# Plasma-enhanced atomic layer deposition of nickel nanotubes with low
resistivity and coherent magnetization dynamics for 3D spintronics

M. C. Giordano  École Polytechnique Fédérale de Lausanne, School of
Engineering, Institute of Materials, Laboratory of Nanoscale Magnetic
Materials and Magnonics, 1015 Lausanne, Switzerland  K. Baumgaertl  École
Polytechnique Fédérale de Lausanne, School of Engineering, Institute of
Materials, Laboratory of Nanoscale Magnetic Materials and Magnonics, 1015
Lausanne, Switzerland  S. R. Escobar Steinvall  École Polytechnique Fédérale
de Lausanne, School of Engineering, Institute of Materials, Laboratory of
Semiconductor Materials, 1015 Lausanne, Switzerland  J. Gay  École
Polytechnique Fédérale de Lausanne, School of Engineering, Institute of
Materials, Laboratory of Nanoscale Magnetic Materials and Magnonics, 1015
Lausanne, Switzerland  M. Vuichard  École Polytechnique Fédérale de Lausanne,
School of Engineering, Institute of Materials, Laboratory of Nanoscale
Magnetic Materials and Magnonics, 1015 Lausanne, Switzerland  A. Fontcuberta i
Morral  École Polytechnique Fédérale de Lausanne, School of Engineering,
Institute of Materials, Laboratory of Semiconductor Materials, 1015 Lausanne,
Switzerland  D. Grundler  École Polytechnique Fédérale de Lausanne, School of
Engineering, Institute of Materials, Laboratory of Nanoscale Magnetic
Materials and Magnonics, 1015 Lausanne, Switzerland<EMAIL_ADDRESS>
###### Abstract

We report plasma-enhanced atomic layer deposition (ALD) to prepare conformal
nickel thin films and nanotubes by using nickelocene as a precursor, water as
the oxidant agent and an in-cycle plasma enhanced reduction step with
hydrogen. The optimized ALD pulse sequence, combined with a post-processing
annealing treatment, allowed us to prepare 30 nm thick metallic Ni layers with
a resistivity of 8 $\mu\Omega$cm at room temperature and good conformality
both on the planar substrates and nanotemplates. Thereby we fabricated several
micrometer-long nickel nanotubes with diameters ranging from 120 to 330 nm. We
report on the correlation between ALD growth and functional properties of
individual Ni nanotubes characterized in terms of magneto-transport and the
confinement of spin wave modes. The findings offer novel perspectives for Ni-
based spintronics and magnonic devices operated in the GHz frequency regime
with a 3D device architecture.

###### keywords:

American Chemical Society, LaTeX

École Polytechnique Fédérale de Lausanne, School of Basic Sciences, Institute
of Physics, 1015 Lausanne, Switzerland École Polytechnique Fédérale de
Lausanne, School of Engineering, Institute of Microengineering, 1015 Lausanne,
Switzerland +41 21 693 38 52

##  1 Introduction

Magnetic nanostructures find applications in data storage devices 1, magnetic
sensors 2, 3 and biomedical applications 4. The demand of high-density
technologies 5 and the need of new applications are driving the expansion of
nanomagnetism towards three-dimensional (3D) device architectures6, 7. Here,
novel physical phenomena arising from complex spin textures become possible.
However, the implementation of high-quality 3D nanomagnetic structures on
large scales is lacking. A promising path is to magnetically coat 3D
nanotemplates, e.g. polymeric scaffolds8, self-assembled9 or bottom-up
grown10, 11 nanostructures. Therefore, a technique for conformal coatings of
high-quality ferromagnetic metals is of uttermost importance to advance 3D
nanomagnetism. Atomic layer deposition (ALD) 12, 13 indeed offers a great
potential. Here the material is deposited through self-limiting reactions
between vapor phase metal-organic precursors (and co-reactants) and the
exposed substrate surface, enabling conformal coatings and thickness control
on the atomic scale. For the deposition of numerous oxides, ALD is a well-
established technique. For metallic thin films, including ferromagnetic
metals, ALD still faces challenges due to the limited number of suitable
precursors, the difficulty in reducing the metal cations, and the tendency of
metals to agglomerate to islands 14, 15. Previous attempts to deposit the
ferromagnetic metal nickel with ALD included one-step and two-step processes.
In the first scenario, the nickel films were deposited directly through the
reaction of the nickel precursor and a co-reactant. Amongst them there were
molecular hydrogen 15, molecular ammonia16, plasma-activated hydrogen17 and
ammonia 17, 18, as well as tert-butilamine19. The use of plasma offered the
possibility to lower the deposition temperature and to increase the purity of
the deposited material 20. However, the recombination of plasma species
limited the conformality of plasma-enhanced ALD (PE-ALD) on high aspect ratio
(AR) nanotemplates 21, 20. For the two-step process, the first step was based
on nickel oxide deposition, where the oxidant co-reactant was ozone10, 22, 23
or water24, and the second step relied on a reduction process, which
transformed the oxide into pure nickel. The reduction was performed either
within the ALD cycle with hydrogen plasma 24, or as post-deposition annealing
in molecular hydrogen 10, 11, or by exploiting both processes. The two-step
process ensured, in principle, good conformality on high aspect ratio
nanotemplates due to the greater homogeneity of nickel oxide deposition via
the thermal ALD growth. The usage of water as co-reactant resulted in Ni
coatings with good conformality on nanotemplate substrates with an AR of about
3:1 24. Metallic Ni coatings achieved with ozone as the co-reactant exhibited
appreciable surface roughness and inhomogeneities as Ostwald ripening was
potentially induced by the annealing temperatures above 400$~{}^{\circ}$C 10,
25.
Nanotubes (NTs) prepared from ferromagnets represent prototypical 3D
nanomagnetic structures 7. They raised significant attention as they support
stable flux-closure magnetic states 26 and avoid the Bloch point structure
along the central axis. This results in a fast and controllable reversal
process when compared to magnetic stripes 27. In addition, previously
unforeseen dynamic effects are possible. Domain walls moving in nanotubes are
predicted to avoid Walker breakdown and give rise to Cherenkov-like spin wave
emission 28. Their curvature generates non-reciprocal spin-wave dispersion 29.
Furthermore, the three geometric parameters, namely the length, the inner
radius and external radius, offer the possibility to tailor spin-wave
confinement. These characteristics make ferromagnetic NTs promising functional
objects for 3D integrated spintronics and magnonics 30. Experimental studies
on individual NTs focused so far on the determination of their static magnetic
properties, magnetization states and reversal magnetization mechanisms, by
techniques such as SQUID magnetometry 3132, cantilever magnetometry 33,
anisotropic magnetoresistance measurements 10, 11, 34 and x-ray based magnetic
imaging 35, 32. Still, the dynamic magnetization of individual ferromagnetic
NTs with the radius on the order of 100 nm has not yet been the subject of
comprehensive experimental investigations. So far studies were limited to
spin-wave resonances of large ensembles of ferromagnetic tubular structures 36
and of rolled up ferromagnetic layers on semiconductor membranes with
micrometric radii37, 38, where azimuthal interference of long-wavelength
magnetostatic spin waves was found. Lenz et al.39 recently reported the
magnetization dynamics of an individual Fe nanorod embedded in a carbon
nanotube, studied by ferromagnetic resonance (FMR) and Brillouin light
scattering (BLS). Dynamics of individual nanotubes 28, 29 remain to be
explored experimentally.
Atomic layer deposition (ALD) has selectively been used for the fabrication of
ferromagnetic Ni NTs by coating GaAs nanowires 10, 11. Weber et al. 33
demonstrated three stable magnetization states in individual ALD-grown Ni
nanotubes measured by cantilever magnetometry. Rueffer et al. 11 measured an
anisotropic magnetoresistance (AMR) effect of up to 1.4 $\%$ on such Ni
nanotubes when applying fields of 1 T at room temperature. The presented NTs
exhibited a rough surface after annealing and their magnetization dynamics
were not reported. Wang et al. 18 reported PE-ALD to coat high aspect ratio
(AR$\sim 13:1$) Si nanopillars with Ni. They achieved a low resistivity of
11.8 $\mu\Omega$cm in planar Ni layers after post-annealing enhancement, but
did not report on functional properties of the high aspect ratio coatings.
In this work we present the fabrication of Ni NTs by PE-ALD on single-
crystalline GaAs nanowires as 3D nanomagnetic structures. Adapting the two-
step ALD process of Ref. 24 we employed nickelocene as the metallorganic
precursor and water as the oxidant agent to first obtain nickel oxide.
Subsequently, an in-cycle plasma enhanced reduction process with hydrogen was
performed to get metallic nickel. We optimized the ALD sequence to enhance the
functional properties and conformality of Ni in the nanotubular shape. We
identified the deposition conditions that led to identical coverage on planar
substrates and on the nanotemplates with an AR above 15:1. The optimized Ni
NTs exhibited, after post-annealing enhancement, smooth surfaces and showed
both a low resistivity of roughly 8 $\mu\Omega$cm and a prominent AMR effect.
We characterized the magnetization dynamics of individual NTs under microwave
irradiation in the GHz frequency regime. We detected a series of resonances by
micro-focused BLS which we attributed to spin-wave modes with discrete wave
vectors $k$ due to azimuthal confinement. To the best of our knowledge, this
is the first experimental investigation of resonant modes in nanoscale
ferromagnetic NTs. Our results prove the high quality of the achieved 3D
ferromagnetic coatings, promising to advance the realization and exploration
of 3D spintronic device architectures such as the proposed racetrack memory.5

##  2 Results and discussion

###  2.1 Growth of nickel thin films and nanotubes

We start by reporting on the PE-ALD growth of Ni thin films on planar Si (100)
substrates and high aspect ratio nanotemplates. The planar wafers and GaAs
nanowires with hexagonal cross-sections (Figure S1), respectively, were coated
with a few nanometers of alumina with the aim of growing Ni on the same
surface, while focusing on the effect of different substrate geometries and
aspect ratios. In order to optimize the composition and morphology we varied
the chamber temperature $T$ and the number of metal oxide deposition steps c
prior to hydrogen-plasma reduction. Figure 1a shows the dependence of the
growth rate of nickel on planar substrates on $T$. Here the number of nickel
oxide was set to $c=1$ and the ALD cycle was repeated $n=1000$ times (inset of
Figure 1b). The chamber temperature was varied from 155 ∘C to 175 ∘C. From
$T=155~{}^{\circ}$C to $T=165~{}^{\circ}$C we observe that the growth rate
increases from a value of 0.12 Å/(NiCp2 pulse) to a value of 0.33 Å/(NiCp2
pulse). It stays constant at 0.33 $\pm$ 0.01 Å/(NiCp2 pulse) for temperatures
ranging from 165 to 172 ∘C. We attribute this temperature regime to the ALD
window. The growth rate is at 0.47 Å/(NiCp2 pulse) for $T=175~{}^{\circ}$C.

Figure 1: (a) Dependence of the growth rate of Ni on the chamber temperature.
Each sample was prepared with $n$ = 1000 cycles. (b) Dependence of growth rate
on the number $c$ of NiCp2/H2O steps in the ALD cycle (inset) for samples
deposited at $T=170~{}^{\circ}$C.

To explore how the growth rate varies at $T=170~{}^{\circ}$C with changing the
number of nickel oxide steps c in the ALD cycle, we varied the parameter $c$
from 1 to 10, while the ALD cycle $n$ was changed consistently in order to
keep the product $c\times n=1000$. In Figure 1b we report the growth rate
estimated on the the number of total $\ce{NiCp2}$ pulses ($c\times n$) as
function of the parameter c, number of steps leading to nickel oxide
formation. We observe that the growth rate drops from 0.33 Å/(NiCp2 pulse) at
$c=1$ to $0.13$ Å/(NiCp2 pulse) at $c=5$ and $7$. It increases again to 0.17
Å/(NiCp2 pulse) for $c=10$. Thin films with a thickness of about 30 nm were
prepared with $c$ set to $1,5,7$ and 10, taking into account the values of
growth rate reported in Figure 1b. They were then annealed under hydrogen flow
(details in Methods). In the following we discuss the properties of annealed
Ni films. Atomic force microscopy (Figure 2a-2d) and X-ray diffraction (Figure
2e) were used to assess the roughness and the crystallographic structure,
respectively, of the four planar Ni films. We extracted rms roughness values
of 2.1 nm for $c=1$ (Figure 2a), $1.3$ nm for $c$ = 5 (Figure 2b), 0.9 nm for
$c=7$ (Figure 2c) and 1.2 for $c=10$ (Figure 2d). Accordingly, the sample
prepared with $c=7$ showed the smoothest surface (compare Figure 5a).

Figure 2: Atomic force microscopy performed on a $1~{}\mu$m$\times 1~{}\mu$m
area of samples with (a) $c=1$, (b) $c=5$, (c) $c=7$, and (d) $c=10$ (scale
bars: 200 nm). (e) X-ray diffractograms of $30~{}\pm 2$ nm thick nickel films
prepared with $c$ ranging from 1 to 10 (see labels). (f) The bars give the
positions of the (111) and (200) reflections of cubic nickel (PDF 00-004-0850)
. The thin films were deposited at 170 ∘C on a silicon substrate coated with 5
nm of Al2O3 and annealed at 350$~{}^{\circ}$C.

In Figure 2e we report the X-ray diffractograms acquired in a glancing
incident angle configuration for the four annealed thin films. The measured
peaks are consistent with the (100) and (200) reflections of cubic nickel (PDF
00-004-0850) shown in Figure 2f. The broadening and positions of peaks of the
samples prepared with $c=1,5$ and $7$ suggest polycrystalline Ni films, with
no preferential grain orientation.
The four depositions were performed simultaneously on Si wafers and GaAs
nanowires (both coated with alumina). We employed different set of nanowires,
with lengths between 7 and 15 $\mu$m and diameters ranging from 100 to 270 nm.
The aspect ratios of different nanotemplates were calculated following the
convention for squared pillars 21 and range from 15:1 to 31:1 (Table S1). The
nanotemplates allowed us to obtain arrays of millions of vertically oriented
nickel nanotubes (Figure 3).

Figure 3: Scanning electron microscopy (SEM) micrographs of (a) a large
ensemble of vertical nanowires with ALD-grown Ni shells (scale bar: 1 $\mu$m)
and (b) selected central regions of Ni NTs (scale bar: 200 nm). Figure 4:
(a-d) HRTEM images (scale bar: 10 nm) of interfaces between GaAs nanowires and
the $\ce{Al2O3}$/Ni shells. The insets show SAED patterns (scale bars: 5
nm-1). (e-h) HADDAF images (scale bars: 100 nm). (i-l) Distribution of Ga and
Ni. (m-p) Elemental analysis of Ni nanotubes on GaAs nanowires. Each column of
panels represents one sample. The four samples were prepared with $c$ set to
$1,5,7$ and $10$ as indicated on top of each column. The chamber (annealing)
temperatures were 170 ∘C (350∘C).

To investigate the chemical composition and the structure of the Ni NTs we
performed transmission electron microscopy (TEM). Figure 4a to 4d shows high-
resolution (HR) TEM micrographs with insets of the corresponding selective
area electron diffraction (SAED) patterns of the Ni nanotubes prepared with
$c=1,5,7$ and $10$, respectively. The images depict the interface between GaAs
nanowires and the ALD-grown Al2O3/Ni shells. The Ni coating with $c=1$ (Figure
4a) exhibits a relatively large peak-to-peak roughness value of about 7 nm.
The Ni shells deposited with $c=5,7,10$ (Figure 4b to d) show a continuous
coverage and form ferromagnetic nanotubes with small roughness. The roughness
variation of the NTs is qualitatively consistent with the characteristics
found for the planar Ni films (Figure 5a). No preferential growth direction is
observed in any of the samples, all made up of grains of different sizes and
orientations. An epitaxial relationship between core and shells was not found.
All shells exhibited a surface layer of nickel oxide of a few nanometers,
attributed to oxidization in air after the annealing treatment. SAED images
contain the typical ring patterns of a polycrystalline texture, which we
attribute to the nickel shell, and a single crystal dot pattern of a zinc-
blende structure, that we assign to the GaAs core, which was grown along the
[111]B directions as previously reported 40. Starting from the center, the
diffraction rings were assigned to the (111), (200), (220) and (311) planes
respectively, belonging to the face-centered cubic lattice of Ni crystals.
Using scanning TEM (STEM) we performed high-angle annular dark-field (HAADF)
and energy dispersive X-ray spectroscopy (EDS) analysis to study the
composition and elemental distribution of the nanostructures. HADDAF images in
Figure 4e to 4h display the edges of NTs which were cleaved and close to the
Si(111) substrate. Corresponding element maps of Ga and Ni, as representative
elements of inner core and outer shell, are shown in Figure 4i to 4l (and
Figure S1). Figure 4m to 4p show line scans of the Ni (red curve) and Ga
(green curve) elemental distributions as a function of position along the NT
diameter. The signals were extracted in the area enclosed by the dotted line
and averaged along the NT long axis. These curves allowed us to quantify the
thickness of the $\ce{Al2O3}$/Ni shells on the nanotemplates (Methods).
Thicknesses of the top and bottom edge of the cross section are highlighted by
yellow and gray shaded regions in Figure 4m to 4p. Similar analysis where
acquired at random spots of five NTs of each deposition, giving averaged
thicknesses of 11.5 $\pm$ 2.4 nm, 14.0 $\pm$ 2.1 nm , 29.4 $\pm$ 4.1 nm and
12.2 $\pm$ 1.5 nm for depositions prepared with $c=1,5,7$ and $10$,
respectively. Only for $c=7$ did we observe that the Ni shell had a thickness
that was comparable with the one deposited on the planar Si (100) substrate.
In Figure S2 we analyzed the step coverage as the ratio between sidewall and
top-surface film thickness for two samples. The values were $\geq 88~{}\%$.

The growth rates of shells were found to vary as a function of $c$. We explain
the different growth rates and surface roughnesses by different growth
mechanisms. The higher the number of nickel oxide steps $c$ in the ALD
sequence, the more the surface is oxidized, to which the hydrogen plasma pulse
is applied. The drop in growth rate for $c>1$ in Figure 5a might indicate that
we moved away from the PE-ALD regime, for which the growth rate is enhanced by
reaction mechanisms involving plasma species, and entered a thermal ALD
regime, which often exhibits a smaller growth rate 41, 42. Furthermore, the
decomposition of the metalorganic precursor is expected to be more pronounced
on nickel than on nickel oxide 43. Consequently the Ni-rich surface realized
by $c=1$ promotes a larger growth rate. The discrepancies in growth rates for
planar films and NT shells are attributed to the recombination of plasma
species by collisions on the enlarged surface area provided by the
nanotemplates. The recombination probability of hydrogen plasma species is
higher on metallic nickel 44, 45 than on an oxidized surface46, 47. For $c=7$
we observed the lowest roughness in planar thin films and the highest growth
rate on the high aspect ratio nanotemplates. We assume that in this deposition
process a scenario was reached where, at each ALD cycle, a monolayer of nickel
oxide was formed before the reducing hydrogen plasma step was inserted. For
$c=10$, the increase in growth rate for the planar film (Figure 1b) suggested
that a larger nickel oxide thickness was realized compared to samples with
$c=5$ and $c=7$, meaning that the duration of the hydrogen plasma pulse was no
longer optimized to reduce the deposited oxide layer. For the same reason, we
speculate that the subsequent annealing treatment in hydrogen may have been
less effective on the reduction-induced crystallization process of sample
$c=10$ and thereby explain the less intense peaks observed in the X-ray
diffractogram of Figure 2e.

###  2.2 Magnetic properties of planar Ni thin films

In Figure 5a to 5e (and Figure S3a and S3b) we summarize different physical
properties of Ni thin films measured as a function of the growth parameter c:
the rms roughness measured by atomic force microscopy (AFM) (Figure 5a), the
magnetization measured at room temperature by SQUID (Figure 5b), the room
temperature resistivity measured by either a four-probe configuration or a van
der Pauw configuration (Figure 5c), the AMR effect (Figure 5d) and the
linewidth of the ferromagnetic resonance (FMR) (Figure 5e). In Figure 5f we
depict the magnetoresistance measured on planar Ni thin films when an applied
in-plane magnetic field $\mu_{0}H=80$ mT was rotated (the angle $\theta$ is
defined between the current direction and $\mathbf{H}$). We display the
specific resistivity $\rho(\theta)$ in terms of
$\dfrac{\Delta\rho(\theta)}{\rho}$ =
$\dfrac{\rho(\theta)-\rho(90^{\circ})}{\rho(90^{\circ})}$. We observe a
$cos^{2}(\theta)$ dependency as expected for the AMR effect. Depending on $c$
we find maximum values of the relative AMR effect
$\dfrac{\Delta\rho(\theta)}{\rho}$ ranging from 4.7 down to 2.1 (Figure 5d).
From field-dependent FMR measurement (Figure 5e) we extracted the effective
magnetization $M_{\rm eff}$ by fitting the Kittel formula 48 to the observed
resonance frequencies (white branch in the color-coded spectra shown in Figure
5g). The spectra were taken by inductive measurements using a vector network
analyzer (VNA) and reflected the scattering parameter $S_{21}$ (Methods).
Values of $M_{\rm eff}$ are summarized in Figure 5b. In Figure 5h we compare
individual spectra taken at the same field of -80 mT for thin films deposited
with different $c$. We observe a resonance frequency of 5.5 GHz and a narrow
linewidth for $c=7$. Extracted linewidths are summarized in Figure 5e.

Figure 5: Physical properties of the Ni thin films as a function of $c$: (a)
roughness, (b) effective and saturation magnetization $M_{\rm s}$ extracted
from $M(H)$ curves (Figure S3b), (c) specific resistivity $\rho$ extracted by
two independent experiments, (d) relative AMR effect, and (e) linewidth
$\Delta f$. (f) Magnetoresistance of four thin-film samples (see labels)
measured at room temperature for a rotating in-plane field of 80 mT. (g)
Magnitude of the scattering parameter $S_{21}$ measured in an in-plane
magnetic field by VNA-FMR on a thin film deposited with $c=7$. White indicates
large absorption. (h) Spectra taken at -80 mT on Ni thin films prepared with
different $c$ (see labels). For $c\geq 5$ the resonance frequencies are found
to be the same for different $c$ but still the linewidth $\Delta f_{\rm mag}$
(indicated by horizontal arrows) varies as summarized in (e).

The narrowest FMR linewidth of $\Delta f$ = 1.23 GHz is registered for a thin
film prepared with $c=7$. The same film exhibits the smallest rms roughness of
0.9 nm. We assume that the small roughness reduces the inhomogeneous linewidth
broadening expected from two-magnon scattering induced by inhomogeneities and
surface irregularities49, 50. Values of $M_{\rm eff}$ are found to be in good
agreement with $M_{\rm s}$ measured by SQUID (Figure S3b) in an in-plane
magnetic field for $c\geq 5$. We measured a maximum effective magnetization
$M_{\rm eff}$ of 380 kA/m for the sample deposited with $c=7$. Within the
error bar it agrees with an $M_{\rm s}$ of 400 kA/m measured by SQUID on the
same sample. Consistent values indicate that additional magnetic anisotropies
(magnetocrystalline anisotropy, surface anisotropy) do not play a significant
role. The experimentally observed values are about 20$~{}\%$ smaller than the
reported saturation magnetization of 490 kA/m of single-crystalline nickel 51.
For $c=1$, $M_{\rm eff}$ is found to be smaller than Ms by about 130 kA/m. The
discrepancy indicates an out-of-plane anisotropy present in this sample. The
lowest values of resistivity $\rho$ that we measured on thin films with $c=1$
and $c=7$ amounted to about 7 to 8.4 $\mu\Omega$cm. They were only slightly
larger than the resistivity of 6.9 $\mu\Omega$cm reported for bulk Ni at room
temperature 52. For $c=10$ we found $\rho$ $\sim$ 20 $\mu\Omega$cm. In the
same sample we find the smallest AMR value of 2.1 $\%$. We attribute the
increased resistivity and the small AMR value for $c=10$ to inhomogeneities
due to non-optimized ALD deposition.

###  2.3 Transport properties of Ni nanotubes

Nickel NTs of the different sets of depositions were extracted from the
nanotemplate substrates and transferred to either Si or glass substrates for
further nanolithography and the integration of metallic leads. Then we
performed magneto-transport and spin-dynamic experiments on individual NTs. In
the following we report on data obtained on samples S1 and S2 (Tab. 1).

Table 1: Geometrical parameters of NTs prepared with $c=7$. S1 (S2) was used of magneto-transport (spin-dynamics) measurements. Sample |  $L$ ($\mu$m) |  $D_{out}$ (nm) |  $t$ (nm)
---|---|---|---
S1 | 14.5 $\pm$ 0.1 | 285 $\pm$ 10 | 29.4 $\pm$ 4.1
S2 | 14.5 $\pm$ 0.1 | 280 $\pm$ 20 | 29.4 $\pm$ 4.1
Figure 6: Magnetoresistance of nanotube S1 measured as a function of a static
magnetic field applied along the long NT axis. The inset shows an SEM
micrograph of S1 contacted by 5 metallic leads made from a Au thin film via
lift-off processing. Voltage ($V$) and current ($I$) probes are indicated.

In Figure 6 and Figure S4, we report the field-dependent resistance $R$
measured on different nanotubes. The data of nanotube S1 deposited with
$\textit{c}=7$ is displayed in Figure 6. The two voltage probes were separated
by 7 $\mu$m. The field was applied long the long axis. While decreasing the
field from 80 mT (red curve), the magnetoresistance remains constant at about
27.4 $\Omega$ until $\approx 30$ mT. From $R$ we calculate a specific
resistivity of 7.8 $\pm$ 2.8 $\mu\Omega$cm. This value is only about 13 $\%$
larger compared to the resistivity of bulk nickel. For further decreased $H$,
the resistance decreases to a minimum value of 27.05 $\Omega$ at -2 mT. The
maximum resistance is restored again for $\mu_{0}H\leq-30$ mT and remains
constant until $-80$ mT. The blue curve is acquired with increasing field and
shows hysteretic behavior near zero field. Considering the long axis of the NT
to be the easy axis, a field of 80 mT is sufficient to saturate the spins
along the current direction. This configuration corresponds to the scenario of
maximum electron scattering for the AMR effect and the resistance is at its
maximum value. For fields below 30 mT, the resistance decreases, which we
attribute to the onset of reversal of the NT’s magnetization. The reduced
resistance suggests that spins tilt away from the current direction. For
nanotubes of similar geometrical parameters the reversal of magnetization is
predicted to occur via vortex domain wall formation 53. Correspondingly spins
would be aligned in an azimuthal direction around the NT, thereby taking an
angle of 90 deg with respect to the current flow direction. This configuration
corresponds to the scenario of minimum electron scattering for the AMR effect
and is consistent with the observed reduction of $R$ in the reversal regime.
The relative variation of $R$ corresponds to an AMR effect of about 1.3 $\%$
at room temperature in small fields. This value is larger by about 30 $\%$
compared to the best low-field AMR effect of Ni NTs reported previously (see
the supplementary information of Ref. 11).

###  2.4 Spin dynamics in Ni nanotubes

Figure 7:  (a)Sketch of the excitation-detection scheme to probe dynamic
magnetization precession in individual NTs. The NT is covered by a coplanar
waveguide (CPW) used for excitation. Laser light is focused through the
transparent substrate onto the NT and the back reflected beam is analyzed by
an interferometer. The field $H$ is collinear with the long axis of the NT.
(b) SEM micrograph of the gold CPW laying on the investigated Ni nanotube S2.
Inset displays a magnified image of S2 under the signal line. (c) BLS spectra
detected on S2 at different applied static magnetic fields. Black circles
indicate peaks attributed to resonance frequencies (eigenmodes). (d) Extracted
eigenmodes plotted as a function of the field and compared with the resonance
frequency given by the Kittel equation 54 for a nickel rod (dotted gray line)
and a nickel thin film (continuous gray line) using $M_{\rm s}=380$ kA/m and a
gyromagnetic ratio of 28.02 GHz/T. (e) Spin wave dispersion relations
calculated for a 30 nm Ni thin film subjected to 27.5 mT (bottom line) and
72.5 m T (top line) by using the Kalinikos and Slavin formalism 55. The
resonance frequencies measured on S2 (boxes in d) are displayed as filled
circles matching the color code of the calculated lines. The wave vector
$k=k_{n}$ was estimated assuming azimuthal confinement of DE modes with
$n=0,1,2$ as illustrated above.

The spin dynamics of individual NTs deposited with $c=7$ were studied after
depositing them on a separate glass substrate (Figure 7a and Figure S5).
Coplanar waveguides (CPWs) made from 120 nm thick Au were fabricated by lift-
off processing such that a NT was fully covered by the signal line of the CPW
(Figure 7b). When a microwave current was applied to the CPW, a dynamic
magnetic field $h_{\rm rf}$ was generated, which exerted a torque on the spins
of the NT. To detect the induced spin precession we used microfocused BLS
($\mu$-BLS) 56. We focused a laser beam through the glass substrate onto the
NT from the backside. The backreflected light was analyzed by means of a
Fabry-Perot interferometer for energy shifts due to inelastic scattering
between the photon and the spin excitation. The chosen geometry with the NT
covered by Au and the laser focused through the substrate prevented the NT
from degrading due to the ambient atmosphere when heated by the focused laser
beam. BLS measurements were performed in the central position of the NT.
Figure 7c shows spectra obtained by $\mu$-BLS while sweeping the signal
generator frequency from 1.5 to 9 GHz in steps of 0.1 GHz. A static magnetic
field was applied along the long axis and varied in steps from 27.5 (bottom
curve) to 94.9 mT (top curve). At each increment we detect more than one peak
in the spectrum, indicating multiple resonant magnon modes (black circles)
whose eigenfrequencies vary systematically with $H$. In Figure 7d we summarize
the resonance frequencies as a function of field. For comparison we show with
a gray dotted and solid line the field-dependencies of FMR frequencies of a
cylindrical nickel rod and a nickel film, respectively, which were calculated
according to the Kittel formula 57 with a gyromagnetic ratio of 28.02 GHz/T
and $M_{\rm eff}=380$ kA/m, i.e. the value measured on the planar ALD-grown Ni
film. Considering the shape-dependent demagnetization effects, we expect that
the resonance frequency for uniform precession of a tube (FMR) resides below
the FMR frequency of the rod. Indeed, measured eigenfrequencies of the NT fall
below the dashed curve. The observed field-dependencies of measured eigenmodes
are similar to the slopes of the calculated curves (broken and dashed lines).
In the following, we speculate on the origin of the multiple observed
eigenmodes. We assume that the peculiar tubular shape of the NT gives rise to
standing spin waves with wavelengths $\lambda_{n}$ fulfilling the constructive
interference condition $n\times\lambda_{n}=3\times D$ along the azimuthal
direction, where $n$ is an integer number ($n=0,1,2,...$) and $3\times D$ is
the circumference $C$ of the hexagonal tube 10. The allowed wave vectors
amount to $k_{n}=2\pi/\lambda_{n}=n\times 2\pi/C$. The mode with $n=0$
represents uniform spin precession. The wave vectors $\mathbf{k}_{n}$ with
$n\geq 1$ point in azimuthal directions. They are orthogonal to the direction
of magnetization $\mathbf{M}$ of the NT which was found to be aligned with
$\mathbf{H}$ for $\mu_{0}H\geq 30$ mT (Figure 6). In such a configuration the
eigenfrequency of a spin wave is expected to increase with increasing $k$
(Damon-Eshbach mode configuration). In Figure 7e we plot the peak frequencies
observed at 27.5 mT and 72.5 mT in that for each field value we attribute
wavevectors $k_{n}$ with increasing $n$ to modes with increasing resonance
frequency. The solid lines shown in Figure 7e reflect the spin wave dispersion
relations f(k) calculated for a 30 nm-thick Ni film according to the formalism
provided by Kalinikos and Slavin 55. The experimental values follow the slope
of $f(k)$ expected for a thin film. However, the observed eigenfrequencies are
smaller than the calculated ones at a given field. The remaining discrepancy
could most likely be attributed to the different dynamic demagnetization
effect of the tubular geometry compared to the planar film and the increased
temperature of the NT in the laser focus. As the magnetization of Ni decreases
with increasing temperature, lower resonance frequencies would indeed be
expected. The results depicted in Figure 7 show that the Ni NTs form spin-wave
nanocavities which impose discrete wave vectors and confine GHz microwave
signals on the nanoscale. They can be produced in large ensembles of
vertically aligned 3D nanomagnetic structures (Figure 3).

###  2.5 Conclusions

Combining cycles of thermal ALD of nickel oxide with a single pulse of plasma-
assisted hydrogen reduction, we both obtained the conformal growth typical of
thermal ALD and exploited the low temperature reduction realized by a plasma
treatment. The optimized ALD sequence resulted in conformal coating of
vertically aligned nanotemplates with Ni where the growth rate was similar to
the growth rate of the planar films. We achieved a specific resistivity of Ni
of about 8 $\mu\Omega$cm on both planar substrates and arrays of
nanotemplates. The low-field AMR effect observed in the hysteretic regime of a
nanotube amounted to 1.3 $\%$ at room temperature. The spin-wave damping was
low allowing us to detect several standing spin wave modes which fulfilled the
constructive interference condition in azimuthal direction of a nanotube. The
Ni thin films and nanotubes thereby exhibited physical properties which make
them promising for functional spintronic elements and magnonic applications in
3D device architectures.

##  3 Methods

###  3.1 Plasma enhanced atomic layer deposition

Si (100) wafers were cleaved in pieces of circa 2 cm x 1cm and employed as
substrates for the deposition of planar Ni thin films. The Si substrates were
cleaned using isopropanol, DI water and dried by blowing nitrogen. GaAs NWs
were grown in a moelcular-beam epitaxy reactor as previously reported 40 and
used as nanotemplates for nanotubes depositions. They were inserted in the ALD
chamber without previous treatment. Nickel growth experiments were performed
in a hot wall Beneq TFS 200 ALD reactor, operated at a pressure of 4-5 mbar,
under a 100 sccm costant flow of ultrahigh purity nitrogen, used both as
carrier and purge gas. Prior to Ni growth, the Si (100) wafer pieces and GaAs
nanowires were coated with 5 nm of ALD Al2O3 in the same ALD chamber, by
alternating trimethylaluminium (TMA) and H2O pulses for 65 cycles. The growth
of nickel was performed using nickelocene (NiC5H5, NiCp2)as a precursor, water
as the oxidant agent and an in-cycle plasma enhanced reduction step with
hydrogen. Nickelocene and water were stored in stainless steel containers
respectively at 80 ∘ C and room temperature to exploit their vapor pressure.
The plasma was generated in a RF parallel plate system where one electrode
with shower head was powered at 150 W, the other was grounded and hosted the
sample substrates. An ALD cycle consisting of $c$ steps of NiCp2/H2O leading
to the formation of nickel oxide, followed by one step of plasma hydrogen,
reducing nickel oxide to metallic nickel, was repeated $n$ times. The ALD
sequence can be summarized as follow: $[[\ce{NiCp2}/{\rm purge}/\ce{H2O}/{\rm
purge}]\times c+\ce{H2}~{}{\rm plasma/purge}]\times n$. The pulse duration of
nickelocene and water in the NiCp2/H2O part of the sequence was set as 2 and
6s, respectively, and were followed by purge steps of 4 and 8 s. Hydrogen
plasma was supplied for 4 s, followed by a nitrogen purge of 8 s. The
thickness of the deposited thin films was measured imaging the films in cross
section by SEM. The growth rate was calculated dividing the thickness by the
total number of NiCp2 pulses $c\times n$ . The first set of depositions was
performed at different chamber temperatures of the range 160 - 175 ∘ C,
keeping c = 1 and n = 1000. A second set of samples was deposited at
$T=170~{}^{\circ}$ C to assess variations of growth rate on the number of
NiCp2/H2O steps $c$, by keeping $c\times n=1000$. To achieve thin films with
thickness of around 30 nm, a third set of depositions was carried out where
the ALD cycle number $n$ was set as 1000, 475, 350, 175 for c set as 1, 5, 7,
10, respectively. After nickel deposition, the samples were annealed in the
same ALD chamber at 350 ∘ C, for 5 hours, under a mixture of hydrogen and
nitrogen gas. No significant thickness variations were registered in the
planar thin films as a consequence of the annealing treatment.

###  3.2 Structural characterization and chemical analysis of the thin films
and nanotubes

We report the properties of the annealed thin films. Their morphologies and
thicknesses were investigated by scanning electron microscopy (SEM) and atomic
force microscopy (AFM). X-ray diffraction spectra were recorded in the
glancing incidence mode on a Malvern Panalytical(Empyrean model)
diffractometer with incidence angle of 2∘. The morphologies of the annealed
nanotubes was investigated by both SEM and transmission electron microscopy
(TEM), chemical element distribution was examined by scanning transmission
electron microscopy (STEM) combined with energy dispersive X-ray spectroscopy
(STEM-EDS). The TEM and STEM experiments were carried out using an FEI Talos
electron microscope operated at 200 kV. The thicknesses of the NTs were
extracted from EDS elemental 2D maps of Ga and Ni using Velox, which were
analyzed along specific lines cutting the NT perpendicularly to the length.
Line scans were smoothened by applying a Gaussian blur (sigma $=1.0$) on the
maps. The width of the line was 15-20 nm. Moving from the inner GaAs core
towards the external NT radius, we assumed the ALD Al2O3/Ni shell to start at
the point where the counts of the Ga element distribution (green curve)
vanished to zero and to end at the point where the counts of the Ni element
distribution (red curve)vanished to zero. We subtracted a nominal thickness of
5 nm ascribable to the Al2O3 layer to obtain the Ni thickness.

###  3.3 Investigation of physical properties

The static magnetic properties of the thin films were assessed using a SQUID
system (MPMS-5, Quantum Design) operated at room temperature (300 K).
Broadband spectroscopy of the deposited thin films was performed using a a
vector network analyzer (VNA)to sweep the frequency and record the
ferromagnetic resonance (FMR) absorption spectra. The thin films were
positioned on top of a CPW connected by microwave tips to a vector network
analyzer (VNA). The 2-port VNA allowed us to generate a microwave magnetic
field with frequencies ranging from 10 MHz to 10 GHz. The applied microwave
current generated an in-plane rf-magnetic field perpendicular to the long axis
of the CPW. The microwave with a power of $-15~{}$dBm was applied at the port
1 of the CPW in order to excite magnetization precession. The precession-
induced voltage was detected at port 2 via reading the scattering parameter
$S_{21}$ where the numbers 2 and 1 in the subscript denote the detection and
excitation port. An external magnetic field $\mu_{0}H$ was swept from 90 mT to
$-90~{}$mT along the CPW’s long axis. The linewidth of the resonant trace
acquired at each field was divided by $\sqrt{3}$ to obtain, with good
approximation, the imaginary component of the FMR response. Resistivity
measurements were performed with a KLA Tencor OmniMap RS75 four-point
resistivity meter. Anisotropic magneto resistance (AMR) measurements were
carried out in Van der Pauw four-point configuration 58. The samples were
bonded into a chip-carrier adapted for a room temperature probe station,
equipped with a custom built 2D vector magnet assembly that allowed us to vary
the in-plane applied field $H$ under an angle $\theta$. AMR measurements were
performed at room temperature applying a current of 2 $\mu$A and a static in-
plane magnetic field of 80 mT applied at an angle varying from 0 to 360 deg.

###  3.4 Experiments on nanotubes with integrated leads and waveguides

The Ni nanotubes were transferred trough an isopropyl alcohol solution on a
4-inch Si(100) wafer covered with 200 nm thick SiO2 for the fabrication of
metallic leads and on a 4-inch fused silica wafer for the fabrication of
integrated coplanar waveguides (CPWs). On both wafers pre-patterned gold
alignment markers were fabricated by photolithography via a custom developed
software previously reported 59 to precisely localize randomly oriented
nanotubes and generate the electron beam lithography pattern for integrated
leads and CPWs. The metallic leads (CPWs) were prepared by electron beam
lithography and a following sputtering (evaporation) of 5 nm Ti/ 400 (120) nm
Au film. The CPWs’ dimensions were chosen to enable impedance matching. The
signal line, having a width of 2.0 $\pm$ 0.1 $\mu$m was separated by gaps of
1.1 $\pm$ 0.1$\mu$m width from the 2$\mu$m wide ground lines. Magneto-
transport properties of nanotubes were investigated at room temperature
employing a probe station. The resistances were measured by 4 probe
measurements with a magnetic field applied in parallel to the tube axis.
Resistivities were calculated as $\rho=\dfrac{R\times A}{l}$ , where $l$ is
the separation between the voltage leads and $A$ is the cross-sectional area
$A$ of the shell around a NW. Correspondingly, $A$ was calculated as the
difference between an external hexagonal area with the hexagon’s long diagonal
$D_{\rm out}$ (measured by SEM) and an internal hexagonal area with diameter
$D_{\rm in}$ (taken as $D_{\rm out}$ \- $2\times t$, with $t$ being the
thickness of the Ni shell).
Spin wave eigenmodes were detected via Brillouin light scattering (BLS)
microscopy at room temperature 56. A coplanar microwave waveguide (CPW) was
fabricated on the glass substrate on top of the NT and oriented in a way that
the signal line covered the NT along its long axis. The end of the CPW was
electrically bonded to a printed circuit board, which was connected to a
signal generator (Anritsu MG3692C) applying a microwave current. The
corresponding magnetic microwave field excited spin precession in the NT at a
fixed frequency. The frequency was swept from 1.5 to 9 GHz. A monochromatic
laser with a wavelength of 473 nm and power of 0.8 mW was focused through the
sample backside to a diffraction limited spot using a specially corrected
$100\times$ objective lens with a large numerical aperture of NA = 0.85. The
recorded BLS signal is proportional to the square of the amplitude of the
dynamic magnetization at the position of the laser spot. The sample was
mounted on a closed loop piezo stage which allowed a precise localization of
the NT spots addressed in the experiments. The power was such that spin
precession was excited in the linear regime. A magnetic field was applied
parallel to the NT long axis via a permanent magnet.

Funding by the German Science Foundation DFG via GR1640/5-2 in SPP1538 “Spin
caloric transport” and SNF via grants 163016, BSCGI0_157705, and NCCR QSIT is
gratefully acknowledged. We thank Gözde Tütüncüoglu and Didier Bouvet for
excellent experimental support.

Further information and data are contained in the supporting information. Data
are available on zenodo.

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See pages - of suppinfo.pdf