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A Mg-In Alloy Interphase for Mg Dendrite
Suppression
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J. Electrochem. Soc.
171 010513
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A Mg-In Alloy Interphase for Mg Dendrite Suppression
Brian C. Lee
and Kimberly A. See
*
,
z
Division of Chemistry and Chemical Engineering, California Institute of Technology, Pasadena, California 91125, United
States of America
Mg metal batteries have attracted much attention as an alternative to Li-ion technology due to the high abundance and volumetric
capacity of Mg metal. Further, early reports show that Mg is less prone to dendritic growth compared to Li, thereby improving the
safety and long-term reversibility of Mg metal anodes. However, dendritic growth of Mg can be observed in various conditions,
causing cell shorting and capacity loss. Herein, we report a chemically-formed Mg-In alloy interphase that suppresses nonuniform
Mg growth during electrochemical reduction. Ex-situ X-ray diffraction shows that upon reduction, Mg alloys into the Mg-In
interphase with no evidence of Mg deposition on top of the surface during initial cycles. Interestingly, further reduction results in
Mg depositing underneath the interphase, which con
fi
rms Mg mobility through the interphase. However, the alloying reaction is
kinetically limited, leading to signi
fi
cant Mg deposition on top of the interphase at high current densities. Thus, alloys on Mg can
affect deposition morphologies, but are limited by the kinetics of Mg conduction through the alloy.
© 2024 The Electrochemical Society (
ECS
). Published on behalf of ECS by IOP Publishing Limited. [DOI:
10.1149/1945-7111/
ad1c13
]
Manuscript submitted November 15, 2023; revised manuscript received December 12, 2023. Published January 17, 2024.
Supplementary material for this article is available
online
Mg batteries have been studied as a promising alternative to
current Li-ion battery technology due to high natural abundance and
well-dispersed deposits of Mg-containing precursors.
1
However, Mg
batteries are only competitive from a performance perspective if a
Mg metal anode is used. Mg metal anodes have a very high
volumetric capacity of 3833 mAh cm
3
and a reasonable gravimetric
capacity of 2295 mAh g
1
, but face a host of challenges that prevent
commercialization. One of the biggest challenges facing Mg metal is
the tendency for Mg metal to passivate by reaction with water,
oxygen, and other electrolyte components to form an insulating
layer, shutting down deposition and stripping of the metal.
2
6
As
such, research into Mg metal anodes initially consisted primarily of
investigation into electrolytes to control the interface and prevent the
passivation of the Mg metal surface. To that end, several classes of
electrolytes, such as those based on Grignard reagents,
5
,
7
12
halide-
containing electrolytes,
13
18
tri
fl
uoromethanesulfonimide (TFSI)
salts,
19
21
fl
uorinated weakly coordinating anions,
22
,
23
and boron
hydride-based anions
24
,
25
were developed to varying degrees of
success.
A potential attractive feature of Mg metal anodes is the smooth
deposition morphology compard to Li metal. Matsui observed that at
the same current densities (2.0 mA cm
2
), Li forms dendrites while
Mg deposits in plate-like morphologies.
26
The lower self-diffusion
barrier of Mg compared to Li was hypothesized to result in smoother
morphologies.
27
Control over metal deposition morphology is
crucial as dendrites can have many detrimental effects, such as
capacity fade as a result of dead metal or shorting of the cell when
the dendrites pierce through the separators, posing safety hazards
associated with thermal ruanway in the presence of a
fl
ammable
electrolyte.
28
However, despite the low self-diffusion barrier in-
herent to Mg metal, the kinetics of the electron transfer coupled with
mass transport limitations can outpace the self-diffusion at high
enough current density. During deposition, the metal cations
nucleate and deposit until the ions at the surface of the negatively
charged electrode deplete.
29
31
The anions are repelled from the
cathode, while the cations are consumed, causing a large space
charge region at the electrode. As a result of the space charge region,
the cations deposit in such a way to maximize growth toward the
anode to minimize the space charge build up, causing dendritic
morphologies.
30
As such, dendritic growth of Mg has been reported
in literature. Davidson et al. observed Mg dendrites at
0.921 mA cm
2
, which is a lower current density than Matsui
reported.
32
However, the growth was mostly observed at the edges
of a Mg ribbon, where higher local
fl
ux would occur. Eaves-Rathert
et al. also observed nonuniform growth of Mg at 0.2 mA cm
2
in a
coin cell geometry using a polymer separator. While the Mg deposits
are not dendritic in the classic sense that the deposits do not form
classic branching structures, hemispherical islands form using the
separator as scaffolding and stack on top of each other to form
deposits that cause shorts through the polymer separator.
33
The
deposits are referred to as 3D growth but cause the same harmful
effects as a classic dendrite would. The separator scaffolding effect
is also observed by Hebié et al. and Ding et al.
34
36
A strategy to suppress Mg dendrites is through modi
fi
cation of
the electrode surface. The formation of arti
fi
cial interphases on Mg
surfaces has been commonly employed to improve Mg cycling by
protecting the anode from passivation.
37
,
38
However, some inter-
phases, particularly alloy-based systems, have been noted in
literature to minimize dendrites as well. Several different mechan-
istic explanations have been given for the dendrite-reducing ability
of alloy interphases. In many alloy interphase systems, the alloy
interphase decreases insulating passivation with the electrolyte,
which prevents preferential deposition at localized points with low
electrical resistance. As such, a smoother Mg morphology on the
electrode is observed.
39
42
Of particular interest are two alloy interphase systems where the
Mg is hypothesized to conduct through the alloy interphase and
deposit underneath. A Mg
Mg-Sn electrode was shown to decrease
Mg stripping and deposition overpotential while maintaining fast
kinetics, exhibiting
<
1 V overpotentials at a very high current
density of 6 mA cm
2
in symmetric cells. The Mg deposits at the
Mg
Mg-Sn interface is observed by cross-section scanning electron
microscopy (SEM). The bulk composition of the alloy remains
unchanged through cycling by X-ray photoelectron spectroscopy
(XPS), indicating mobility of Mg through the Mg-Sn interphase and
depositing underneath at the Mg
Mg-Sn interface. However, the
study focused mainly on cycling stability, and the mechanism of the
electroreduction process in the interphase was not determined.
43
In a Mg-Ga alloy interphase system, shorts when cycling in a
symmetric cell were prevented when using the alloy interphase.
Calculations suggest Mg thermodynamicaly prefers plating at the
Mg
Mg-Ga interface rather than the Mg-Ga
electrolyte interface.
Experimentally, Mg metal deposition on top of the interface is not
observed by X-ray diffraction (XRD) or XPS and the interphase
composition experiences minimal changes. However, kinetic limita-
tions are present in the system, with the symmetric cell galvanostatic
cycling at a relatively low current density of 0.1 mA cm
2
with a
concentrated electrolyte (0.8 M Mg(TFSI)
2
in glyme) at moderately
z
E-mail:
ksee@caltech.edu
*Electrochemical Society Member.
Journal of The Electrochemical Society
, 2024
171
010513
1945-7111/2024/171(1)/010513/7/$40.00 © 2024 The Electrochemical Society (
ECS
). Published on behalf of ECS by IOP Publishing Limited
elevated temperatures (40 °C). In addition, the Mg deposition
underneath the interphase was not con
fi
rmed by cross-section
SEM.
44
Meng et al. evaluated a Mg-Bi interphase, but its impact
on Mg morphology was not discussed.
45
Inspired by the initial work on alloy interphases, here we show an
arti
fi
cial interphase based on In metal and its alloys with Mg to
reduce Mg dendrites and cell shorting. We then investigate the
mechanism of the alloy interphase during reduction. The alloy-based
interphase is prepared by a simple chemical redox reaction of Mg
metal with InBr
3
in solution. The Mg-In interphase suppresses Mg
dendrites, as observed by SEM. In a symmetric cell with glassy
fi
ber
separators, the Mg-In interphase results in a signi
fi
cant increase in
the cycling lifetime before shorting. Ex-situ XRD shows the Mg
content in the interphase increases as Mg
2
+
is reduced, indicating
that magnesiation of the alloy phases occurs upon reduction instead
of Mg metal plating. Cross-section SEM images show Mg deposi-
tion underneath the interphase, con
fi
rming Mg mobility through the
interphase. Electrochemical characterization reveals that the electro-
alloying of Mg into the Mg-In interphase is kinetically slow,
resulting in a low threshold current density where exceeding said
current density results in Mg deposition on top of the Mg-In
interphase. The kinetic limitation is an important consideration to
understand the behavior and bene
fi
ts of the alloy interphase.
Experimental
General considerations.
All manipulations were performed in
aN
2
-
fi
lled glove box (MBraun,
<
1 ppm H
2
O and O
2
) unless
otherwise stated. Tetrahydrofuran (THF, Fisher Scienti
fi
c) was dried
on a solvent puri
fi
cation system then dried over 4
sieves before
use. Anhydrous MgCl
2
(99.9%, Fisher Scienti
fi
c), anhydrous AlCl
3
(99.999%, Sigma Aldrich), anhydrous InBr
3
(99.99%, Thermo
Scienti
fi
c), and anhydrous hexanes (99%, mixed isomers, Sigma
Aldrich) were used as received. Magnesium hexamethyldisilazide
(Mg(HMDS)
2
, 97%, Sigma Aldrich) was recrystallized in hexanes at
20 °C before use. Magnesium foil (99.9%, MTI Corporation,
0.1 mm thick) was cleaned with 0.1 M acetic acid in air, then
brought inside the glove box, where it was further polished with
320 then 1500 grit silicon carbide sanding paper (3M).
Mg-In electrode and electrolyte preparation.
A 20 ml scintil-
lation vial was charged with 17.8 mg of InBr
3
, which was then
dissolved in 5 ml of THF. The Mg foil was punched with a 6 mm
hammer-driven punch or cut into appropriate pieces, then placed in
the reaction vessel for
approx.
2 h. The foil was then washed
thoroughly in THF before use. The magnesium-aluminum chloride
complex (MACC) electrolyte was prepared in an N
2
-
fi
lled glove
box according to Barile et al. with the addition of Mg(HMDS)
2
as in
Kim et al. in 5 ml batches.
15
,
17
Solutions of MACC
+
Mg(HMDS)
2
(30 mM AlCl
3
+
60 mM MgCl
2
+
10 mM Mg(HMDS)
2
) were
prepared by adding 2.5 ml of chilled THF (cooled to approximately
0 °C on a Peltier plate) dropwise to anhydrous AlCl
3
(20 mg). THF
(2.5 ml) was added to anyhdrous MgCl
2
(28.5 mg) and allowed to
stir for 1 min. The AlCl
3
was completely dissolved in THF to yield a
colorless solution. The AlCl
3
and MgCl
2
solutions were combined,
and the resulting solution was stirred at 420 rpm until it turned clear
and colorless (
approx.
6 h). The electrolyte was subsequently
conditioned by adding Mg(HMDS)
2
(17.5 mg) and allowed to stir
until the solution turned clear.
Electrochemical measurements.
The galvanostatic cycling and
long term deposition experiments were conducted on a VMP3
potentiostat (Bio-Logic) in a two-electrode geometry with the
reference and the counter electrodes shorted. Cells were assembled
in a 0.25 in inner diameter polytetra
fl
uoroethylene (PTFE) Swagelok
cells with Mo current collectors. GF/D (Whatman)
fi
lters were used
as separators for the symmetric cell galvanostatic cycling experi-
ments with 0.1 ml of Mg(HMDS)
2
electrolyte, at 0.025, 0.05, and
0.5 mA cm
2
, with 1 h half-cycles. The GF/D
fi
lters were punched
with a 0.25 in diameter hammer-driven punch and dried at 80 °C
under reduced pressures before use. The long-term deposition
experiments for ex-situ characterizations were done using a
Mg
Mg-In working electrode and a Mg foil counter electrode.
PTFE donut separator (0.25 in outer diameter, 0.125 in inner
diameter) to eliminate scaffolding effects of a physical separator.
The wells of the separator were
fi
lled with 0.05 ml of
MACC
+
Mg(HMDS)
2
electrolyte.
Tandem deposition-cyclic voltammogram experiments for ki-
netics.
The experiments were concuted on a VMP3 potentiostat
(Bio-Logic) in a three-electrode geometry. A Ag/Ag
2
S reference was
used,
46
which is well-characterized in the MACC
+
Mg(HMDS)
2
system.
47
Cells were assembled in a glass four-neck heart cell using
electrode holders (Gamry Li BMC 1.5 mm substrate holder).
The experiment consists of an initial deposition step for
approx.
4
μ
Ah cm
2
, then a 5 s open circuit hold, then a cyclic voltammetry
(CV) scan from open circuit potential to
1 V then back to
1.7 V
vs Ag/Ag
2
S at 5 mV s
-1
. This sequence is repeated for different
current densities, at 0.01, 0.025, 0.05, 0.075, 0.1, and 0.5 mA cm
2
.
Physical characterizations.
All characterization was completed
on working electrodes after rinsing with 1 ml of THF and drying in a
N
2
glove box. SEM images were taken with a ZEISS 1150 variable
pressure
fi
eld emission scanning electron microscope with a 15 kV
accelerating voltage and an in-lens secondary electron detector.
Energy dispersive X-ray spectroscopy (EDS) data were collected
using an Oxford X-Max Silicon Drift Detectors X-ray energy
dispersive spectrometer with a 15 kV accelerating voltage.
Samples were brie
fl
y exposed to air during transfer into the
instrument. Cross sections of electrode samples were taken by
cooling the electrode with liquid nitrogen, then cutting with a
scalpel. XRD patterns of the Mg-In electrodes were collected using
the Rigaku SmartLab diffractometer equipped with a HyPix-3000
detector and a Cu K
α
X-ray source with a 40 kV accelerating
voltage, using a Rigaku airtight sample holder. The patterns were
collected from 10 to 50° 2
θ
at 3° per minute in a Bragg-Brentano
geometry. The grazing incidence XRD pattern of the pristine
Mg
Mg-In electrode was collected in a parallel beam geometry
with an
ω
of 0.3°.
Results and Discussion
First we discuss the preparation of the Mg-In alloy interphase.
The modi
fi
ed electrode is prepared by a chemical reaction of Mg
metal with 50 mM InBr
3
in tetrahydrofuran (THF). Mg is a stronger
reducing agent compared to In with an E
0
of
2.36 V vs the standard
hydrogen electrode (SHE)
48
compared to
0.34 V
49
vs SHE for In.
The Mg metal can therefore chemically reduce the In
3
+
in solution.
To determine the fate of the Br
, EDS is measured on the Mg metal
surface after reaction (Fig. S3). Negligible quantities of Br are
detected in the interphase via EDS (roughly estimated at
<
1% by
weight), which suggests the MgBr
2
precipitates out into solution
rather than on the interface. After reaction with InBr
3
, the foil turns
gray. If the foil is allowed to react for a longer period, on the order of
20 h, then the surface layer
fl
akes off revealing reactive Mg metal.
Thus, the reaction is not self-limiting. To evaluate the morphology of
the solid reaction product on the Mg surface, SEM images are taken
of the Mg metal surface before and after a 2 h reaction with InBr
3
.
The SEM images is shown in Figs.
1
a,
1
b. The Mg foil before the
reaction with InBr
3
is smooth with some grooves from the polishing
process with sandpaper. After the reaction, the surface layer is rough
and uneven, but covers the entire exposed Mg surface.
The XRD pattern measured in Bragg-Brentano geometry of the
Mg foil after the reaction with InBr
3
is shown in Fig.
1
c. The (002)
re
fl
ection of Mg metal is the strongest peak, suggesting that the
X-rays can penetrate through the surface
fi
lm to the Mg metal
substrate (Fig. S1). However, additional peaks are observed sug-
gesting the surface layer has crystalline domains. The additional
Journal of The Electrochemical Society
, 2024
171
010513
peaks can be assigned via a four phase Rietveld re
fi
nement to In
metal and two different Mg-In compounds in the binary phase space:
β
and
β
′′
.
50
The
β
phase is a disordered
fcc
structure that forms at
high In content at lower temperatures, and
<
50% In content at
higher temperatures.
51
The
β
′′
phase is an ordered tetragonal
structure that forms near 50% In content. The
β
′′
phase forms
upon slow cooling between In contents of 35 to 50%, suggesting the
phase is thermodynamically favorable at the speci
fi
ced In contents.
52
Both the alloys have a moderately wide solid-solubility window near
room temperature (
approx.
10% in composition), which may enable
the electroalloying of Mg upon reduction. Both In and Mg metals
have some degree of solubility as well. From the quality of the
diffraction pattern collected, we were unable to determine the exact
Mg-In contents within each phase. We can estimate the phase
fractions of each crystalline phase within the Mg-In interphase from
the Rietveld re
fi
nement by excluding the Mg fraction due to the
visible contributions from the Mg foil substrate. The phase fractions
of the phases given in weight percent are 62% In, 15%
β
′′
, and
23%
β
. Because multiple phases are formed after reaction with
InBr
3
, we will refer to the surface
fi
lm as the Mg-In interphase
hereafter. Mg foil electrodes with the Mg-In interphase will be
referred to as Mg
Mg-In electrodes.
To investigate the effects of the Mg-In interphase on the
electrochemical stripping and deposition of Mg
2
+
, the Mg
Mg-In
electrode is galvanostatically cycled in a symmetric cell and
compared to a bare Mg control. The galvanostatic cycling data are
shown in Fig.
2
. The electrolyte of choice for the electrochemical
characterization of the electrode is a MACC electrolyte chemically
conditioned with Mg(HMDS)
2
.
53
MACC
+
Mg(HMDS)
2
is chosen
for its high anodic stability compared to Grignard-based electrolytes
and its potential susceptibility for dendrites due to its low Mg
2
+
concentration (60 mM) that causes low conductivity. A PTFE
Swagelok union cell with Mo current collectors are used to avoid
stainless steel corrosion with the chloride-containing electrolyte. At
a current density of 0.025 mA cm
2
, bare Mg metal electrodes show
a rapid increase in the plating and stripping overpotential until erratic
spikes are observed after about 9 cycles, which is attributed to
electrical soft shorts in the cell (Fig.
2
a). Upon disassembly of the
cell, black deposits are observed on both sides of the porous glassy
fi
ber separator which we assume to be Mg. In contrast, the Mg
Mg-
In electrodes exhibit a much slower and gradual increase in over-
potential, and cell shorting is not observed until after 144 cycles
Figure 1.
SEM image of the Mg metal foil (a) before and (b) after reaction
with InBr
3
in THF for 2 h. A rough surface
fi
lm is observed after reaction
across the foil. (c) XRD pattern of the Mg foil after reaction with InBr
3
with
a four phase Rietveld re
fi
nement to Mg, In, and the Mg-In alloys
β
and
β
′′
,
collected in a Bragg-Brentano geometry.
Figure 2.
Galvanostatic cycling of Mg metal and Mg
Mg-In electrodes at (a)
0.025 mA cm
2
and (b) 0.5 mA cm
2
using MACC
+
Mg(HMDS)
2
electrolyte with a glassy
fi
ber separator. Each half-cycle is 1 h. A higher
overpotential and faster shorting behavior is observed with bare Mg metal
electrodes when compared to Mg
Mg-In electrodes at low current densities.
When cycling at 0.5 mA cm
2
, there are no meaningful differences between
the galvanostatic behavior of bare Mg and Mg
Mg-In electrodes.
Figure 3.
SEM images of the (a), (b) bare Mg and (c), (d) Mg
Mg-In
electrode after applying
0.05 mA cm
2
for 10 h. A PTFE donut separator is
used to prevent the separator from affecting the morphology. A hetero-
geneous deposit is clearly observed in (a), (b) while absent in (c), (d).
Journal of The Electrochemical Society
, 2024
171
010513
(Fig.
2
a). The increasing overpotentials upon galvanostatic cycling
suggests an instability in the Mg
Mg-In electrode upon extended
cycling, likely the result of an insulating interphase forming on top
of the Mg-In alloy, or a mechanical breakdown of the Mg-In alloy.
However, compared to galvanostatic cycling results of bare Mg
electrodes, the overpotential increase is much slower, suggesting an
improvement in the electrode stability. The galvanostatic cycling
behavior at 0.05 mA cm
2
is shown in Fig. S4, and similar trends
hold. However, 0.05 mA cm
2
is not a commercially relevant
current density. As such, the symmetric cells were cycled at elevated
current densities to study the kinetic limitations of the Mg
Mg-In
electrodes. The galvanostatic cycling performance of the bare Mg
and Mg
Mg-In electrodes at a current density of 0.5 mA cm
2
are
shown in Fig.
2
b. No meaningful differences in electrochemical
performance between bare Mg and the Mg
Mg-In electrode are
observed, suggesting a possible kinetic limitation with the inter-
phase.
First, we probe the low current reduction behavior to investigate
the changes to morphology. SEM images of the electrode surfaces
are taken after applying
0.05 mA cm
2
for 10 h. To eliminate the
effect of the separator scaffolding on the deposition morphology, a
donut PTFE separator is used in a Swagelok cell. The well of the
separator is
fi
lled with the electrolyte solution, thus providing a
planar electrode surface with no scaffolding to seed nucleation or
affect growth morphology. The SEM image of the bare Mg electrode
after reduction is shown in Figs.
3
a and
3
b. Small spherical Mg
clusters are found scattered across the electrode. This morphology is
similar to the Mg deposits observed by Eaves-Rathert and cow-
orkers, arising from hemispherical deposition of Mg under moderate
overpotentials.
33
The SEM of the Mg
Mg-In electrode after reduc-
tion is shown in Figs.
3
c and
3
d. In contrast to the bare Mg electrode,
the Mg
Mg-In electrode morphology changes minimally with the
rough surface becoming marginally less fractal. No hemispherical
growths are observed suggesting Mg is not plated on the surface. To
determine if the Mg-In alloy remains intact following an oxidation,
the Mg
Mg-In electrode is oxidized at 0.05 mA cm
2
and imaged.
The SEM and corresponding EDS maps are found in the SI in
Fig. S18. In is still present homogenously dispersed on the electrode
surface following oxidation, suggesting that at least for short term
cycling, the Mg-In alloy remains intact.
However, the signi
fi
cantly different galvanostatic cycling beha-
vior of the Mg
Mg-In electrode at a higher current density suggests
different processes are occuring at high current densities. To study
the surface morphology after a high current reduction step, SEM
images of the Mg
Mg-In electrode surface after a reduction at
0.5 mA cm
2
are collected. The SEM image and the EDS
elemental maps of the Mg
Mg-In electrode surface after the
reduction are shown in Figs.
4
a
4
c. Contrary to the morphology at
low current density, deposits are observed on top of the electrode
with no In present in the deposits, indicating the deposits are Mg
metal. Thus, there appears to be a rate at which Mg will deposit on
top of the interphase.
To electrochemically probe if Mg is deposited at the electrode
surface, i.e. on top of the Mg-In surface, a two-stage electrochemical
experiment is devised. First, the Mg
Mg-In electrode is reduced at a
constant current density. Then, a positive sweep cyclic voltammo-
gram (CV) is collected. The experiment is repeated with increasing
current densities. Two CVs from the experiment are presented in
Fig.
4
d, one after a deposition at 0.025 mA cm
2
and another at
0.5 mA cm
2
. The CV after reduction at 0.5 mA cm
2
exhibits a
depletion effect where the current density decreases as the potential
is swept positively. This depletion effect in the CV is attributed to
stripping of freshly-deposited surface Mg on top of the Mg-In
interphase, such as the deposit imaged in Fig.
4
a. Once all the
freshly-deposited Mg is stripped off, the current density decreases. A
similar phenomenon was observed by Melemed et al. where a
depletion effect was observed via CVs after a deposition of fresh Ca
on top of a passivation layer.
54
However, no such depletion effect is
observed in the CV after reduction at 0.025 mA cm
2
, which
indicates no such deposited Mg is available on the surface of the
electrode. A clear kinetic limitation in the Mg-In interphase is
therefore observed where past a critical current density, the Mg-In
interphase acts as a deposition substrate instead of suppressing
dendrites.
To gain more insights into the reduction process in the Mg
Mg-In
electrode, we investigate the mechanism by characterizing the
Mg
Mg-In electrode after reduction ex-situ. First, the XRD pattern
of the Mg
Mg-In electrodes are collected after reduction at low
current densities to observe the changes in the crystalline domains in
the interphase. In Fig.
5
, the XRD patterns of the Mg
Mg-In
electrode before and after reduction are shown. The patterns can
be
fi
t via a Rietveld re
fi
nement to the four phases identi
fi
ed in
Fig.
1
c: Mg, In, and the binary Mg-In alloys
β
and
β
′′
. The largest
peak is the Mg(002) re
fl
ection with a preferential orientation for the
(002) plane, which suggest that the X-rays penetrate through the
entire Mg-In interphase to the Mg foil substrate. Therefore, we
assume that all the crystalline phases of the Mg-In interphase are
captured by XRD. All the In-containing phases can be
fi
t well due to
isolated peaks. The intensities of the two XRD patterns are normal-
ized to the Mg(101) re
fl
ection to better compare the relative
intensities of the re
fl
ections of different phases. After reduction of
the Mg
Mg-In electrode, an increase in the intensity of the re
fl
ec-
tions attributed to the
β
′′
phase is observed. The increase in intensity
is most evident in the
β
′′
(100) re
fl
ection near 27° 2
θ
. A corre-
sponding decrease in the intensity in the In re
fl
ections is observed,
easily observable in the In(101) and In(002) re
fl
ections near 33° 2
θ
and 36° 2
θ
, respectively. The calculated phase fractions from the
Rietveld re
fi
nement are shown in Table
I
. Upon reduction of the
Mg
Mg-In electrode, a clear decrease in the In phase fraction is
observed, as well as a slight decrease in the In-rich
β
phase alloy. A
signi
fi
cant increase in the Mg-rich
β
′′
phase is observed. The XRD
patterns thusly indicate that the crystalline domains of the Mg-In
interphase magnesiates upon reduction. Based on the diffraction
patterns, two possible mechanisms can be posed. The Mg-In alloy
could act essentially as an alloy electrode deposited on a conductive
substrate (Mg), where the redox occurs only on the deposited alloys.
Figure 4.
(a) SEM image, (b) Mg EDS map, and (c) In EDS map of the
Mg
Mg-In electrode after a deposition of 1.25 mAh cm
2
at a current density
of 0.5 mA cm
2
. Large Mg deposits are observed on the surface. (d) Cyclic
voltammograms of the Mg
Mg-In electrode swept at 5 mV s
1
after a
deposition of 3 mC, at a current density of either 0.025 mA cm
2
or
0.5 mA cm
2
. The small anodic peak observed after higher current density
deposition is an indicator of Mg deposits on top of the Mg
Mg-In electrode.
Journal of The Electrochemical Society
, 2024
171
010513
Indeed, In and the
β
′′
phase alloy have been shown in literature to be
capable of acting as Mg bulk alloy electrodes with changing Mg
contents as the electrode reduces and oxidizes.
55
A similar phenom-
enon is observed in the Mg-Hg amalgam surface system, where upon
reduction of the surface modi
fi
ed electrode, an increase in magne-
siation in the Mg-Hg alloy on the surface of the electrode is
observed. Based on the increase in magnesiation upon reduction in
the alloy interphase, the authors suggest that the amalgam surface is
not active as a protective surface, but rather as a simple Mg-Hg alloy
electrode deposited on top of an inert Mg metal substrate.
56
However, an alternate mechanism can be hypothesized, wherein
Mg mobility through the interphase can occur alongside alloying
into the interphase, allowing Mg to plate underneath the interphase.
To determine whether Mg mobility through the Mg-In interphase
is possible, ex-situ cross-section SEM images of the Mg
Mg-In
electrodes before and after reduction are collected. The cross-section
SEM image of the Mg
Mg-In electrode before reduction with the
corresponding EDS map is shown in Figs.
6
a
6
d. The EDS map
shows a mix of both Mg and In in the interphase layer, which is
approx.
5
μ
m thick. Some O is found on the surface layer, likely due
to air exposure during sample preparation and transfer.
The cross-section SEM image of the Mg
Mg-In electrode after
reduction is shown in Figs.
6
e
6
h. Mg deposits of a different
morphology from the Mg foil substrate are observed beneath the
Mg-In interphase. Interestingly, the deposits do not form a uniform
layer underneath, instead forming localized veins. We hypothesize
that Mg nucleation occurs preferentially underneath the interphase
where the interphase is the thinnest, due to the kinetic limitations of
the interphase. Upon nucleation, further Mg deposition preferentially
occurs at these sites. The deposited Mg also exhibits more oxidation
during sample preparation compared to the Mg substrate as
evidenced by the O EDS map (Figs.
6
d and
6
h). The increased
surface area of the deposited Mg compared to the densely packed
Figure 5.
XRD patterns of the Mg
Mg-In electrode in a (a) pristine
condition and (b) after a 0.025 mA cm
2
reduction for 100 h. The patterns
are
fi
t with a four-phase Rietveld re
fi
nement that includes In, Mg, and two
Mg-In phases
β
and
β
′′
. Upon reduction, a growth of the Mg-rich
β
′′
phase is
observed with a corresponding decrease in In metal.
Table I. Phase fractions of the In-related phases present in the
interphase layer, as calculated from the Rietveld re
fi
nement.
Phases
a
)
Pristine (wt%)
Reduced (wt%)
In
62%
16%
β
′′
15%
64%
β
23%
20%
a) The Mg phase is excluded as the X-ray penetrates through the alloy
layer into the Mg substrate. The remaining three phases are renormalized.
Figure 6.
(a) Cross-section SEM image of the pristine Mg
Mg-In electrodes and corresponding EDS maps for (b) Mg, (c) In, and (d) O. (e) Cross-section SEM
images of the Mg
Mg-In electrode after a 100 hrdeposition at 0.025 mA cm
2
and corresponding EDS maps for (f) Mg, (g) In, and (h) O. Compared to the
pristine Mg
Mg-In electrode, the Mg underneath the Mg-In layer is clearly heterogeneous, with freshly deposited veins of Mg visible in the Mg and O maps.
Journal of The Electrochemical Society
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171
010513
commercial Mg foil likely causes faster oxidation in the deposited
Mg. Some Mg deposits are also observed on top of the interphase,
which is attributed to the depletion effects associated with the
extreme length of the reduction necessary to deposit enough Mg
underneath the interphase to observe with an SEM (100 h rat
0.025 mA cm
2
) (Fig. S17). No Mg deposits are observed when a
shorter reduction time was used, as seen above in Figs.
3
c
3
d.
Discussion
Next we discuss the results in the context of metal deposition
models. Chazalviel and coworkers described a model in which the
dendritic growth of metals is caused by the formation of a space
charge region at the surface of the cathode, due to repulsion of
anions and consumption of cations.
29
,
57
The model is used to
describe the diffusion-limited aggregation (DLA) that causes the
classic branching dendritic morphology. However, the deposits
observed in Figs.
3
a and
3
b do not resemble a DLA morphology.
Instead, the deposits observed on a bare Mg electrode are better
described as hemispherical growth. Hemispherical growth occurs
under mixed diffusion and kinetic control at intermediate current
densities. Unlike a DLA regime, full mass transport limitation is not
yet achieved.
58
A cartoon of the reduction process with a bare Mg
electrode is shown in Fig.
7
a. The kinetics of the deposition overtake
the kinetics of the self-diffusion of Mg across the surface. As such,
instead of forming a smooth surface and minimizing the surface
energy, the cations reduce as soon as they arrive on the surface.
Based on our experimental observations, we theorize that the
electroreduction process at the Mg
Mg-In electrode is different.
Processes occurring on the Mg
Mg-In electrode are governed by
different rates. In Fig.
7
b, the same reduction process is shown
instead with an alloy interphase. In the Mg
Mg-In electrode, upon
reduction of the Mg
2
+
cation at the surface of the electrode, the
cation alloys into the Mg-In interphase. The alloying process appears
more facile than the self-diffusion of Mg, indicated by the suppres-
sion of dendrites at same current densities. As such, even at
moderate current densities, the alloying process can kinetically
keep up with the cation
fl
ux. The concentration gradient in the
interphase then causes Mg to diffuse, depositing underneath the
interphase at the Mg
Mg-In interface. At elevated current densities
however, the diffusion through the Mg-In interphase becomes
kinetically limiting, resulting in Mg deposition on top of the
interphase.
The kinetic limitation of the alloy interphase controlling the
deposition location of the metal is also observed in literature in Li
alloy interphase systems. In a Li-Sn interphase system, Li plates
underneath the interphase below the exchange current density of the
Li
Li-Sn electrode. However, upon increasing the deposition current
density past the exchange current density, Li begins plating on top of
the Li-Sn interphase.
59
In another study investigating a Li-Sn
interphase system, computed surface energies of the electrode
surface suggest that at low current densities, plating underneath
the interphase is energetically favorable, but at high current densities
Li accumulates at the surface and delaminates.
60
Conclusions
Herein, we demonstrate more uniform reduction with a Mg
electrode with a Mg-In alloy interphase compared to Mg metal. The
Mg
Mg-In electrode is easily prepared by a chemical reaction of a
Mg foil with a solution of InBr
3
in THF. The Mg
Mg-In electrode is
capable of undergoing more deposition and stripping cycles before
evidence of soft shorts are observed compared to Mg metal. While
Mg electrodes grow heterogeneous Mg deposits upon reduction, the
Mg
Mg-In electrodes do not under the same conditions and instead
undergo alloying reactions followed by Mg deposition occurring
below the alloy. However, the Mg
Mg-In electrode is kinetically
limited by the alloying process. At higher current densities, the
alloying kinetics cannot keep up with the rate of Mg
2
+
reduction and
Mg deposits on top of the Mg-In interphase. Improved alloy
interphases thus must have higher Mg mobility, again highlighting
the crucial role of divalent solid-state ion conductivity in Mg-based
electrochemical cells. Higher mobility could be achieved by
targeting a new phase or leveraging defect engineering strategies.
Nevertheless, the Mg-In system serves as a proof-of-concept system
for the use of alloy interphases to leverage alloy solid solubility to
reduce Mg deposition and cell shorting, an issue that is likely to
impact commercial cells due to the use of separators.
Acknowledgments
This research was supported by the Packard Fellowship for
Science and Engineering. KAS acknowledges support from the
Alfred P. Sloan Foundation and the Camille and Henry Dreyfus
Foundation. X-ray diffraction data were collected at the X-ray
Crystallography Facility in the Beckman Institute of the California
Institute of Technology. SEM and EDS data were collected at the
GPS Division Analytical Facility of the California Institute of
Technology. The authors thank Dr. Forrest Laskowski and Dr. Chi
Ma for insightful discussions.
ORCID
Brian C. Lee
https://orcid.org/0000-0002-0898-0838
Kimberly A. See
https://orcid.org/0000-0002-0133-9693
References
1. U.S. Geological Survey, Mineral Commodity Summaries 2019; 2019.
2. E. Peled and H. Straze,
The kinetics of the magnesium electrode in thionyl
chloride solutions.
J. Electrochem. Soc.
,
124
, 1030 (1977).
3. L. P. Lossius and F. Emmenegger,
Plating of magnesium from organic solvents.
Electrochim. Acta
,
41
, 445 (1996).
4. C. Liebenow,
Reversibility of electrochemical magnesium deposition from
grignard solutions.
J. Appl. Electrochem.
,
27
, 221 (1997).
5. Z. Lu, A. Schechter, M. Moshkovich, and D. Aurbach,
On the electrochemical
behavior of magnesium electrodes in polar aprotic electrolyte solutions.
J. Electroanal. Chem.
,
466
, 203 (1999).
6. J. G. Connell, B. Genorio, P. P. Lopes, D. Strmcnik, V. R. Stamenkovic, and
N. M. Markovic,
Tuning the reversibility of mg anodes via controlled surface
passivation by h2o/cl
in organic electrolytes.
Chem. Mater.
,
28
, 8268 (2016).
Figure 7.
Cartoon illustrations of the Mg
2
+
reduction process at (a) bare Mg
and (b) Mg
Mg-In electrode surfaces. In bare Mg, the rate of self-diffusion
(k
s
) of Mg across the surface is in competition with the rate of deposition
(k
dep
), while in Mg
Mg-In electrodes, the rate of alloying (k
alloy
) competes
against the rate of reduction (k
red
) to determine the morphology.
Journal of The Electrochemical Society
, 2024
171
010513
7. L. W. Gaddum and H. E. French,
The electrolysis of grignard solutions.
J. Am.
Chem. Soc.
,
49
, 1295 (1927).
8. W. V. Evans and R. Pearson,
The ionic nature of the grignard reagent.
J. Am.
Chem. Soc.
,
64
, 2865 (1942).
9. J. D. Genders and D. Pletcher,
Studies using microelectrodes of the mg(ii)/mg
couple in tetrahydrofuran and propylene carbonate.
J. Electroanal. Chem. Interfac.
Electrochem.
,
199
, 93 (1986).
10. D. Aurbach, H. Gizbar, A. Schechter, O. Chusid, H. E. Gottlieb, Y. Gofer, and
I. Goldberg,
Electrolyte solutions for rechargeable magnesium batteries based on
organomagnesium chloroaluminate complexes.
J. Electrochem. Soc.
,
149
, A115
(2002).
11. O. Mizrahi, N. Amir, E. Pollak, O. Chusid, V. Marks, H. Gottlieb, L. Larush,
E. Zinigrad, and D. Aurbach,
Electrolyte solutions with a wide electrochemical
window for rechargeable magnesium batteries.
J. Electrochem. Soc.
,
155
, A103
(2007).
12. Y. Vestfried, O. Chusid, Y. Goffer, P. Aped, and D. Aurbach,
Structural analysis
of electrolyte solutions comprising magnesium-aluminate chloro-organic com-
plexes by raman spectroscopy.
Organometallics
,
26
, 3130 (2007).
13. R. E. Doe, R. Han, J. Hwang, A. J. Gmitter, I. Shterenberg, H. D. Yoo, N. Pour, and
D. Aurbach,
Novel, electrolyte solutions comprising fully inorganic salts with high
anodic stability for rechargeable magnesium batteries.
Chem. Commun.
,
50
, 243
(2014).
14. T. Liu, Y. Shao, G. Li, M. Gu, J. Hu, S. Xu, Z. Nie, X. Chen, C. Wang, and J. Liu,
A facile approach using mgcl 2 to formulate high performance Mg 2. electrolytes
for rechargeable mg batteries.
J. Mater. Chem. A
,
2
, 3430 (2014).
15. C. J. Barile, E. C. Barile, K. R. Zavadil, R. G. Nuzzo, and A. A. Gewirth,
Electrolytic conditioning of a magnesium aluminum chloride complex for
reversible magnesium deposition.
J. Phys. Chem. C
,
118
, 27623 (2014).
16. J. H. Ha, B. Adams, J.-H. Cho, V. Duffort, J. H. Kim, K. Y. Chung, B. W. Cho,
F. L. Nazar, and S. H. Oh,
A conditioning-free magnesium chloride complex
electrolyte for rechargeable magnesium batteries.
J. Mater. Chem. A
,
4
, 7160
(2016).
17. S. S. Kim, S. C. Bevilacqua, and K. A. See,
Conditioning-free mg electrolyte by
the minor addition of Mg(HMDS)2.
ACS Appl. Mater. Interfaces
,
12
, 5226
(2020).
18. D.-T. Nguyen, A. Y. S. Eng, M.-F. Ng, V. Kumar, Z. Sofer, A. D. Handoko,
G. S. Subramanian, and Z. W. Seh,
A high-performance magnesium tri
fl
ate-based
electrolyte for rechargeable magnesium batteries.
Cell Rep. Phys. Sci.
,
1
, 100265
(2020).
19. S.-Y. Ha, Y.-W. Lee, S. W. Woo, B. Koo, J.-S. Kim, J. Cho, K. T. Lee, and
N.-S. Choi,
Magnesium(II) bis(tri
fl
uoromethane sulfonyl) imide-based electrolytes
with wide electrochemical windows for rechargeable magnesium batteries.
ACS
Appl. Mater. Interfaces
,
6
, 4063 (2014).
20. I. Shterenberg, M. Salama, H. D. Yoo, Y. Gofer, J.-B. Park, Y.-K. Sun, and
D. Aurbach,
Evaluation of (CF3SO2)2N- (TFSI) based electrolyte solutions for
Mg batteries.
J. Electrochem. Soc.
,
162
, A7118 (2015).
21. M. Salama et al.,
Unique behavior of dimethoxyethane (DME)/Mg(N(SO2CF3)2)
2 solutions.
J. Phys. Chem. C
,
120
, 19586 (2016).
22. J. T. Herb, C. A. Nist-Lund, and C. B. Arnold,
A
fl
uorinated alkoxyaluminate
electrolyte for magnesium-ion batteries.
ACS Energy Lett.
,
1
, 1227 (2016).
23. Z. Zhao-Karger, M. E. G. Bardaji, O. Fuhr, and M. Fichtner,
A new class of non-
corrosive, highly ef
fi
cient electrolytes for rechargeable magnesium batteries.
J. Mater. Chem. A
,
5
, 10815 (2017).
24. R. Mohtadi, M. Matsui, T. S. Arthur, and S.-J. Hwang,
Magnesium borohydride:
from hydrogen storage to magnesium battery.
Angew. Chem. Int. Ed.
,
51
, 9780
(2012).
25. T. J. Carter, R. Mohtadi, T. S. Arthur, F. Mizuno, R. Zhang, S. Shirai, and
J. W. Kampf,
Boron clusters as highly stable magnesium-battery electrolytes.
Angew. Chem. Int. Ed.
,
53
, 3173 (2014).
26. M. Matsui,
Study on electrochemically deposited Mg metal.
J. Power Sources
,
196
, 7048 (2011).
27. M. Jäckle and A. Groß,
Microscopic properties of lithium, sodium, and magnesium
battery anode materials related to possible dendrite growth.
J. Chem. Phys.
,
141
,
174710 (2014).
28. C. Hendricks, N. Williard, S. Mathew, and M. Pecht,
A failure modes,
mechanisms, and effects analysis (fmmea) of lithium-ion batteries.
J. Power
Sources
,
297
, 113 (2015).
29. J.-N. Chazalviel,
Electrochemical aspects of the generation of rami
fi
ed metallic
electrodeposits.
Phys. Rev. A
,
42
, 7355 (1990).
30. V. Fleury, M. Rosso, J.-N. Chazalviel, and B. Sapoval,
Experimental aspects of
dense morphology in copper electrodeposition.
Phys. Rev. A
,
44
, 6693 (1991).
31. R. M. Brady and R. C. Ball,
Fractal growth of copper electrodeposits.
Nature
,
309
, 225 (1984).
32. R. Davidson et al.,
Formation of magnesium dendrites during electrodeposition.
ACS Energy Lett.
,
4
, 375 (2019).
33. J. Eaves-Rathert, K. Moyer, M. Zohair, and C. L. Pint,
Kinetic- versus diffusion-
driven three-dimensional growth in magnesium metal battery anodes.
Joule
,
4
,
1324 (2020).
34. S. Hebié, H. P. K. Ngo, J.-C. Leprêtre, C. Iojoiu, L. Cointeaux, R. Berthelot, and
F. Alloin,
Electrolyte based on easily synthesized, low cost triphenolate
-
borohydride salt for high performance Mg(TFSI)2-glyme rechargeable magnesium
batteries.
ACS Appl. Mater. Interfaces
,
9
, 28377 (2017).
35. S. Hebié, F. Alloin, C. Iojoiu, R. Berthelot, and J.-C. Leprêtre,
Magnesium
anthracene system-based electrolyte as a promoter of high electrochemical
performance rechargeable magnesium batteries.
ACS Appl. Mater. Interfaces
,
10
, 5527 (2018).
36. M. S. Ding, T. Diemant, R. J. Behm, S. Passerini, and G. A. Gif
fi
n,
Dendrite growth
in Mg metal cells containing Mg(TFSI)2/glyme electrolytes.
J. Electrochem. Soc.
,
165
, A1983 (2018).
37. S.-B. Son, T. Gao, S. P. Harvey, K. X. Steirer, A. Stokes, A. Norman, C. Wang,
A. Cresce, K. Xu, and C. Ban,
An arti
fi
cial interphase enables reversible
magnesium chemistry in carbonate electrolytes.
Nat. Chem.
,
10
, 532 (2018).
38. Z. Liang and C. Ban,
Strategies to enable reversible magnesium electrochemistry:
from electrolytes to arti
fi
cial solid
electrolyte interphases.
Angew. Chem. Int. Ed.
,
60
, 11036 (2021).
39. Y. Zhao, A. Du, S. Dong, F. Jiang, Z. Guo, X. Ge, X. Qu, X. Zhou, and G. Cui,
A
bismuth-based protective layer for magnesium metal anode in noncorrosive
electrolytes.
ACS Energy Lett.
,
6
, 2594 (2021).
40. S. Shin, J. H. Kwak, S. H. Oh, H.-S. Kim, S.-H. Yu, and H.-D. Lim,
Reversible
Mg-metal batteries enabled by a Ga-rich protective layer through one-step interface
engineering.
ACS Appl. Mater. Interfaces
,
15
, 28684 (2023).
41. B. Yang, L. Xia, R. Li, G. Huang, S. Tan, Z. Wang, B. Qu, J. Wang, and F. Pan,
Superior plating/stripping performance through constructing an arti
fi
cial inter-
phase layer on metallic Mg anode.
J. Mat. Sci. Technol.
,
157
, 154 (2023).
42. B. Wan, H. Dou, X. Zhao, J. Wang, W. Zhao, M. Guo, Y. Zhang, J. Li, Z.-F. Ma,
and X. Yang,
Three-dimensional magnesiophilic scaffolds for reduced passivation
toward high-rate mg metal anodes in a noncorrosive electrolyte.
ACS Appl. Mater.
Interfaces
,
12
, 28298 (2020).
43. R. Lv, X. Guan, J. Zhang, Y. Xia, and J. Luo,
Enabling Mg metal anodes
rechargeable in conventional electrolytes by fast ionic transport interphase.
Natl.
Sci. Rev.
,
7
, 333 (2020).
44. C. Pechberty et al.,
Alloying electrode coatings towards better magnesium
batteries.
J. Mater. Chem. A
,
10
, 12104 (2022).
45. Z. Meng, Z. Li, L. Wang, T. Diemant, D. Bosubabu, Y. Tang, R. Berthelot,
Z. Zhao-Karger, and M. Fichtner,
Surface engineering of a mg electrode via a
new additive to reduce overpotential.
ACS Appl. Mater. Interfaces
,
13
, 37044
(2021).
46. C. Horwood and M. Stadermann,
Evaluation of a Ag/Ag2S reference electrode
with long-term stability for electrochemistry in ionic liquids.
Electrochem.
Commun.
,
88
, 105 (2018).
47. F. A. L. Laskowski, S. H. Stradley, M. D. Qian, and K. A. See,
Mg anode
passivation caused by the reaction of dissolved sulfur in Mg
S batteries.
ACS
Appl. Mater. Interfaces
,
13
, 29461 (2021).
48. A. J. Bard and L. R. Faulkner,
Electrochemical Methods
(John Wiley & Sons) 2nd
ed. (2001).
49. A. J. Bard, R. Parsons, and J. Jordan,
Standard Potentials in Aqueous Solution
(CRC Press, New York City, NY) 1st ed. (1985).
50. A. A. Nayeb-Hashemi and J. B. Clark,
The In-Mg (indium-magnesium) system.
Bull. Alloy Phase Diag.
,
6
, 149 (1985).
51. G. V. Raynor,
The constitution of the magnesium-indium alloys in the region 20 to
50 atomic per cent of indium.
Trans. Faraday Soc.
,
44
, 15 (1948).
52. K. Schubert, F. Gauzzi, and K. Frank,
Kristallstruktur einiger Mg-B3-Phasen.
Z.
Metallkde.
,
54
, 422 (1963).
53. H. S. Kim, T. S. Arthur, G. D. Allred, J. Zajicek, J. G. Newman, A. E. Rodnyansky,
A. G. Oliver, W. C. Boggess, and J. Muldoon,
Structure and compatibility of a
magnesium electrolyte with a sulphur cathode.
Nat. Commun.
,
2
, 427 (2011).
54. A. M. Melemed and B. M. Gallant,
Electrochemical signatures of interface-
dominated behavior in the testing of calcium foil anodes.
J. Electrochem. Soc.
,
167
, 140543 (2020).
55. F. Murgia, E. T. Weldekidan, L. Stievano, L. Monconduit, and R. Berthelot,
First
investigation of indium-based electrode in mg battery.
Electrochem. Commun.
,
60
,
56 (2015).
56. C. Pechberty, J.-B. Ledeuil, J. Allouche, R. Dedryvère, L. Stievano, and
R. Berthelot,
Surface amalgam on magnesium electrode: protective coating or
not?
Energy Technol.
,
11
, 2201098 (2023).
57. V. Fleury, J. N. Chazalviel, M. Rosso, and B. Sapoval,
The role of the anions in the
growth speed of fractal electrodeposits.
J. Electroanal. Chem. Interfacial
Electrochem.
,
290
, 249 (1990).
58. L. Guo and P. C. Searson,
On the in
fl
uence of the nucleation overpotential on
island growth in electrodeposition.
Electrochim. Acta
,
55
, 4086 (2010).
59. G. Whang, Q. Yan, D. Li, Z. Wei, D. Butts, P. Sautet, J. Luo, and B. Dunn,
Avoiding dendrite formation by con
fi
ning lithium deposition underneath Li
Sn
coatings.
J. Mater. Res.
,
36
, 797 (2021).
60. A. Hagopian, J. Touja, N. Louvain, L. Stievano, J.-S. Filhol, and L. Monconduit,
Importance of halide ions in the stabilization of hybrid sn-based coatings for
lithium electrodes.
ACS Appl. Mater. Interfaces
,
14
, 10319 (2022).
Journal of The Electrochemical Society
, 2024
171
010513