated Rh nanosheets. Reproduced from (A–K) ref. 29.
24
Unexplored binary alloys
(A)
(n:m) = (1:1)
A1B1
(2:1)
A2B1
(3:3)
A3B3
(3:1)
A3B1
(4:1)
A4B1
[001]
[111]
L10-type
[001]
[111]
[110]
[110] L11-type
B2-type
β2-type
(B)
Pd
In
Fe
In/Fe
Fe/Pd
Fe/Pd/In
: Pd
: Fe
: In
1 nm
Pd/In
Figure 3. First synthesis of pseudo-Z3 structure in uninvestigated binary alloys. (A)
Various ordered binary alloys based on the fcc framework. (B) High-angle annular dark-
field scanning transmission electron microscopy and atomic-resolution energy-dispersive
X-ray spectroscopy images, and the model of the unit cell in Z3-type Fe(Pd,Ind)3, where
the superscript refers to the Wyckoff letter. Reproduced from ref. 34.
25
(B)
Eform (eV atom–1)
EL12−EZ3 < 0
EL12−EZ3 (eV atom–1)
(A)
L12
Z3
EL12−EZ3 (eV atom–1)
(C)
EL12−EZ3 > 0
(D)
A1-PdInx@FeOy
Hg
In
Zn
Ga
Ge
Z3-type Fe(Pd,Ind)3
Pd
Tl
Cd
Sn
Fe
miscible immiscible
Pb
Pd
Fe
miscible miscible
In powder
(mp: 156 ºC)
Pd@FeOx
Reductive
50 nm annealing
Reductive
annealing
50 nm
Figure 4. Key factors for forming the Z3-type Fe(Pd,Ind)3 structure. (A) Formation
energies (Eform) of Z3-type and L12-type FeaPdbInc [(a, b, c) = (2, 6, 0), (1, 6, 1), and (2,
5, 1)] obtained from first-principles calculations, corresponding to E[L12- or Z3-type
FeaPdbInc] – (aE[Fe] + bE[Pd] + cE[In]), where E[X] is equivalent to the total energies of
X at the ground states. (B) Change of EL12 and EZ3 dependent on the quantity of In, where
EL12 and EZ3 are equal to x×E[L12-(Fe1, In1)Pd6] + (1–x) ×E[L12-Fe2Pd6] + x×E[Fe] and
x×E[Z3-Fe2(Pd5, In1d)] + (1–x) ×E[Z3-Fe2Pd6] + x×E[Pd] (0 ≤ x ≤ 1), respectively. (C)
Difference of EL12 and EZ3 (x = 1) in the case of substituting M instead of In (M = Zn, Ga,
10
Ge, Cd, Sn, Hg, Tl, and Pb). (D) Schematic indicating that nanoscale-homogeneous
26
nanoparticulate precursor powder is a key factor for forming Z3-type Fe(Pd,Ind)3.
Reproduced from ref. 34.
27
(A)
Interface of core/shell
(111) in L10type structure Anisotropic lattice mismatch
ei = (δi – dPt)/dPt (i = 1, 2)
Pt(111)
4 layers
[Atomic distances]
δ2 (≠ δ1)
dPt
2.82 Å (Unstrained Pt)
L10-type
structure
-2
-4
-6
|e1| ≤ |e2|
-8
-6
-4 -2
e1 %
-8
0.1
0.1
0.0
0.0
0.0
(D)
G at U = 1.23VRHE (eV)
: Ni
: Cu
: Zn
e2 %
: none
: Mn
: Fe
Increasing activity
(C)
ΔGPt – ΔGStrained Pt
(B)
: none : Ni
: Cu
: Mn
: Fe
: Zn
2H2O + *
*OOH
O2 + *
*O2
*OH
Unstrained Pt
*O
0 0 1 2 3 4
(H+ + e–) transferred
-4
-6
|e1| ≤ |e2|
-8
-6
-4 -2
e1 %
-8
e2 %
-2
0.2
0.1
0.1
0.0
0.0
0.0
G at U = 1.23VRHE (eV)
ΔGPt – ΔGStrained Pt
Increasing activity
(E)
H2O + *
1/2 O2 + *
1/2 *O2
*OH
*O
1 2
(H+ + e–) transferred
Figure 5. Oxygen reduction reaction (ORR) activity of Pt enhanced by anisotropic strain.
(A) Model of L10-alloy@Pt4-layer core@shell structure and anisotropic lattice mismatch
introduced on the {111} planes of the Pt shell (Pt{111}) induced by the {111} planes of
the L10-type structure. (B and D) Strain–ORR activity relationship for associative (B) and
dissociative (D) mechanisms, where GPt – ΔGStrained Pt corresponds to the difference in
these activation barriers at the reactions [*O2 + H+ + e– → *OOH (B) and *OH + H+ + e–
→ H2O + * (C), where * refers to the state of adsorption on the catalysts]. (C and E) Gibbs
28
free energies for the ORR on unstrained and strained Pt via associative (C) and
dissociative (E) mechanisms. Reproduced with permission from (B–E) ref. 85. Copyright
2020, American Chemical Society.
29
Keywords
Nanomaterials, ordered stacking fault, ordered point defect, inter-element miscibility,
ordered alloy
Glossary
Bulk energy: internal energy when the total energy of a substance is divided into internal
and surface energies
Cohesive energy: energy required to separate each atom from a solid
Density of states (DOS): electron energy distribution formed by orbital hybridizations
between all of the atoms in a solid. The shape of the DOS is determined by the symmetry
of the structure, the inter-atom distance, and the species of the constituent elements. In
particular, the DOS for electrons with a maximum energy near the Fermi level is used to
describe various physical and chemical properties, such as electrical conductivity and
catalytic properties.
First-principles calculations: method of solving the kinetic energy of electrons in a
substance by numerical calculations in accordance with quantum theory. In many cases,
an approximate solution is obtained by expressing the electrons in terms of their density;
i.e., by using density functional theory.
Formation energy: difference in energy between the bulk energy of the alloy and the
bulk energy of each constituent element; i.e., the energy obtained by alloying
Intermetallic binary alloy: alloy structure in which the constituent elements in a binary
alloy are arranged at specific atomic positions
Kirkendall effect: phenomenon in which hollows are formed in a substance because of
differences in the rate of atomic diffusion for each element in the substance. This cavity
formation indicates that atoms diffuse by using defects in the material.
Solid–solution alloy: alloy structure in which multiple elements are randomly arranged
based on the crystal structure of a mono-metal
Wyckoff letter: nonequivalent atomic positions in the unit cell based on the space group.
In addition, atomic positions are often expressed by adding the multiplicity of equivalent
atomic positions. In the case of Z3-type Fe(Pd,Ind)3, the space group is P4/mmm and the
atomic positions are Fe1a (0, 0, 0), Fe1c (0.5, 0.5, 0), Pd4i (0.5, 0, 0.23), Pd1b (0, 0, 0.5),
and In1d (0.5, 0.5, 0.5).
Highlights
An investigation of unprecedented crystal structures enables development of new
functions and enhancement of well-known properties.
Although an infinite number of crystal structures are geometrically possible, the crystal
structures of metal nanoparticles depend on thermodynamics.
Mono-metal nanoparticles with ordered stacking faults and ordered point defects, as well
as unprecedented ordered alloy nanoparticles, are generated by stabilization.
Stabilization is by (1) transformation from unfavorable to more-favorable structures
during growth of the nanoparticles and epitaxial growth of other metals on ordered
stacking faults structures, (2) suppression of the atomic diffusion rate, and (3) substitution
of third elements based on the inter-element miscibility.
Box 1. Formation of metastable alloy NMs with well-known crystal structures
Metastable alloy NMs with well-known crystal structures stabilized by nano-size effect
have been synthesized by kinetic chemical synthesis methods. For example, solidsolution alloys between immiscible elements have been obtained by a simultaneous
reduction of multiple metal precursors with different redox potentials [1, 2, 9-15].
Moreover, control over the crystal structure can be also done by fine-tuning of the
reduction rate in the simultaneous reduction procedure [1]. Interestingly, the kinetically
formed metastable alloy NMs can potentially transform into different metastable phases
by an appropriate external stimulus. For example, Pd–Ru alloy composed of immiscible
elements transformed from A1 to A3 structures by introducing hydrogen atoms [2].
Therefore, the modification of nucleation process is important for the alloying of
immiscible elements and the formation of metastable crystal structures.
Box 2. Phase stability of Au nanoribbons with 4H structure
4H-Au nanoribbons, which are used as templates for forming 4H structures of various
mono-metals and solid solution alloys, are stable under high temperature and high
pressure. According to in situ TEM observation under high temperature, 4H structure was
kept until < 800 °C [90]. As a result of the pressurization experiment at room temperature,
it was confirmed that the 4H structure was maintained up to 1.2 GPa and the 4H structure
was maintained as an fcc/4H heterostructure up to about 26 GPa [91]. Surprisingly, in situ
TEM under 1 mbar of CO gas observed the transformation of fcc (stable phase) to 4H
(metastable phase) structures of Au nanospheres on 4H-Au nanoribbons [92]. Firstprinciples calculations and experiments strongly support that this phase transition is
driven by surface energy gain that exceeds bulk energy loss. These results indicate that
4H-Au nanoribbon is an effective material as a template for the 4H phase formation of
other metals.
Box 3. Formation of intermetallic compound NMs
A large difference in redox potentials is a serious problem in synthesizing intermetallic
compound NMs, because the simultaneous reduction method tends to form a phasesegregated structure. Then, a step-by-step chemical synthesis method is effective for
alloying such an element pair. For example, after the growth of metal oxides (low redox
potential) on noble metal NMs (high redox potential) with monodisperse size and shape,
the reductive annealing for the nanoparticulate precursor powders is conducted at high
temperature (>500 °C), by which highly ordered intermetallic compound NMs are formed
[34, 93]. Recently, in order to avoid the inter-particle fusion happening at such a high
temperature, the synthesis of intermetallic compounds NMs by the introduction of a third
element [94] and alloying noble metal NMs with base metals in a solution that excludes
oxygen [51, 52] have been reported. In both methods, relatively high temperature around
300 °C in a solution system allows the atomic diffusion within particles and the formation
of intermetallic compounds. These approaches facilitate the control of particle size and
shape, and the investigation on the phase stability of intermetallic compounds including
or excluding nano-size effects, respectively.
Outstanding Questions
Various metastable phases have been discovered in the nanoscale regime. These phases
are metastable or kinetically stable in bulk. Can we regard them to be thermodynamically
stable, considering nano-size effects?
4H-Au nanoribbons are formed from 2H-Au nanowires. Is it possible to form 4H-Au
nanoplates from 2H-Au nanosheets? Because 4H-Au nanoplates exhibit different surfaces
compared with 4H-Au nanoribbons, can the different metastable phases be stabilized by
epitaxial growth of other mono-metals or solid–solution alloys on 4H-Au nanoplates?
To reveal the contribution of C atoms in hcp-type Rh nanoparticles with short-range
ordered point defects to the phase stability, is it possible to remove only the C atoms yet
maintain the crystal structure?
Inter-element miscibility of In, which is miscible with Pd but immiscible with Fe, restricts
substitution sites of In in L12 and Z3-FePd3 structures to sites where Fe and In are not
adjacent. Does such an inter-element miscibility restrict diffusion paths until forming Fe–
Pd–In alloy phases? Can the Z3-type Fe(Pd,Ind)3 structure be formed for nanoparticles
smaller than 10 nm by the nano-size effect?
...
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