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Controlling diffusion in Ni/Al reactive multilayers by Nb-alloying
Volker Schnabel, Alla S. Sologubenko, Stefano Danzi, Güven Kurtuldu, and Ralph Spolenak
Citation: Appl. Phys. Lett. 111, 173902 (2017);
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Published by the American Institute of Physics
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APPLIED PHYSICS LETTERS 111, 173902 (2017)
Controlling diffusion in Ni/Al reactive multilayers by Nb-alloying
€ven Kurtuldu,2
Volker Schnabel,1,a) Alla S. Sologubenko,1 Stefano Danzi,1 Gu
and Ralph Spolenak
Laboratory for Nanometallurgy, Department of Materials, ETH Z€
urich, Vladimir-Prelog-Weg 5, 8093 Z€
Laboratory of Metal Physics and Technology, Department of Materials, ETH Z€
urich, Vladimir-Prelog-Weg 5,
8093 Z€
urich, Switzerland
(Received 4 September 2017; accepted 10 October 2017; published online 23 October 2017)
Metallic reactive multilayers are known as high energy-density storage systems. Conventionally,
these multilayers are tailored for high reaction rates with the purpose to achieve high maximum reaction temperatures and explosive-like behavior upon mixing. However, in some instances such as neutralization of biological hazards or chemical energy-storage systems, a low heat flow rate is desired.
In the present work, we show that Nb-alloying presents an efficient approach to stabilize the asdeposited state and to form a diffusion barrier in situ, effectively reducing the heat flow rate by more
than 50%. The validation of the concept is carried out by a comparative study of thermally induced
phase reactions in Ni/Al and (Nb-Ni)/Al reactive multilayers. Kinetics of the phase reactions in these
systems were followed by differential scanning calorimetry, analytical scanning transmission electron microscopy, and in situ electron diffraction analysis. The results confirm alloying as a design
strategy for tailoring reaction kinetics of reactive multilayers. Published by AIP Publishing.
Reactive multilayer thin films are chemical energy storage systems.1,2 These heterogeneous structures are composed
of alternating thin layers of constituents with a negative
enthalpy of mixing.1–3 Through external stimuli such as an
increase in temperature,4 a spark discharge,5 or mechanical
shock,6 an exothermic mixing between the constituents is
triggered. The released energy can be used, e.g., to solder
materials.7,8 The ternary (Ni-V)/Al reactive multilayer system,8–10 with high reaction rates and low ignition thresholds,
is one of the well-known examples for such applications. To
utilize the reactive multilayers as materials for long-term
thermal batteries11 or neutralization of biological hazards,12
the duration of the heat flow interval has to be increased,
requiring a decrease in the heat flow rate d2 H=ðdTdtÞ, where
H; T; and t are the enthalpy, temperature, and time of the
reaction, respectively. Thermal batteries built from reactive
multilayers show a high potential to replace more traditional
Fe/KClO4 pellet based systems.13 Moreover, compared to
the conventional thermal batteries, Ni/Al multilayer systems
do not produce toxic reaction residue while maintaining similar mass specific heats of reaction.1,13
The reaction heat, temperature, and reaction velocity
have been controlled by arranging three single element
layers, introducing different reaction interfaces.14 To
decrease the heat flow rate, low-density compacts of multilayer reactive particles15 and inert-mediated reactive multilayers16 have been proposed. In both cases, the heat flow rate
is critically dependent on the volumetric system density, detrimentally affecting the amount of energy that can be stored
by the system. In the present work, we propose an alloying
concept that has the potential to overcome this conflict.
Author to whom correspondence should be addressed: volker.schnabel@
Our approach is based on a comparative study of (NbNi)/Al and a reference Ni/Al reactive multilayer. We have
chosen Ni/Al as a reference reactive multilayer because it is
one of the most studied materials.1,17,18 In our work, differential scanning calorimetry (DSC) and analytical transmission electron microscopy (TEM) were complemented by
in situ heating TEM experiments. The scanning transmission
electron microscopy (STEM) operation mode of TEM using
atomic number-sensitive high-angle annular dark field
(HAADF) imaging was accompanied by high-resolution
energy-dispersive X-ray spectroscopy (EDS) analyses. For
EDS elemental distribution maps and the chemical composition evaluations, the K-series for Al and Ni and L-series for
Nb were used. The HAADF-STEM and EDS-STEM studies
were also complemented by electron diffraction analyses of
the system phase configurations.
Two reactive multilayer systems, Ni/Al and (Nb-Ni)/Al,
were synthesized by direct current magnetron sputtering on
Si-wafers. Nb was introduced only in the Ni-layer by cosputtering from two targets, Nb and Ni, simultaneously. The
total thickness of the films was 1200 nm with the bilayer
thickness of 60 nm. For deposition, circular targets with a
diameter of 76 mm were used. The set powers for Ni, Nb, and
Al were 250, 300, and 300 W, respectively. Figures 1(a)–1(d)
show the scanning electron microscopy (FEI Magellan 400)
micrographs and elemental distribution maps of the (Nb-Ni)/
Al and Ni/Al reactive multilayers in the as-deposited state,
respectively. The SEM images demonstrate the morphological similarity of the Ni/Al and (Nb-Ni)/Al multilayers. The
elemental distribution maps of two systems in the asdeposited state were acquired in the same conditions (STEM
probe size, duration, dwell-time, and resolution of the EDSSTEM scans). The evaluation of the chemical composition of
an Nb-Ni layer of the (Nb-Ni)/Al system in the as-deposited
state was performed by the “interactive TEM” evaluation
111, 173902-1
Published by AIP Publishing.
Schnabel et al.
FIG. 1. (a)–(d) SEM micrographs and elemental distribution maps of (NbNi)/Al (a) and (c) and Ni/Al (b) and (d) reactive multilayer systems in the
as-deposited state. (e) A dark field TEM image of a Nb-Ni-layer. The inset
shows an electron diffraction pattern from the (Nb-Ni)/Al multilayer. We
positioned the objective aperture on the wide diffuse ring, marked with the
red circle to obtain the dark field image. (f) The heat flow dH=dT of the
(Nb-Ni)/Al and Ni/Al multilayer in the upper two panels and the heat flow
rates d2 H=ðdTdtÞ of both systems in the bottom panel within a temperature
range of 100 to 500 C. A summary of the electron diffraction analysis presented in Fig. 2 is also provided here. The nano-crystalline Nb-Ni is labeled
with nc-Nb-Ni.
routine based on the Cliff-Lorimer algorithm of the ESPIRIT
software package of the SuperX-EDX. Figure 1(e) shows a
dark field (DF) TEM micrograph of a Nb-Ni layer acquired
with the objective aperture positioned on the broad diffuse
ring in the corresponding electron diffraction pattern (shown
in the inset). The occurrence of the contrast modulations
within the layers in DF TEM images infers that the Nb-Ni
layers are nano-crystalline with the crystallite size of about
1–2 nm.
Differential scanning calorimetry (DSC, METTLER
TOLEDO, DSC 1 STAR system) was employed to evaluate
the effect of Nb-alloying on the heat flow rate of Ni/Al reactive multilayers. Figure 1(f) shows the temperature dependence of heat flow for (Nb-Ni)/Al (top) and for Ni/Al
(middle) and the heat flow rate for both systems (bottom)
within a 100–500 C temperature range. A summary of the
phase identification analysis by in situ electron diffraction is
also included in Fig. 1(f).
The (Nb-Ni)/Al system exhibits a nearly monotonous
segment on the heat flow rate curve from 209 to 390 C [Fig.
1(f), bottom], followed by prominent rate modulations within
the 390 to 430 C temperature interval. At variance, the Ni/
Al system shows three distinct exothermic peaks within the
209 to 390 C temperature interval, which agrees well with
literature data19 on Ni/Al multilayers of 50 nm bilayer thickness. In the Ni/Al system, the highest rate increase of 0.22 W
g1 K1 occurs within the temperature interval of 202
to 243 C [bottom graph in Fig. 1(f)]. The maximal heat
flow rate transition of the (Nb-Ni)/Al reactive multilayer
within the 209 to 390 C temperature range corresponds to a
Appl. Phys. Lett. 111, 173902 (2017)
0.1 W g1 K1 heat flow rate. The total energy release in both
systems can be obtained by integrating the heat flow over time,
which yields 840 and 1070 J g1 for the (Nb-Ni)/Al and Ni/Al
reactive multilayer systems, respectively. The addition of Nb
to the Ni/Al reactive system leads therefore to only about 21%
reduction in total energy storage capacity while decreasing the
maximum of the heat flow rate by more than 50%.
To clarify the effect of phase transformations on the
observed difference in the heat flow rate, we performed electron diffraction studies during in situ TEM heating experiments on both Ni/Al and (Nb-Ni)/Al reactive multilayers.
The thermally induced evolution of the phase configurations
in both, Ni/Al and (Nb-Ni)/Al, reactive multilayers is presented in Figs. 2(a) and 2(b). The in situ electron diffraction
studies were performed within the temperature range from
20 to 500 C at a heating rate of 20 K min1 with 10 min isothermal segments for electron diffraction data acquisition.
For the diffraction patterns analysis, the 2D diffraction patterns were integrated into radial averages using the FIT2D
software.20 The logarithm of the electron diffraction intensity profile plotted as a function of a spatial frequency (Å1)
allowed the identification of the set of the peaks in the intensity profiles and assigning the set to a particular phase. The
temperatures discussed are unevenly spaced and guided by
the DSC curves presented in Fig. 1. As a reference, we present Ni/Al structural evolution in Fig. 2(a). In the temperature range from 20 to 170 C, only Ni (PDF 04-0850) and Al
(PDF 04-0787) phases were observed. At 190 C, the formation of Al9Ni2 (PDF 06-0699) was detected, which agrees
well with the literature data on Ni/Al reactive multilayers
with a bilayer thickness larger than 25 nm.19 The 200 C isotherm shows in addition to Al, Ni and Al9Ni2, also the occurrence of Al3Ni (PDF 04-007-0402). Furthermore, it was
found that the Al-fraction has decreased in comparison to
that at 20 C. Considering that at 210 C pure Al was not present anymore, we conclude that at this temperature, the
whole volume fraction of the Al-phase has fully reacted into
the intermetallics Al9Ni2 and Al3Ni, which agrees with the
results of Blobaum et al.19 who reported a complete transformation of Al to intermetallic phases at temperatures lower
than 242 C in Ni/Al reactive multilayers. From 210 to
240 C, corresponding to the end of the first exothermic peak,
the Al3Ni phase fraction increases. Moreover, no trace of the
Al9Ni2 phase was found at this isotherm, and only peak sets
belonging to Al3Ni and Ni were present at 275 C. The Al3Ni
phase has entirely transformed to Al3Ni2 at 350 C, and
finally, in the temperature range of 350–500 C, only Al3Ni2
and Ni were detected. The phase reaction sequence from
Al9Ni2 to Al3Ni and finally to Al3Ni2 observed here is in
agreement with literature data19 for Ni/Al multilayers with
the bilayer thickness ranging from 25 to 200 nm.
In Fig. 2(b), we present the phase analysis of (Nb-Ni)/Al
reactive multilayers as a function of temperature. At 20 C, a
broad peak between 2 and 4 Å1, corresponding to the nanocrystalline (nc-) Nb-Ni layers, and a set of peaks, attributed to
Al can be observed. In the temperature range of 20–200 C,
no structural changes are detected. The 210 C isotherm
shows in addition to the Al-peaks and the broad feature of the
nc-Nb-Ni layers, yet another set of peaks revealing the presence of the Al3Ni phase. The onset of Al3Ni phase formation
Schnabel et al.
Appl. Phys. Lett. 111, 173902 (2017)
FIG. 2. (a) and (b) The in situ electron
diffraction structural analysis for the
Ni/Al and (Nb-Ni)/Al reactive multilayers, respectively. Within the (NbNi)/Al reactive multilayers, we
observe an initial formation of Al3Ni,
whereas in the Ni/Al system, the metastable Al9Ni2 phase forms initially.
within the temperature interval of 200–210 C corresponds to
the heat flow onset observed by DSC in Fig. 1. From 210 to
320 C, the peak intensities corresponding to the Al phase
decrease. The decrease in the Al phase volume fraction coincides with the heat flow, explored by DSC. At 320 C, the
temperature corresponding to the maximum of the continuous
heat flow DSC peak, no trace of the Al face centered cubic
phase is present. For, the 390 C isotherm corresponding to
the end temperature of the continuous heat flow segment of
the DSC curve [Fig. 1(f) top], electron diffraction analyses
revealed not only the presence of the Al3Ni phase but also the
formation of Al3Ni2 (PDF 01-083-3987). The 425 and 500 C
isotherms show Al3Ni2 and Al3Nb (PDF 13-0146) phases.
However, the Al3Ni phase, which was observed at 320 C, is
not present anymore. Hence, from the comparison of the
intensity profiles of the in situ TEM electron diffraction patterns, we conclude the presence of a difference in the phase
formation sequence between Ni/Al and (Nb-Ni)/Al reactive
multilayers. For the (Nb-Ni)/Al reactive multilayers, we
observe the Al phase up to 270 C, whereas in the case of the
Ni/Al reactive multilayer, the Al phase has been consumed
already at 210 C. Therefore, the complete consumption of Al
by the intermetallic phases in the (Nb-Ni)/Al reactive multilayers is slower compared to that in the Ni/Al reference system. This observation is in agreement with the lower heat
flow rate measured by DSC (Fig. 1). Furthermore, in the comparative in situ electron diffraction analysis, we observe a difference in initial phase formation. Although in the Ni/Al
system, Al3Ni forms subsequently to the formation of the
metastable Al9Ni2 phase,19 we observe a direct formation of
the stable Al3Ni phase in the (Nb-Ni)/Al system. We propose
that the difference in the phase formation sequence is due to a
difference in the initial compositional gradient at the reactive
multilayer interface. With the smaller grain size in the Nb-Ni
rich layer of the (Nb-Ni)/Al reactive multilayer compared to
the grain size of Ni within the Ni/Al system, we expect a fast
diffusion of Ni towards the multilayer interface. With a faster
diffusion towards the interface, we expect a faster Ni saturation within Al at the interface. We propose that the faster
saturation favors the formation of the stable Al3Ni in the (NbNi)/Al reactive multilayers, which is in agreement with the
evaluation of d’Heurle et al.21
In Ni/Al based reactive multilayers, a large part of the
energy release is due to chemical mixing.22 To understand
the difference in heat flow rates observed, it is essential to
complement the structural analyses presented in Fig. 2 by the
evaluation of chemical intermixing as a function of temperature. We evaluated the extent of thermally induced interdiffusion of annealed and subsequently quenched free-standing
Ni/Al and (Nb-Ni)/Al reactive multilayers at temperatures of
350 C and 390 C, respectively. For the elemental distribution analyses of the multilayer cross-sections, TEM lamellas
from the annealed and as-deposited reactive multilayers
were prepared using a focused ion beam (FIB-SEM) dualbeam station (Zeiss NVision40) operating with a Gaþ source
at 30 and 5 kV. Figures 3(a) and 3(b) present the Ni- and Alelemental distribution maps of the Ni/Al lamella in the asdeposited state. The maps clearly identify the Ni-rich and
Schnabel et al.
Appl. Phys. Lett. 111, 173902 (2017)
FIG. 3. (a) and (b) The elemental distribution maps for Ni and Al on the
Ni/Al reactive multilayer in the asdeposited state, respectively. (c) and
(d) The elemental distribution maps for
Ni and Al of the Ni/Al reactive multilayers, which were annealed at 350 C.
(e) The chemical composition for the
as-deposited (dotted line) and annealed
(solid line) Ni/Al reactive multilayers
measured across to the layered structures. (f) The results of line scan measurements for the (Nb-Ni)/Al reactive
multilayers, whereas the elemental distribution maps in the as-deposited and
annealed state are presented in [(g)–(i)]
and [(j)–(l)], respectively.
Al-rich layers. Figures 3(c) and 3(d) present the Ni- and Alelemental distribution maps of the Ni/Al reactive multilayers
annealed at 350 C. At variance to the as-deposited state, one
observes a less pronounced separation between the Ni- and
Al-rich layers, indicating interdiffusion of the constituents.
Additional to elemental distribution maps, the elemental
intensity profiles across the layers [indicated by the white
arrows in Figs. 3(a)–3(d)] were carried out. The length of the
profiles is 120 nm, which corresponds to the thickness of two
bilayers. Figure 3(e) presents the elemental content profiles
of the Ni/Al reactive multilayer measured across the layers.
Here, dotted and solid lines represent the elemental content
profiles measured in the as-deposited and annealed states,
respectively. In the as-deposited state, a higher Al content
within the Ni-rich layer than the Ni content within the Al
layer is measured. This could be due to interdiffusion of the
constituents during the multilayer deposition. The evaluation
routine yields the maximum Ni content of 82 at. % in the
Ni-rich layer, whereas the maximum Al content within the
Al-rich layer is almost 100%. The same analyses were performed also for the (Nb-Ni)/Al reactive multilayers [Figs.
3(f)–3(j)] in the as-deposited and annealed states. In this
case, 390 C as the annealing temperature is chosen, which
corresponds to the end of the continuous heat flow segment
[Fig. 1(f), top]. Our results show that in the cross-sections of
the (Nb-Ni)/Al reactive multilayer, one can clearly identify
the Nb-Ni and Al rich layers in the as-deposited state. The
maximum of the Ni content is 61 at. %. Hence, the maximum
Ni content within the Ni-rich layer is 21 at. % lower compared to the Ni/Al reactive multilayer, which is compensated
by a maximum Nb content of 23 at. %. The comparison of
the elemental content profiles in the as-deposited and
annealed states reveals an increase in the Ni content within
the Al layer similar to the Ni/Al reactive multilayers.
Compared to the amount of Ni, the increase in the Nb content within the Al rich layer upon annealing is insignificant.
However, annealing increases the Al content within the Ni
rich layer by about 20 to 25 at. %. Furthermore, in contrast to
the Ni/Al system, a non-monotonous transition between the
Nb-Ni and Al-rich layer is observed at 390 C. Instead, we
observe a transition layer of about 3 nm thickness at the
native (Nb-Ni)/Al interface. In this transition layer, the Ni
content exhibits a minimum. This minimum in the Ni content
infers the formation of an Al-Nb layer at the native (Nb-Ni)/
Al interface, which is in agreement with the Al3Nb phase
formation observed by in situ electron diffraction at temperatures above 390 C (Fig. 2). The DSC analysis in Fig. 1(f)
(top) shows two separate exothermic peaks at 416 and
447 C, respectively. In Nb/Al reactive multilayer systems,
such peaks have been connected to a two stage process,23,24
namely, the propagation of the initial reaction phase laterally along the reactive interface and a subsequent growth
normal to the interface. The formation of a 3 nm thin Al-Nb
layer, as observed in Fig. 3, along with the occurrence of
two separate heat-flow peaks at 416 and 447 C is in agreement with the two stage process reported in the literature.23,24 In case of the ternary (Nb-Ni)/Al multilayer
system, this two-step process is likely to be superimposed
by the formation of Al3Ni2.
For diffusion-controlled growth in layered systems, the
activation energy can be determined using the Kissinger formalism.25–28 Both the Ni/Al and (Nb-Ni)/Al reactive multilayers are evaluated by DSC at heating rates of 10, 20, 40,
and 80 K/min to obtain the activation energies for the initial
growth process. A dynamic shift in peak temperature as a
function of the heating rate allowed calculation of the activation energy of the corresponding initial reaction. The following values were obtained for the reactions just above 200 C:
1.0 eV for (Nb-Ni)/Al and 1.1 eV for Ni/Al. These values are
at the lower end of the reported activation energies of 1.0 to
1.9 eV for the formation of Al3Ni and Al9Ni2.19,29 Battezzati
et al.29 proposed grain-boundary interdiffusion as a possible
rate-controlling mechanism in Ni-Al reactive multilayers.
The value of 1.1 eV for the Ni/Al reactive multilayer is
within the range of 1.1 to 1.35 eV reported for grain boundary diffusion in Ni/Al.30 Therefore, we propose grainboundary interdiffusion as an initial rate-controlling process
for the reactive multilayers studied here.
Schnabel et al.
Appl. Phys. Lett. 111, 173902 (2017)
Interdiffusion can be quantified by the diffusion flux j as19
j ¼ D~ ;
where D~ is the chemical diffusion coefficient, and dc is the
compositional change along an infinitesimal distance dx.31,32
From the elemental content analyses, we know that the initial
Ni content within the (Nb-Ni)/Al layer is 21 at. % lower
compared to the maximum Ni content within the Ni/Al
system. This, according to Eq. (1), causes a compositional
induced decrease in diffusion flux, effectively reducing
chemical intermixing and heat flow rate.
Taking into account the influence of the chemical potential gradient dl on the chemical diffusion coefficient
D~ ¼ B
d ln c
where B is a mobility parameter,32 we conclude that Nballoying stabilizes the Ni-rich layer,33 reducing the chemical
potential gradient, and further decreases the diffusion flux.
Both the compositional and chemical induced decreases in
diffusion flux may explain the lower heat flow rate observed.
Moreover, the formation of an Al-Nb layer at the native (NbNi)/Al interface may further reduce the effective interdiffusion.
Hence, from the comparative analyses of (Nb-Ni)/Al
and the reference Ni/Al reactive multilayers, we learn that
Nb-alloying reduces the maximum heat flow rate by more
than 50% while only reducing the total energy-storage
capacity by about 21%. Through a combined in situ electron
diffraction and EDS-STEM analysis, we identify a reduction
in chemical potential driving force and the formation of an
Al-Nb layer at the native (Nb-Ni)/Al interface as a reason for
the reduced heat flow rate of (Nb-Ni)/Al reactive multilayers.
As a result of the study, we conclude that Nb-alloying stabilizes the as-deposited state and facilitates the formation of a
diffusion barrier. We propose this approach as a microstructural design concept enabling the control of diffusion in reactive multilayer systems.
A further reduction in the heat flow rate, while maintaining a large energy-storage capacity, could be achieved by
nanoscale microstructural design in the as-deposited state.
Furthermore, we propose Mo as a promising alloying element.
Similar to Nb-Ni, Mo-Ni forms intermetallics on the Ni rich
side of the phase diagram.34 In addition, Mo-Ni in contrast to
Nb-Ni exhibits a positive enthalpy of mixing as a regular solution.34 This should lead to a lower reduction in energy-storage
capacity. Moreover, Al-Mo exhibits large enthalpy of mixing
similar to Al-Nb,35 which is also beneficial for energy-storage
capacity and the formation of a diffusion barrier.
Film deposition and characterization were carried out at
the FIRST Center for Micro- and Nanoscience and the
Scientific Center for Optical and Electron Microscopy
(ScopeM) at ETH Zurich, respectively.
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