# A review of uranium-based thin films

R. Springella, E. Lawrence Brighta,b, D. A. Chaneya,b, L. M. Hardinga, C.
Bella, R. C. C. Wardc and G. H. Landera,d Corresponding author email:
<EMAIL_ADDRESS><EMAIL_ADDRESS>a School of Physics,
University of Bristol, Tyndall Avenue, Bristol BS8 1TL. b European Synchrotron
Radiation Facility, 38040 Grenoble, France. c Clarendon Laboratory, Oxford
Physics, Parks Road, Oxford OX1 3PU, UK. d European Commission, Joint Research
Centre (JRC), Directorate for Nuclear Safety and Security, Postfach 2340,
D-76125 Karlsruhe, Germany.

###### Abstract

Thin films based on silicon and transition-metal elements dominate the semi-
conducting industry and are ubiquitous in all modern devices. Films have also
been produced in the rare-earth series of elements for both research and
specialized applications. Thin films of uranium and uranium dioxide were
fabricated in the 1960s and 1970s, but there was little sustained effort until
the early 2000s. Significant programmes started at Oxford University
(transferring to Bristol University in 2011), and Los Alamos National
Laboratory (LANL) in New Mexico, USA. In this review we cover the work that
has been published over the last $\sim$20 years with these materials.
Important breakthroughs occurred with the fabrication of epitaxial thin films
of initially uranium metal and UO2, but more recently of many other uranium
compounds and alloys. These have led to a number of different experiments that
are reviewed, as well as some important trends. The interaction with the
substrate leads to differing strain and hence changes in properties. An
important advantage is that epitaxial films can often be made of materials
that are impossible to produce as bulk single crystals. Examples are U3O8,
U2N3 and alloys of U-Mo, which form in a modified bcc structure. Epitaxial
films may also be used in applied research. They represent excellent surfaces,
and it is at the surfaces that most of the important reactions occur in the
nuclear fuel cycle. For example, the fuel-cladding interactions, and the
dissolution of fuel by water in the long-term storage of spent fuel. To
conclude, we discuss possible future prospects, examples include bilayers
containing uranium for spintronics, and superlattices that could be used in
heterostructures. Such applications will require a more detailed knowledge of
the interface interactions in these systems, and this is an important
direction for future research.

###### keywords:

Uranium, actinides, thin films, epitaxy

††articletype: REVIEW ARTICLE

###### Contents

  1. 1 Introduction
    1. 1.1 Early efforts (before $\sim$ 2000) on uranium-based films
  2. 2 An overview of the growth of uranium-based films
    1. 2.1 Deposition Techniques
    2. 2.2 Characterisation Techniques
  3. 3 Uranium Metal Phases and Alloys
    1. 3.1 Introduction
    2. 3.2 Production of metallic uranium films
      1. 3.2.1 Uranium containing multilayers
      2. 3.2.2 Achieving single crystal films of $\alpha\mbox{-U}$
      3. 3.2.3 The hunt for hexagonal close packed uranium
      4. 3.2.4 Achieving single crystal films of $\gamma\mbox{-phase}$ alloys
    3. 3.3 Science with uranium metal thin films
      1. 3.3.1 Induced magnetism in uranium-ferromagnetic multilayers
      2. 3.3.2 Uranium metal bilayers for future spintronic applications
      3. 3.3.3 Manipulating the CDW States in $\alpha\mbox{-uranium}$ using epitaxial strain
      4. 3.3.4 Correlated disorder in pseudo-bcc uranium alloys
      5. 3.3.5 Thin films of transuranium metals
  4. 4 Uranium Oxide Systems
    1. 4.1 Introduction
    2. 4.2 Surface studies of UO2
    3. 4.3 Growth of oxides
    4. 4.4 Science with UO2 thin films
      1. 4.4.1 Photoemission experiments on UO2 and PuO2 epitaxial films
      2. 4.4.2 Antiferromagnetism of UO2 epitaxial films
      3. 4.4.3 Enhanced paramagnetism in strained epitaxial UO2 films
      4. 4.4.4 Search for exchange bias using UO2 thin films
      5. 4.4.5 Dissolution studies of UO2
      6. 4.4.6 Studies of phonons with irradiated UO2
      7. 4.4.7 Studies of irradiated films
      8. 4.4.8 Use of films to assure high-quality surfaces
      9. 4.4.9 Studies of higher oxides
  5. 5 Uranium Hydrides, Nitrides, and Silicides
    1. 5.1 Introduction
    2. 5.2 Growth of thin films of hydrides, nitrides, and silicides
    3. 5.3 Science with thin films of hydrides, nitrides, and silicides
      1. 5.3.1 Thin films of uranium hydrides
      2. 5.3.2 Magnetism and electronic structure of uranium nitride thin films
      3. 5.3.3 Reactivity studies of UN
      4. 5.3.4 Thin films of uranium silicides
  6. 6 Conclusions and future prospects
  7. 7 Acknowledgments

##  1 Introduction

Thin films are ubiquitous in modern technology. They form the basis of the
semiconductor industry: from light emitting diodes to the millions of
transistors in every single computer central processing unit. Digital memory
technologies are similarly underpinned by thin films, from the spin-valve
heterostructures in hard drive read heads to ferroelectric random-access
memory. More recently, the rapid developments of several high profile quantum
computing architectures are based on thin films of superconducting aluminium.
It is not unreasonable to argue that thin films have transformed the
technology of the late 20th century, and continue to do so to this day. At the
same time, for many years thin films have provided researchers with key
insights into fundamental condensed matter physics. Examples in this area
include the discovery of the integer and fractional quantum Hall effects in
GaAs heterostructures [1] and oscillatory exchange coupling in magnetic/non-
magnetic multilayers [2, 3]. These discoveries earned Nobel prizes, but there
are a host of other novel effects in thin film layers and heterostructures.

Sitting at the bottom of the periodic table, the actinides are defined by the
presence of 5$f$ electrons which give rise to a plethora of weird and
wonderful physical properties [4] that vary significantly across the series as
the nature of the 5$f$ electrons changes from largely itinerant in Th and U,
to almost fully localized in Am and beyond, with the notoriously complex Pu
separating the two sub-series. The elements up to and including Pu exhibit a
vast range of isomorphs, whereas Am and beyond crystalise into a double
hexagonal close-packed (dhcp) structure and behave akin to the heavier rare-
earth metals. Likewise, superconductivity gives way to magnetism across the
series. Plutonium and uranium also display two properties unique to single
element materials: negative thermal expansion in Pu and ambient pressure
charge-density modulations in U. The phenomena found in actinide containing
compounds are no less fascinating and include heavy-fermion behavior [5, 6,
7], spin fluctuation states, large spin-orbit coupling, Jahn-Teller
distortions, quadrupolar ordering [8] piezoelectricity, and magnetorestriction
[9], to name but a few.

The marriage of thin films and actinides provides a vast parameter space of
experimental and theoretical exploration and there are some distinct areas of
study where real advances have been made, and there are still many exciting
opportunities for the future; in fundamental research and investigations on
applied nuclear fuel and waste materials, driven by a renewed global appetite
for advanced fuel materials for modern 21st century nuclear reactor fleets.
Fig.1 highlights the range and connectivity of the materials, properties and
phenomena that can be found in the current body of literature. This approach
offers some practical experimental and theoretical advantages over more
traditional bulk materials, as well as opening new scientific avenues for
study.

Figure 1: Image showing range of uranium-based materials; properties and
phenomena that have been investigated in thin film form.

These sample systems have macroscopic surface areas (typically of the order 1
cm2), with typical masses of 100’s of micrograms. This provides a basic
advantage for active work in that many facilities and institutions become
accessible that otherwise would be restricted, and transport, handling, and
storage of these samples is significantly easier than for their bulk
counterparts. For example, a typical 1000 Å film of UO2 would have an activity
of only $\sim$1.5 Bq and would contain a comparable uranium mass to that found
in the human body ($\sim$100 $\mu$g).

Sample synthesis is rapid compared to typical bulk methods, not by mass, but
by sample type. This means that one can study a range of compositions quickly,
which is ideal for surveys of phase diagrams, alloys, additives etc. These
techniques are good for controlling growth parameters in situ, which means
that one can mimic surfaces and interfaces of relevant nuclear materials, or
design multilayers/interfaces to explore fundamental 2D physics of 5$f$
states. One can engineer the structure, composition, phase, and stoichiometry,
one can grow amorphous materials, polycrystalline materials, controlling the
grain size, single crystals: modifying the crystal quality, strain fields etc.
It allows us to experimentally model many complex effects, (such as radiation
damage) in a much simpler way by more careful control of variables and removal
of complexity from bulk systems. This connects much more directly with
theoretical modelling in nuclear materials, where these idealised systems are
able to feed into computational models and vice versa.

Epitaxial matching to a substrate crystal allows fine control of the
crystalline growth. This raises the question of whether different allotropes
of the films can be stabilised by using different substrates, which act as
”guiding templates”. As we shall see, this use of substrates to apply strain
and alter the properties has already been done for $\alpha$-uranium metal.
Some success, also with uranium, has been achieved in preparing a hexagonal-
close packed (hcp) structure using templating. Notably here the hcp structure
does not exist in the bulk. If the interface between film and substrate is
fully understood there is the possibility that new allotropes of bulk
materials can be produced with unknown properties. The first challenge, of
course, is to find the right substrate and method to prepare an epitaxial
sample, but this search for new structures could clearly open exciting
scientific possibilities.

Broadly, the power of thin films for fundamental physics or materials science
can be categorised in terms of (a) surface and dimensional effects; (b)
proximity effects; (c) strain effects. Some of these are already being
exploited. The interface between the film and the substrate provides not only
flexibility in a geometric sense, but also provides a pathway for electronic
interactions. For example, the actinide elements have a large spin-orbit
parameter, which is roughly dependent on Z4, where Z is the atomic number, and
this parameter is important for spintronic applications. Preparing and
exploiting such samples clearly is a considerable challenge, but understanding
the interfaces is the primary one.

New physics can be expected, such as topological ground states. Such states
have already been predicted in 2012 [10] for Pu and Am compounds. This comes
from the fact that the 5$f$ states in Pu are almost exactly at the boundary
between itinerant (before Pu) and localized (after Pu), but this condition can
also be found for uranium compounds, e.g. UNiSn, [11] and UTe2 [12].
Initially, of course, the effort on actinide films has been confined almost
entirely to using the elements Th and U, which can be handled without
difficulty in most laboratories. However, in the longer term, facilities that
can handle actinides up to at least Cm can be envisaged. There are huge
advantages working with small quantities of these elements, and much that
could be discovered, especially about Pu, an element that has six allotropes
in the solid state before it melts [13]. In this respect it is worth noting
that the only transuranium epitaxial films produced so far are those of NpO2
and PuO2 at Los Alamos National Laboratory [14]. Some of the science deduced
from experiments on these samples are discussed below.

The 5$f$ electrons are also the key features of many heavy-fermion compounds,
including ferromagnetic superconductors containing uranium [7, 15, 5], where
it is commonly assumed that it is the hybridization of the conduction and 5$f$
states that lead to the peculiar properties of these systems. In the famous
case of URu2Si2 the physics of the system remains unresolved, despite a vast
amount of both theory and experiment [6] since the discovery of
superconductivity in this material in 1986. More recently, UTe2 has attracted
wide attention as a possible topological triplet superconductor [12]. The
prize for the highest superconducting Tc (18 K) of any heavy-fermion material
still goes to PuCoGa5, which is mainly a mystery 20 years after its discovery
[16]. All the studies referenced above have been performed on bulk samples.
Usually, but not all, the experiments reported used single crystals. If
epitaxial films of these materials, especially the heavy fermions, were
available, further experiments could be easily envisaged.

The basic research motivation for producing actinide thin films is thus
abundantly clear. But there is another motivation, equally important. Beyond
their undoubted fundamental interest, understanding actinide compounds has
substantial importance due to the property that all are radioactive and from
the 15 elements, one typically finds 7 fissile isotopes. From this subset,
fissile uranium-235 in the chemical form of UO2 powers the large majority of
the world’s operational nuclear reactors, generating $\sim$10% of world power,
which is more than a quarter of the world’s low carbon electricity production
[17].

Although much is understood about UO2 [18], there are still questions to be
answered, particularly relating to the surface and interface reactions and
properties of UO2. For example, the synthesis of fuel/clad interfaces opens up
the possibility of designing experiments to test pellet-clad-interaction, and
uranium metal/oxide interfaces can be used to investigate the behaviour of
stored metal wastes. UO2 surfaces can be used to investigate interactions with
aqueous environments, simulating ’leakers’ (split fuel pins during operation,
giving rise to high temperature water and steam exposure) and intermediate
longer term spent fuel storage scenarios. Doped UO2 systems could pave the way
for studies of modified fuel types to improve thermal conductivity, to improve
structural degradation during operation, or to improve end-of-life behaviour.

In addition, there is intense interest in developing alternative actinide
compounds to fuel reactors in a safer and more efficient way. These so-called
“advanced technology fuels (ATF)” include uranium silicides, nitrides,
metallic uranium alloys, thorium compounds, as well as other more exotic fuel
designs. A campaign of any new fuel composition in bulk form, and then
proceeding studies on radiation behaviour, thermal properties, or interaction
with coolant/storage media, are understandably, intensive operations. Using
thin films can shortcut many of the typical hurdles and provide a great deal
of supporting information in a much shorter time. It is possible for example,
to synthesise a new fuel design, simulate corrosion behaviour in long-term
storage to assess its feasibility, without ever having to embark on a full in-
reactor fuel performance review.

Hopefully, it should now be clear that thin films, and particularly epitaxial
films, have an important and irreplaceable role to play in advancing our
understanding of the actinides and their compounds both from a fundamental
aspect as well as those compounds of great practical importance to meeting our
ongoing energy needs in a decarbonising world. In this review we will cover
the growth methods and considerations for various uranium-based films, detail
many of the key experiments conducted to date and what they have taught us
already, before laying out a roadmap for the future, highlighting the
scientific areas we feel hold the most promise and would benefit most from a
thin film approach.

###  1.1 Early efforts (before $\sim$ 2000) on uranium-based films

Probably the first recorded use of thin films was by Steeb [19] who
demonstrated in 1961 that vapour deposition onto heated substrates such as MgO
produced an epitaxial film of UO2 with a thickness of $\sim 100\,\mbox{\AA}$
that could be further oxidized to U4O9. Further work on the structure of UO2+x
was done by electron microscopy at Stuttgart by Steeb et al.. This was
followed by work with electron microscopy by Navinsek [20] and Nasu et al.
[21] using different substrates - the best being identified as LiF and NaF.
They also observed fission tracks after the samples were irradiated in a
reactor. These efforts seem to have reduced once suitable bulk single-crystal
samples were produced and the quantitative study of the structure of UO2+x
using neutron diffraction was demonstrated. The use of neutrons allowed the
positions of the light oxygen atoms to be deduced, and became a major tool in
characterizing such systems [22].

For uranium metal, the first production of thin films was reported by T.
Gouder in 1993 [23] in Karlsruhe who deposited monolayers of uranium onto
various substrates to explore localisation effects in the uranium overlayer.
The first effort to produce epitaxial thin films was reported by Molodtsov et
al. in 1998 [24] in Dresden. The main objective of their work was to measure
resonant photoemission from the surface of uranium [24, 25], and scanning
tunneling spectroscopy [26]. As discussed later, in Section 3.2.3, a key
difficulty in the interpretation of these studies is that there was no X-ray
characterization of the samples, as it was not possible to cap the samples and
remove them from the preparation chamber. Nonetheless, interest in different
structural forms of uranium, as well as the surface layers, was stimulated by
these experiments. Theoretically Hao et al. [27], had earlier predicted that
the surface 5$f$ states in the bcc form would be more localized than in the
$\alpha$ (orthorhombic) form. Later Stojic et al.[28] predicted that such
localization, and ordered magnetism, would even occur at the surface of
$\alpha$-U, but no evidence for this has been found.

In a series of experiments, a group in first Darmstadt and then Mainz in
Germany grew epitaxial films of the heavy fermion compounds UPd2Al3 and
UNi2Al3, which have hexagonal symmetry, with molecular-beam techniques and
used heated (111) oriented LaAlO3 (LAO) substrates [29]. Their interest was in
transport measurements, as both systems show antiferromagnetic order with
superconductivity at lower temperatures, but with the material remaining
antiferromagnetic. Tunneling spectroscopy was used to demonstrate the crucial
role of the antiferromagnetic fluctuations in inducing the superconductivity
in UPd2Al3 [30], and measurements of the optical conductivity were also made
[31]. Rather similar measurements were made on epitaxial films of UNi2Al3 [32,
33]. At the same time one of the thin films of UPd2Al3 was used in a series of
synchrotron experiments to show how the coherence of the x-ray beam, together
with the large absorption at the uranium $M_{4}$ edge, allows information to
be obtained on the spatial position of the scattering volume [34]. This
technique is also discussed in Section 4.4.2, for more recent experiments on
UO2 epitaxial films.

In the 1990’s there was also considerable interest into whether memory systems
could be based on the magneto-optical Kerr effect (MOKE), and many different
systems were studied. This effect requires a bulk ferromagnetic signal.
Samples consisting of multilayers of amorphous UAs and elemental Co were
produced and the MOKE measurements showed that the uranium had a magnetic
moment at room temperature [35]. A more detailed experiment later took place
[36] to measure the XMCD signal at the uranium $M_{4}$ edge in a sample of the
form [UAs80/Co20]12. The XMCD data confirmed a moment of $\sim 0.80$
$\mu_{\mathrm{B}}$ per U atom at low temperature, but this rapidly declined at
higher temperatures. This study showed relatively poorly defined interfaces,
with diffusion between the layers.

##  2 An overview of the growth of uranium-based films

###  2.1 Deposition Techniques

Thin films, in a research sense, typically range from the Angstrom (Å) to the
micron scale, and involve the controlled deposition of the material of
interest onto a prepared surface of a chosen substrate. There is a range of
chemical or physical processes that one can employ, and many other reviews and
textbooks have dealt with this subject comprehensively [37, 38, 39]. Here we
will focus on just the subset of those techniques that have been used for
U-based deposition. It is worth noting that deposition of U has some specific
considerations, which depend on materials restrictions of a particular nation,
and are centred around the basic radioactivity of the starting material of
depleted U.

The choice of deposition method depends on the final application, whether a
fundamental study of basic physics, or an applied nuclear materials
investigation, this will influence the choice of material, metal, oxide,
intermetallic etc. and the required physical and chemical structures. This
Section will try and provide a strategic roadmap for new and existing research
groups who wish to utilise uranium deposition, by comparing and contrasting
the most successful examples in the literature.

Physical vapour deposition (PVD) is by far the most frequently adopted
technique in this field and can be generalised as the vaporisation of a
starting material that is then condensed onto a substrate [38]. PVD methods
are flexible, they can be used for metal, compound, and multilayer deposition
and one can control stoichiometry, phase and crystalline quality, to some
degree. The drawbacks are the need for large apparatus, high or ultra-high
vacuum, and bulk solid starting materials, and that the deposited material is
often highly energetic, so some additional thermalisation energy at the
substrate position is often required.

Figure 2: External (left) and internal (right) photographs of the actinide DC
magnetron sputtering system at Bristol.

Of the many PVD options, sputtering is the most prevalent and there are many
examples of research groups employing this approach [40, 41, 42, 43, 44]. DC
magnetron sputtering is commonly a UHV ($<10^{-9}$ mbar base vacuum)
technique, which consists of discs/ingots of starting material (in these
cases, either U metal or UO2 ceramic, typically many centimetres in diameter)
that are bombarded by a plasma of ionised noble gas, such as argon (pressures
typically $<10^{-2}$ mbar). Note that ceramic targets require the use of
pulsed-DC or RF sputtering techniques. Typical deposition rates vary from 0.1
to 2 Å/s. This has been used for U metal deposition [41], bilayer spintronics
[45], multilayers [46, 47] and intermetallics [48].

A sophisticated modification to this technique is the use of triode
sputtering, which employs a tungsten filament for electron emission to
stabilise the plasma at the source [40, 44]. This has the advantage of using
much smaller starting quantities, so more exotic starting materials are
accessible, it is a more efficient use of material; the deposition ‘racetrack’
produced by DC magnetron sputtering can yield efficiencies in the range of
only 5%. However, the lateral homogeneity at the substrate position is not as
good. Both of these methods can be adapted for reactive sputtering by feeding
a small partial pressure of reactive gas into the chamber (pressures typically
$<10^{-4}$ mbar), and this has been used to successfully grow oxides [49, 50,
51], nitrides [52, 53], oxynitrides [54] and hydrides of U [55, 44].

For the simplest polycrystalline films, one has control over magnetron power,
sputter gas pressure, target to substrate distance and substrate temperature.
For reactively grown compounds, we have the gas partial pressure as an added
lever, and for binary and even tertiary systems, the relative powers of the
co-depositing magnetrons become the crucial control mechanism for the
formation of specific phases. Even then, it may be difficult to achieve phase
pure materials. However, here one can utilise epitaxial matching to ‘lock-in’
desired phases, which prove elusive, even in bulk systems: the line compounds
of the U-Si system for example [48].

There are important differences with well-known literature examples of
epitaxial thin film systems, which have extremely close lattice matches and
result in high-quality crystals, often of single domains with mosaics below
0.05∘ (where the mosaic width is the rocking curve full width at half maximum
[FWHM]) [56]. Most U-based substrate matches are far from ideal, but are
required to stabilise epitaxy of the many compounds, phases and orientations
described in this review, hence they can have more complicated
crystallographic domains and have mosaics from 0.1 - 2∘ [41, 49, 53]. Also,
the multilayers are not of the same quality as some of the famous rare-earth
or tunnel junction heterostructures and superstructures [57, 58], although
these standards could be possible for UO2/ThO2 on CaF2, for example.

Pulsed laser deposition (PLD) is also a vacuum-based technique, but where the
vaporisation is typically performed by a high-power pulsed excimer laser. A
plasma ‘plume’ then carries energetic material towards a substrate, which can
then be thermalised to aid the crystalline growth of the film. It is also
possible to deposit reactively, to make oxides, nitrides etc. The most notable
PLD work at Los Alamos National Laboratory, used a KrF excimer laser
($\lambda=248$ nm, repetition rate 1-5 Hz) in varying partial pressures of
oxygen, employing substrate heating and rotation, to stabilise UO2, U3O8 and
UO3 oxides of uranium [59, 60]. Although there is little work in the
literature, using PLD for U-based deposition, it has many of the same
attributes as sputtering. It could be used for metal deposition, and for
multilayer and heterostructure synthesis. It typically requires bulk starting
materials; the deposition rates are similar and crystalline quality is
comparable. However, binary systems, such as silicides and carbides might be
more complicated, whereby sputtering or an evaporation technique could be more
suitable. Also, the energetics of the deposition process must be controlled
sufficiently to prevent the insertion of defects into the growing film.

Molecular beam epitaxy (MBE) is a UHV-based technique that uses Knudsen
effusion cells or direct e-beam heating (for the more refractory materials
such as U) onto small quantities of starting material to provide gradual
sublimation. The energetics of this process are lower, and the atoms have
longer mean free paths. This results in controlled deposition rates of
fractions of an Å per second and near layer-by-layer growth. MBE has an
advantage in the ability to monitor the growing surface in real time using
electron diffraction (RHEED – Reflection High Energy Electron Diffraction)
without the requirement for differential pumping. Gas sources can be added to
grow oxides, etc, often with cracker stages to produce confined beams of
highly-reactive atomic oxygen, or ozone. This is an expensive technique that
is focussed on the synthesis of high-quality epitaxial films and is more
commonly found in the manufacture of semiconductor devices and magnetic
memory; GMR and magnetic tunnel junctions, for example [58], although MBE has
also been used extensively to grow epitaxial rare-earth layers and
superlattices [57]. This technique can be used to deposit metals, oxides and
more complex ternary/quaternary intermetallic materials. Some of the early
work at Darmstadt, investigated the $\mbox{UPd}_{2}\mbox{Al}_{3}$and UNi2Al3
systems [30, 32] as epitaxial thin films. MBE requires a great deal of
investment and is not good for high throughput studies, i.e. for fast
exploration of phase diagrams or a wide range of compositions, but is useful
for particular studies where epitaxial quality is crucial.

Aside from PVD there are also chemical methods of deposition, which generally
avoid the requirement for bulk solid U-metal or U compounds, and although the
scope of these has been limited in terms of the range of U-based materials,
they feature prominently in the literature. Chemical vapour deposition (CVD)
is a vacuum deposition technique that encompasses an enormous range of
materials; it involves the reaction or decomposition of volatile precursors
onto a substrate wafer. In terms of U-based materials, groups in Cologne [61]
and UC Berkeley [62] have successfully used CVD and Magnetic field-assisted
CVD to make thin films of UO2, employing the decomposition of U(IV) amidate
and reduction of uranium hexakis-tert-butoxide, respectively. These methods
are yet to yield epitaxial films but could be useful for investigating UO2
grain morphologies in polycrystalline samples. The deposition rates are higher
than most PVD techniques and these techniques could be used for efficient
growth of $>\mu$m thick layers. Although studies have so far been focused on
uranium oxides, it may also be possible to modify the methods to deposit
metals, nitrides and carbides.

The sol-gel process has also been used to prepare UO2 films [63]. This is most
prevalent in metal oxide fabrication, such as TiO2, where a colloidal solution
or ‘sol’ is deposited onto a substrate, then becomes a two-phase wet gel, and
liquid is removed slowly to allow for densification and eventual film
synthesis. In the case of UO2, uranyl acetate in methanol and acetic acid were
heated together to form the precursor sol, which was then dropped onto
substrates that were spun to coat evenly. A final heating stage was used to
drive off remaining liquid and form a dense film. This method has some
drawbacks in terms of crystal structure control and synthesis of epitaxial
films. However, it is possible to synthesise $>\mu$m thick layers, and to dope
the uranium oxide with typical semiconductor dopant concentrations, which can
be very difficult to achieve in most PVD methods.

Polymer assisted deposition (PAD) is another chemical solution method, which
is common for metal oxide thin film growth [64]. Precursors are made using
metal ion-coordinated polymers. In this way, the polymer properties, such as
viscosity, can be modified to control the metal ion distribution to form
homogeneous films. U-based PAD has been successfully used by groups at Los
Alamos National Laboratory [64, 14] for more than a decade. They have reported
epitaxial synthesis of a number of U-oxides, UN2 and UC2 [14], as well as the
growth of PuO2 epitaxial films [65, 66]. In the case of uranium oxides, for
example, an aqueous solution of UO2(NO3)2 is added to a polymer, which is then
spin coated onto carefully chosen and prepared substrates with the desired
lattice matches. These are then annealed in the presence of oxygen at 1000 ∘C
to form epitaxial films. This has similar advantages to other chemical
processes, where film thicknesses are routinely larger than those made with
PVD, but clear progress has been made with the PAD technique for uranium, such
that epitaxial films of similar quality to those produced by the best PVD
methods are possible.

There are a number of options when embarking on a new research programme in
this field. In terms of overall strategy, one needs to consider some important
questions. Firstly, what sort of material(s); pure metal, compound,
heterostructure etc. as this will be the first limiting step in terms of
synthesis choice. What amount of starting material is required/is accessible?
Does the starting material need to be a compound first, or is it better to
make the composition during the deposition process? Often, even if the correct
composition is present in the starting material, it is necessary to adjust one
or more of the components during growth. For example, a UO2 target material
will require additional oxygen to reach stoichiometry. How many samples are
required? Some techniques are more suited to high throughput than others.

For metallic systems, especially those containing U, oxidation of the surface
(which can be the entire depth of the film in some cases) is a major problem.
This means that a capping layer is necessary – typical materials are
aluminium, chromium, tungsten, etc., however, the authors recommend niobium,
as this layer develops a thin passivating oxide layer of approximately 20 Å.
Capping layers can sometimes be necessary for oxides also, as these will
become hyper-stoichiometric over time and this can affect physical properties.
In some cases, buffer layers are used to provide a chemical barrier and to
mediate the lattice mismatch between substrate and film. The requirement of
extra layers then has consequences for the choice of synthesis method.

Typical research projects/programmes involve preparation of starting material
and substrate surfaces; this could be chemical cleaning, sonication, Ar plasma
cleaning, or a combination of these steps. Compositions of desired materials
are usually tested first on substrate standards, such as glass, or silicon,
and where polycrystalline samples are enough then this can provide the basis
for the remainder of the synthesis. However, where epitaxy is required, then
the strategy is more refined, as careful substrate matching is required,
considerations of temperature, thermal expansions of the different materials
and possible interfacial mixing. It is not always easy to predict which
substrate to use.

###  2.2 Characterisation Techniques

For most chemical-based synthesis routes the characterisation takes place ex-
situ, once the films have been made. However, for all of the PVD techniques
described, some in-situ characterisation is routinely used during the
deposition process. For more detailed investigations of the physical structure
of the films it is more likely that an ex-situ measurement will be employed,
and this will depend on the length-scale of interest and whether lateral or
longitudinal information is important. Here, we present the most common
techniques in more detail with particular recommendations for U-based
materials.

The deposition rate is in general the primary characterisation parameter for
thin films as this determines the layer thicknesses, and there are many
instruments and techniques that are found in the literature. The method of
choice will depend on whether in- or ex-situ measurements are needed, the
thickness regime, and the required precision; of course, it is often the case
that a combination of methods is preferred. For thickness determination within
deposition chambers/vessels, quartz crystal microbalances (QCM) are commonly
employed, which exploit the Sauerbrey equation [67], relating the frequency of
oscillation of a piezoelectric crystal (quartz for example) with the mass
deposited, i.e. as the thickness increases and more mass is deposited, the
frequency decreases. The resolution is typically 1 Hz for resonant frequencies
in the MHz range, which means that this technique has approximately, monolayer
sensitivity. One major advantage is that this can be used during the growth
process. However, this means that it has to be mounted in the deposition
system itself, which can be complicated, and for a precise measurement it must
be at the same position as the substrate, which can be spatially restrictive.
Less common methods of thickness determination during growth include laser
interferometry [68], RHEED, utilising the oscillating intensity of the
specular diffraction spot [69, 70], and ellipsometry [71].

Figure 3: The left hand image shows a RHEED pattern from a (001) LAO
substrate, [110] azimuth. Kikuchi streaks are visible, and diffraction maxima
on arcs where the 2D reciprocal lattice rods intersect the Ewald sphere,
indicating a smooth surface. The right hand images are from a 1900 Å UO2 film
grown on LAO. The surface is rougher and 3D spot patterns are observed, formed
by transmission through 3D islands. The [100] and [110] UO2 azimuths are
shown, respectively. RHEED images have been adapted from Ref. [72]

Both low energy electron diffraction (LEED) [73] and reflection high energy
electron diffraction (RHEED) [74, 75] are used to identify the crystal
structure. It is possible to distinguish between polycrystalline, highly
textured, and single crystal systems, and in the most advanced cases, even
monitor strain as a function of growth. For MBE, electron diffraction can be
acquired during growth, however, for PLD and sputtering, typically the
synthesis process has to be paused to view the diffraction image, unless a
double differentially pumped electron beam path is employed [69, 70]. Fig. 3
shows typical RHEED images from a single crystal lanthanum aluminate
($\mbox{LaAl}\mbox{O}_{3}$ or LAO) substrate and UO2 epitaxial film [72].

X-ray reflectivity (XRR) is used to probe the electron density profile of the
film [76], which gives information about the sample morphology; the thickness,
the roughness at each surface/interface [77, 78], and the value of the
electron density itself, which can be used to infer the composition of the
film. The geometry is typically in a specular or longitudinal mode, where the
incident and reflected angles are equal (incorporating any offset angle due to
sample surface misalignment). Therefore, the wave-vector momentum transfer
(usually written Q or qz) is along the surface normal, with no sensitivity to
lateral features. There are many freely available resources for modelling this
reflectivity spectrum [79, 80, 81], which use Parratt’s recursive method [82],
and then employing a range of fitting algorithms in order to explore the
parameter space and find the best global minimum.

Fig. 4 shows two reflectivity curves for thin Si and U films, with data shown
as the open black circles and the fits as solid orange and blue lines,
respectively [83]. The oscillations are known as Kiessig fringes, which depend
on the film thickness [76]; an approximate relationship between film thickness
and fringe separation is highlighted in the figure. Models of the reflected
intensity are often sensitive to changes in the layer thickness on the order
of a fraction of an Å. However, one should note that for films greater than
$1000\,$Å thick, these fringes become very small and are eventually too
difficult to resolve. For a perfectly smooth film the intensity decays as a
function of $1/\mathrm{Q}^{4}$, surface and interface roughness cause this
intensity to decay more rapidly [76]. The roughness is usually approximated as
the root mean square of the thickness variation of a layer and appears as a
Gaussian spread of electron density at an interface [77]. This means that
interfacial roughness and interdiffusion manifest in the same way for specular
XRR.

Figure 4: Example reflectivity data for typical U and Si thin layers, taken
from Ref. [83]. The experimental data are represented by the open black
circles and the fitted models of the reflected intensity are shown as blue and
orange solid lines, respectively. Total thickness oscillations, Kiessig
fringes, are visible for both samples. The U film also displays a longer
wavelength periodicity, more visible at low Q, due to a thin oxide layer. The
electron density dependence of the critical angles, is also obvious here as
the U film $\theta_{C}$ is at a much higher Q than the silicon.

The roughness also affects the amplitude of the fringes, as does the electron
density differences between the layers. The overall electron density of the
topmost layers in a sample affects the position of the critical angle,
$\theta_{C}$ (shown in Fig. 4), where total external reflection ends, and the
x-rays first start to penetrate the film. There is often a reduction in
intensity associated with the smallest angles, below $\theta_{C}$, and this is
known as a footprint effect (also visible in Fig. 4, where the incident beam
size becomes so large that a sizeable fraction of the photons are no longer
incident on the surface of the sample.

XRR can also be used for more complex systems and in more complex geometries
in order to extract even more information. For instance, bilayers, trilayers
etc. will exhibit multiple periodic intensity oscillations that result from
coherent reflections from each of the interfaces [76]; see the U metal film in
Fig. 4, where a thinner oxide layer appears as a much longer wavelength
oscillation, more prominent at low Q values. Heterostructures and multilayers
with repeat units will result in the appearance of sequential peaks of
intensity, known as Bragg peaks, whose position will depend on the bilayer
thickness and intensity on the number of repeat units [84]. The technique
itself can be extended further by allowing the incident and reflected angles
to vary at various positions in Q (along the specular ridge) to measure
intensity as a function of Qx. Modelling this intensity is more difficult and
one has to employ the distorted wave born approximation, which maps height-
height correlations to generate lateral coherence lengths, and a jaggedness
factor, which describe the distribution of height variation and the smoothness
of this variation as a function of lateral dimension in the sample [85, 76].

In summary, XRR is an extremely powerful and versatile technique, which non-
destructively investigates the physical structure of a thin film sample. Due
to the often sizeable parameter space, and extreme variation in intensity,
fitted models may look convincing, exhibiting excellent figures of merit, but
they can often be misleading, so it is wise to synthesise a series of samples,
where only one or two growth parameters are varied systematically.

X-ray diffraction (XRD) has been used to study the crystalline nature of
materials for over a hundred years and has some very special considerations
when used to study thin films and heterostructures [86, 87]. Thin films
inherently have a small sample volume, and the majority of the photons will
pass directly through the sample without scattering. However, the atomic form
factor and therefore the scattering amplitude vary as a function of atomic
number (Z), which means that the observed intensity varies as a function of Z2
. This is an important advantage when considering U-based thin films, because
U is such a strong scatterer that even films of just a few tens of Å have
measurable intensities on a standard laboratory x-ray source (Cu
K${}_{\alpha}\sim 1.54\,$Å, for example).

Typical measurements to determine phase and structure, are in a similar
specular/longitudinal geometry to that described for XRR, however they use
2$\theta$ angles in ranges from 15 - 140∘. The positions of the peaks are the
first indication of the crystal structure of the materials in the film,
although the spectra can often be dominated by intensity from the substrate.
At the two crystallographic extremes; glass gives an amorphous signature,
which results in a long, damped periodic intensity background, which is
relatively weak overall, but persists at all angles, whereas single crystal
substrates only exhibit extremely strong intensity peaks at specific angles,
relating to the $d$-spacing along the unique growth axis out of the plane.

Polycrystalline films are grown when there is no obvious lattice match between
substrate and film, or when no thermalisation has been used to aid epitaxy.
Usually, all of the reflections that one would expect for a powder are
visible, and it is even possible to use the Scherrer equation [88] to
determine grain size, which analyses the peak widths in the same way as for
bulk samples; this works well if the grains are smaller than the film
thickness, otherwise it is just a measure of the film thickness itself and
does not indicate a lateral grain dimension (could be plate-like in shape for
example). Normally, for bulk materials, it is standard practice to measure in
a longitudinal geometry, and that is true in the first instance for thin-film
measurements, however, if a researcher wants to improve their measured signal
to be more surface sensitive, this can be achieved by fixing a small incident
angle (few degrees in 2$\theta$) to fully illuminate the sample, and then scan
the detector angle to achieve the desired 2$\theta$ range.

Figure 5: Typical XRD spectrum for a $500\,$Å UO2 epitaxial thin film,
deposited on yttria stabilised zirconia (YSZ), taken from [89]. The data are
represented by the open black circles and a fitted model of the data is shown
as a solid red line. The principal (002) reflection from the [001]-oriented
YSZ substrate is very intense and has a mixture of Gaussian and Lorentzian
contributions, and is well-modelled by a pseudo-Voigt function. The main (002)
reflection of the UO2 film is predominantly Gaussian, which is typical for
thin films and there are Laue fringes, relating to the film thickness, visible
either side.

In many cases, the polycrystalline film will have a preferred orientation, or
texture [88], which manifests as a preferential intensity for particular
reflections, which deviates strongly from the intensities expected from a
theoretically ideal powder pattern [88]. For high symmetry structures, this is
typically with the closest packed plane flat, facing upwards along the surface
normal, so the [110] and [111] orientations for bcc and fcc crystals,
respectively, for example. Due to the inherent energetics of most deposition
processes and thermalisation, impurities and other structural defects, most
polycrystalline samples will exhibit some form(s) of microstrain or residual
stress. Microstrain depends on grain orientation, as local lattice spacing
variations may vary from grain to grain and can be analysed using the
Williamson-Hall analysis [90, 91, 88], relating peak position, width and grain
size to the microstrain. Residual stress results in an average change in
lattice spacing, using $\sin^{2}\psi$ analysis [88, 92, 93], where lattice
parameters are calculated from the $d$-spacing of a selected family of planes,
measured as a function of sample inclination.

Figure 6: The left hand panel shows typical rocking curve data and models for
the calcium fluoride substrate and the epitaxial USi3 film, in purple and
blue, respectively. The right hand panel shows an example $\phi$-scan, which
is an azimuthal rotation about the specular direction at off-specular {115}
and {113} reflections of CaF2 and USi3, respectively. Taken from [83]

Epitaxial films have a unique axis along a crystallographic surface normal and
they are deposited onto single crystal substrates, which also exhibit a unique
axis, see Fig. 5. For some sample systems these axes might not be coincident,
and an angular offset can be measured between the rocking curve peaks of the
film and substrate. Usually, there will already be a predicted lattice match
and the first high-angle scan is a survey scan with wide open diffracted beam
slits to allow maximum intensity with low resolution. In this way, it is
usually possible to observe all of the reflections aligned close to the
surface normal of the film. Remember that the crystallographic directions will
be different to the flat surface, since the polished substrate surface will
not be perfectly aligned along a crystal plane. At this point, the slits can
be narrowed to improve resolution and a more detailed measurement of the
position of film and substrate peaks allow the determination of respective
lattice parameters. Fig. 5 also highlights another inherent feature in a high-
angle diffraction spectrum, which are unique to thin films, the phenomenon of
Laue fringes. These are due to a beating frequency in the diffracted signal
due to added interference of x-rays reflected at interfacial boundaries (not
too dissimilar from the Kiessig fringes in XRR). The main peak for thin layers
is also a lot more Gaussian in shape as the number of monolayers decreases. In
fact, in this regime, it becomes possible to model the whole diffraction
spectrum using discrete numbers of lattice planes to generate the observed
peak widths [87].

Once the orientation has been determined, a rocking-curve measurement [56,
88], which varies incident and reflected beam angles at a fixed $2\theta$, is
used to align at the maximum of the peak. The full width at half maximum
(FWHM) of this rocking curve is also a standard measure of crystal quality
(for films and bulk crystals) and is known as the mosaicity. Literature values
for single crystals of U-based studies can be from 0.1 - 2∘, see the left hand
panel of Fig. 6. For many epitaxial thin film systems an unusual phenomenon is
observed in the rocking curves [56]. They consist of two peak shapes, a sharp
component, together with a wider contribution. There are a number of theories
as to why this is present, and a recent paper by Wildes [94] gives a good
discussion of such profiles. It is usual to take the wider component as a
reflection of the bulk mosaicity of the deposited film.

To confirm that a film is indeed epitaxial and then to relate the rotational
orientation of the substrate to the film, it is necessary to measure off-
specular reflections and probe their azimuthal dependence, by rotation around
the surface normal, a so-called phi scan [56]. Fig. 6 shows a phi-scan for a
[001]-oriented USi3 film grown on a [001]-oriented CaF2 single crystal
substrate [83]. The first point to note is that if this were simply textured
then there would be no discrete azimuthal dependence. Second, the number of
reflections in the 360∘ range indicates how many domains are present, as there
may be more than one in-plane direction that results in a lattice match; in
fact, this is very common for cubic materials. Finally, the azimuthal angular
offset between substrate and film determines the in-plane lattice match
relationship.

One could go further and measure truly in-plane diffraction, where the whole
system is set at a slight tilt angle and the detector is moved out of the
scattering plane, until the Q-vector is pointing almost along the surface of
the sample, this is known as grazing incidence x-ray diffraction (GIXRD) [95,
96]. In this regime, the surface is severely truncated, and the diffraction
spots become elongated and more rod-like, since the Fourier transform of a
Dirac $\delta$ function (analogous to lattice for a single layer) is a
constant (hence, rod of scattering). These are known as crystal truncation
rods and can be modelled to give detailed in-plane information [95, 96].

To summarise, there are many ways that XRD can be used to give a good
understanding of the crystal structure(s), lattice parameters, lattice
matches, stresses and strains etc. However, most x-ray spot sizes at the
sample position will be several mm2, which means that these techniques are not
suited for local features, individual grain information, stresses, strains
etc., and a microscopy method may be better.

Scanning electron microscopy (SEM) uses a focused beam of electrons to raster
across the sample, which produces secondary and backscattered electrons, which
can be used for imaging or diffraction imaging, achieving magnifications of
approx. $\times$250,000 [97]. For epitaxial films, which are often smooth (rms
roughness of $<$10’s Å), the SEM image can look featureless, although it is
possible to magnify the edge of a sample to observe the film/substrate
interface and estimate the film thickness. This only works for thick films
$>$1000 Å and the resolution is no better than 100 Å. However, there are other
operational modes of an SEM, which can be used to give higher resolution
images, compositional and crystallographic information. If the SEM system has
an additional focussed ion beam column, then it is possible to cut thin cross-
sectional foils through the surface and across the film/substrate interface
and image them in transmission mode [98].

Compositional information can be gathered using energy-dispersive x-ray
spectroscopy (EDX), where the electron beam stimulates the emission of
characteristic x-rays, which relate to specific electron shell transitions in
given elements [97]. This is not highly accurate quantitatively, but it can
give a good approximation of the elemental composition. Moreover, the rastered
electron beam gives lateral compositional data, so that a map of elemental
composition can be constructed, which is especially useful for heterogeneous
systems.

Figure 7: EBSD data collected for a 1000 Å UO2 film deposited on an annealed
polycrystalline YSZ substrate, showing an average grain diameter of about 40
$\mu$m. Each colour represents a particular grain orientation and these
exhibit a random distribution, expected from a polycrystalline material. Taken
from [99].

Finally, electron back-scatter diffraction (EBSD) can be used to investigate
the structure, phase and crystal orientation [100]. For bulk materials, the
surface has to be extremely smooth, but for thin film synthesis the resulting
surface is often far better, in terms of rms roughness, than any mechanically
prepared material, as long as the lateral features are larger than the lateral
resolution ($\sim$100 Å). For single crystals, EBSD would just give a single
orientation and a map of the surface would be uninteresting. However, for
polycrystalline samples it is possible to generate a map of the grain
structure and individual orientations, see Fig. 7 for an example UO2
polycrystalline film, deposited on polycrystalline YSZ [99].

It should be noted that SEM is largely non-destructive for low resolution
imaging and ESBD. However, the sample can undergo significant damage if images
are taken at high resolution, or it may need to be coated if non-conducting.
Also, the probing depth of the EDX process is on the order of microns, which
means that signals are often dominated by the substrate materials.

High-resolution Transmission electron microscopy (HRTEM) images a thin cross-
section of a sample, using a magnetically focussed, highly collimated, high
energy ($\sim 200\,\mbox{keV}$) electron beam in a transmission geometry
[101]. This is typically a UHV set up, although specialist systems are able to
deliver gas environments [102]. Typical magnifications can reach
$\times$1,000,000 so can image in the $10\,\mbox{\AA}$ regime, and can operate
in a variety of different acquisition modes where the image contrast can be
simply due to thickness variation, atomic number or crystallographic
orientation. It is also possible to generate diffraction images and analyse
the crystal structure at a very local level. Fig. 8 shows an example HRTEM
high resolution image and diffraction image from an epitaxial UO2 film
deposited on lanthanum aluminate (LAO) as an example [49].

Figure 8: Left hand panel is a high resolution TEM image of a UO2/LAO cross-
section taken from Bao et al. [49] and the right hand panel shows diffraction
images for, A, the [001]-oriented UO2 film and B, the [001]-oriented LAO.

Electron energy loss spectroscopy (EELS) [103], which uses an electron
spectrometer to measure the energy lost by electrons due to inelastic
scattering to probe core shell states, much in the same way as EDX, can be
used to make high resolution elemental composition maps. Note that HRTEM
usually relies on initial preparation via SEM and focussed ion beam milling. A
cross-sectional lift out is prepared by careful etching and final platinum
adhesion to a movable needle to remove the sample, before attaching it to a
standard TEM grid. Specific consideration has to be made for U-based
materials, as the electron density of U poses a particularly difficult
challenge for the transmission of an electron beam. This means that foils of
$<$1000 Å are preferable, and this requires added preparation time and care,
compared with more standard materials.

Ion beam analysis (IBA) [104] is a less common, but complementary field of
materials characterisation that, particularly for the case of thin films, uses
Rutherford backscattering spectroscopy (RBS) to map the depth-dependent
profile of sub-micron layers with element specificity. This technique
typically uses the energy profile of backscattered helium ions to build a
model of the component species in a thin layer, where the depth resolution is
limited by the energy resolution of the detector. This can be especially
effective for heavy ions, such as uranium [105, 106, 107], where it is
possible to operate in isotope detection limits of the ppm. The most advanced
high resolution systems can operate with a depth resolution of less than 20 Å
[108], which although not at the same scale as XRR, does provide simultaneous
element-specific information.

X-ray photoemission spectroscopy (XPS) is based on the photoelectric effect
[109, 110], which is the emission of electrons from a given material, due to
an incident light source. Most lab-based facilities use an Al or Mg K-edge
x-ray source with energies of 1486.6 and 1253.6 eV, respectively, which can be
monochromated to improve the final energy resolution of the resulting spectra
[109]. The electrons are emitted from core shells within the different atomic
species in the film and their energies are then determined by an electrostatic
hemispherical analyser. This means that the system has to be in UHV
conditions, and in fact the main chambers are often in the 10-11 mbar range.
Fig. 9 shows a schematic of a typical XPS [83].

Figure 9: Left hand panel shows a schematic of a typical XPS system,
indicating the photelectron path within the concentric hemispherical analyser.
The right hand panel gives a depiction of the electron configuration filling
order for sub-shells for a uranium atom. Here, the principle quantum number,
orbital angular momentum and total angular momentum are labelled, n, l and j,
respectively. Taken from [83]

At a basic level, the spectra of the core levels of the constituent elements
give a fingerprint of the composition, similar in information to EDX, however,
the probing depth here is less than 100 Å, which means that the signal will
result from interactions with the film in all but the thinnest layers. An
analysis of the integrated areas of respective core-level peaks can also give
quantitative compositional information [111]. Fig. 9 shows the expected core
levels for a U atom [83].

One of the most powerful uses of XPS is in characterising the binding state of
the constituent elements, where specific chemical shifts in energy appear due
to the existence of particular valence states [109]. In some compounds the
core levels of the metallic ions have features that are even more sensitive to
the local coordination chemistry. For example, UO2 has ‘so-called’ shake up
satellites that are sensitive to the oxygen stoichiometry, and with careful
measurement it is possible to measure excess oxygen to better than about 3%,
i.e. x to $\pm 0.06$ in UO${}_{2\pm\,x}$ [112, 113].

XPS is clearly a very powerful non-destructive technique for chemical
analysis, but there are possible modifications from the standard that make it
even more useful. Since many of the deposition systems are in UHV conditions,
it is also possible to combine these facilities with XPS. Focussed x-ray
sources, or focussed analysers, can achieve lateral resolutions of 10’s
$\mu$m, so that it is possible to laterally map the chemical states of
samples. Argon plasma sources can be used to gently etch through the sample,
which gives depth profiling information, which could be crucial for
understanding complex interfaces. Finally, the use of a much lower-energy
ultraviolet source (UPS) means that the valence states are accessible with
good resolution, which in conjunction with the XPS spectra, creates a more
complete chemical picture [109].

##  3 Uranium Metal Phases and Alloys

###  3.1 Introduction

Figure 10: Phase diagram of uranium as a function of pressure and temperature,
with structure of $\alpha\mbox{-U}$ shown expanded outside the conventional
unit cell to highlight the corrugated atomic planes along the [010] direction.
For every atom there are four neighbouring atoms in the corrugated (010)
sheets, two at $d_{1}=2.754\,\mbox{\AA{} (AB)}$ and two at
$d_{2}=2.854\,\mbox{\AA{} (AC)}$, which are shown by thick and thin grey
bonds, respectively. A further eight neighbouring atoms at distances
$>3.2\,\mbox{\AA{}}$ are found in the adjacent corrugated planes, examples of
the two types (AD) and (AE) are shown as black dashed lines. Adapted from
[114]

Before discussing the progress in created thin films of U metal and related
alloys, we first discuss some of the fundamentals of bulk U metal, including a
brief discussion of its phase diagram and some important electronic properties
that have been studied over many decades. Much of this discussion will be
focused on the basic physics, to frame the importance and power of studies
with thin films, but it is important not to forget the use of metallic U in
the earliest days of nuclear energy production, as well as of course weapons.

The phase diagram of U as a function of pressure and temperature is shown in
Fig. 10. Similar to many metals U adopts a high-temperature body-centred
cubic, bcc, $\gamma$-phase stable from the melting point of $1132$ ∘C down to
$772$ ∘C. At pressures below approximately $3.5\;\mbox{GPa}$ this temperature
marks the transition to a tetragonal $\beta$ phase, stable between 772 ∘C and
662 ∘C, with a complex atomic structure and 30 atoms in the unit cell. Efforts
were made to solve the $\beta\mbox{-U}$ structure as early as the 1950’s
[115], but it was only in 1988 with the advent of modern neutron powder
diffraction, together with Rietveld analysis, that the the structure was
eventually solved [116].

Significantly before this is in 1937 Jacob and Warren [117] solved the atomic
structure of the room-temperature stable, orthorhombic $\alpha$-phase which is
adopted below 662 ∘C at atmospheric pressure or, above approximately
$3.5\;\mbox{GPa}$, directly from the bcc phase at around $800$ ∘C. The
$\alpha$-phase allotrope, defined within the Cmcm space group and shown in
Fig. 10 is a highly open, consisting of series of corrugated atomic chains
nested along the $[010]$ direction. Such a structure produces a series of
highly anisotropic interatomic distances reflecting the complex role of the
5$f$ electrons in stabilizing the structure [118, 119, 120]. At ambient
pressure this low symmetry, highly open orthorhombic character is unique among
the elements, however it has been found as a common stable structure for the
high-pressure forms of the light rare-earths and is also closely related to
the high-pressure forms of the heavier actinides [121]. For a more in-depth
description of the structures of the three bulk allotropes the reader is
directed to Ref. [115] and for a detailed account of extensive structural
studies conducted on bulk $\alpha\mbox{-U}$ in the 1980’s and early 1990’s the
reader is directed to the review by Lander et al. [114].

Given the complexity of the U phase diagram and the unusual structures
contained within it, it is perhaps unsurprising that large single crystals of
U metal are difficult to prepare. Indeed, until the production of single
crystals of $\alpha$-U from a molten salt process in the late 1990’s [122],
only the few crystals made by E. Fisher in the 1950’s (known to contain
measurable impurities of Si and Fe) were available for use [114]. Obtaining
single crystals of the two high temperature allotropes has proved even more
elusive. Long lived metastable single crystals of the $\beta$ phase were first
synthesised by A. N Holden in 1951 by quenching a $1.36$ at.%U-Cr alloy from
the $\beta$-phase in water [123]. However these crystals were not truly single
phase and contained a likely Cr rich precipitate. Later, the same authors
produced phase pure single crystals by reducing the Cr content to 0.5 at.%
however such crystals transform to the $\alpha$-phase within a few hours at
room temperature [123]. To date, bulk single crystals of the bcc $\gamma$
phase have never successfully been produced despite significant efforts.

Despite the challenges in crystal growth, subsequent decades saw great numbers
of detailed and varied studies into the structural, electronic and
thermodynamic properties of the bulk U allotropes. For the room temperature
$\alpha$-phase, many of the experiments focused on pursuing the various
possible long range ordered states that could exist at low temperature. The
major findings of most studies these studies (pre-1994) are encompassed by the
review article in Ref. [114]. Perhaps the greatest experimental effort was
expended on probing the unexpected and mysterious phase transition near 43 K,
first observed in temperature dependant measurements of the elastics constants
in 1961 [114]. Key to understanding general properties of $\alpha\mbox{-U}$,
as well as the CDW state it hosts, was a pioneering experiment in 1979 [124]
to measure the phonon dispersion in a bulk single crystal, and a subsequent
experiment by Smith et al. in 1980 [125] that extended the phonon studies to
low temperature, and showed significant phonon softening near
$q_{\mathrm{CDW}}$ highlighting the transition as soft-phonon driven. Robust
ab-initio calculations that accurately modelled the observed dispersion were
not developed until 2008 by J. Bouchet [126].

At even lower temperatures a superconducting state was established,
$T_{\mathrm{c}}=0.7\;\mbox{K}$ [127] but there is still much ongoing debate as
to the exact nature of the superconductivity, bulk or filamentary [128, 129],
and the importance of sample purity considerations. The coexistence of these
two ground states is unique amongst the elements, and reminiscent of more
exotic highly correlated electron systems such as the high $T_{\mathrm{c}}$
cuprate family of superconductors. However, the exact relationship between the
ordered states is still unclear. It has been determined that pressure
initially suppresses the CDW and enhances the superconductivity, increasing
$T_{\mathrm{c}}$ to over 2 K, however further increase in pressure to
$10\;\mbox{GPa}$ suppressed both phenomena [114, 130]. Additionally, magnetism
has long been discussed and sought in connection with U metal, early neutron
work on thermal expansion even claimed to have found extra peaks [131].
Although later it was established that these came from multiple scattering
processes and to date no magnetism has been found in bulk $\alpha$-U [114].

Considering the relative difficulty of obtaining bulk single crystals for the
various allotropes and the obvious applications for epitaxial strain to modify
the different types of long range order present, it becomes clear that the
synthesis of U thin films, especially epitaxial films, is of great interest.
The following Sections will lay out the general synthesis considerations for
the key metallic systems before detailing a number of important case studies
where metallic U thin films have been employed to provide scientific insight
into this fascinating element and its many allotropes that would not have been
possible by relying on bulk synthesis routes alone. In terms of
superconductivity, there is ongoing work on elemental, and alloy thin films,
but the situation at the time of writing is currently unclear in thin films.

###  3.2 Production of metallic uranium films

The number of laboratories world wide that are sufficiently equipped to
deposit metallic U films is understandably small. Aside from the issues of
sourcing appropriate purity starting material, the base vacuum within any
potential deposition system must be sufficiently low to avoid the immediate
formation of $\mbox{UO}_{2}$ or other compounds. As a result the ratio of
metal film producing laboratories to oxide or other U compound producing
laboratories is also small. As discussed in Section 1.1 the first U deposition
for the explicit purpose of forming metallic films was conducted in the 1990’s
by T. Gouder in an effort to induce localisation [23] after which there was
also early work on heavy fermion films [29], and attempts to stabilise the
metastable hcp form of U in thin layers [24, 25]. Following these early
examples, we will now discuss a number of the key U thin film systems that
have been developed to date and are detailed in Table 1. The majority these
systems were first fabricated using the PVD method DC magnetron sputtering,
see Section 2.1 and the optimum growth conditions discovered by an iterative
process using many of the characterisation methods discussed in Section 2.2.

Figure 11: X-ray reflectivities (Cu $\mathrm{K}_{\alpha}$ source) for (a)
[U(32 Å)/Fe(27 Å)]30, and (b) [U(89 Å)/Gd(20 Å)]20. Data are shown as black
open circles, and the red lines are best fit models. Adapted from Ref. [46].

####  3.2.1 Uranium containing multilayers

The deposition of polycrystalline or amorphous U layers is substantially
easier than the epitaxial systems described below. However, there has been
significant and continued interest in such systems in the context of
multilayers and bilayers to investigate the possibility of induced behaviour
between U and other metals as a direct result of their proximity across the
interface. The first work on uranium containing multilayers was conducted by
Beesley et al. [132, 133], which spurred significant further work detailed in
Sections 3.3.1 and 3.3.2. The precise growth details for each system can be
found in the references included in table 1 however in general the U layers
are grown by sputtering techniques with minimal substrate cleaning and little
or no heating during growth. As induced effects typically exist over
relatively short length scales, a key parameter in such films is the
interfacial quality between the layers, and this can vary dramatically from
system to system. Note that similar growth considerations are valid for both
multilayer and bilayer systems; however, significantly more characterisation
has been performed on the multilayer systems and thus they will form the focus
of the following discussion. As described in Section 2.2, XRR is an invaluable
tool for characterising thin film systems. Typical XRR curves from U/Fe and
U/Gd multilayers are shown in Fig. 11.

Analysis of these and related data demonstrate that the interfacial properties
of the U/Fe, U/Co, and U/Ni multilayers are quite different from that of the
U/Gd multilayers. In the case of the transition metal systems, the interfaces
are not as chemically sharp, and there is an interdiffused region of thickness
$\sim$ 15 Å at each interface. In contrast, for the U/Gd multilayers, the
interfaces are much sharper, and there appears no significant interdiffusion
between Gd and U [46]. This conclusion is supported by transmission electron
microscopy studies - a typical image of one such film is shown in Fig. 12 .
Here the layers are well defined with relatively low roughness, and the Gd
layers are strongly crystalline, whereas the U layers have small nano-
crystallites or are amorphous. In many cases the brighter Gd crystallites
extend across the layers suggesting sizes of as much as 50 Å in the vertical
growth direction.

Figure 12: TEM images of a [U(50 Å)/Gd(50 Å)]20 multilayer. The mass contrast
between the U and Gd can be easily distinguished; the lighter Gd scatter the
electron beam less than the U atoms and thus appear brighter. The layers are
labelled in the high-resolution image (b); (a) shows a lower resolution. Taken
from Ref. [134].

Regarding the growth directions of the layers, in all cases the structure of
the 3$d$ transition element follows the expected preferred orientation, i.e.
bcc for Fe, hcp for Co, and fcc for Ni. The U is in a poorly defined $\alpha$
structure. Surprisingly however, in the U/Gd multilayers, although the Gd
forms in the expected hcp phase, the U forms also in a hcp symmetry, with a
unique $c$ (hexagonal axis) growth axis. There is no ordering that could be
found within the hexagonal planes, so these consist of a number of random
domains of the hexagonal basal plane [46]. The $c_{\mathrm{Gd}}$ is 5.84 Å,
which is close to the bulk value of 5.78 Å, and the $c_{\mathrm{U}}$ = 5.60 Å.
If we assume that the atomic volume is the same as that of the of the
$\alpha$-U form, then the $a_{\mathrm{U}}\sim 2.91$ Å, giving a
$c_{\mathrm{U}}/a_{\mathrm{U}}$ ratio of 1.92, which is much larger than that
expected for a hard sphere model 1.633. The only element close to this is Cd
with a ratio $c/a$ = 1.89. We return to the topic of hcp-U in Section 3.2.3.

####  3.2.2 Achieving single crystal films of $\alpha\mbox{-U}$

The key breakthroughs in the synthesis of epitaxial metallic U layers came
with the realisation of the importance of unreactive epitaxial buffer layers
deposited prior to the U layer, and the addition of substrate heating to allow
sufficient mobility of deposited atoms. Refractory metal seed/buffer layers
are routinely implemented in the growth of rare-earth systems to prevent
interactions with the substrate and/or to bridge a large mismatch in lattice
parameters between the substrate and overlayer in order to achieve successful
adhesion of a crystalline thin film. As many substrates contains oxygen, which
reacts readily with U metal, the addition of these (nominally) non-reactive
buffers opens up a new region of phase space. It is fortunate that the
refractory metals that prove good chemical buffers also demonstrate excellent
epitaxial matches with U. As we will see below, however, the details of these
epitaxial matches are not obvious in many cases. To date high quality
epitaxial layers of, orthorhombic and hexagonal U metal have been achieved
[41, 135, 136], as well as pseudo-bcc U-Mo alloys [137, 138]. As with the
multilayer systems described above each system will be briefly explored here
and further details for each system can be found in the references provided in
Table 1.

Utilising this approach of high temperature growth and epitaxial refractory
metal buffers the epitaxial growth of $\alpha\mbox{-U}$ was first achieved in
2008 via deposition onto thin, single crystal buffer layers of either Nb(110)
or W(110) at 600 ∘C [41]. The optimal growth temperature was later refined to
450 ∘C [135]. In these cases, the substrates were single crystals of sapphire,
Al2O3, epi-polished parallel to the (11.0) plane. This excellent, if non-
intuitive, epitaxial match had already been identified [139]. The resulting
films were capped with a thin layer of one of the two refractory metals to
protect the U layer from atmospheric degradation. It was found that Nb serves
as a better capping layer as the passivating oxide, Nb2O5, is limited to
approximately 20-30 Å and can be stable for many years. Despite the use of
similar $bcc$ buffer layers ($a_{\textrm{Nb}}=3.300$ Å and
$a_{\textrm{W}}=3.165$ Å), there are significant differences in the
orientations, domain structures and strains induced in the U overlayers for
these two buffer materials.

(a)

(b)

Figure 13: Two-dimensional projection illustrating the growth of the
$\alpha$-U(110) epitaxial film on (a) Nb(110) (a=3.300 Å) with Nb atoms shown
in dark grey and uranium atoms in light blue. The close packing direction of
uranium is shown by light grey, solid lines, dashed lines are guides for the
eyes and (b) W(110) (a = 3.165 Å) where U atoms are shown in blue and tungsten
in red. The single-crystal sapphire ($\alpha$-Al2O3) substrate atoms are shown
in either projection. Adapted from Ref. [41]

Firstly, for the Nb(110)/$\alpha\mbox{-U}(110)$ epitaxial match the overlayer
grows with the orientation relationship illustrated in Fig.13a. The Nb buffer
has a growth axis [110], and one of the in-plane [1-11] axes is aligned
parallel to the sapphire [00.1] [139]. The deposited U atoms self organise in
order to align the close-packed rows of each layer, indicated by grey diagonal
lines in Fig. 13a. The epitaxy is driven by the close match between the
distances $d_{\textrm{Nb}}=a_{\textrm{Nb}}=3.311$ Å and
$d_{\textrm{U}}=\frac{1}{2}(a_{\textrm{U}}^{2}+b_{\textrm{U}}^{2})^{\frac{1}{2}}=3.264$
Å thus circumventing conventional wisdom that two systems with almost 6%
maximum lattice mismatch should never form an epitaxial system. As the system
cools back to room temperature from the elevated growth temperature, the in-
plane $c$-axis is locked into a state of tensile strain due to the positive
thermal expansion coefficient of $\alpha$-U along the $[001]$-axis ($+30\times
10^{-6}$ K-1). It is expected that the $a$-axis, which is closest to the
surface normal, should then contract slightly in order to maintain the unit
cell volume. The low mosaic spread of both the Nb and U layers
($\Delta\omega=0.15^{\circ}$) in Ref. [41] indicated high quality epitaxial
matches between the substrate, buffer and U overlayer.

Secondly, the W/U epitaxial system results in layers of $\alpha$-U(001) with
multiple domains, and an $a$-axis that is held in a state of tensile strain
[41, 135]. The complexity arises partly because the epitaxy of Al2O3/W(110)
results in the formation of two domains that are related via a 70.5∘ in-plane
rotation about the W[110] axis. The epitaxial match for one of eight predicted
U domains [41] is shown in Fig. 13b, such that the distances
$d_{U}=d_{110}=2.556\,\mbox{\AA{}}$ and $d_{W}=2d_{211}=2.584\,\mbox{\AA{}}$
are aligned producing rows of U and W atoms in register. The eight possible
domains arise since the condition
$\mathrm{U}[110]\parallel\mathrm{W}(1\bar{1}1)$ can be satisfied in eight
equivalent ways. However, the associated experimental data reported by Ward et
al. [41] was inconclusive and showed only four domains, further reduced to two
strong and two weak domains by the underlying sapphire orientation. It is
noteworthy that these domains form a pseudo-hexagonal symmetry with angles
between domains of either 52 or 63∘. These unusual angles arise because of the
orthorhombic symmetry of U. However, this pseudo-hexagonal symmetry could have
played a role in the earlier experiments reporting hcp-U [24, 25, 26, 140] as
X-ray characterisation of the films was not performed.

####  3.2.3 The hunt for hexagonal close packed uranium

Figure 14: Possible epitaxial matches of of hcp-U on Gd. (i) Domain 1, which
has $[a]_{\mathrm{U}}$ parallel to $[a]_{\mathrm{Gd}}$ and a repeat motif
every nine lattice spacings along $[a]_{\mathrm{Gd}}$. (ii) Domain 2 is
oriented at 30 to domain 1 and has $[b]_{\mathrm{U}}$ parallel to
$[a]_{\mathrm{Gd}}$ and a repeat motif every seven lattice spacings along
$[a]_{\mathrm{Gd}}$. Reprinted from [136].

As explained above, and shown in Fig. 10 the phase diagram of U does not
contain the hexagonal close packed (hcp) form, despite it being a common
structure in elemental metals. However, as was first pointed out by Axe [141],
the single soft phonon mode that relates the bcc and orthorhombic
$\alpha$-forms passes first through a hypothetical hcp structure. As such one
could imagine that the hcp structure could be stabilised in thin film form if
an appropriate epitaxial match and growth conditions could be identified. On
cooling these materials might well develop magnetic ordering, super-
conductivity and/or CDWs; indeed, some of these phenomena have already been
predicted [142].

There are reports that U may have been stabilized in such a hcp phase [24,
25]. The first paper in 1998 [24] describes the fabrication of the films and
their characterization with low-energy electron diffraction (LEED). After
evaporation of the U metal onto a W(110) substrate, the films were annealed at
$1127\,^{\circ}$C, and the subsequent LEED pattern suggested a close-packed
structure of with an interatomic spacing of 3.2(1) Å. They then state that:
“The LEED data, however, do not allow us to decide definitively whether the
close-packed pattern relates to a cubic structure like $fcc\mbox{-Th/Ce}$, or
to a hexagonal arrangement, as for most rare-earth metals.” Because the films
were fabricated in a closed system, and strongly oxidise if exposed to air,
they were not able to perform an X-ray analysis on these films. This would, in
any case, be difficult, as the films were only $\sim$ 80 Å thick, which is
close to the limit for analysis with a laboratory-based X-ray diffractometer.
Subsequent resonant photoemission experiment were performed on these samples
to show that the material had considerable 5$f$ weight at the Fermi level
[25], and scanning tunnelling spectroscopy [26], again showing the predominant
feature of the 5$f$ states near the Fermi energy. In Ref. [26] the authors
report lattice parameters of $\bm{a}=3.5(3)\,\mbox{\AA{}}$ and
$\bm{c}=5.4(2)\,\mbox{\AA{}}$ , giving a $c/a\sim 1.61(8)$ , which is not far
from the close-packed ideal value of 1.633. More recently a 70 Å film was also
grown by Chen et al. [140] who reported (from LEED) a close-packed structure
with a U-U spacing of 3.15(10) Å. The $\bm{c}\mbox{-axis}$ lattice parameter
was not stated, but using angular-resolved photoemission they reported
evidence for the 5$f$ states hybridising with the conduction-electron states,
as well as $f–f$ hybridisation. Again, a difficulty in this work is the
absence of comprehensive X-ray structural characterisation.

Figure 15: Specular scans of a $1400\;\mbox{\AA}$ hexagonal-close-packed U
film grown onto a polycrystalline $\mbox{Cu}{111}$ surface over the first
seven days following growth. Arrows indicate the direction of growth/decay of
the Bragg peaks. The dashed line represents the position of bulk
$(002)_{\alpha}$. These data have not been corrected for absorption of the
x-ray beam. Reprinted from [143].

Regarding capped systems - that can be removed from UHV conditions for X-ray
characterisation - as discussed in Section 3.2.1, it was reported that the U
layers (which were usually $<100\,$ Å in thickness) in U/Gd multilayers formed
in a hexagonal structure with $c=5.60(1)\,$ Å. If an atomic volume similar to
the $\alpha$-U phase is assumed, then this gives $a=2.91(1)\,$ Å and a ratio
$c/a$=1.92, a very large value. As reported in [136], an attempt was made
extend this result and form an epitaxial hcp layer by depositing a $500\,$ Å U
film onto a $500\,$ Å epitaxial buffer of Gd grown on Nb. The $c$-axis was
found to be $5.625(5)\,$ Å, and using the off-specular family of (10.4)
reflections $a$ was to found to be $2.96(2)\,$ Å, giving an atomic volume
close to the $\alpha$ phase and $c/a=1.90(2)$, consistent with the value found
in the multilayers. Some diagrams of possible epitaxy are shown in Fig. 14,
but these were not experimentally confirmed and attempts to increase the U
layer thickness resulted in exfoliation. Furthermore, in the successful,
$500\,\mbox{\AA{}}$ thick system, the hcp-U layer displayed poor mosaicity of
1.5∘, compared to $<0.2^{\circ}$ for the $\alpha$-U films on Nb.
Interestingly, theory [136], has predicted a value for $c/a=1.84$ with a
similar atomic volume as $\alpha$-U. A more recent theory [142], has found
$a=3.01\,$ Å, and a $c/a=1.82$. When these authors fixed their $a=2.96\,$ Å to
agree with the experiment [136], they calculated $c/a=1.864$. Thus, the large
$c/a$ ratio for the hcp phase seems to be a strong feature of both experiment
and theory.

Recently an extensive study of preparing hcp-U has been reported by Nicholls
et al. [143] to try and understand whether this form of U can be made stable
and thick enough to investigate other material properties e. g. electronic
transport, or lattice dynamics. In addition to using the substrates discussed
above, an effort was also made on Cu(111) and Ir(111) faces, as these fcc
materials present a hexagonal face in this orientation that has lattice
parameters close to those discussed above for hcp U. Fig. 15 shows that the
initial phase deposited was indeed hcp, but that a short time after growth
($\sim 20$ minutes for a film of $1400\,\mbox{\AA{}}$) there is a rapid
decomposition of the hcp phase, which transitions to the stable orthorhombic
structure with the the main grain orientations (020) and (110) [143] . Thinner
samples of hcp films are stable for somewhat longer times, but in all cases
observe phase decomposition of the hcp phase was observed as a function of
time. It was also observed in the same study that the specific decay path hcp
to orthorhombic U varies depending on film thickness.

####  3.2.4 Achieving single crystal films of $\gamma\mbox{-phase}$ alloys

Figure 16: Nominal bcc lattice parameter (measured at room temperature) as a
function of Mo content as determined by lattice parameter and calculated using
the empirical fit shown in red, data shown as open circles [144]. $\alpha$,
$\gamma$ and phase separation fields are highlighted in red, white and blue,
respectively. The first two contain all metastable phases with the indicated
parent structure. The latter contains metastable $\gamma$ along with the
$\gamma^{\prime}$ phase, the tetragonal compound U2Mo. For detailed discussion
of the various metastable structures the reader is referred to Ref. [145].
3000 Å and 700 Å films are shown as blue and green solid points, respectively.
(Insert) Schematic of film structure. Adapted from from Ref. [138]

It has been known since the mid 1950’s that $\alpha\mbox{-U}$ makes for a poor
nuclear fuel as, due to its orthorhombic structure, it displays both
anisotropic thermal expansion and poor dimensional stability under irradiation
[114, 146]. As such, the high temperature, $bcc$, $\gamma\mbox{-phase}$ was
highlighted early on as an enticing prospect for increased U density for low-
enriched fuel solutions [147]. Long lived $\gamma\mbox{-like}$ metastable
samples have been obtained through the alloying of various transition metals
combined with rapid cooling techniques since the 1960’s, up to the present day
[148, 149, 150]. However, a direct result of such fabrication methods is the
inability to produce single crystal samples thus severely limiting the
experimental probes that can be brought to bear when attempted to understand
the properties of this allotrope. Adamska et al. [151] were the first to
demonstrate that epitaxial matching could provide a mechanism to avoid rapid
cooling and instead lock in the metastable $\gamma\mbox{-like}$ phase using
epitaxial strain. This procedure makes use of the
$\mathrm{Nb}/\mathrm{Al}_{2}\mathrm{O}_{3}$ match detailed in section 3.2.2 to
provide a $[110]$ surface onto which to co-deposit U and Mo at
$800^{\circ}\mathrm{C}$, inside the stable $\gamma$ region, and the similarity
in $\bm{a}_{\mathrm{Nb}}=3.33\,$ Å and $\bm{a}_{\gamma\mbox{\tiny-UMo}}\approx
3.41\,$ Å allows a simple “cube-on-cube” epitaxial match that produces
sufficient strain to preserve the $\gamma^{\mathrm{s}}$ phase upon cooling to
room temperature. This work was built on further by Chaney et al. [138] and
refined to span the large majority of the $\gamma\mbox{-like}$ region in the
metastable UMo phase diagram as shown in Fig. 16. Mo was chosen as proof of
concept as it provides the strongest stabilising effect of all the transition
metals and is also the most likely future fuel candidate, however this growth
approach can reasonably be extended to any $\gamma\mbox{-like}$ U-transition
metal alloy system with high probability of success.

Table 1: U allotrope, alloy, and multilayer thin films that have been produced, with references of their first mention in publication. *This structure was not confirmed with structural characterisation. Material | Form | Substrate | Deposition method | Reference
---|---|---|---|---
Allotropes |  |  |  |
hcp* U |  (001)? |  (110) W | vapour deposition | [24]
$\alpha$ U |  (110) |  (110) Al2O3 with (110) Nb buffer | DC mag. sputtering | [41]
$\alpha$ U |  (001) |  (110) Al2O3 with (110) W buffer | DC mag. sputtering | [41]
$\alpha$ U |  (021)/(110) |  (110) Al2O3 with (110) Nb and (001) Gd buffer | DC mag. sputtering | [41]
hcp U |  (001) |  (110) Al2O3 with (110) Nb and (001) Gd buffer | DC mag. sputtering | [46]
$\alpha$ U |  (001) |  (110) Al2O3 with (110) Nb and (110) W buffer | DC mag. sputtering | [135]
Alloys |  |  |  |
U-Mo | poly | glass | DC mag. sputtering | [137]
U-Mo |  (110) |  (110) Al2O3 with (110) Nb buffer | DC mag. sputtering | [137]
U-Zr | poly | glass | DC mag. sputtering | [151]
U-Zr |  (110) |  (110) Al2O3 with (110) Nb buffer | DC mag. sputtering | [151]
Multilayers & Bilayers |  |  |  |
UAs/Co | amorphous | glass | mag. co-sputtering | [35]
U/Fe | poly | polyimide (Kapton) and glass | DC mag. sputtering | [132]
U/Fe | poly |  (110) Al2O3 with (110) Nb buffer | DC mag. sputtering | [46]
U/Gd | poly |  (110) Al2O3 with (110) Nb buffer | DC mag. sputtering | [46]
U/Co | poly |  (110) Al2O3 with (110) Nb buffer | DC mag. sputtering | [46]
U/Fe | poly | glass | DC mag. sputtering | [45]
U/Ni | poly | glass | DC mag. sputtering | [45]

###  3.3 Science with uranium metal thin films

This Section discusses: the early attempts to make multilayers containing one
layer of pure U metal; the important breakthrough to achieve epitaxial
$\alpha$-U metal; the discovery that interfacial strain changes the properties
of the $\alpha$-U film; experiments attempting to stabilize hcp-U films; and
the first efforts to make U/ferromagnet heterostructures for spintronics
experiments. We end with a short description of experiments on thin films (not
epitaxial) of plutonium and important photoemission experiments showing the
delocalisation of the Pu 5$f$ electrons as a function of film thickness.

####  3.3.1 Induced magnetism in uranium-ferromagnetic multilayers

The fundamental question in multilayer samples which combine ferromagnetic and
non-magnetic materials is whether the latter has any magnetic moment induced
on it due to the proximity of the former. Early work by Beesley et al. on U/Fe
multilayers [132, 133] employed Mössbauer spectroscopy and polarized-neutron
reflectivity to examine the magnetic moments at the Fe atomic site, and
determined the easy axis of magnetisation was in the plane of the layers.
Later, the element specific synchrotron-based techniques of X-ray magnetic
circular dichroism (XMCD) [152] and X-ray resonant reflectivity (XRRR) [153]
were developed allowing information on the magnitude and profile of the much
smaller induced moment at the U atoms sites. Notably it was found that the U
moments are indeed induced, but are only significant close to the interface
for multilayer systems containing transition metal atoms. As mentioned in
Section 3.2.1 the interface sharpness was relatively poor for multilayers
containing 3$d$ transition metals with approximately 15 Å of interdiffusion
between the layers. Furthermore, Mössbauer experiments [154] suggest that, at
least in the case of Fe, the structure of the Fe in this region was amorphous,
and it is in this region that the largest induced U moments are found. This
sharp decrease in the U moment as a function of distance away from the
interface has been supported by theory [155]. The magnetic moments of the
transition metals were slightly smaller than the bulk values, and the easy
direction of magnetisation was in the plane of the films [47].

The work with the transition-metal ferromagnets was summarized in Ref. [156]
in 2008. It was noted that that the maximum U moment in the U/Fe system was
$\sim$ 0.12 $\mu_{\mathrm{B}}$ and reduced for both the U/Co and U/Ni systems.
This behaviour is ascribed to the position of the 3$d$ band relative to the
Fermi level ($E_{\mathrm{F}}$). Given that the 5$f$ band is centered close to
the Fermi level, $E_{\mathrm{F}}$, a maximum in $3d\mbox{-}5f$ band overlap,
and thus maximised hybridization, occurs when the transition metal $3d$ band
also lies close to $E_{\mathrm{F}}$. This criteria is maximised for Fe and
decreases for the higher elements, Co and Ni.

(a)

(b)

Figure 17: (a) Charge diffuse scattering (black) and the magnetic diffuse
scattering (blue) at the Gd $M_{5}$ edge of a U/Gd multilayer of composition
[U(50 Å)/Gd(50 Å)]20. Data are shown as open symbols and lines are modeled
fits. Inset shows the magnetic diffuse scattering from the two X-ray
polarization states. (b) Schematic of columnar growth of Gd in the layers. The
red atoms have the full moment of $\sim$ 7 $\mu_{\mathrm{B}}$ parallel to the
applied field, whereas the atoms at the edge have their moments in some random
direction that averages to zero over the whole layer. The reduction of the
moment as a function of $D=(A+B)$ may be easily calculated. Both figures
reprinted from Ref. [134].

In contrast, considering now the lanthanide/uranium multilayer system, U/Gd
[134], there is no overlap in energy between that of the U 5$f$ band and that
of the localised Gd 4$f$ electrons, since the latter are well removed from
$E_{\mathrm{F}}$. Hence the induced U moment is much smaller, $\sim$ 0.02
$\mu_{\mathrm{B}}$ per atom, but it does oscillate through the U layer, as
might be expected if this is driven by a RKKY-type interaction mediated via
the conduction electrons. Notably the Gd magnetic moments are much reduced to
$\sim$ 4 $\mu_{\mathrm{B}}$ from the elemental value of 7.5 $\mu_{\mathrm{B}}$
known for pure Gd.

This largeapparent reduction of the magnetic moment in U/Gd multilayers has
been investigated using various techniques [134]. One possible scenario to
explain the reduced magnetisation is that the particular type of columnar
growth observed by TEM, and discussed in Section 3.2.1, leads to the pinning
of Gd moments at the boundaries of the columns. To investigate this requires
tools that probe the length scales laterally within the layers, rather than
the vertical growth direction. By tuning polarized X-rays to the $M_{5}$ edge
of Gd at 1187 eV it was possible to measure sufficiently far away from the
specular ridge to learn something about length of both the chemical and
magnetic correlations [134]. When data are obtained with the X-ray circularly
polarized in the two opposite directions the magnetic components change sign.
Hence the charge diffuse scattering, which represents the chemical
configuration, can then be obtained by adding the two contributions and
dividing by two, whereas the magnetic contribution is obtained by subtracting
the two contributions. The results are shown in Fig. 17a, and it can be easily
seen that the intensities fall abruptly at the same value of Qx, the component
of the moment transfer in the plane of the layers. This shows that the
chemical and magnetic diffuse scattering define a similar length scale over
which these correlations occur, which is not surprising given the proposed
link to the columnar growth. A detailed analysis shows that this distance is
$120\pm 20$ Å, whereas the interlayer distance is $\sim$ 50 Å. This
information then allows a simple model to be constructed based on the idea
that the columnar growth resembles the model in Fig. 17b, which explains the
strong reduction of the Gd saturated moment, and confirms the columnar growth
of Gd [134].

####  3.3.2 Uranium metal bilayers for future spintronic applications

Figure 18: In-plane angular dependence of the coercive field $\mu_{0}H$ and
normalized remnant moment $M_{r}^{*}$ for $\mathrm{FM}=\mathrm{Fe}$ and
$d_{\mbox{U}}=65\mbox{\AA}$. Lines are a guide to the eye. Taken from Ref.
[45].

In the field of spintronics, where information is carried by the spin magnetic
moment of the electron rather than charge, one of the key parameters is the
spin-orbit coupling parameter of the material, which can have a significant
impact on both the generation of spin currents (via the spin Hall effect for
example), and the spin lifetime of carriers in non-magnetic systems. Since the
spin-orbit interaction increases in the periodic table as approximately
$Z^{4}$, it is of interest to see whether U can be used in such devices, and
be competitive with other materials. The first experiments with U [157] were
done on a bilayer consisting of a 125 Å of permalloy
$(\mbox{Ni}_{0.8}\mbox{Fe}_{0.2})$ ferromagnetic layer deposited on glass with
a layer of 30 Å of uranium metal (non-magnetic layer) on top, capped with a 30
Å film of Nb to prevent oxidation. The key parameter to determine is the spin
Hall angle $\theta_{\mathrm{H}}$, where this is defined as the ratio of the
injected spin current and resulting charge current. It is this quantity which
should increase strongly with heavier elements. The method applied uses
dynamical spin pumping [158]. In spin pumping, the precessing magnetisation of
an externally excited ferromagnet (FM) undergoing ferromagnetic resonance
(FMR) is dynamically coupled to the charge carriers in an adjacent non-
magnetic (NM) system, resulting in a net transfer of spin angular momentum
across the ferromagnetic/non-magnetic interface. The measured value of
$\theta_{\mathrm{H}}$ for U is positive and has a value of +0.004. This value
is smaller than the value measured of +0.006 [159] for Pt, which is somewhat
surprising. Since both the $5f$ and $6d$ shells of U are less than half
filled, one would, to first approximation, expect a negative
$\theta_{\mathrm{H}}$ for U, as is found for transition metals. In detail,
however, the situation is more complex: the extrinsic and intrinsic
contributions to the spin Hall angle need to be isolated, which depends on the
presence of impurity scattering in the samples [160]. The crystal structure
and resultant electronic band structure play an important role in the
magnitude of the intrinsic contribution, and can be calculated in terms of the
Berry phase curvature. Theoretical work by Wu et al. examined the $\alpha$,
$\gamma$, and hcp phases of U, and suggested that the hcp phase shows the
largest spin Hall conductivity near to $E_{\mathrm{F}}$. In the same work they
also showed that the nature of any 3$d$ transition metal ion impurities (which
might be present at a disordered interface) can significantly affect the
magnitude and even the sign of $\theta_{\mathrm{H}}$ [161]. In more recent
work the same authors show the peculiar result that spin accumulation in
uranium films is highest at the side of the film opposite to the impurity
position [162]. Clearly, there is more to do on this subject to understand the
value of U in spintronic applications.

In another study the anisotropic magnetic properties of thin U/Fe and U/Ni
bilayers [45] were investigated. Here the transition metal layers of $\sim$
100 Å were deposited on glass, and the U layer thicknesses ($d_{U}$) were
varied from 0 to 80 Å. The growth was carried out at room temperature, and the
layers were polycrystalline, as expected. The U appeared to be mostly [001]
textured as the preferred growth axis. Magnetisation measurements at room
temperature were made of the coercive field as a function of angle in the
plane, to investigate the impact of the heavy U on the magnetic anisotropy of
the transition metal layer: since magnetocrystalline anisotropy is inherently
linked to spin-orbit coupling, the proximity of enhanced spin-orbit effects in
the U layer might be expected to have a significant impact. Some data are
shown in Fig. 18.

From these results the uniaxial anisotropy coefficient $K_{\mathrm{eff}}$ may
be calculated and this is found to have a maximum in the U/Fe system at
$d_{U}\sim 55$ Å. Further examination suggests that quantum-well states in the
U may be affecting the changes in $K_{\mathrm{eff}}$, suggesting some impact
of the electronic structure of neighbouring U layer on the magnetic layer.
Once again, the interface probably needs further investigation, especially
since an induced moment on the U may be present, and as discussed in Section
3.3.1, an interdiffused layer of at least 15 Å, might be present.

####  3.3.3 Manipulating the CDW States in $\alpha\mbox{-uranium}$ using
epitaxial strain

Figure 19: (a) shows the U/Nb growth axis as [110] with the plane of the film
in grey containing the [001] axis. In the bulk material [114] the CDW is
formed along the propagation vectors, Q1, Q2, Q3, and Q4 so they all lie
slightly away from the yellow plane containing [100] and [010]. (b) shows the
$d$–space of the (220) (specular) charge peak as a function of temperature for
both the bulk and from a 250 Å epitaxial film. Taken from Ref. [135].

As explained in Section 3.1, significant research efforts were expended to
understand the low temperature charge ordering in $\alpha\mbox{-U}$ through
which it was established that there exists a series of CDW transitions
starting at 43 K and ending at 22 K with a fully commensurate CDW state
described by the wavevector
$q=\frac{1}{2}\bm{a}^{*}+\frac{1}{6}\bm{b}^{*}+\frac{5}{27}\bm{c}^{*}$ [114].
It was also determined that one of the key parameters determining CDW
behaviour is the length of the $\bm{a}\mbox{-axis}$, as this is the direction
along which the primary component of the CDW distortion lies [114]. It is
known from work on the phonons under pressure that on compression the soft-
phonon mode in the [100] direction hardens (i.e. increases in energy)
increasing the energy required for the CDW to form, suppressing the CDW state
and enhancing $T_{\mathrm{c}}$ from $\sim 0.5$ K at ambient pressure to a
maximum of $\sim 2$ K at 1.2 GPa [130]. In a naïve picture of competing CDW/SC
ground states, tensile strain along the $a$-axis should therefore enhance the
CDW stability and suppress superconductivity. Given this, the difference in
both the magnitude and sign of the strain exerted on the $\alpha$-U $a$-axis
in the epitaxial Nb/U and W/U systems described in Section 3.2.2 opened the
door to the potential to achieve something impossible in bulk, manipulating
the CDW state via epitaxial strain.

Figure 20: (a) The lattice parameters of the strained uranium film as a
function of temperature. The size of the symbols indicates approximately how
much of the sample has that lattice parameter (or volume) at the given
temperature. The bulk values are given by stars. (b) Temperature dependence of
the intensity (filled circles) and full-width at half maximum (FWHM) (open
triangles) for the (1.5 0 3). (c) Theoretical phonon dispersion for bulk vs
thin film calculated at 0 K using methods reported in ref. [126]. Experimental
points at 300 K [124] are shown as open circles for comparison. Insert shows
the effect of compression (negative) or expansion (positive) of the $a$-axis
on the unstable mode. All panels taken from Ref. [135].

The initial exploratory studies in this area were conducted in 2008 [136] and
found that the two systems do indeed display vastly different CDW states.
Starting with the Nb/U films system, the CDW reflections were found in the
bulk positions [136] but with a dramatic change in the relative reflection
intensities compared to bulk where all satellites are equal in intensity
[114]. Fig. 19b gives the clue as to why the CDW is so close to that found in
the bulk, but is in fact “weaker”, as measured by the smaller change of the
$d_{220}$ in the film compared to the bulk. The $d_{220}$ spacing, which
includes a component from the $\bm{a}$-axis, is smaller for the film than for
the bulk, i.e. $\bm{a}$-axis is under compression in the film. This implies a
weaker CDW interaction, hence less intensity in the CDW peaks, and less change
of lattice parameter on cooling below $T_{\mathrm{CDW}}$ although the latter
stays at about the same value. Since the film was grown on a Nb buffer, which
is superconducting at $\sim 9\,\mathrm{K}$, no effort was made to see whether
the film was superconducting, since a simple resistive measurement would not
detect the lower $T_{\mathrm{c}}$. Regarding the change in relative reflection
intensities; for each layer of satellites, $h\pm\frac{1}{2}$, the CDW
reflections are defined by the vectors Q1 to Q4, shown in Fig. 19a, and it was
found that for $h+$ the satellites could be divided into two clear pairs of
comparable intensity, a strong pair Q1/Q3 and a weak pair Q2/Q4. The matching
intensity within a pair is expected as the only difference is $\pm k$ which is
a true mirror plane, however to explain the intensity variation between pairs
the authors note that the stronger pair is closely aligned with the growth
axis [110]. It can thus be inferred that a major effect of epitaxial strain is
to promote or suppress the formation of CDW domains depending on their
relative alignment with the growth axis. The variation of the CDW wavevector
was also measured as a function of temperature and it was found that, unlike
bulk, there were no lock-in transitions of the individual components,
$\bm{q}_{x}$, $\bm{q}_{y}$, and $\bm{q}_{z}$ [136].

The situation in the U/W system is quite different to that in the U/Nb system,
and the complex domain structure discussed in Section 3.2.2 needs to be
suppressed in order to simplify reciprocal space for measurements. By
depositing a thin (100 Å) layer of niobium between the substrate and tungsten
layer, the W(110) buffer is limited to a single domain and the total number of
observable $\alpha$-U domains is reduced from eight to four. The domains are
unequally populated and only two strong reflections (separated by an in-plane
rotation of $\sim 56^{\circ}$) tend to be observed [135]. As shown in Fig.
20a, at 250 K the atomic volumes are comparable, but the $a$ and $c$ axes have
expanded, and the $\bm{b}$ is contracted. The CDW occurs at $\sim$ 120 K
(three times that of the U/Nb film and the bulk) and results in a large
further expansion of the $\bm{a}$-axis; at the same time the $\bm{c}$-axis
(which is not constrained, as this is the growth direction) contracts to allow
the atomic volume to approach that of the bulk sample. Further to the dramatic
increase in $T_{\mathrm{CDW}}$, the most remarkable result of the tensile
epitaxial strain is to fundamentally change the CDW character, such that it
becomes fully commensurate with the lattice, and changes from three- to one-
dimensional losing both $\bm{q}_{y}$ and $\bm{q}_{z}$ components retaining
only the $\bm{q}_{x}=0.50$ component. Fig. 20b shows the development of this
peak, and its width, as a function of temperature.

This work showed conclusively that the expansion of the $\bm{a}$-axis is the
key to the marked increase in $T_{\mathrm{CDW}}$. In the U/W films the
$a$-axis is under tension, allowing its expansion. It is also confirms the
theory developed first in Ref. [130] that the soft-mode phonon behaviour, and
the associated strong electron-phonon interaction, are crucial to the
development of the CDW. In the context of these different types of long range
order, further strain tuning of epitaxial films could provide invaluable
information regarding the interplay between these two low temperature ground
states.

####  3.3.4 Correlated disorder in pseudo-bcc uranium alloys

As introduced in Section 3.2.4 there has been growing interest in bcc uranium
– transition metal alloys over the last 20 years with a re-assessment of their
potential role as advanced nuclear fuels [163]. However, as traditional
synthesis methods are incapable of producing single crystals, detailed
structural analysis has been severely limited. The first seminal work on this
problem was presented by Yakel in 1969 [149, 164] who obtained diffraction
from single grains at the tip of a polycrystalline needles. He observed
significant diffuse scattering in the $\gamma^{\mathrm{s}}$ phase, and
attributed it to local structural modifications which preserved the global bcc
structure as a configurational average refuting previous work suggesting
chemical order [148]. However, his proposed solution was both complex and
largely ignored by the majority of the U-alloy community who, for the most
part, have considered $\gamma^{\mathrm{s}}$ as fully stabilised bcc. This is
understandable as the characteristic diffuse intensity is at least eight
orders of magnitude weaker than the Bragg reflections and therefore when
working with polycrystalline samples and standard powder diffraction
refinement techniques is impossible to detect [165, 166, 167].

Figure 21: Diffuse X-ray scattering patterns taken with $\lambda=0.98$ Å at
room temperature from the three different Mo concentrations as marked outside
quadrants. (Main) Reciprocal space reconstructions of the $(hk2)_{\mathrm{s}}$
plane produced using the superstructure unit cell with
$b_{\mathrm{s}}\parallel b_{\mathrm{p}}$. Reflections are categorized and
indexed in the bottom right quadrant, with indices in superstructure notation.
Reflections are indicated in black (Mo parent Bragg peaks), blue (N – domain
1), green (N – domain 2) and orange (H). The Nb buffer reflections are next to
the Mo parent peaks. Powder rings arise from the polycrystalline Nb cap,
whereas narrow intense peaks correspond to substrate Bragg reflections, none
of which are included schematically. The ($\bar{1}32$) reflection, marked with
an asterisk, is highlighted in all data sets to show mirror symmetry
relations. (Insert) Reconstructions for the $(h1l)_{\mathrm{s}}$ plane
highlighting the evolution of H type signal. The H point is marked with a
circle. 31 at% peaks are indexed with $\mathrm{U}_{2}\mathrm{Mo}$ with c
parallel to $a_{\mathrm{p}}$ (red) and $c_{\mathrm{p}}$ (blue). Adapted from
Ref. [138].

As explained in Section 3.2.4 the single crystal synthesis problem was
recently overcome using epitaxial matching as a stabilising force in
substitute of rapid cooling [138, 151] allowing the $\gamma^{\mathrm{s}}$
phase to be studied in detail using modern synchrotron techniques for the
first time. Using the diffuse x-ray scattering diffractometer at the ESRF
[168] Chaney et al. confirmed the bcc structure is not preserved locally and
instead a three-dimensional structural modulation gives rise to a well defined
diffuse X-ray pattern, shown in Fig. 21. Further, the exceptional flux and
resolution provided by modern synchrotrons allowed for another key
breakthrough, the separation of the diffuse signal into two distinct types of
different origin. The first group (H) were re-indexed as a coherent precursor
structure of the intermetallic $\mbox{U}_{2}\mbox{Mo}$ allowing the remaining
diffuse alloy reflections (N) to be explained by a substantially simplified
model. The unique structural solution was given by a local orthorhombic
structure, formed by atomic displacements along the $\langle 001\rangle_{s}$
with no requirement for chemical ordering. The superstructure is C-face
centered and defined with four atoms in the unit cell at
$(0,\nicefrac{{1}}{{4}}\pm\delta,\nicefrac{{1}}{{4}})$ where
$\left|\delta\right|=0$ corresponds to the bcc structure. Not only is this a
simpler solution, it also removes the unusual condition present in the
previously suggested structure, whereby there are two atomic sets with
distinctly different coordination and no chemical ordering.

Figure 22: Experimental and theoretical phonon dispersion curves for the U 23
at. % Mo alloy at room temperature. The transverse (longitudinal) acoustic
modes are shown as blue squares (red circles). The VCA theory is shown as
dashed (TA) and dotted (LA) white lines. The light shading shows the widths
(from theory) after allowing for the alloying; note this is not on a linear
scale. The soft mode $\mathrm{TA}_{1}[110]_{\mathrm{p}}$ predicted by theory
at the N point could not be observed, as the diffuse (elastic) scattering
dominates the signal. The panel below the main figure shows the intrinsic
widths (after deconvoluting with the instrumental resolution function and the
theoretical width) of the phonon modes, which have both a directional, as well
as an energy dependence. Adapted from Ref. [138].

Given all atoms are displaced from their bcc positions, the individual
displacements, which transform from the parent to superstructure, can be
viewed as a frozen phonon. The responsible modes being of general type
$\mathrm{TA}_{1}[110]_{\mathrm{p}}$, with polarization along the fourfold axis
$[001]_{\mathrm{p}}$ and are the exact modes predicted to destabilize the
high-temperature bcc phase [126, 169]. Twelvefold degeneracy generates six
equivalent and equally occupied superstructure domains, however, given the
diffuse nature of the reflections it is clear that the identity of any one
choice of domain is only preserved over a relativity short distance, $\sim 30$
Å along $\bm{b}_{\mathrm{s}}/\bm{c}_{\mathrm{s}}$ and $\sim 22$ Å along
$\bm{a}_{\mathrm{s}}$. As such, Chaney et al.. [138] suggest that the local
superstructure is better thought of, not as local order, but as correlated
displacive disorder that lowers the local symmetry with three dimensional
correlations between nearby atoms governed by rules represented as a frozen
phonon. Twelvefold degeneracy maintains the higher average symmetry, while
allowing anisotropic neighbour distances reminiscent of
$\alpha\mbox{-uranium}$ to be recovered. This situation can be understood
intuitively as U has a desire, rare among the elements, to occupy highly open
structures that produce extreme anisotropic local environments [119, 120].
Thus, by attempting to stabilise U onto a highly symmetric bcc lattice the
mismatch in preferred symmetry creates an intrinsic conflict within the
system. In the absence of sufficient thermal energy, this conflict creates a
structural instability resolved by the formation of correlated disorder. The
same experiments also showed this correlated disorder was intrinsically
tunable. The authors use a qualitative metric combining correlated volume and
magnitude of atomic displacement to capture the “correlated disorder strength”
and show strong tunability with alloy composition, as evidenced by the
variation of intensity and FWHM of the diffuse reflections in Fig. 21 as a
function of minor alloy content.

Building on the discovery of correlated displacive disorder Chaney et al.
[138] explored the phonon dispersion for the 23 at. % Mo system using the
grazing incidence inelastic X-ray scattering technique pioneered at the ID28
beamline [170], ESRF [171, 172]. It is well understood that correlated
disorder can couple to other periodic phenomena such as lattice dynamics [173]
and as mentioned in Section 3.1 the uranium phonon dispersion has been subject
to many notable theoretical efforts including high-temperature modelling [169,
174] and more recently calculations for the UMo alloy itself [175]. However,
due to the lack of suitable single crystal samples no experimental data
existed making this a significant gap in understanding. The first dispersion
was published in 2019 by Brubaker et al. [167] who extracted a single enlarged
grain from an annealed polycrystalline 20 at. % sample by laser cutting. This
work, along with the studies by Chaney et al. [138], agree well with each
other as well as with theoretical predictions. The full dispersion as measured
by Chaney et al. [138] including room temperature ab initio calculations is
shown in Fig. 22. Both experiments also observed extraordinary phonon lifetime
suppression away from zone centre. As an alloy some intrinsic phonon
broadening is expected, however this contribution was simulated and determined
to be $<2\,\mbox{meV}$ for all $\bm{q}$, thus, the authors concluded that
almost all phonon lifetime suppression in the system is attributable to
disorder-phonon coupling.

Figure 23: Valence-band spectra of pure Mg substrate and increasing coverage
of Pu. The thickness is indicated in the number of monolayers. Spectra were
obtained with the photon energy $hv=40.8\;\mbox{eV}$. The vertical bars
indicate the position of the $5f\mbox{-localized}$ peak and the triplet of
narrow features. Reprinted from [176].

####  3.3.5 Thin films of transuranium metals

The number of laboratories that can perform experiments on transuranium films
is, of course, extremely limited. To our knowledge the only laboratory to
produce epitaxial thin films is Los Alamos National Laboratory, where
epitaxial transuranium oxide films have been grown and will be covered in
Section 4.4.1. However, the controversy over the electronic structure of the
light actinides has been such that various efforts in photoemission and
related spectroscopies have been undertaken over the past 20 – 30 years, many
of which are discussed in a recent Plutonium Handbook[13]. For our purposes in
this review we highlight one particular measurement that was able to monitor
the localization of the 5$f$ electrons in Pu metal films by monitoring the
photoemission spectra as a function of sample thickness [40, 176].

The key spectra are reprinted in Fig. 23 and show an appreciable shift for the
thinnest sample. Of course, in this work no X-ray characterization could be
made of the samples, so the assumption is that the Pu metal was in the
$\alpha$-Pu form and was polycrystalline. These spectra have been often cited
in work on the electronic structure of Pu metal. The 3-peak (A, B & C)
structure are characteristic of 5$f$ localization, together with the peak D
that represents the position of the 5$f$ states below $E_{\mathrm{F}}$.

##  4 Uranium Oxide Systems

###  4.1 Introduction

As the difficulties of manufacturing pure U metal fuel elements, as well as
the safety issues raised by low-melting temperature of 1132 ∘C, became more
evident in the 1950’s, uranium dioxide was introduced as an alternative
nuclear fuel in the early 1960’s. It remains today the fuel of choice of a
majority of power reactors. A complete industry is devoted to the manufacture
of such nuclear fuels and their treatment post-irradiation. The properties of
UO2 have thus been studied extensively in both a fundamental, as well as an
applied, sense [18]. A major issue is related to the surface properties of
UO2, as it is at the surface that reactions will initially occur, and, with
irradiated fuels, the potential danger arises from material escaping from the
surfaces of such fuels. Films in the thin limit are essentially simply
surfaces, so it is reasonable to expect that they can help with the
understanding the behaviour of UO2 surfaces. We shall show one example of
research in this area. Another key aspect is related to interfaces involving
UO2, for example the interface between the fuel and the cladding, and here
too, thin bilayers should offer new insights.

From a more fundamental point of view, UO2 is a semiconductor with a band gap
of $\sim$ 2 eV, which is almost twice that of silicon. It has garnered
interest in applications (especially in satellites) requiring high reflectance
in the wavelength range range of 40 - 100 Å. Both UO2 and UN films reflect
significantly more [177, 178] than all known materials such as Au and Ir, and
UO2 is also cheaper. Additionally, there has been some effort to consider UO2
for solar cells, as, if doped with oxygen, UO2 is a p-type semiconductor [177,
178, 42].

###  4.2 Surface studies of UO2

An excellent review of the chemical reactions at the surface of UO2 was given
by Idriss [179] in 2010. Those studies (so far) have not used thin films, but
are done on specially prepared surfaces of single crystals from bulk samples.
Many of the properties deduced are relevant to thin-film research, and
frequently can be extended with suitable films. On the chemical structure of
the surface, earlier work was reported in the 1980s [180, 181]. An
understanding of the chemical structure of the surfaces is important because
with the fluorite CaF2 structure of UO2 the only primary surface that is non-
polar is the (110). The other two surfaces, (100) and (111), are polar, so
that with these latter two surfaces there will always be an atomic
rearrangement at the surface to resolve the polar discontinuity with the
vacuum [182]. We will need some of these results to understand later studies
of the dissolution of UO2.

Upon further oxidation the chemical formula can be generalised as UO2+x, where
$x$ = 0 at stoichiometry, and proceeds through many different phases, U4O9
($x$ = 0.25), U3O8 ($x$ = 0.67), and finally arrives at UO3 ($x$ = 1). As
oxidation starts at the surface, various experiments have been performed by
Stubbs and collaborators using the technique of x-ray surface truncation rods
to understand the initial oxidation process [183, 184]. Other notable
experiments on surfaces have been reported by Seibert [185], and Spurgeon et
al. [186] on UO2+x. There has also been a series of theoretical investigations
on the surface structure of UO2 [187, 188, 189], with the last two
concentrating on the interaction of water with the surface.

Another area that has been studied is the magnetic structure and phase
transition at the surface of UO2. In the bulk this aspect has been studied
since the 1960s [8, 190], but the first observation of the magnetic structure
from the surface was with resonant magnetic scattering [191, 192]. This work
showed an unusual aspect of the phase transition at 30 K from the paramagnetic
to antiferromagnetic state, and later experiments clearly defined this as a
so-called “surface transition”, i.e. a transition that is induced directly by
the presence of the reduced symmetry at the surface [193]. Perhaps counter-
intuitively, we shall show an example where a study of the magnetism, has been
able to illustrate some important information about the interface of UO2 with
the substrate.

###  4.3 Growth of oxides

Given the abundant work reported on UO2 surfaces in Ref. [179], it is perhaps
surprising that it took so long before epitaxial films of UO2 were made
systematically. As reported in Sec. 1.1, epitaxial films were already made in
the 1960’s [19] and 1970’s [20, 21], but relatively little was done with them.

The first recorded epitaxial UO2 films were made at Los Alamos National
Laboratory and reported in 2007 [64]. The substrates were LAO, which can be
bought commercially. Since that time a whole series of papers have reported
the growth of epitaxial UO2 and also higher oxides with a wide variety of
methods on many different substrates [50, 59]. These are given in Table 2.

Epitaxial UO2 films were grown by DC-magnetron sputtering at Oxford in 2010
[72] using substrates of LAO following the success of Burrell et al. [64], but
a paper on these samples was not published until 2013. In the meantime, a
paper by Strehle et al. [43] appeared in 2012, also using DC-sputtering, and
reporting epitaxial samples of UO2 and U3O8. Strehle et al. [43] published
much useful information on the growth of UO2-based epitaxial films, and they
also demonstrated that epitaxial films could be grown on yttrium-stabilized
zirconia (YSZ), even though the mis-match was a large 6%. This is important in
UO2-film research as YSZ substrates can be readily obtained in the three
important orientations (100), (110) and (111), allowing epitaxial films of UO2
to be grown in all three principal orientations. Strehle et al. also showed by
chemical analysis with RBS and XPS that the (1–1.2) sapphire substrate
material is unstable with respect to Al transport into the uranium-oxide
overlayers.

Table 2: Uranium oxide thin films that have been produced, with references of their first mention in publication. *No substrate given in the paper. Material | Form | Substrate | Deposition method | Reference
---|---|---|---|---
UO2+x |  (001) |  (001) MgO | vapor deposition | [19]
UO2+x |  (001) |  (001) NaCl | DC sputtering | [20]
UO2+x |  (001) |  (001) NaCl, KCl, KBr, NaF, LiF | vapor deposition | [21]
UO2+x | poly | * | pulsed laser ablation | [194]
UO2, UO2-x | poly | Si wafer | reactive DC sputtering | [195]
doped UO2 | poly | Al2O3, MgO | sol-gel | [63]
UO2 |  (011)? |  (100) LaAlO3 | polymer-assisted deposition | [64]
U3O8 |  (100) |  (001) Al2O3 | polymer-assisted deposition | [64]
UO2 |  (001) |  (001) YSZ | reactive DC mag. sputtering | [43]
UO2 |  (001) |  (012) Al2O3 | reactive DC mag. sputtering | [43]
$\alpha$-U3O8 |  (001) |  (012) Al2O3 | reactive DC mag. sputtering | [43]
UO2 |  (001) |  (001) LaAlO3 | reactive DC sputtering | [49]
UO2 |  (001) |  (001) CaF2 | reactive DC mag. sputtering | [49]
$\alpha$-U3O8 |  (001) |  (012) Al2O3 | polymer-assisted deposition | [14]
$\beta$-U3O8 |  (001)? |  (001) Al2O3 | polymer-assisted deposition | [14]
UO2+x | poly | Ni | electrodeposition | [196]
UO2 |  (110) |  (100) TiO2 | reactive DC mag. sputtering | [50]
UO2 |  (111) |  (001) ZnO | reactive DC mag. sputtering | [50]
UO2 |  (001) |  (001) SrTiO3 | reactive DC mag. sputtering | [171]
UO2 |  (011) |  (011) YSZ | reactive DC mag. sputtering | [51]
UO2 |  (111) |  (111) YSZ | reactive DC mag. sputtering | [51]
UO2 |  (001) |  (001) LSAT | pulsed laser deposition | [59]
$\alpha$-U3O8 |  (001) |  (001) Al2O3 | pulsed laser deposition | [59]
$\alpha$-U3O8 |  (001) |  (001) LSAT | pulsed laser deposition | [59]
$\alpha$-U3O8 |  (001) |  (001) MgO | pulsed laser deposition | [59]
$\alpha$-UO3 |  (001) |  (001) MgO | pulsed laser deposition | [59]

###  4.4 Science with UO2 thin films

####  4.4.1 Photoemission experiments on UO2 and PuO2 epitaxial films

We shall highlight here the ARPES results obtained from UO2 and PuO2 epitaxial
films. Whereas polished single crystals of UO2 exist and were already used for
ARPES measurements, there are no sizable single crystals of PuO2, so that
effort is unique and important. The central question posed by the theory
presented also in Ref. [14] is to what extent are the actinide 5$f$ electrons
hybridizing with the oxygen 2$p$ electrons? The hybrid density functional
theory discussed predicts that in UO2 the mixing between these two orbitals is
present, but relatively small, whereas it should be much stronger in PuO2.
Below in Figs. 24 and 25 we show the UO2 and PuO2 results from the ARPES
measurements.

Figure 24: The small dispersion of $\sim$ 100 meV across the zone in the 5$f$
band for UO2. Adapted from Ref. [14] Figure 25: ARPES work on PuO2 showing
that the dispersion in PuO2 is $\sim$ 3 eV. The black bar indicates
approximately the position and width of the similar U 5$f$ states found in
UO2. Taken from Ref. [65]

Clearly, the much greater dispersion ($\sim$ 3 eV) in the 5$f$ level in PuO2
[65] than found in UO2 of $\sim$ 0.1 eV is a key finding confirming the theory
used in this investigation of the actinide dioxides.

####  4.4.2 Antiferromagnetism of UO2 epitaxial films

The work published on the films from Oxford/Bristol about the low-temperature
antiferromagnetic (AF) structure of UO2 films illustrates the challenges of
working with epitaxial films, which must be grown on substrates. The films
were produced at ITU, Karlsruhe, as well as at Bristol. All initial films used
LAO substrates and the depositing temperature was 650 ∘C. A careful
examination [49] of the UO2 films showed tetragonal symmetry, even for films
over 2000 Å thick. The $a$ and $c$ axes are shown in Fig. 26 below.

Figure 26: Evolution of the out-of-plane, $c$, and the in-plane, $a$, lattice
parameters determined from reciprocal space mapping measurements with UO2
thickness (deposited on LAO substrates). The dashed line is the bulk value of
$a$ and $c$. The upper panel shows the strain, determined as
$\frac{2(c-a)}{(c+a)}$. Taken from Ref. [49]

The UO2 films had a rocking curve width that exceeded 1∘, (2∘ for the thinnest
samples) although the substrate peaks were very sharp. Some of the thinner
samples showed no signs of AF magnetic ordering. This certainly is related to