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PHYSICAL REVIEW B 96, 195103 (2017)
Electronic properties and bonding in ZrHxthin films investigated by valence-band
x-ray photoelectron spectroscopy
Martin Magnuson,*Susann Schmidt, Lars Hultman, and Hans Högberg
Thin Film Physics Division, Department of Physics, Chemistry and Biology (IFM), Linköping University, SE-58183 Linköping, Sweden
(Received 24 May 2017; revised manuscript received 20 September 2017; published 1 November 2017)
The electronic structure and chemical bonding in reactively magnetron sputtered ZrHx(x=0.15, 0.30, 1.16)
thin films with oxygen content as low as 0.2 at.% are investigated by 4d valence band, shallow 4p core-level,
and 3d core-level x-ray photoelectron spectroscopy. With increasing hydrogen content, we observe significant
reduction of the 4d valence states close to the Fermi level as a result of redistribution of intensity toward the H
1s–Zr 4d hybridization region at ∼6 eV below the Fermi level. For low hydrogen content (x=0.15, 0.30), the
films consist of a superposition of hexagonal closest-packed metal (αphase) and understoichiometric δ-ZrHx
(CaF2-type structure) phases, while for x=1.16, the films form single-phase ZrHxthat largely resembles that
of stoichiometric δ-ZrH2phase. We show that the cubic δ-ZrHxphase is metastable as thin film up to x=1.16,
while for higher H contents the structure is predicted to be tetragonally distorted. For the investigated ZrH1.16
film, we find chemical shifts of 0.68 and 0.51 eV toward higher binding energies for the Zr 4p3/2and 3d5/2
peak positions, respectively. Compared to the Zr metal binding energies of 27.26 and 178.87 eV, this signifies
a charge transfer from Zr to H atoms. The change in the electronic structure, spectral line shapes, and chemical
shifts as a function of hydrogen content is discussed in relation to the charge transfer from Zr to H that affects
the conductivity by charge redistribution in the valence band.
DOI: 10.1103/PhysRevB.96.195103
I. INTRODUCTION
Zirconium hydrides are important in many applications
in the nuclear industry [1], in getter vacuum pumps, as
hydrogen storage, in powder metallurgy, and as zircaloy [2].
The hardness, ductility, and tensile strength of ZrHxalloys can
be controlled by varying the hydrogen content [3]. Increasing
the H content results in substoichiometric δ-ZrHxphase
(CaF2-type structure, x=∼1.6−2[4] as shown in Fig. 1)
and a body-centered tetragonal εphase (x=1.75−2) with
aThH
2-type structure [5] has also been observed [4]. The
δ-ZrHxstructure originates from incerting H atoms to occupy
all or parts of the tetrahedral interstitials in the CaF2structure.
Hydrogen acts as a hardening element that prevents dislocation
movements and the hydride material becomes a ceramic that
is harder, but less ductile than Zr metal [6].
Traditionally, bulk Zr hydrides are synthesized by
annealing Zr metal in hydrogen gas for a time period of days
to a few weeks at high temperatures between 400 and 900 ◦C
for a homogeneous diffusion process at different H contents
[7]. However, defects in the bulk of the Zr metal result in
variations in the diffusion rate for hydrogen that cause ZrHx
alloys with composition gradients and less well-defined
polycrystalline structure including grain boundaries. At room
temperature in air, ZrHxquickly forms a nanometer-thin
surface oxide layer that prevents further oxidation into the
bulk. With increased annealing temperature, oxygen proceeds
deeper into the bulk of the material, in particular along grain
boundaries [8], and a few percent of oxygen cannot be avoided
with the annealing-diffusion hydration synthesis method.
Consequently, the above conditions hinder adequate bonding
determination in hydrides by spectroscopy.
Previous experiments on bulk ZrHxusing ultraviolet pho-
toelectron spectroscopy showed significant changes in the
*Corresponding author: Martin.Magnuson@liu.se
electronic structure for high hydrogen contents (x=
1.63−1.94) both at the states near the Fermi level (EF) and
in deeper-lying states around −7eV[9]. These changes were
associated with a phase transition from a cubic (fcc) structure
to a face-centered tetragonal (fct) structure as a result of
spontaneous symmetry breaking that removes the degeneracy
by a distortion known as the Jahn-Teller effect [10] that lowers
the total energy of the 4d25s2electron configuration in the Zr
valence band.
For high hydrogen contents (x=1.63−1.9), core-level
x-ray photoelectron spectroscopy (XPS) studies of bulk ZrHx
materials indicated a significant shift of the Zr 4p and 3d core
levels by 1.0 and 0.7 eV [7], respectively. For x=1.52−1.68,
valence-band XPS also showed an interstitial Zr–H bonding
peak around 6.4 eV below EF[11]. Although these studies
indicate significant changes in the Zr–H bond for different
hydrogen contents, the spectra are affected by superimposed
O2pstates that occur in the same energy region (5–8 eV) as
the Zr–H bonding structures [12].
For the ZrHxthin films in this work, we apply XPS at the
Zr 3d,4p core levels and the 4d valence band to investigate
the electronic structures to determine the chemical bonding
and conductivity properties as a function of relatively low
hydrogen content (x=0.15, 0.30, and 1.16) and compare the
result to bulk α-Zr as reference. These hydride compositions
were chosen to obtain a spread in the type and amount
of Zr–H bonds. Deposited δ-ZrHxfilms with CaF2-type
structure are found to be stable outside the homogeneity
region determined for bulk δ-ZrH2with compositions ranging
from about ZrH1.5to stoichiometric [13]. To control the
structure and hence the electronic properties of the materials
and to overcome the problems with oxidation, we deposited
homogeneous thin films of ZrHxusing reactive direct current
magnetron sputtering (rDCMS). The thin films were grown
without external heating and contamination of oxygen, and
other contaminates were very low. Thin films are particularly
2469-9950/2017/96(19)/195103(7) 195103-1 ©2017 American Physical Society
MAGNUSON, SCHMIDT, HULTMAN, AND HÖGBERG PHYSICAL REVIEW B 96, 195103 (2017)
FIG. 1. Unit cell of stoichiometric δ-ZrH2(space group 225) with
Zr-atoms (red spheres) in (0,0,0), (½,½,0), (½,0,½), and (0,½,½)and
with H (white spheres) in all tetrahedral sites (¾,¼,¼), (¼,¾,¼),
(¼,¼,¾), (¾,¾¾), (¼,¼,¼), (¾,¾,¼), (¼,¾,¾), and (¾,¼,¼).
favorable for correct characterization of the material’s
fundamental bonding properties and to identify different
phases in understoichiometric transition-metal hydrides.
II. EXPERIMENTAL
A. Thin-film deposition and XRD characterization
The Zr-H films studied were deposited on Si(100) sub-
strates by rDCMS. ZrHxfilms with x=0.15, 0.30, and 1.16
were deposited in an industrial high-vacuum coating system
(CemeCon AG, Würselen, Germany). Here, a zirconium target
wassputteredinAr(99.9997%)/H2(99.9996% mixtures using
a fixed Ar partial pressure of 0.42 Pa with 5, 10, and 20% H2.
The substrate bias was set to −80 V and no external substrate
heating was applied. The films were deposited for 120 s,
resulting in films with thicknesses ranging between ∼800
and ∼840 nm. More detailed information on the deposition
conditions is presented in Ref. [14]. The investigated α-Zr
reference sample was a commercial zirconium target with a
purity of 99.9% (Kurt J. Lesker Company, Clairton, PA, USA).
B. XPS measurements
Valence-band (VB) and XPS measurements of the Zr 3d,
4p,4dvalence band and O 1score-level regions were per-
formed in a surface analysis system (AXIS UltraDLD, Kratos,
Manchester, UK) using monochromatic Al-Kα(1486.6 eV)
radiation with an incidence angle of 54◦and a spot size
of 300 ×800 μm. The electron energy analyzer detected
photoelectrons perpendicular to the sample surface with an
acceptance angle of ±15◦. The spectra were recorded with
a step size of 0.1 eV and a pass energy of 10 eV, which
provided an overall energy resolution better than 0.5 eV. The
binding-energy scale of all XPS spectra was referenced to the
EF, which was set to a binding energy of 0 eV [15]. The ZrHx
thin films in this work were examined before and after Ar+
sputtering for 600 s using an Ar+incident angle of 20◦at 4 keV
rastered over an area of 2 ×2mm
2.
C. Structural model and density-functional theory calculations
The geometry relaxation was performed using the
Perdue-Burke-Ernzerhof (PBE) functional [16] including the
Grimme–van der Waals density-functional theory (DFT)-D2
scheme [17]. The first-principle calculations were carried out
using DFT implemented in the Vienna ab initio simulation
package (VASP)[18] with an exchange-potential functional
using the general gradient approximation (GGA). A hydrogen-
potential suited for short bonds (hydrogen projected aug-
mented wave (PAW) potential H_h) with an energy cutoff
of 1050 eV was used. For the self-consistent calculations,
the PAW [19] method was used with the PBE and the
Heyd-Scuseria-Ernzerhof functionals of the GGA. The HSE06
functional [20] is a hybrid functional including a linear
combination of short and long-range PBE exchange terms and
a short-range Hartree-Fock term that improves the formation
energies and band gaps and peak positions. The screening
parameter of HSE06 was 0.2˚
A−1with a plane-wave cutoff
energy of 700 eV. A 29 ×29 ×29 kgrid was used for the
standard PBE functional for the structure relaxation and an
11 ×11 ×1kgrid was used for the HSE06 functional as
the computational cost is significantly higher for the HSE06
in comparison to the PBE functional. Using the HSE06
exchange-correlation functional, the energy positions of the
bands are shifted by ∼1.8 eV toward higher energy relative to
the EF, and are thus much more realistic, in comparison to the
corresponding energy positions, using the PBE functional. The
charge-transfer calculations between the different elements
were made using standard Bader analysis [21].
III. RESULTS AND DISCUSSION
Figure 2shows Zr 3dcore-level XPS spectra of the
ZrHxthin films in comparison to α-Zr metal. As observed,
there is a significant chemical shift toward higher binding
energies for the film grown with hydrogen in the plasma
compared to the pure metal film, which is a consequence of the
higher electronegativity of H (χ=2.2) in comparison to Zr
(χ=1.33) [22]. For x=1.16, the Zr 3d5/2peak position was
determined to 179.38 eV. This is in good agreement with the Zr
3d5/2binding energy of 179.375 reported for bulk δ-ZrH1.74
[23] as well as 179 eV determined for bulk ZrH1.9material
[24] and close to the 180-eV binding energy determined for
bulk ZrH1.64 [25] and ZrH1.9[7]. Furthermore, the position
of the 3d5/2peak is more or less identical to the binding
energy determined for the reactively sputtered ZrH1.64 film in
Ref. [14] with 179.44 eV. Note the small energy shift toward
lower binding energy for the less hydrogen-rich film with
x=1.16 (179.38 eV) that is consistent with a smaller charge
transfer from Zr to H in this film in comparison to the x=1.64
film (179.44 eV). In addition, for the ZrH1.16 film, there is a
high-energy shift of +0.51 eV while between x=0.15 and
0.30, the Zr 3dcore-level high-energy shift is much smaller
(+0.06 eV) and is an indication of smaller average charge
transfer from Zr toward H in these samples.
The metallic reference sample has the Zr 3d5/2and Zr
3d3/2peak positions at 178.87 and 181.28 eV (2.41-eV peak
splitting) in good agreement with literature values [26]. The Zr
3d5/2-3d3/2spin-orbit splitting is the same in the films.Prior
to sputtering, the spectra showed structures of ZrO2following
air exposure, but at the Zr 3dedge after sputtering there are
no features connected to oxygen.
195103-2
ELECTRONIC PROPERTIES AND BONDING IN ZrHx. . . PHYSICAL REVIEW B 96, 195103 (2017)
Intensity (arb. units)
184 182 180 178
Binding Energy (eV)
x
=0.00
x
=0.15
x
=0.30
x
=1.16
3d5/2
3d3/2
Zr 3d XPS
FIG. 2. Zr 3dcore-level XPS spectra of the ZrHxthin films in
comparison to α-Zr metal. The 3d5/2,3/2spin-orbit splitting (2.41 eV)
is indicated by the horizontal arrow.
To evaluate the asymmetric tails toward higher binding
energies in each spectrum, a Doniac-Sunjic function [27]
corresponding to the electron-hole pair excitations created at
the Fermi level to screen the core-hole potential was applied.
The Doniac-Sunjic profile is essentially a convolution between
a Lorentzian with the function 1/E(1 −α), where Eis the
binding energy of each peak and αis a parameter known as
the singularity index that is related to the electron density of
states at the Fermi level [28]. The intensities of the tails are
related to the amount of metallicity in the system due to the
coupling of the core-hole with collective electron oscillations.
For α-Zr, the singularity index is large (α=0.15) and the
3d5/2/3d3/2branching ratio is higher (1.75) than the statistical
value of 1.67 (5/3), signifying high conductivity. For x=0.15
and 0.30, the singularity index and branching ratio remain
almost the same as for α-Zr, but for x=1.16, the tail becomes
significantly smaller with a singularity index of only α=0.04
with a branching ratio that is smaller (1.61) than the statistical
ratio (1.67). This shows that the number of bands crossing the
EFare reduced and thus the expected conductivity.
Figure 3shows Zr 4pshallow core-level XPS spectra of
the ZrHxthin films in comparison to α-Zr metal. As in the
case of the Zr 3dlevel, the largest chemical shift of +0.68 eV
from 27.26 eV in the metal to 27.94 eV is found for the film
with x=1.16, while a smaller shift of 0.1 eV occur for x=
0.30 with a similar shift in the composition range x=0.15.
The binding energy determined for the Zr 4p3/2peak of the
Intensity (arb. units)
40 35 30 25 20
Binding Energy (eV)
ZrHx4p-edge
x=1.16
x=0.30
x=0.15
x=0.00
Zr
4p
4p1/2
4p3/2
O 2s
FIG. 3. Zr 4pcore-level XPS spectra of the ZrHxthin films in
comparison to α-Zr metal. The 4p3/2,1/2spin-orbit splitting (1.3 eV)
is indicated by the vertical lines, from Ref. [26].
hydride with x=1.16 is slightly lower than the 28.1- and the
28.0-eV values reported in Refs. [9] and [12], respectively. For
both these studies, the spectra were recorded for oxidized bulk
samples, where the interaction of Zr atoms with O causes the
Zr 4ppeak to shift to a higher binding energy. Finally, the
4p3/2peak position of the metal reference of 27.26 eV is in
agreement with the reported value of 27.1 eV in Ref. [26].
The observed binding energies of the α-Zr spin-orbit split
4p3/2,1/2levels are 27.26 and 28.60 eV (1.3 eV splitting). This
is consistent with tabulated values of 27.1 and 28.5 eV [26].
At the shallow Zr 4plevel, we observe both a significant
broadening of the spectra and a high-energy shift in the
binding energy relative to metallic α-Zr. Only for the film
with x=0.30 some oxygen is left even after sputtering.
This was also observed by time-of-flight energy elastic recoil
detection analysis (ToF-ERDA) showing a slightly higher
oxygen content of 0.7% for this film [14] compared to 0.2%
for the films with x=0.15 and 1.16.
Figure 4shows a set of high-resolution VBs XPS spectra
of the ZrHxthin films (0.15, 0.30, and 1.16) in comparison
to α-Zr metal. In the region 0–2 eV from the EF, the spectra
are dominated by the Zr 4dvalence states with a 4d5/2−3/2
spin-orbit splitting of ∼1.3 eV, most clearly observed in
pure α-Zr (x=0). For x=0.15, the shape of the VB-XPS
spectrum is very similar to that of pure α-Zr with a larger
broadening and appears to be very little affected by the small
hydrogen content. On the contrary, for x=1.16, a major
spectral redistribution of intensity changes of the shape of
the Zr 4dband is further broadened so that the spin-orbit
splitting is no longer resolved. The most prominent new
feature appears around 6 eV binding energy and is due to the
195103-3
MAGNUSON, SCHMIDT, HULTMAN, AND HÖGBERG PHYSICAL REVIEW B 96, 195103 (2017)
Intensity (arb. units)
12 10 8 6 4 2 0 -2
Binding Energy (eV)
ZrHx Valence band
x=1.16
x=0.30
x=0.15
x=0.00
Zr
4d
Zr-H
hybr.
EF
4d3/2
4d5/2
O 2p
FIG. 4. Valence-band XPS spectra of the ZrHxthin films in
comparison to α-Zr metal.
strong H 1s–Zr 4dhybridization and bonding. For x=0.30, a
significant broadening of the Zr 4dvalence states at ∼1eVis
observed. The largest change is in the double-peaked feature
between 6 and 7.8 eV. This feature is consistent with previous
observations in bulk material [23] with δ-ZrH1.52,δ-ZrH1.65,
and δ-ZrH1.74, where the height of the Zr 4dpeak close to EF
decreases and the Zr–H bonding peak increases with increasing
hydrogen content. These observations suggested a donation
of Zr 4delectrons toward the Zr–H bond, while the Zr–Zr
bond weakens. For higher hydrogen content (x=1.94) the
states closest to the EFwere enhanced due to a fcc-fct phase
transition [9]. However, bulk materials are often hampered by
the appearance of superimposed O 2pstates between 5 and
10 eV, i.e., in the same energy region as the Zr–H bonding
peak. As ZrO2is known to have prominent peaks of O 2p
states between 5 and 10 eV and O 2sstates around 22.5 eV
from the EF[12]. Thus, in bulk materials, the contribution of
the Zr–H bond peak cannot be fully distinguished or separated
(disentangled) from the O 2pstates due to oxygen impurities.
Figure 5shows calculated density of states (DOS) for
crystalline δ-ZrHxfor x=0, 0.5, 1.0, 1.5, and 2.0. As
observed, for x=0, the states within 4 eV from the EFare
dominated by Zr 4dstates while strong Zr 4pstates occur at
29–30 eV from EF. This is also the case for increasing xbut a
prominent peak in the energy region 5–10 eV is mainly due to
H1sstates that does not exist in pure Zr. At 5–10 eV below EF,
bonding H–H s-sσ,Zr–Hd-sσ, and Zr–Zr d-dσ-interactions
occur while antibonding H–H s-sσ∗interactions appear above
EF. We find that the H 1scharacter in δ-ZrHxstrongly depends
on the hydrogen content and increases with x. Notably, a deep
DOS minimum at 2–4 eV occurs for x=0.5 when hydrogen
is introduced. Closer to the EF,Zr4d-egstates dominate at
Density of States (DOS/eV/atom)
-35 -30 -25 -20 -15 -10 -5 0
Energy (eV)
Zr s
Zr p
Zr d
H s
H p
ZrH0.5
Zr
ZrH1.0
ZrH1.5
ZrH2.0
Zr 4p
FIG. 5. Calculated DOS for α-Zr and δ-ZrHx,wherex=0.5,
1.0, 1.5, and 2.0. To enhance the valence bands, the intense Zr 4p
states at −30 eV were scaled down by a factor of 5.
0–2 eV while H 1sstates are negligible. For δ-ZrH2, there is
an intense peak at the EFof Zr 4d-t2gcharacter that dominates
and causes instability of the crystal structure at 0 K. At finite
temperature, phonon interactions and tetragonal (Jahn-Teller)
distortion of the structure leads to a splitting of the intense
t2gpeak at the EFand a pseudogap develops. This leads to
lowering of the total energy of the system and restabilization
of the structure. Thus, the t2gstates at EFare reduced at
finite temperatures and likely reduce the conductivity and
other transport properties. This is consistent with previous
DFT studies including the δ- and εphases [29–32]ofZrH
x
while comparisons to experimental electronic structure studies
of the Zr–H bonds of thin films have been lacking. Thus, the
order and bonding of the hydrogen atoms in the octahedral
sites has not been well understood.
Figure 6shows the structures (left) and calculated total
energies (right) of ZrHxfor x=0.5, 1.0, 1.5, and 2.0 (top
to bottom). Here, Jahn-Teller distortion is predicted for high
hydrogen content (x>1), where the cubic δstructure becomes
unstable and lowers the symmetry. On the contrary, for low
hydrogen content, the cubic structure is stable (x=0.5) or
metastable (x=1.0). For x=2.0, the minimum of c/a is
located at 0.886, for x=1.5 at 0.868, while for x=1 there is
a metastable state for c/a =1.0 and a lowest stable state for
c/a =0.792. Thus, the cubic δ-ZrHxphase is metastable as
thin film up to x=1.16, while for higher H contents the ZrHx
structure is predicted to be distorted by the Jahn-Teller effect.
195103-4
ELECTRONIC PROPERTIES AND BONDING IN ZrHx. . . PHYSICAL REVIEW B 96, 195103 (2017)
FIG. 6. (Left panel): Model structures. (Right panel): Calculated
total energies for ZrHx,wherex=0.5, 1.0, 1.5, and 2.0.
To evaluate the charge-transfer dependence on x,we
calculated the Bader charges that are listed in Table Itogether
with the lattice parameters for different compositions. Unit
cells of δ-ZrHx(Fig. 6, left panels), where the H atoms
were replaced by vacancies, were used as model systems for
different x. Among the hydrides, the structure and stability
properties of ZrH2have been of particular interest where the
attention has been focused on the δ-εphase transition induced
by stress, and the influence of the Jahn-Teller distortion to
stabilize the structure [30,33]. As the electronegativity of Zr
(1.33) is lower than for H (2.2), the formal oxidation states of
H and Zr in ZrH2are −1 and +2, respectively. Generally, the
predicted charge transfer from Zr increases with increasing
xand lattice parameter. In comparison to the pure Zr and H
TABLE I. Calculated lattice parameters and charge transfer for
ZrHxfor different x. Experimental values for α-Zr are given in
parenthesis [34].
System a-lattice parameter c-lattice parameter Zr H
α-Zr 3.2577 (3.2327) 5.1358 (5.1493) 0 –
ZrH0.54.6484 4.6607 +0.439e−0.878e
ZrH1.04.7457 4.7457 +0.845e−0.845e
ZrH1.54.7554 4.75547 +1.175e−0783e
ZrH2.04.7811 4.78117 +1.491e−0.745e
elements, the Bader charge on the Zr atoms in, e.g., ZrH2,
increases significantly by +1.49e, i.e., significant charge is
withdrawn from the 4d25s2valence orbitals of Zr. According
to the Bader analysis of ZrH2, the charge is transferred
to the lowest unoccupied orbital, of the hydrogen atoms
(1s1→1s1.75), where the Bader charge decreases by −0.75e
on each atom. A distortion of the cubic δ-ZrHxstructure
causes a slight reduction of the charge transfer.
The combined 3dand 4pcore-level and 4dvalence-band
studies show several interesting effects. For the hydride
system, we observe a significant reduction of the Zr 4dband
within ∼3 eV from the EFinfluencing the shallow 4pcore level
by a broadening and an asymmetric redistribution of intensity
toward higher binding energy. There is no trace of oxygen on
the samples, except for x=0.30, where small signs of O 2p
and O 2sintensity can be distinguished. For the hydrides, a
bonding band at 6 eV below the EFdoes not occur in pure Zr
and is related to strong H 1s–Zr 4dhybridization. This is also
accompanied by a chemical shift with binding energies that
are 0.6 eV higher than pure α-Zr indicating significant charge
transfer toward H, in particular for x=1.16. In our x-ray
absorption near-edge structure study [34], a chemical shift
toward higher energies in comparison to Zr metal was found
to be due to changes in the oxidation state that depend on
the structure and the formation of Zr–H bonding at sufficient
hydrogen loading. Previous valence-band XPS experiments
have shown a dominant Zr 4d peak at 1 eV and another peak
around 6.4 eV corresponding to a Zr–H bond [11], consistent
with DOS calculations [29–31]. The chemical shift of the Zr
edges toward higher bonding energies confirms a significant
charge transfer from Zr toward H with increasing hydrogen
content. It has been argued that a tetragonal distortion of the
cubic δ-ZrHxstructure in combination with phonons make the
structure stable [30]. However, as shown in Fig. 6, this is only
the case for high H content, while at low H content, the cubic
structure is stable or metastable and does not distort. Thus, for
low H content, the cubic δstructure with tetrahedral Zr sites
favorably affects the materials properties in comparison to a
tentative octahedral coordination that for hydrogen is inhibited
according to the Hägg rules on size, coordination number, and
electronegativity.
For ZrHxfilms deposited by rDCMS, we have previously
shown that it is possible to grow films with tailored composi-
tions by altering the concentration of H2in the plasma [14].
The properties of the films resemble those of the parent Zr
metal where increasing hydrogen content in the films yields
higher resistivity seen from the values ranging from ∼70 to
∼120 μ cm in comparison to 42.6μ cm for α-Zr [35].
The measured hardness values are around 5.5 GPa, which
is harder than the ∼3 GPa determined for bulk α-Zr, using
nanoindentation [36].
In studies of bulk ZrHx, there have been problems with
oxide, as Zr is known to slowly form stable ZrO2on the
surface and at the grain boundaries [12]. Previous studies have
thus been limited to more or less oxidized polycrystalline
bulk materials that contain randomly ordered crystallites
with grain boundaries that affect the results of electronic
structure characterization. Moreover, investigations of the
electronic structures of understoichiometric transition-metal
hydrides have been scarce. Thin films synthesized by reactive
195103-5
MAGNUSON, SCHMIDT, HULTMAN, AND HÖGBERG PHYSICAL REVIEW B 96, 195103 (2017)
magnetron sputtering offer an alternative route to grow metal
hydrides such as ZrHxwith well-defined properties and at
a fraction of time (deposition time of about 2 min for an
800-nm film) compared to traditional hydration/percolation
processes that normally takes 1–2 days or weeks depending
on the H content [7]. A further advantage of the films
is the low level of contaminants, in particular, the oxygen
content that is of the range of 0.2−0.7% as determined
by ToF-ERDA [14]. Although 3dcore-level shifts in ZrHx
[14] have previously been observed, indicating charge-transfer
from Zr to H, a deeper analysis including branching ratio and
singularity index is necessary for the understanding how the
electronic properties influence the conductivity. For low H
content (x=0.15 and 0.30), the films contain a mixture of
metallic α-Zr and δ-ZrHxphases and the metallicity remains.
For higher H content (x=1.16), a homogeneous oriented,
textured, and understoichiometric δ-ZrHxphase is formed
with a lower singularity index and smaller branching ratio
that overall indicate more ceramic properties than for the
lower H contents. Producing highly stoichiometric and phase
pure materials is thus important for optimizing electrical and
hardness properties in a vast number of applications.
IV. CONCLUSIONS
Early transition-metal hydrides are a new class of thin-film
materials where the cubic δ-ZrHxphase is metastable up to
x=1.16, while for higher H contents the ZrHxstructure is
predicted to be distorted by the Jahn-Teller effect both in thin
films and bulk synthesized materials. By combining valence-
band and shallow core-level x-ray photoelectron spectroscopy
with electronic structure calculations, we have investigated the
electronic structure and chemical bonding of ZrHxthin films
with x=0.15, 0.30, and 1.16 in comparison to α-Zr metal. For
the understoichiometric δ-ZrH1.16 film, there are significant
chemical shifts of 0.68 and 0.51 eV for the 4p3/2and 3d5/2
peak positions toward higher energies compared to the metal
values of 27.26 and 178.87 eV. We find that even though there
is a significant charge transfer from the Zr 4dstates toward
the H 1sstates in δ-ZrH1.16, there are still states crossing the
Fermi level that signify metallicity. This is due to the fact
that there is a significant asymmetric redistribution of spectral
intensity from the shallow Zr 4pcore level toward the Zr 4d
valence band that accompanies the charge-transfer. There are
important changes at the Fermi level with a splitting of the 4d
band into a pseudogap where the Fermi level is located in a
local minimum that stabilize the structure and minimize the
total energy of the system. For the hydrides, a bonding band
around 6 eV below the Fermi level does not occur in pure Zr
and is related to strong H s–Zr 4dhybridization. This is also
accompanied by a chemical shift with binding energies that are
0.68 eV higher than pure α-Zr, indicating significant charge
transfer toward H, in particular for x=1.16.
ACKNOWLEDGMENTS
The research leading to these results has received funding
from the Swedish Government Strategic Research Area in
Materials Science on Functional Materials at Linköping Uni-
versity (Faculty Grant SFO-Mat-LiU No. 2009-00971). M.M.
acknowledges financial support from the Swedish Energy
Research (Grant No. 43606-1), the Swedish Foundation for
Strategic Research (SSF) (Grant No. RMA11-0029) through
the synergy grant FUNCASE, and the Carl Trygger Foundation
(Grants No. CTS16:303 and No. CTS14:310). Dr. Grzegorz
Greczynski is acknowledged for assistance with the measure-
ments of the Zr 3dcore-level XPS spectra.
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