A Review of Grain Boundary and Heterointerface Characterization in Polycrystalline Oxides by (Scanning) Transmission Electron Microscopy

Abstract

Interfaces such as grain boundaries (GBs) and heterointerfaces (HIs) are known to play a crucial role in structure-property relationships of polycrystalline materials. While several methods have been used to characterize such interfaces, advanced transmission electron microscopy (TEM) and scanning TEM (STEM) techniques have proven to be uniquely powerful tools, enabling quantification of atomic structure, electronic structure, chemistry, order/disorder, and point defect distributions below the atomic scale. This review focuses on recent progress in characterization of polycrystalline oxide interfaces using S/TEM techniques including imaging, analytical spectroscopies such as energy dispersive X-ray spectroscopy (EDXS) and electron energy-loss spectroscopy (EELS) and scanning diffraction methods such as precession electron nano diffraction (PEND) and 4D-STEM. First, a brief introduction to interfaces, GBs, HIs, and relevant techniques is given. Then, experimental studies which directly correlate GB/HI S/TEM characterization with measured properties of polycrystalline oxides are presented to both strengthen our understanding of these interfaces, and to demonstrate the instrumental capabilities available in the S/TEM. Finally, existing challenges and future development opportunities are discussed. In summary, this article is prepared as a guide for scientists and engineers interested in learning about, and/or using advanced S/TEM techniques to characterize interfaces in polycrystalline materials, particularly ceramic oxides.

Full text

Crystals 2021, 11, 878. https://doi.org/10.3390/cryst11080878 www.mdpi.com/journal/crystals
Review
A Review of Grain Boundary and Heterointerface
Characterization in Polycrystalline Oxides by (Scanning)
Transmission Electron Microscopy
Hasti Vahidi 1, Komal Syed 1, Huiming Guo 1, Xin Wang 1, Jenna Laurice Wardini 1, Jenny Martinez 1
and William John Bowman 1,2,*
1 Department of Materials Science and Engineering, University of California, Irvine, CA 92697, USA;
vahidih@uci.edu (H.V.); ksyed1@uci.edu (K.S.); huimingg@uci.edu (H.G.); xinw15@uci.edu (X.W.);
jwardini@uci.edu (J.L.W.); jennypm@uci.edu (J.M.)
2 Irvine Materials Research Institute, Irvine, CA 92697, USA
* Correspondence: will.bowman@uci.edu
Abstract: Interfaces such as grain boundaries (GBs) and heterointerfaces (HIs) are known to play a
crucial role in structure-property relationships of polycrystalline materials. While several methods
have been used to characterize such interfaces, advanced transmission electron microscopy (TEM)
and scanning TEM (STEM) techniques have proven to be uniquely powerful tools, enabling quan-
tification of atomic structure, electronic structure, chemistry, order/disorder, and point defect dis-
tributions below the atomic scale. This review focuses on recent progress in characterization of pol-
ycrystalline oxide interfaces using S/TEM techniques including imaging, analytical spectroscopies
such as energy dispersive X-ray spectroscopy (EDXS) and electron energy-loss spectroscopy (EELS)
and scanning diffraction methods such as precession electron nano diffraction (PEND) and 4D-
STEM. First, a brief introduction to interfaces, GBs, HIs, and relevant techniques is given. Then,
experimental studies which directly correlate GB/HI S/TEM characterization with measured prop-
erties of polycrystalline oxides are presented to both strengthen our understanding of these inter-
faces, and to demonstrate the instrumental capabilities available in the S/TEM. Finally, existing chal-
lenges and future development opportunities are discussed. In summary, this article is prepared as
a guide for scientists and engineers interested in learning about, and/or using advanced S/TEM
techniques to characterize interfaces in polycrystalline materials, particularly ceramic oxides.
Keywords: oxides; ceramics; polycrystalline; grain boundary; heterointerface; transmission electron
microscopy; scanning transmission electron microscopy; energy dispersive X-ray spectroscopy;
electron energy-loss spectroscopy; scanning electron nanodiffraction; 4D-STEM
1. Introduction
1.1. Background and Motivation
Solid-solid interfaces are ubiquitous in materials science and engineering with wide-
ranging properties and applications [1,2]. In polycrystalline bulk materials and thin films,
interfaces directly impact mechanical, optical, thermal, magnetic, electrical and (elec-
tro)chemical properties due to existence of local heterogeneity in structure, composition,
chemistry, and electronic structure down to the atomic scale [1–16]. Ceramics such as ox-
ides [4,17] are particularly prone to GB effects as annealing the GBs out by coarsening the
grains usually requires ≥1000 °C, which is energy-intensive, costly, and deteriorates de-
vice components. Therefore, understanding the basics of interfaces is key in optimization
of ceramics for a wide range of applications including electrochemical energy conversion
and storage, optical, magnetic, and mechanical applications, thermal applications includ-
ing thermal/environmental barrier coatings, refractories, etc. [5–7,18–25]. This effect
Citation: Vahidi, H.; Syed, K.; Guo,
H.; Wang, X.; Wardini, J.L.; Mar-
tinez, J.; Bowman, W.J. A Review of
Grain Boundary and Heterointerface
Characterization in Polycrystalline
Oxides by (Scanning) Transmission
Electron Microscopy. Crystals 2021,
11, 878. https://doi.org/10.3390/
cryst11080878
Academic Editors: Xialu Wei, Diletta
Giuntini and Yanhao Dong
Received: 28 May 2021
Accepted: 19 June 2021
Published: 28 July 2021
Publisher’s Note: MDPI stays neu-
tral with regard to jurisdictional
claims in published maps and insti-
tutional affiliations.
Copyright: © 2021 by the authors. Li-
censee MDPI, Basel, Switzerland.
This article is an open access article
distributed under the terms and con-
ditions of the Creative Commons At-
tribution (CC BY) license (http://crea-
tivecommons.org/licenses/by/4.0/).

Crystals 2021, 11, 878 2 of 58
becomes more significant in nanocrystalline ceramics with higher volume fractions of in-
terfaces. However, elucidating nanoscopic features of these interfaces is challenging and
requires analytical methods offering high spatial resolution [21,26–28]. Among available
techniques, modern S/TEMS are uniquely capable of qualification and quantification of
structure, composition, and electronic structure down to the atomic scale. For example,
cation and anion sublattices can be resolved using high angle annular dark field (HAADF)
and annular bright field (ABF) STEM imaging, respectively. HAADF images are formed
by thermal diffuse scattering of electrons by nuclei and allow locating cation columns.
ABF images [27] are formed by electrons channeled by light atoms, which forward focus
the primary beam enabling detection of anions [29,30]. Core-loss electron energy-loss
spectroscopy (EELS) can quantify elemental concentrations (sensitivity ~1 at%) and probe
electronic structure and bonding, giving information on ion valence, coordination, and
defect concentration [7,31–34]. Grain orientation and strain can be mapped in atomic-res-
olution images, or by precession electron nanodiffraction [7,31]. While there are books [35]
and review papers [36–39] in the literature describing the role of GBs and interfaces in
ceramics, a comprehensive review on the power of S/TEM in characterizing these inter-
faces is needed. Thus, this article reviews recent progress of interface characterization us-
ing S/TEM techniques with a focus on ceramics oxides. In addition, major challenges and
future perspectives are discussed.
1.2. Types of Interfaces
An interface is a planar feature separating two phases or domains of matter and can
have properties different from those of the “bulk” material on either side. Here, interfaces
between crystalline ceramics, particularly metal oxides, in either single or multi-phase
form are discussed. In this context, a grain boundary (GB) is an interface where the two
phases comprise the same material (i.e., equivalent stoichiometry and crystal structure).
GBs have been studied since at least the 1940s, described by Burgers [40] as a “transition
surface”—what we today call a dislocation array (Figure 1).
Figure 1. Schematic diagram of a transition surface between two domains, formed by a set of par-
allel lines of edge dislocations, all situated in the plane x = 0 [40]. Reprinted from Proceedings of
the Physical Society (1926–1948), Geometrical considerations concerning the structural irregulari-
ties to be assumed in a crystal, J. M. Burgers, Proc. Phys. Soc. 1940, 52, 23. © IOP Publishing. Re-
produced with permission. All rights reserved.

Crystals 2021, 11, 878 3 of 58
1.2.1. Heterointerface
A heterointerface (HI) describes an interface between two phases that differ structur-
ally and/or chemically [4]. Figure 2 contains schematic diagrams and BF-TEM micro-
graphs of GBs and HIs in a single phase and a multiphase polycrystalline oxide. GBs/HIs
comprise an approximately two-dimensional “core” or “structural” region where the
atomic structure deviates from that of both adjacent crystals over a distance of ~ 1 unit cell
in the direction normal to the interface plane. A diffuse three-dimensional “space charge
zone” or “chemical” region extends up to several nm into adjacent crystals depending on
the point defect concentration [40].
Figure 2. Schematic diagrams of interface between (a) same material (GBs) and (b) different mate-
rials (HIs) at atomic scale (c) BF-TEM micrograph of GBs in a polycrystalline Gd/Pr co-doped CeO2
(d) BF-TEM micrograph of HI and GBs in a polycrystalline MgAl2O4, YSZ and Al2O3 multiphase
oxide. Grains in red and blue in (a) and (b) are depicting different materials with different chemis-
try and crystal structures for demonstration [2,8,41]. Figure 2a, b reprinted by permission from
Springer Nature Customer Service Centre GmbH: Springer Nature, Journal of Electroceramics,
when two become one: An insight into 2Dconductive oxide interfaces, Pryds, N.; Esposito, V.; Dk,
N.; Dk, V., Copyright © 2021, Springer Science Business Media New York. Figure 2c reprinted
from Solid State Ionics, 272, William J. Bowman, Jiangtao Zhu, Renu Sharma, Peter A. Crozier,
Electrical conductivity, and grain boundary composition of Gd-doped and Gd/Pr co-doped ceria,
9–17, Copyright (2014), with permission from Elsevier. Figure 2d reprinted from Acta Materialia,
14, Komal Syed, Mingjie Xu, Kenta K. Ohtaki, David Kok, Keyur K. Karandikar, Olivia A. Graeve,
William J. Bowman, Martha L. Mecartney, Correlations of grain boundary segregation to sintering
techniques in a three-phase ceramic, 100890, Copyright (2020), with permission from Elsevier.

Crystals 2021, 11, 878 4 of 58
1.2.2. Coherent/Incoherent
When two solids share the same or very similar crystal structures, the atomic col-
umns will be continuous along the interface plane. In this case, the interface will have
minimum strain possible and is called coherent. If the mismatch strain associated with a
coherent interface is high enough, particularly in presence of large atomic misfits, the total
energy of the system increases until the surface is replaced with a more energetically fa-
vorable semi-coherent interface, where the excess energy is compensated by generation of
misfit dislocations. When atomic structures on both sides of the interface are completely
different, there is no possibility of a continuous lattice at the interface. Even when the
atomic structures match, the interatomic distances may differ by around 25% or more. In
both cases the interface is called incoherent [42,43].
1.3. Additional GB Terminology
1.3.1. GB Character
There are five macroscopic degrees of freedom to describe crystallographic space of
GBs, three of which determine relative rotation of two grains, including direction of a
rotation axis and rotation angle relative to the axis, and two of them define the interface
plane [44–46]. The GB character distribution (GBCD) function (Δ,  ) with misorienta-
tion Δ and plane normal  statistically defines the GB character in polycrystals [46]. Ad-
ditionally, considering atomic arrangements in crystals, three microscopic degrees of free-
dom are employed to describe grain translation at the in-planar boundary plane and at
the normal GB plane [47,48]. The GB energy and properties depend on the degrees of
freedom, and it is crucial to link GB energy/properties to degrees of freedom through GB
engineering and design [49,50]. GBs can be classified into symmetric and asymmetric GBs,
depending on whether the GB is a mirror plane (Figure 3a,b) [51]. Tilt, twist and mixed
GBs depend on relative position between the rotation axis and the GB plane. For instance,
the rotation axis of tilt GBs is normal to the GB plane (Figure 3a,b), while the rotation axis
of twist GBs is parallel to the GB plane (Figure 3d). The rotation axis of mixed GBs is
neither normal nor parallel to the GB plane (Figure 3c), but commonly they can be re-
solved into several symmetric special GBs, like symmetric twist boundaries, for analysis
[52]. Polycrystalline materials comprise grains in a wide range of crystallographic orien-
tations and therefore characters. Researchers have synthesized and studied bicrystals to
manually change the GB character and study its influence on properties [36,53,54]. We
will discuss experimental examples of bicrystals in the following sections.

Crystals 2021, 11, 878 5 of 58
Figure 3. Schematic of macroscopic GB geometry (a)symmetric tilt GB (STGB); (b) asymmetric tilt
GB (ATGB); (c) mixed GBs (MGB); (d) twist GBs (twin GB); (e) facet structure of (b); (f) facet struc-
ture of (c) [51]. Figure 3 reprinted from Progress in Materials Science, 98, Jian Han, Spencer L.
Thomas, David J. Srolovitz, Grain-boundary kinetics: A unified approach, 386–476, Copyright
(2018), with permission from Elsevier.
1.3.2. Low Angle GBs
GBs are categorized based on the orientations of two adjacent grains. The two grains
are related to one another by a rotation axis, a rotation angle and meet on a plane. When
the misorientation angle is less than 10–15°, the GB is called a low angle boundary [35,40].
Tilt and Twist boundaries are in this category and most real GBs are of a mixed type of tilt
and twist.
1.3.3. High Angle GBs
High angle GBs have misorientation angle is higher than 10–15° [35,40]. They are
more disordered and open in structure comparing to low angle GBs. They usually occur
in ceramic materials with two or more types of ions which result in perfect dislocations
on the oxygen sublattice, as well as stacking faults on anion sublattice [55].
1.3.4. Special, Twin, and Coincident Site Lattice (CSL) GBs
The term “special” GB has been used to describe GBs of unique structure, properties,
or both. Such special boundaries are often of interest in the field of GB engineering [56].
One example of a special GB is the coincidence-site lattice (CSL). The coincidence-site lat-
tice (CSL) is a well-known model for GBs that was primarily developed from studies of
metals by considering the relative misorientation of the adjoining grains. In CSL GBs, two
grains are in a relative orientation such that some atomic sites of one grain coincide with
the atomic sites of the other grain [35,56]. ∑ is defined as the degree of fit between the
structures of the two grains. The inverse fraction of coincidence sites, ∑−1, is an important
parameter to characterize CSL boundaries and a lower ∑ value is an indication of higher
order symmetry [40]. Coherent twin boundaries are unique CSL boundaries where the
structure mirrors across the GB plane.

Crystals 2021, 11, 878 6 of 58
1.3.5. Bicrystals
Bicrystals are composed of two specific single crystals with controllable interfacial
structure. Because many single crystals are and have been commercially available, bicrys-
tals are crucial tools for research on interfaces and have been extensively explored in un-
doped/doped single phase polycrystalline systems [6,7,54,57–69]. They are often used to
analyze individual interfaces at atomic scale, such as CSL boundaries and their relation to
properties [6,33]. While bicrystals can be used as models on for structure property analy-
sis, they do not fully represent GBs and heterointerfaces in polycrystalline materials,
which contain grains with many different orientations [57,60].
1.3.6. GB Complexions
Complexions can be defined as quasi-2D “phases” that primarily occur at interfaces
and surfaces [56,61]. These interfacial structures can be classified into different categories
depending on their composition, structure, and thickness. These categories are: (i)
clean/undoped GBs, (ii) sub-monolayer segregation, (iii) bilayer segregation, (iv) multi-
layer segregation, (v) nanoscale intergranular films, and (vi) wetting films [62].
1.4. How Do GBs and HIs Form in Polycrystalline Oxides?
The formation of GBs is accompanied by an increase in Gibbs free energy [40]. The
thermodynamic equilibrium is defined as the state with the lowest energy of the crystal,
where all GBs are eliminated. Kinetic and geometrical conditions prevent this from hap-
pening in reasonable times and therefore, a local equilibrium assumption is defined when
analyzing GBs [62]. HIs may exist under thermodynamic equilibrium as the result of a
phase transformation. A non-exhaustive list includes four common formation mecha-
nisms: (i) Polycrystalline ceramic sintering yields GBs and HIs; the relative stability of the
interface during densification and grain growth dictates which boundaries are likely to
form. (ii) Phase transformations, reactions and corrosion can cause the formation of HIs.
(iii) Coating processes where one material is applied to a substrate’s surface to yield a
thick layer (>1 μm) or thin film (<1 μm), yielding a HI between the materials. (iv) Thick
layer coating processes often involve a slurry (ceramic precursor particles suspended in
liquid), and include a sintering step, which creates GBs/HIs within the coating material
via (i). In thin films—produced by methods such as sputtering [63], molecular beam epi-
taxy [64], or pulsed laser deposition (PLD) [65–69], GBs/HIs can form within the thin film
during growth depending on growth conditions. This review is concerned only with
GBs/HIs in polycrystalline oxides formed via (i), (ii), and (iv), and ignores substrate-coat-
ing HIs formed via (iii), as there is a very substantial literature on HIs of type (iii) moti-
vated by research and development of oxide thin film and semiconductor devices.
1.5. How Are GBs and HIs Characterized?
As mentioned earlier, characterizing individual GBs and HIs is a considerable chal-
lenge due to the requisite structural and chemical sensitivity on the length of single unit
cells. Moreover, sampling a statistically relevant number of GBs/HIs often requires char-
acterizing many tens, hundreds, or thousands of individual GBs/HIs. The former chal-
lenge relates to the spatial resolution of characterization techniques, which limit the suit-
able techniques to electron microscopy (EM), atom probe tomography (APT), X-ray pho-
toelectron spectroscopy (XPS) [70] on intergranular fracture surfaces, electrochemical
strain microscopy (ESM) [71], scanning tunneling microscopy (STM) where the GB inter-
sects the surface, and atomic force microscopy (AFM) for studying GB energy using
grooves [72]. As useful as the current characterization techniques are, they are still suffer-
ing from a tradeoff between spatial resolution and analyzed sample volume within a rea-
sonable time. Among these techniques, EM techniques have for decades offered signifi-
cant contributions to interfacial characterization [73,74]. For example, TEM phase contrast
imaging was the key to the identification of ion-blocking amorphous phases at the GBs,

Crystals 2021, 11, 878 7 of 58
affecting oxygen ion conductivity [75,76]. STEM annular bright field (ABF) imaging up-
graded with the help of negative Cs STEM aberration correction was used to directly im-
age oxygen and oxygen vacancies. This information helps shed light on the underlying
mechanisms driving GB/HI defect chemistry, paving the way for modeling GBs and their
properties in future [8].
In addition to EM, APT is useful for quantitative probing of GBs/HIs, owing to its
sufficient spatial resolution and chemical sensitivity. APT is a combination of field ion
microscope and mass spectrometer, capable of accurately reconstructing the morphology
and chemical composition of the specimen in three dimensions at the nanoscale [77]. The
detection efficiency (i.e., the ratio of signals detected to signals produced) of APT is
around 30–80%, which is reasonable among other techniques (up to 80% at low count rates
for STEM-EELS and <1% for STEM-EDXS) [78].
The APT technique has been used to measure segregations of defects in oxide GBs in
addition to learning about charge density and electrostatic potential [79]. This technique
is particularly beneficial in ceramics including oxides due to the combination of surface
band bending under high electric field and/or surface defects, which make accumulation
of charge on the surface possible, and photons of lower energies can be absorbed [80].
However, data reconstruction can be a limiting factor, as it can potentially create artifacts
[81]. APT has been coupled with STEM imaging and spectroscopy to provide a comple-
mentary understanding of GB structure and chemistry in Hf and La doped-alumina and
Y-doped ceria [82,83]. Just like S/TEM, extensive optimization of specimens is necessary
to obtain reliable APT results [84]. In a work by Diercks, et al. [79], GB segregation was
observed in a 10% and 30% Nb-doped ceria (NDC10 and NDC30) using APT. APT is an
inherently destructive method, as illustrated by the overlay of TEM micrograph before
and after APT shown in Figure 4a. The Nd% map is shown in Figure 4b, where fraction
of Nd ions is illustrated for clarity. There is an evident increase in the Nd concentration at
the GB compared with the grain bulk. Figure 4c shows the same volume as in Figure 4b,
except regions containing at least 66 at% O are highlighted instead (the boundaries of
these regions are referred to as iso-concentration surfaces). This figure illustrates a change
in chemistry at the GB, this time a clear decrease in the O concentration.
Figure 4. (a) Overlay of transmission electron microscope images of an atom probe NDC10 speci-
men before and after atom probe analysis. A GB is observed in the before image. The after image
indicates that the atom probe analysis encompassed the entire GB region. Atom probe data recon-
structions showing (b) Nd, and (c) a 66 at% oxygen iso-concentration surface [79]. Reproduced
from [79] with permission from The Royal Society of Chemistry.

Crystals 2021, 11, 878 8 of 58
Quantitative 3-D maps of the elements as well as segregation of Nd, Al and Si cou-
pled with O depletion at the GB are presented in Figure 5. This information was used to
model the interfaces and correlate them with electrochemical impedance spectroscopy
data to yield 3-D potentials at the GBs. At 30% Nb, local gaps in the 3-D potentials suggest
conduction pathways through the potential barrier at the GB [79].
Figure 5. (a) Atom probe data reconstruction for Nd0.30Ce0.70O2 (NCD30) showing Nd ions and a 3-
D region of interest around the GB extracted for further analysis. (b) 2-D projections of the concen-
tration of each species in at% from the depicted region of interest. The GB is especially apparent in
the O deficiency and Nd enhancement. There is also evidence of Al and Si enrichment at the GB
[79]. Reproduced from [79] with permission from The Royal Society of Chemistry.
GBs in polycrystalline oxides have also been investigated using X-ray photoelectron
spectroscopy (XPS) [70], a widely used surface analytical method. Researchers used XPS
to analyze the fracture surfaces in Mg-doped alumina to assess Mg GB segregation. The
concentrations of Mg and Ca at and near the GB are measured and shown in Figure 6.
While Mg has much higher concentration in the bulk comparing to Ca, the segregation of
Ca to the GB is more significant.

Crystals 2021, 11, 878 9 of 58
Figure 6. (a) concentration of Mg and Ca (at%) as a function of distance (nm) from the GB [70], (b) Electrochemical strain
microscopy (EMS) detecting space charge in a polycrystalline Sm-doped ceria. ESM map of response magnitude for ±3V
perturbation, revealing increased response in the GB [71]. Figure 6a reprinted from Journal of the American Ceramic So-
ciety, 57, Taylor, R.I., Coad, J.P. and Brook, R.J., Grain Boundary Segregation in Al2O3, 539–540, Copyright (2006), with
permission from John Wiley and Sons. Figure 6b reprinted from Progress in Materials Science, 89, Giuliano Gregori, Ro-
traut Merkle, Joachim Maier, Ion conduction and redistribution at grain boundaries in oxide systems, 252–305, Copyright
(2017), with permission from Elsevier.
Scanning probe techniques such as ESM [71] and AFM measure local changes in the
surface topography of polycrystalline oxides and have proven useful for quantifying GB
energies and mapping functional properties. GB energy plays an important role in segre-
gation phenomena. GB segregation generally occurs to lower the energy of the GBs. Both
experiments [85] and simulation [86] studies have shown that GB energies can change
with temperature and have a strong effect on microstructure development. A common
scanning probe technique that allows measurement of relative GB energy (ϒGB/ϒS) is AFM.
When a solid surface is polished and etched, the interfaces are revealed due to preferential
removal of material near the GBs [72]. The geometry of the grooves formed by etching can
then characterized by AFM and used to measure average GB energy [87]. Many studies
have been conducted on single phase oxides for GB energy measurements [88–91].
2. Characterizing GBs and HIs in S/TEM
2.1. A Brief Introduction to TEM and STEM
The transmission electron microscope (TEM) and scanning TEM (STEM) are power-
ful tools for directly imaging atomic arrangements, and quantitative chemical analysis
down to the atomic scale. Books and review articles are written focusing on TEM [92–97],
STEM [98–101], and the various analytical techniques available in both [93,102–105]. A
basic introduction to these techniques as they pertain to GB/HI characterization in poly-
crystalline oxides is provided here. In general, TEMs can be configured to operate in mul-
tiple modes. In “TEM mode” (also called conventional TEM) a broad electron beam pro-
vides parallel illumination to the specimen. In “STEM mode” the electron beam is con-
verged to a point (or “probe”) by pre-specimen electromagnetic focusing lenses and
scanned in a raster pattern over the specimen.
Modern S/TEMs routinely demonstrate atomic resolution imaging and chemical
analysis in two dimensions, and more recently in three dimensions [79,80,106]. TEM de-
velopment dates to the era of de Broglie overcoming the resolution of a light microscope
by electrons [107]. Later, TEM was built and demonstrated by Max Knoll and Ernst Ruska
in 1933 [108]. The developments of spherical and chromatic aberration correctors for TEM
and STEM in 1998 [109] and 2002 [110], respectively, extended the resolution limit to sub-

Crystals 2021, 11, 878 10 of 58
Angstrom, creating novel opportunities for the accurate characterization of material de-
fects, such as GBs and His [111]. In addition, enhanced electron and spectroscopic signal
detection capabilities [112] have increased the ease of use by removing technical barriers
such as resolution and signal-to-noise-ratio improvement [113]. Today, a variety of anal-
ysis techniques are available to probe important characteristics and functionalities of or-
ganic and inorganic specimens in solid, liquid, and gas phases. This review, highlights
imaging, diffraction, and spectroscopy techniques suited to characterize interfaces within
polycrystalline ceramics, such as atomic resolution imaging, microstructure analyses (e.g.,
grain orientation mapping), and chemistry mapping (e.g., elemental concentration and
chemical segregation).
2.2. Imaging and Selected Area Electron Diffraction in the TEM
Due to the wavelike nature of electrons and periodic arrangement of atoms within a
crystalline specimen, electrons scattered from the specimen generate a diffraction pattern
that contains rich crystal structure information. From the position and symmetry of the
diffraction spots, it is possible to determine the crystal structure, unit-cell parameters, lat-
tice type, and defects such as twinning, while the intensities of the diffraction spots are
related to the arrangement of atoms within the unit cell [114]. Selected-area electron dif-
fraction (SAED) is the most used method to acquire this information, and it requires the
microscope to operate in “diffraction mode”. The ray diagram presented in Figure 7 ex-
plains the differences between diffraction and imaging modes in a TEM [92]. As in light
optics, ray diagrams are often used in TEM to show how electron lenses control the elec-
tron beam by depicting the path electrons take as they traverse the electron column. As
can be seen from Figure 7, the information contained in the diffraction pattern is also con-
tained in the image However, the TEM image displays this information in real space while
the diffraction pattern represents this information into reciprocal space, the two are re-
lated through a Fourier transform. Depending on the application, one or the other repre-
sentation may prove to be more useful.

Crystals 2021, 11, 878 11 of 58
Figure 7. Ray diagrams demonstrating two basic operation modes in a conventional TEM. In (a)
diffraction mode, the DP is projected onto the viewing screen and (b) image mode, projecting the
image onto the screen [92]. Switching between the modes requires changing the strength of the
intermediate lens. Figure 7 reprinted by permission from Springer Nature Customer Service Cen-
tre GmbH: Springer Nature, Springer eBook, Diffraction in TEM, David B. Williams, C. Barry
Carter, Copyright © 2021, Springer Science Business Media New York.
A SAED pattern is acquired by isolating a small area of the specimen for diffraction
analysis with an area-limiting aperture (usually 0.1 um or greater). In addition to crystal
structure determination, the SAED pattern also contains information about overall sample
crystallinity and can be used to establish orientation relationships between multiple dif-
fracting crystals (e.g., between precipitates and a host matrix, or at an interface like GB).
Further structural information is encoded into the diffraction pattern through the pres-
ence of features such as extra spots, the spot splitting, satellite spots, streaks and diffuse
scattering that can indicate the presence of superlattice [115], point defect ordering, or
extended defects [3], or can indicate particle shape [92]. Widely used techniques such as
bright-field (BF) and dark-field (DF) TEM imaging, demonstrated in Figure 8, rely on fil-
tering information from the SAED pattern to either exclude (BF-TEM) or preferentially
select (DF-TEM) diffracted beams for image formation, generating contrast based on
whether a particular region does or does not meet the Bragg condition such as examples
in [116]. In the case of apertureless imaging, the transmitted and diffracted intensities re-
combine, and the diffraction contrast is suppressed, while the BF and DF TEM images
show much better diffraction contrast due to the coexistence of large intensity in transmit-
ted beam and large loss of intensity in the diffracted beam, and vice versa. For example,
heavier phases tend to be darker in BF and brighter in the DF mode. While the BF-TEM

Crystals 2021, 11, 878 12 of 58
helps learn about the morphology of the sample, DF can provide information about nano-
crystal size and distribution in addition to defects such as dislocations, stacking faults and
twinning.
.
Figure 8. Ray diagrams demonstrating the formation of BF, DF, and centered DF images by selecting specific beams in the
diffraction plane with the objective aperture and reforming the filtered image at the imaging plane. (a) BF images are
produced by selecting the direct beam while both (b) DF and (c) centered DF images are formed by filtering out the direct
beam, using only one or more diffracted beams to form the image. Centered DF images are formed by tilting the incident
beam so that the diffracted beam aligns with and is centered on the optic axis, which reduces aberrations. The bottom row
shows the region of the diffraction pattern that is selected by the aperture in each case [92]. Figure 8 reprinted by permis-
sion from Springer Nature Customer Service Centre GmbH: Springer Nature, Springer eBook, Diffraction in TEM, David
B. Williams, C. Barry Carter, Copyright © 2021, Springer Science Business Media New York.
Another useful technique is weak beam imaging that allows the formation of diffrac-
tion contrast in BF and DF modes using a weakly excited beam. The weak beam dark-field
imaging (WBDF) approach is popular as it offers advantages over BF, such as having a
stronger contrast that is easier to understand using simple physical models. WBDF is rou-
tinely used for highlighting the detailed defect structure, and is hence very suitable for
lattice defect observation, including dislocations, stacking faults and grain boundaries
[117]. Furthermore, high resolution TEM (HRTEM) images demonstrate the atomic ar-
rangement at interfaces and enable observation of individual defects, which can be used
to analyze secondary phases or impurities and can probe local crystallinity (degree of
structural ordering) at the GBs [118]. One key consideration is that high resolution images
require complementary image simulation before the image contrast can be conclusively
assigned to an atomic arrangement.
2.3. Imaging Techniques in the STEM
In STEM mode, the electron beam is converged to form a probe and scanned across
the specimen [98,99,119]. Imaging in the STEM is done primarily using elastically and in-
elastically scattered electrons which have interacted with, and passed through a thinned

Crystals 2021, 11, 878 13 of 58
section of material, usually <300 nm. The scatter collection angle is critically important for
image interpretation. The main imaging modes in STEM are presented in Figure 9 and
explained below in the order of high to low scattering collection angles. They are often
collected simultaneously and can be coupled with EELS and EDXS analysis.
When high angle scattered electrons are collected which are incoherent, the image
would be a high-angle annular dark-field (HAADF) STEM image. The contrast is gener-
ated based on the efficiency of the Rutherford scattering process at that angle of collection,
and so is roughly proportional to the square of the atomic number. In addition, atomic
positions are very clear since there are no contrast reversals. This technique is the most
widely used but is not ideal for imaging light elements since the Rutherford scattering
processing is inefficient for light atoms with smaller atomic nuclei.
Medium and or low-angle annular dark-field (MAADF or LAADF) STEM is another
commonly used condition that collects lower angle coherent scatter. These images retain
some of the characteristics of HAADF images (e.g., no contrast inversion, interpretability
of atomic columns over a large focus range) and is also good at showing features that
diffract strongly (strained areas such as dislocations, nanosized coherent or semi-coherent
precipitates etc.).
Annular bright field (ABF) is when the inner part of the collection disc is collected
which is dominated by phase contrast (thickness and defocus dependent). This method
works very well with the aid of simulations and seems to show that the outer part of the
BF disc produces images that are more incoherent and reveal all atoms as dark shadows,
including light atoms such as oxygen [120] or hydrogen [121]. ABF pairs well with ana-
lytical STEM techniques where localized chemical information is desired, typically pre-
ferred to operate in STEM mode for chemical analysis.
Bright-field (BF) STEM is analogous to bright-field TEM images attained in C-TEM.
The angle of scatter is small, and the electrons are coherent. Just as in HRTEM, the contrast
is critically dependent on sample thickness, microscope defocus and can show contrast
inversions. The benefit of BF-STEM compared with HR-TEM is that it can often be col-
lected simultaneously with ADF-STEM images of varying collection conditions (HAADF,
MAADF/LAADF).
A comparison of STEM BF, STEM DF and HAADF STEM is given in Figure 10 for
Hexoloy SiC.
Figure 9. Schematic configuration of STEM detection set-up. (a) annular DF (ADF) detector under ABF, ADF and HAADF
conditions; and (b) circular BF detector under LABF, MABF, and HABF conditions. The camera length is shortened in the

Crystals 2021, 11, 878 14 of 58
order of ABF, ADF, HAADF in DF-STEM imaging, and LABF, MABF, HABF in BF-STEM imaging. (c) Schematic diagram
of double-detector STEM imging. The top annular detector and the bottom circular detector are placed with a separating
distance [117]. Figure 9 reprinted from Applied Physics Letters, 101, Yasutoshi Kotaka, Direct visualization method of the
atomic structure of light and heavy atoms with double-detector Cs-corrected scanning transmission electron microscopy,
133107, Copyright (2012), with permission from AIP Publishing.
Figure 10. Hexoloy SiC. (a) STEM BF, (b) STEM DF, and (c) STEM HAADF images showing a graphite inclusion (arrow)
containing a few smaller oxide inclusions. Figure 10 reprinted from [122] with permission.
2.4. Specimen Preparation and Requirements for S/TEM
As discussed in the introduction, S/TEM specimens are in the form of a thin film or
particles typically around <300 nanometers. There are several methods for preparing such
thin specimens in oxide; the most useful ones are briefly explained below.
2.4.1. Focused Ion Beam (FIB) Lift-out and Milling
A common way of preparing TEM specimens for materials—particularly ceramics
and oxides—is using focused ion beams (FIB) [123]. In this method, a thin lamella of a
specific thickness is etched and extracted from a specific area of the sample. First, a plati-
num or carbon layer is deposited on the surface to protect the material during ion milling.
Then, an ion beam is used to mill two parallel trenches on both sides of the sample (Figure
11). The FIB is usually equipped with an internal micromanipulator needle for extracting
the lamella after milling and transferring to a TEM gird. This process can take place exter-
nally using a fine electrostatically charged glass needle. The thin lamella extracted from
the bulk is then ready for final thinning and cleaning, before the TEM analysis [124]. To
prepare an electron beam sensitive specimen such as biological materials, cryogenic-FIB
(cryo-FIB) was developed to allow sample preparation at lower temperatures to reduce
redeposition, thermal damages and preserve materials structure for subsequent analyses.
Recently, cryo-FIB techniques have been extended to oxides for lithium-ion batteries, for
example to characterize lithium dendrites within a frozen liquid electrolyte [125,126].

Crystals 2021, 11, 878 15 of 58
Figure 11. Focused ion beam TEM specimen preparation (a) layout positioning of the ion beam, secondary electron (SE)
detector and the metallic source (b) milling trenches with low voltage and weak current. SEM image in secondary electron.
(c) welding platinum between the slice and the needle (micromanipulator) [124]. Reprinted from [124] with permission
from Springer Science & Business Media.
2.4.2. Mechanical Polishing
This conventional method is still widely used today due to its simplicity and ability
to create large areas of electron-transparent specimens. In most cases, mechanical polish-
ing is used to obtain a flat surface in bulk materials, compacts, thin films and even fine
embedded particles [124]. However, mechanical polishing needs to be followed by other
final thinning methods such as Ar ion milling [127] and twin-jet electropolishing [128] to
produce an electron transparent specimen. First, a small disc of ~3 mm in diameter (i.e.,
the standard size accepted by TEM specimen holders) is extracted from the material of
interest using a tool called a “dimpler”. The disc is mounted and polished to a thickness
of ~ 100 μm (perpendicular to the disc radius) using a combination of rotary polisher and
grinding wheel and turntable (Figure 12). This is followed by either twin-jet electropolish-
ing [128] or Ar ion-milling [127] to produce damage-free, vanishingly thin and electron
transparent regions. Twin-jet electropolishing (an extension of electropolishing) is a fast,
damage-free and cheap method for preparing TEM specimens from electrically conduc-
tive bulk materials such as metals and alloys. In this method the material is electrochem-
ically removed from both sides of a mechanically thinned disc until sample perforation
occurs [129].
Ion-milling involves bombarding the sample with relatively low energy (e.g., <1–2
keV) ions or neutral atoms to remove material from the surface until a very thin electron
transparent specimen is formed. Variables including accelerating voltage, ion energy and
incident angle of irradiation can be changed depending on the specimen requirements
and polishing rate. Ar is widely used in ion milling due to its inert and heavy nature,
though other ions are utilized. While ion milling is a common and useful technique, it is
good to be aware of its artifacts and possibilities of redepositing material from one side to
the other side of your specimen.

Crystals 2021, 11, 878 16 of 58
Figure 12. (a) Rotary grinder/polisher, (b) dimpling diagram [124]. Reprinted from [124] with permission from Springer
Science & Business Media.
2.4.3. Ultramicrotome
The ultramicrotome technique uses a sharp diamond or glass knife to produce a na-
noscale cross-section through manipulating six degrees of freedom (X-axis, Y-axis, rota-
tional angle and clearance angle of the knife holder, the pitch and rotation of the block
sample). Thereby, the knife-block relative position can be accurately regulated and ad-
justed according to the observation from the stereomicroscope attached at the ultramicro-
tome [130]. To obtain desired a TEM specimen, the sample requires to be sheared off by a
clean single-side razor blade to form a trimmed surface (generally <1 mm2). Then, the well
sheared sample goes through trimming and sectioning at high and/or low speeds, to ac-
quire ultra-thin slices (40–100 nm) which are easily transferred to TEM grids by a loop
[131]. Figure 13 presents a schematic of ultramicrotome and images showing the sample
cross sectioning. Due to the need to mechanically slice the specimen, it is not as practical
in hard materials, where slicing can potentially lead to breaking and failure.
Figure 13. (a) Schematic of ultramicrotome and (b) images illustrating the ultramicrotome cross-sectioning and slices (the cutting
direction is parallel to the desired surface and the slice is with nanoscale thickness and wide dimension) [132]. Figure 13a reprinted
from Journal of Power Sources, High-resolution chemical analysis on cycled LiFePO4 battery electrodes using energy-filtered trans-
mission electron microscopy, Joshua D.Sugar, Farid El Gabaly, William C. Chueh, Kyle R. Fenton, Tolek Tyliszczak, Paul G. Kotula,
Norman C. Bartelt, 246, 512–521, with permission from Elsevier (2013). Figure 13b reprinted from [130] Methods in Cell Biology,
152, Valentina Baena, Richard Lee Schalek, Jeff William Lichtman, Mark Terasaki, Serial-section electron microscopy using auto-
mated tape-collecting ultramicrotome (ATUM), 41–67, Copyright (2019), with permission from Elsevier.
2.4.4. Nanomilling
Nanomilling is an effective tool for final thinning and cleaning of TEM specimens
following other preparations by low voltage Ga+ FIB milling or mechanical polishing. Na-
nomilling uses a focused Ar beam to polish both sides of the TEM specimen [54,133]. The
spot size of an Ar beam is as small as 1 μm—smaller than the size of TEM lamella—so

Crystals 2021, 11, 878 17 of 58
there will be no milling or redeposition of the specimen grid material (e.g., copper or mo-
lybdenum) if correct milling procedures are used [134]. Nanomill operates at sub-kV ac-
celeration voltage and low current (<150 pA) under liquid N2 environment to reduce ion
beam damage and achieving effective thinning and cleaning [135]. Therefore, as a post-
FIB processing step, nanomilling can not only remove a surface amorphous layer and im-
planted Ga+ layer effectively but avoids additional damage and contamination [136]. The
final TEM specimen will be clean, thin (<100 nm), and flat [137,138]. Figure 14 shows a
clear comparison between two TEM specimens with and without the nanomilling as post-
FIB processing step. The specimen that went through 500 V nanomilling exhibited clear
lattice fringes even at the edge of the sample, while the specimen that went through 5 kV
FIB polishing showed obvious redeposition and damaged layers which impeded imaging
and energy spectrum analysis [134].
Figure 14. HRTEM images of La0.29Sr0.7Al0.65Ta0.35O3 substrate with nickelate and aluminate heterostructures below the
surface. (a) specimen polished with 5 kV FIB. (b) specimen went through nanomilling at 500 V as a final treatment step.
The lattice fringes are clear even at sample edge [134]. Reprinted from Ultramicroscopy, Practical aspects of the use of the
X2 holder for HRTEM-quality TEM sample preparation by FIB, Willem van Mierlo, Dorin Geiger, Alan Robins, Matthias
Stumpf, Mary Louise Ray, Paul Fischione, Ute Kaiser, 147, 149–155, Copyright (2014), with permission from Elsevier.
2.5. Specimen Contamination in S/TEM
Residual hydrocarbons in the microscope column coming from pump oil, specimen
air lock or even user contact with holder can be a resource of contamination [92,139]. Car-
bonaceous materials then deposit on the surface of perfectly thinned specimen and signif-
icantly affect high resolution imaging and microanalysis. This can be minimized by heat-
ing the specimen to >100 °C in a heating holder to burn off the carbonaceous contamina-
tion or cooling the specimen to liquid-N2 temperature in a cooling holder. Plasma cleaning
of specimen and holder prior to operation is a very successful approach in reducing con-
taminations [140].
2.6. Analytical Techniques in TEM/STEM
2.6.1. Energy Dispersive X-ray Spectroscopy (EDXS)
EDXS is a powerful qualitative and quantitative analytical technique available in
S/TEM used for chemical analysis of a desired specimen. The technique relies on the X-
rays generated in the microscope due to the interaction between the electron beam and
sample. Two different types of X-ray signals are generated: Bremsstrahlung X-rays and
Characteristic X-rays. Bremsstrahlung or braking radiation results from the deceleration
of primary beam electrons when deflected by the atomic nuclei in the sample. The

Crystals 2021, 11, 878 18 of 58
characteristic X-rays are a result of sample atoms being ionized by the primary electron
beam. Simply, a core-shell electron is excited and ejected from the atom, while an outer-
shell electron replaces it; this energy difference is released as an X-ray. These characteris-
tics X-rays are like fingerprints for each element and as a result, EDXS spectra are very
useful to determine elemental composition of any sample [102]. Figure 15 exhibits a sum-
marized illustration of X-ray signal generation, collection, data display and resulting
quantification [102]. The methods for quantitative analysis of EDXS data is discussed be-
low.
Figure 15. Schematic diagram showing the generation of X-rays in the sample, detection of the X-
rays by the detector (using generation of electron and hole pairs), pulse analysis and quantification
process of the spectrum to derive the composition of the sample [102]. Reprinted from Science of
microscopy, Scanning Transmission Electron Microscopy, Peter D. Nellist, New York, NY, 2007,
65–132. Springer, Copyright (2007), with per-mission from Springer Science+Business Media, LLC.
EDXS is a standard analytical technique in S/TEM for determining elemental concen-
trations in each sample. However, the quantification of low Z elements, including oxygen
and lighter elements, has been widely known to be challenging in EDSX. This has to do
with poor detection efficiency of these elements, severe absorption of X-rays within the
specimen etc. Thus, even with very thin TEM samples, absorption and fluorescence, cor-
rection cannot be ignored for lighter elements. However, with detector advancements in
recent years, particularly use of multiple silicon-drift detectors (SDDs), improved effi-
ciency in signal collection, etc., it has become possible to quantify elements like oxygen
with good accuracy [6,141].
2.6.2. Electron Energy-Loss Spectroscopy (EELS)
EELS is the analysis of the energy distribution of electrons that have lost energy
through inelastic scattering while passing through an electron transparent sample
[98,142–144]. There are two general types of EELS in S/TEM, both of which rely on the

Crystals 2021, 11, 878 19 of 58
facts that inelastically scattered electrons have kinetic energy corresponding to the energy
lost, and that they can be dispersed in space by passing through a magnetic field, i.e., an
energy-loss spectrometer. The first is energy-filtered imaging, known as energy-filtered
TEM (EFTEM) [145], in which electrons having lost a relatively narrow range of energy
(e.g., ~5eV) are filtered out of the transmitted beam and used to form a real space image.
The specific amounts of energy form a real space image. The second is parallel EELS,
where all electrons having lost a relatively large range of energy (e.g., <2000 eV) are dis-
persed onto a detector plane and typically analyzed further in terms of a one-dimensional
spectrum of counts versus energy loss. Both provide spatially resolved information about
chemical composition and electronic structure, though parallel EELS is currently more
commonly used because an entire EEL spectrum can be recorded at specified locations in
a S/TEM specimen, potentially providing a rich set of data containing information about
specimen thickness, chemical composition, and electronic structure of multiple elements
simultaneously.
EELS is now incorporated into S/TEMs that typically operate using high energy elec-
tron (60–300kV) sources. The interactions between electrons and matter are either elastic
(approximately zero energy loss) or inelastic (finite energy loss), providing various infor-
mation about the sample. Elastic scattering involves Coulomb interactions with the atomic
nuclei, while inelastic scattering refers to the interactions between a fast incident electron
and the atomic electrons that surround the nucleus. Inelastic processes can be understood
in terms energy band theory. Figure 16 shows the energy level diagram in a solid and
several signals generated after electron interaction and excitation in the specimen [142].
Figure 16. Energy-level diagram of a solid depicting K and L Shell and the valence band Ef is the Fermi level and Evac is
the vacuum level. The primary process of inner and outer shell excitation is shown on the left, the secondary processes
from phonons and electrons on the right. [142]. Figure 16 reprinted by permission from Springer Nature Customer Service
Centre GmbH: Springer Nature, Springer eBook, An Introduction to EELS, R.F. Egerton, Copyright © 2021, Springer Sci-
ence Business Media New York.
The energy loss spectrum consists of three main components: (i) the zero-loss or
“elastic” peak (ZLP), (ii) the low energy-loss region (∆E <50 eV) where optical information
can be determined by examining excitations into low-lying states above the Fermi energy
[98], and (iii) the high energy-loss or core-loss region, where compositional and electronic
information can be obtained from the electron beam interaction with deep core states (∆E
>50 eV). These regions are indicated in Figure 17 for YBa2Cu3O7.

Crystals 2021, 11, 878 20 of 58
Figure 17. Electron energy loss spectrum (EELS) core loss edges of YBa2Cu3O7 [142]. Figure 17 reprinted by permission
from Springer Nature Customer Service Centre GmbH: Springer Nature, Springer eBook, An Introduction to EELS, R.F.
Egerton, Copyright © 2021, Springer Science Business Media New York.
Electron energy-loss spectrum of YBa2Cu3O7 is a high-temperature superconductor.
The electron intensity is depicted on a logarithmic scale, showing zero-loss, plasmon
peaks and ionization edges arising from each element [142,146].
The ZLP represents electrons that are transmitted without losing measurable energy,
including electrons scattered elastically and those that excite phonon modes, for which
the energy loss is less than the experimental energy resolution (in conventional EELS sys-
tems). While the ZLP does not contain spectroscopic information about the specimen, it is
useful for energy calibration of the loss spectrum, deconvolution of plural scattering in
thicker specimens and determining local thickness [145,147]. The width of the ZLP is re-
lated to the energy spread of the incident beam emitted from the electron source and is
typically 0.25–1eV [148,149]. These spreads are sufficient for applications such as core-loss
elemental mapping, bulk plasmon analysis and analyses of wide-band gap materials but
can mask multiple narrow excitations such as surface plasmons, near-edge fine structures,
or vibrational excitations at very low losses. The tails of the ZLP ultimately define the
spectral resolution and can obscure the optical excitations [150–155]. Spectral resolution
can be improved to less than 10 meV in advanced monochromated electron microscopes
that include specialized equipment to reduce the electron energy spread of incident beam
[148]. In a typical monochromator, the electron beam is dispersed by energy (like a pre-
specimen EEL spectrometer) and a narrow band of the dispersed beam is selected by an
energy-selecting slit which defines the energy spread of the electron beam, which is inci-
dent on the specimen. This technique has proven to be useful in high spatial resolution
vibrational and low-loss spectroscopy in the electron microscopy, with monochromated
STEMs having now provided impressive results for several years [92,149,154–159].
The lower loss energies (<50 eV) are related to the formation and destruction of pho-
nons and electrons of the outermost orbitals, including both single and collective

Crystals 2021, 11, 878 21 of 58
collections, which are plasmons and electrons transmitting from valence band to empty
conduction band, allowing band gap energy measurements (Figure 18) [160,161]. The
most common application of energy losses in this low range is the measurement of sample
thickness, simply by comparing zero loss intensity with inelastic scattering intensity.
Figure 18. Valence EELS and Al-L23 for bulk and GB. Difference between bang gap in the bulk and at the GB [161]. Re-
printed with permission from Copyright © 2021 American Chemical Society.
The high-loss electrons are related to the core binding energy, and therefore provide
both qualitative and quantitative chemical information, as the intensity of each edge cor-
responds to the quantity of electrons generating a specific energy. In addition to the core-
loss edges, the first 30–40 eV beyond each edge is related to the change in density of un-
occupied states, which affects the electron loss near edge structure (ELNES), or fine struc-
ture [162–164].
The detection efficiency of EELS is generally much higher (up to 80% at low count
rates) than EDXS systems (<1%) [78]. This is because a large proportion of inelastically
scattered electrons are forward scattered into the EEL spectrometer, whereas X-ray emis-
sion is isotropic and only a small fraction is intercepted by finite-sized detectors; in recent
years, additional (e.g., 4) X-ray detectors have been introduced to the S/TEM column to
enhance efficiency. Remember that the detector efficiency will be affected by the specimen
thickness, the probe current, and the solid angle of collection of the detector. In EELS, a
thin specimen is preferred as more electrons transmit the specimen and can pass through
the spectrometer entrance aperture, which can be expanded in diameter to increase col-
lection efficiency.
The first step for analyzing EELS data is to subtract the background and any potential
plural scattering effects that occur in a thick specimen when a significant fraction of inci-
dent electrons passing through a sample scatter inelastically more than once [158,165].
When the background is removed, intensities of elements can be converted to concentra-
tions.
STEM-EELS analysis on an interface in a Y-doped ZrO2 bicrystal is shown in Figure
19 [166]. In this case, STEM was used to observe the interface (a) and EELS was used to
generate signal maps of each element relative to the bulk (b–d). The line profile of

Crystals 2021, 11, 878 22 of 58
elemental signals is shown in (e) that shows depletion of 0.5–1 nm for O and Zr and en-
richment of Y at this synthetic interface. Therefore, the combination of STEM with EELS
is a powerful method, allowing the investigation of defect chemistry at atomic level, pav-
ing the way for understanding the relationship between GB and ionic conductivity [166].
Figure 19. STEM-EELS of a YZS bicrystal. (a) Z-contrast image at the GB obtained in Nion UltraSTEM 200 operated at 200
KeV, the yellow dashed box marks the area where an EELS spectrum was acquired. (b–d) atomic resolution, integrated
signal maps of Zr-L23, Y-L23 and O-K edges, respectively, normalized to the bulk concentration. (e) normalized integrated
signal profiles based on the analysis of the O-K, Zr and Y-L23 edges in (d). solid symbols result from quantification per-
formed on a spectrum image including the O-K and Zr and Y-M23 edges instead. The black line represents the stoichio-
metric O content that would be expected from measured Zr and Y signals. (f). averaged Y-L23, and Zr-L23 EELS spectra
from the bulk (black) and the dislocation core (red). Figure 19 reprinted from [166] under the terms of the Creative Com-
mons CC BY license.
2.7. Other Emerging Techniques
2.7.1. Cathodoluminescence (CL) in S/TEM
In an electron microscope, the energy of electrons is transferred to specimen, bringing
it to an exited state. As the excited electrons relax to the ground state, they release both
high and low energy photons. The higher energy core-shell transitions generate X-ray
emission, while the lower energy photons that are associated with valence band energy
states and transitions between free electrons, generate cathodoluminescence (CL) which
has applications in S/TEM (CL-S/TEM), particularly in the semiconductors, where an in-
cident electron beam generates electron-hole pairs during the imaging mode. It is possible
to separate electron hole pairs using an external voltage to generate a signal of electron
charge pulse onto the STEM screen. The strength or weakness of the signal is related to

Crystals 2021, 11, 878 23 of 58
the status of electron-hole pairs. When electron-hole pairs are separated, the signal is
strong and when they recombine in presence of defects such as dislocations, the signal is
weak. If the electron-hole recombines, they give off visible light in form of a spectrum
which contains information about band gaps and dopants of the system. According to a
recent study [167], it is impossible to interpret the emitted light spectrum to the transitions
excited by the electrons directly. This is due to the existence of Cherenkov effect [168] in
addition to ratio-induced transitions in the spectra. Therefore, the spectra need to be cor-
rected for each sample shape, geometry, thickness, and beam energy.
2.7.2. Electron Beam Induced Current (EBIC-S/TEM)
In EBIC, high energy electrons induce a current that is scanned across the specimen
surface and variations in the EBIC induced current are used to map the electronic activity
of the specimen. In this technique, the specimen needs to be coated to create ohmic con-
tacts for applying an external voltage. EBIC has been used to identify defects such as GBs
[169] or minority carrier properties [170] in semiconductors. It has also proven useful in
characterizing oxides such as GBs and structure of zinc oxide varistors [171–173].
2.8. Scanning Electron Diffraction Techniques
2.8.1. Precession-Electron Nano Diffraction (PEND)
PEND was introduced in 1944 by Vincent and Midgley [165]. This technique has be-
come a major crystallographic investigation tool to identify crystal structure and relevant
parameters down to the nanometer fine scale. It is also used to map the local orientation
to investigate crystal texture, rotation and strain and virtual STEM imaging at the na-
nometer scale, in connection with the scanning system of the microscope [166]. In this
system, the incident electron beam is tilted and precessed using the pre-specimen optics
to form a conical electron probe. The result of this process is a higher number of spots in
the diffraction patterns acquired and integrated over the precession cycle (Figure 20).
Compared to the case for an unprecessed beam, these PEND reflections have intensities
that are determined by integrating through the Bragg condition of each reflection. There-
fore, PEND enables better indexing of high order diffraction patterns, allowing comple-
mentary studies to X-ray diffraction, especially in multi-phase samples, disordered crys-
tals or other systems that are not suitable for conventional X-ray crystallographic analysis
[174].

Crystals 2021, 11, 878 24 of 58
Figure 20. (a) A schematic ray diagram for precession electron diffraction (PED), illustrating the rocking/de-rocking action
of the beam before and after the specimen. (b–e) Illustrations of how precession alters the recorded diffracted intensities,
here from the [001] zone axis of Er2Ge2O7: (b) without precession, (c) with a precession angle of 20 mrad, (d) with a
precession angle of 47 mrad. The pattern of diffracted intensities seen in (d) is similar to that seen in the kinematic simu-
lation shown in part (e). Figure 20 reprinted from [174] under the terms of the Creative Commons CC BY license.
PEND is now a powerful tool for looking at the twinning behavior in hexagonal al-
loys [149] and phase changes in oxide structures like Ni-YSZ electrodes during reduction
and oxidation [146]. The sensitivity of this technique to local orientations has been the key
to strain analysis in semiconductors [157].
2.8.2. D-STEM
In this method, the probe is scanned over the specimen in a two-dimensional (2D)
array and at each probe position a 2D diffraction pattern is recorded, thus generating 4D
data that can be further analyzed. This technique enables orientation, strain, and phase
mapping down to the atomic scale, as well as virtual STEM imaging. Due to the volume
of obtained data during 4S-STEM, data analysis can be complicated via requiring capable
algorithms and simulation, though these are expected to become more automated over
the next few years. In addition to providing spatially resolved diffraction patterns, 4D-
STEM data can also be used for STEM imaging using “virtual” detectors. This is accom-
plished in post processing by integrating the scattered intensity within arbitrary segments
of the diffraction pattern at each probe position, thus recreating a STEM image [175]. Fig-
ure 21 [175] shows imaging of Y-doped ZrO2 using 4D-STEM. Two methods of image in-
terpretation are shown: (i) generation of a virtual detector and (ii) selecting diffraction
patterns from various regions of constant contrast over the probe position in the real
space. As all the diffraction patterns for each position in the raster scan are stored, a com-
prehensive analysis of the GB crystallography in addition to local strain mapping can be
performed (Figure 21a). Additionally, virtual STEM images can be created, such as BF,
ADF and HAADF can be reconstructed from a single dataset showing in a by selecting
angular ranges of interest (Figure 21b).

Crystals 2021, 11, 878 25 of 58
Figure 21. (a) Grain map from a Y-doped ZrO2 thin-specimen. (b) Series of STEM images, reconstructed by selecting the following
angles from the STEM dataset in a. below 18 mrad, 37–81 mrad and 18–37 mrad [176]. Figure 21 reprinted by permission from
Cambridge University Press, Development of Diffraction Imaging for Orientation Analysis of Grains in Scanning Transmission
Electron Microscopy, M Watanabe and D Williams, Copyright © 2021.
For this purpose, computer image processing methods are used to classify the crystal
orientations automatically, which is called “automated crystal orientation mapping”
(ACOM). This method allows tracking orientation changes of hundreds of grains simul-
taneously and has been used significantly for in-situ mechanical testing measurements in
nanocrystalline metals [177–179]. In a work by Garner, et al. [180], the orientation map of
monoclinic zirconia reveals a fine level of detail with grains as small as 5 nm in the outer
region. This was also done for tetragonal zirconia grains and therefore can be used for
mapping any type of phase distribution [180].
2.9. In-Situ S/TEM
In-situ S/TEM is a technique where an electron microscope is used to watch and cap-
ture specimen’s response to one or more external stimuli in real time. Operando S/TEM is
an extension of in-situ analysis whereby specimen properties are measured in addition to
simultaneous direct observation, thus shedding light on the specimen functionality. Both
techniques have emerged as a powerful field of nanomaterials characterization. Now, pre-
viously mysterious reactions and processes can be dynamically monitored under realistic
or close-to-realistic conditions. The past decades have witnessed continuous development
in both in-situ technology and the applications in a wide field of research, particularly in
catalysis, energy storage and conversion [181,182], gas-solid interactions, semiconductor
growth, nanomechanical testing, electrical properties and phenomena, and corrosion
[183] (Figure 22).

Crystals 2021, 11, 878 26 of 58
Figure 22. In-situ S/TEM involves direct specimen observation while applying stimuli, and oper-
ando S/TEM includes measuring a performance metric of the material [184]. Reprinted from Con-
trolled Atmosphere Transmission Electron Microscopy, Spectroscopy of Solids, Gases, and Liquids
in the ETEM, Crozier, Peter A., and Benjamin K. Miller, 2016, 95–141, Copyright (2016), with per-
mission from Springer International Publishing Switzerland 2016.
2.10. Electron Beam Damage
The bombardment of thin TEM specimens by intense electron beams can result in
irreversible atomic displacements and chemical reactions generally termed “electron
beam damage”. This can particularly affect any quantitative analysis performed on sam-
ples exposed to such intense, energetic electron beams [185,186]. The beam damage in
S/TEMs can be a result of inelastic scattering caused due to the interaction between the
energetic incident electrons and sample atoms/electrons. This beam damage can thus re-
sult in various phenomena in TEM samples including mass loss, phase decomposition,
crystallization, amorphization, atom transport, etc. [185].
Two main mechanisms of radiation damage in TEMs are knock-on damage (elastic
beam-atom collision) and radiolysis (ionization damage), or a combination of both
[185,186]. In knock-on beam damage, the sample-beam interactions can cause preferential
surface sputtering in oxides, while radiolysis can generate point defects and break chem-
ical bonds, etc. [185] While radiolysis mainly occurs in insulating oxides, knock-on dam-
age can occur in both conducting and insulating oxides. Furthermore, knock-on damage
is significant during prolonged exposure to e-beam and can become predominant at
higher accelerating voltages ≥200 kV. Other mechanisms include thermal effects for pho-
non excitation and electric-field-driven migration [185]. Jiang [185] reviewed that while
majority of the observed TEM/STEM beam damage in reported studies have been inter-
preted as knock-on damage or radiolysis, most observations appeared to involve collec-
tive atomic displacements in which atoms of the same element move together in synchro-
nization.
The electron beam damage phenomena have been studied extensively in different
oxides, both for bulk and GB characterization, such as spinel [187–189], alumina [190,191],
perovskites [192–194], etc. In oxides, electron beam damage phenomena are often notice-
able when using aberration-correction STEMs, which by design provide very small and
intense. With EDXS particularly, longer acquisition times are preferred for better counting
statistics, but this increases the likelihood of e-beam damage, as well as contamination

Crystals 2021, 11, 878 27 of 58
[195]. This makes it difficult to interpret any chemical information obtained from an area
damaged by e-beam. Electron beam damage was shown to influence the oxidation state
of cations such as Ce4+ and reduce it to Ce3+, particularly near GBs, providing a methodol-
ogy for further investigating oxygen-vacancy concentrations in solid oxide solutions [196].
Figure 23 shows the Ce3+ fraction across two GBs in Gd- and Fe-doped ceria. In both plots,
the bulk contains larger than expected amounts of Ce3+ due to the e-beam damage (30%
and 20%, respectively). The Ce3+ drops to zero or negative values, with a slight rise at the
GB plane followed by another drop near the GB on the other side. After that, Ce3+ increases
again to high values in the bulk. It is noteworthy that the slight increase of Ce3+ at the GB
can be due to the beam damage, intrinsic boundary properties or both.
Figure23. Plots of the Ce3+ fraction from STEM-EELS line profiles across grain boundaries in cerium oxide doped with Gd
and Fe with around (a) 30% and (b) 20% Ce3+ in the bulk [196]. Figure 23 reprinted from Journal of the European Ceramic
Society, 34, J.P. Winterstein and C.B. Carter, Electron-beam damage and point defects near grain boundaries in cerium
oxide, 3007–3018, Copyright (2014), with permission from Elsevier.
3. Applications of TEM to Oxide GBs and HIs Characterization
3.1. Imaging Atomic and Nano Structures
Atomic and nanoscale imaging by HRTEM is a common approach to assess GBs in
many oxide systems, such as alumina [197–200], zirconia [74,201,202], spinel [203,204],
ceria [76,205], etc. [206–208] In addition, aberration-corrected TEM is widely used in ox-
ides [27,209–211]. HRTEM micrographs and schematic models of GBs in a rare-earth (RE)
doped alumina and a pure alumina are shown in Figure 24 [195]. Lattice fringes joining at
the interface are observed in the RE-doped GB. This can indicate relaxation in the crystal-
lographic structure near this GB and is consistent with observations of interface broaden-
ing in RE-doped GB, which itself is consistent with observations of interface broadening
in RE-doped GBs and bicrystals compared to pure ones. It is concluded from the micro-
graphs that the segregants are confined in a ~1 nm region at the GB and that the GB struc-
ture is sensitive to GB type/character.

Crystals 2021, 11, 878 28 of 58
Figure 24. HRTEM micrographs of (a) un-doped and (b) La-doped alumina GBs viewed edge on.
Schematic models of (c) un-doped and (d) RE doped GB in alumina [195]. Note the wider GB re-
gion in RE doped alumina shown in (b) and (d). Figure 24 reprinted by permission from Springer
Nature Customer Service Centre GmbH: Springer Nature, Journal of Materials Science, Characteri-
sation of fine-grained oxide ceramics, West, G.D.; Perkins, J.M.; Lewis, M.H., Copyright © 2021.
Spherical aberration (Cs)-corrected scanning transmission electron microscopy
(STEM), with a sub-angstrom-sized electron probe can is another powerful technique for
understanding structures of interfaces and segregation of dopants [212–215]. In a study
by Schustertisch, et al. [215], HR-STEM, complementary EELS analysis and density func-
tional theory (DFT) were used to accurately determine GBs structures in TiO2 at atomic
scale (Figure 25). The existence of a unique nanoscale phase like bulk anatase TiO2 was
discovered in a fabricated Σ13 (221) [11
0] rutile TiO2 bicrystals. This work show that it is
possible to embed regions of TiO2 anatase in rutile TiO2 structure by GB fabrication and
design. In other words, the power of GB analysis and engineering in oxides allows for
further tuning the GB structure and properties in such materials.

Crystals 2021, 11, 878 29 of 58
Figure 25. Atomic resolution HAADF, ABF STEM images and the DFT atomic structure of the 13 bicrystal along the (a–
c) [11
0] and (d–f) [114
], respectively. Image simulations based on the atomic structure of rutile are shown as insets in the
experimental HAADF and ABF STEM images. Gray and red spheres represent Ti and O atoms, respectively. Columns of
Ti + O, Ti, and O atoms are marked by solid green circles, dashed purple circles, and blue ellipsoids, respectively [216].
Figure 25 reprinted from https://pubs.acs.org/doi/full/10.1021/acs.nanolett.0c04564 (accessed on 27 May 2021) under the
terms of the Creative Commons CC BY license. Further permissions related to the material excerpted should be directed
to the ACS.
Most ceramics are composed of a mixture heavy and light element, the latter of which
are hard to detect. Annular bright-field (ABF)-STEM combined with HAADF imaging is
a powerful tool for imaging both light and heavy elements. Such studies shed light on the
origin of structure property relationships in oxides and improve even more through the
incorporation of theoretical calculations [122,217]. This technique was proved to be useful
in imaging cation and oxygen sublattices in a (210) Σ5 CeO2 GB with a tilt axis of [001]
(Figure 26) [33]. In addition, the oxygen non-stoichiometry used to reveal the role of the
oxygen vacancies on the GB atomic structure. Modeling of the GB structure in this case
was done using a static lattice calculation with the GULP program [218]. Figure 26a and
(c) show HAADF and ABF images taken from this GB, along with simulated HAADF and
ABF images of the following in (b) and (d).
The bright spots in the HAADF image correspond to Ce column locations and the
black and gray spots in ABF correspond to Ce and O column locations, respectively. The
presence of O columns in the GB shown by the gray contrast is weaker compared to the
O columns in the bulk. This is reasonable due to lower density of O column in GB com-
pared to the bulk region.

Crystals 2021, 11, 878 30 of 58
Figure 26. (a) HAADF, (b) simulated HAADF, (c) ABF image and (d) simulated ABF of a [001] (210) Σ5 GB in a CeO2 thin film. Sim-
ulated images are from a non-stoichiometric GB model structure. The structural units of each boundary are indicated by the poly-
gons. A noise-reduction procedure was applied to the ABF image by a background subtraction filter [34]. The structural units of
each boundary are indicated by polygons. Figure 26 reprinted with permission from Copyright © 2021, Oxford University Press.
3.2. Measuring GB Character
To determine the GB character in oxides, grain orientation mapping with a spatially
resolved transmission electron diffraction techniques are used. In cases where the charac-
ter of many GBs is of interest, researchers measure the GB misorientation angle distribu-
tion, which indicates the fractional length of boundaries with a given misorientation angle
between adjacent grains. Such technique can employ PEND, 4D-STEM, or atomic resolu-
tion imaging to quantify GB character and orientation by automated acquisition and in-
dexing of the diffraction data sets [7,147,219]. Figure 27 shows a of grain orientation over-
lay map on the BF-TEM image in polycrystalline Pr and Gd doped-ceria. Each grain is
colored based on the crystallographic direction oriented parallel to the electron beam di-
rection (see the color-coded stereographic inset). The misorientation angle distribution
measured agrees well with polycrystalline cubic structures full of randomly oriented
grains [220].

Crystals 2021, 11, 878 31 of 58
Figure 27. (a) Schematics representation of PEND in a TEM that enables grain and GB orientation mapping via precession
of the electron beam. (b) Bright-field TEM image of Pr and Gd-co doped ceria (PGCO) specimen with overlaid inverse
pole figures, grain orientation is indicated via color code corresponding to inset stereographic triangle. (c) Misorientation
angle distribution and Mackenzian distribution [7]. Reprinted from Copyright © 2021 American Chemical Society.
3.3. Chemical Analysis by EDXS and EELS
Analytical electron microscopy (AEM) techniques such as EDXS and EELS are pow-
erful methods for qualitative and quantitative chemical and electronic structural analysis
of the interfaces in oxides down to atomic scale. Modern AEMs readily provide spatially
resolved spectroscopic data from points, lines, and 2D areas (also called “maps”). These
methods have been applied readily for several decades, as described in books and reviews
[99,103–105,142]. STEM-EDXS is used to characterize GB and HI segregations in YSZ-
Al2O3-MgAl2O4 ceramic oxides (Figure 28) and correlate them with thermal properties of
these composites [41].

Crystals 2021, 11, 878 32 of 58
Figure 28. STEM HAADF images and EDXS maps for six (representative) interfaces in a flash-sintered MgAl2O4-Al2O3-
YSZ ceramic. Mg Kα, Al Kα, Y Kα, β, Zr Kα, β and O Kα elemental EDXS maps (net counts) and their corresponding intensity
profiles (net counts normalized by the total counts) are shown. (a) Alumina-alumina GB shows depletion of Al and segre-
gation of Y/Zr, (b) spinel-spinel GB shows depletion of Mg and segregation of Y/Zr, (c) YSZ-YSZ GB shows depletion of
Zr and segregation of Al, (d) spinel-alumina GB shows segregation of Y/Zr, and (e) YSZ-spinel and (f) YSZ-alumina inter-
faces do not show any segregation. Scale bar is 2nm [41]. Figure 28 reprinted from Acta Materialia, 14, Komal Syed, Mingjie
Xu, Kenta K. Ohtaki, David Kok, Keyur K. Karandikar, Olivia A. Graeve, William J. Bowman, Martha L. Mecartney, Cor-
relations of grain boundary segregation to sintering techniques in a three-phase ceramic, 100890, Copyright (2020), with
permission from Elsevier.

Crystals 2021, 11, 878 33 of 58
In a work by Ikuhara et al. [211], STEM-EDXS was used to study segregation behavior
in YSZ bicrystal at atomic scale. Figure 29 represents EDXS maps and line profiles of Zr
and Y, respectively. From this data, chemically ordered segregation around the GB core is
detected. Y atoms are strongly segregated to the immediate left and right planes of the
cation mirror plane (Mc in Figures 29c, d), whereas Y is depleted in the mirror plane and
Zr tends to segregate on the mirror plane (Figure 29b). Such detailed observations in GBs
of oxides can lead to a better understanding about defect chemistry and related properties.
In this case, the interactions between Y and oxygen vacancies govern segregation behavior
by inducing a phase transition near the GB. Such characterization method capable of
atomically mapping elements enables GB design at atomic scale, paving new avenues to-
wards controlling and optimizing properties of oxides, particularly when supported by
theoretical calculations such as Monte Carlo (MC), DFT, or phase field simulations
[18,22,221].
Figure 29. Energy-dispersive X-ray spectroscopy maps and intensity profiles. (a), (b) EDXS elemental maps for a. Zr K and b. Y K
and (c), (d). Intensity profiles by summing the X-ray counts in the maps in the direction parallel to the GB for (c) Zr-K and b. Y-K
(normalized to the total Zr and Y counts). Reprinted from [212] under the terms of the Creative Commons CC BY license.
Chemical doping is fundamental method of control and design GB functionality, and
dopants are often added to ceramics improve sintering and control grain growth. Dopants
or unwanted impurities in the system can segregate and modify GB properties. For in-
stance, intergranular glassy phases are often observed in sintered ceramics [74,200]. It was
elucidated in a work by Ishihara et al. [222], that small differences in the amounts of Si
impurity led to distinct GB atomic structures in a Ti-doped Σ13 α-Al2O3 GB (Figure 30). In
GB type-I, Ti3+ is replacing Al3+ while in type-II, Ti3+ oxidizes to Ti4+ and the Al vacancies
were introduced to the GB to compensate for that leading to three-times larger structure
units (oxidation states confirmed by EELS).

Crystals 2021, 11, 878 34 of 58
Figure 30. ADF-STEM, ABF-STEM and EDXS count maps of Al-K, Ti-K, and Si-K edges, respectively, obtained from (a)
type-I and (b) type-II Ti-doped Σ13 α-Al2O3 GBs viewed along the [12
10] direction. The white boxes on the ADF-STEM
image indicates the size of the unit structure of pristine GB. The arrows indicate the positions of brighter Z-contrast in the
GB core, and corresponding locations are also indicated in respective images. All the scale bars are 1 nm. Reprinted from
[222] under the terms of the Creative Commons CC BY license.
EELS is regularly employed in both TEM and STEM, and in addition to elemental
concentrations can provide information about the electronic structure, bonding environ-
ment, chemistry, and oxidation states of atoms and ions at interfaces. While application
of EELS can be very similar to EDXS, it acquires additional information mentioned, has
higher detection efficiency, making it more promising for analysis and quantifications in
e-beam sensitive materials and light weight elements. In a study by Saito et al. [223], seg-
regation behavior of impurities to an MgO bicrystal was investigated. Some impurities
tend to segregate to GB and interact with native structural defects, dominating the struc-
tures of the GB and hence determining many of its properties. A case study of the Cliff-
Lorimer method for quantitative GB analysis in oxides is a work by Bowman et al. [8] that
analyzes the cation concentrations at the GBs of a polycrystalline, oxygen conducting ma-
terial, Pr0.04Gd0.11Ce0.85O2−δ. Using the quantitative EELS analysis along with other charac-
terization techniques, the authors concluded that the boundaries with higher solute con-
centrations have lower activation energy of ion conduction across GBs.
Yang et al. [224] published a study on GB segregation in Mn-doped SrTiO3. This ma-
terial has promising magnetic and electric properties, which can be improved by engi-
neering the GBs. It was found that Mn2+ segregates inside GBs both in Sr and interstitial
sites, while Mn4+ is found to substitute Ti in bulk SrTiO3. Figure 31 represents the HAADF
image, EELS elemental maps of (a) Sr, (b) Ti and (c) Mn. The line profile in (e) is extracted
by averaging the compositions across the images shows an enrichment in Mn and Sr is
accompanied by Sr deficiency at GB core. It is evident that Sr is roughly 50 at% away from
the GB while the sum of Ti and Mn is also roughly 50 at%. This would suggest that Mn
substitutes for Ti and should, therefore, have a +4 charge state [224]. This interface was
also investigated by EDXS (Figure 31f,g). It is evident that there is significant difference
between qualitative elemental maps (b), (d) and quantitative maps (e), (g) which proves
the importance of quantitative analysis.

Crystals 2021, 11, 878 35 of 58
Figure 31. Atomic resolution mapping of the Mn-doped SrTiO3 GB. EELS compositional maps of (a) Sr, (b) Ti, (c) Mn, and
(d) composite maps were obtained from the MSA with RGB false colors applied. The scale bar is 0.5 nm. A few nm away
from the GB core, Mn has a much lower but detectable concentration, and substitutes the Ti site as indicated by the coin-
cident dashed circles in (b) and (c). The relative concentration profile (at%) across the GB in (e) is obtained by concerting
the vertically summed up EELS signal of the two-dimensional map using the Cliff–Lorimer factors of a 200 kV microscope.
Atomic EDXS mapping of the Mn-doped SrTiO3 GB. (f) shows count maps of Sr, Ti and Mn, respectively. (g) shows the
atomic fraction of elements normalized to1(or 100%) using the Cliff-Lorimer factors. As seen clearly in these maps, Ti
shows reduced counts but increased concentration (or atomic fraction) as the concentration of Sr drops inside the GB.
Reprinted from [224] under the terms of the Creative Commons CC BY license.
3.4. In situ S/TEM of GBs and HIs
In-situ mechanical testing inside electron microscopes has been an effective approach
to investigate interfacial deformation and fracture modes in ceramics. Among different
techniques, in situ nanopillar compression is widely used to explore the mechanical be-
havior and deformation mechanisms, including dislocation motion, stacking fault for-
mation and GB sliding, in ceramic materials at small scale [225,226]. Dynamic interactions
between individual dislocations and well-controlled GBs in SrTiO3 can be directly ob-
served under TEM during nanoindentation [227]. The dislocation impediment effect in-
duced by high-angle or low-angle GBs was found to be related to both the geometric and
structural stabilization effects of the GBs [227]. In another example, Kondo, et al. [228]
conducted nanoindentation at a Zr-segregated ∑13 GB in Al2O3 and observed the crack
nucleation and propagation along the GB inside TEM. Post-mortem HAADF STEM re-
vealed that atomic-scale cleavage path is along the center layer of the triple Zr-rich layers,
pointing to the importance of local oxygen coordination at the dopant-enriched GB core
in determining the fracture path and therefore, toughness, of oxide ceramics. Combining
localized laser heating inside TEM, in-situ mechanical testing at ultra-high temperatures

Crystals 2021, 11, 878 36 of 58
up to the melting point of ceramics could be realized. Using this approach, Grosso et al.
[229], investigated the temperature dependency of yield strength in nanocrystalline Sc2O3-
stablized ZrO2 from room temperature to 2050 °C. Their results cast doubt on the previous
hypothesis that GB diffusion is the reason for Hall-Petch break down in nanocrystalline
ceramics, at least in this material system. This approach has also been used to investigate
GB and surface diffusivities in ZrO2 bicrystal by the same group of authors [230].
In addition to mechanical testing, in-situ characterization of other crystalline defects,
such as oxygen vacancies under thermal or electrical stimulus is of great importance for a
wide range of functional properties in ceramics. Klie, et al. [3] performed in-situ heating
coupled with STEM-EELS on a series of perovskite oxides to capture changes in the struc-
ture, composition, and valence states. An excess of oxygen vacancies was observed at the
GB before heating. At 724 K, the amount of oxygen vacancies increased both in the bulk
and at the GB, even though the GB cation structure did not change upon heating. Accord-
ing to the EELS data collected from the same GB upon cooling, oxygen diffusion occurs
from the bulk into the GB plane which agrees with the increase of oxygen vacancies upon
heating. In other words, the oxygen diffuses to the bulk upon heating and goes back to
the GB upon cooling. The impact of electrical biasing on oxygen vacancy migration in
CeO2 was studied by Gao, et al. [231] using STEM-EELS. The formation and ordering of
oxygen vacancies in the CeO2 thin film within an Au/CeO2/Nb-SrTiO3 junction was di-
rectly observed by applying a positive voltage. These promising results have proven that
the oxygen vacancy migration mechanism is responsible for the resistance switching in
CeO2 and possibly other oxides with similar properties [231]. Such material systems can
be used for next generation memories due to their small power consumption, high den-
sity, and other properties. It is important to understand and control the switching behav-
ior of nanoscale oxide devices.
In-situ electron microscopy combined with electron loss near edge structure (ELNES)
can be used to monitor changes in the electronic structure of many elements such as 3d-
transition and rare earth metal oxides, including Fe and Ce oxide [232,233]. Numerous
experiments have suggested a strong correlation between changes in valence states of cat-
ions and the L3/L2 or the ionization edges (white lines) ratio and O-K edge fine structure.
In addition, the O-K pre-peak, which is caused due to transitions from the O1s state to the
unoccupied hybridized O 2p in transition metal 3d orbitals, helps quantify the transition
metal valence and the concentration of mobile holes and the transition metal spin states
[234–240].
In 2000, Wang, et al. [241], reported the L3/L2 ratio for direct measurement of Co and
Mn valence states. Significant differences between O-K edges in Co3O4 and CoO allowed
to distinguish Co oxidation state in these two structures [232,233].
In an in-situ heating experiment coupled with STEM-EELS, metallic cobalt was dis-
tinguished from cobalt in 2+ and 3+ oxidation states [237]. The ELNES used for this study
are shown in Figure 32. A clear energy shift in the Co-L3 and L2 peaks of CoO (contains
Co2+) compared to Co3O4 (contains Co2+and Co3+) is observed which is related to transition
of 2p3/2 to 2p1/2 core-shell electrons into unoccupied 3d orbitals. In summary, the following
techniques can be used to measure the oxidation state of each cation at the GBs and HIs,
where defect chemistry at such interfaces plays an important role and it directly relevant
to the concentration of such defects at the interface.

Crystals 2021, 11, 878 37 of 58
Figure 32. Comparison of (a) O K-edges and (b) Co-L2,3 edges of metallic Co, CoO, and Co3O4 [237]. Reprinted from in situ
electron energy loss spectroscopy study of metallic Co and Co oxides, Zhao, Y., Feltes, T.E., Regalbuto, J.R., Meyer, R.J., &
Klie, R.F., Journal of Applied Physics, 2010, 108(6), 063704, with the permission of AIP Publishing.
In-situ EM has proved to be beneficial in research of battery materials, particularly
observing material degradation during electrochemical performance and cycling [242].
While lithium-ion batteries with Ni-rich layered oxides such as LiNi0.8Co0.15Al0.05O2 (NCA)
are of interest as cathodes, they suffer from capacity fade hypothesized to be correlated to
microcracks that initiate and grow through GBs [243,244]. These microcracks serve as
pathways for electrolyte penetration into secondary NCA particle and grow with the elec-
trolyte corrosion and mechanical strain during lithiation/delithiation [245]. These hypoth-
eses are validated by the comparison ELNES in the bulk and GB. According to Figure 33b,
Ni shows an increase in the L3/L2 ratio indicates the change of Ni oxidation state from the
bulk to the GB, while the same thing did not occur for cobalt (Figure 33e). In addition, GB
were observed to be preferential cites for in situ exsolution of metallic particles [246]. This
phenomenon is related to the lower formation energy for vacancy and nucleation sites at
GBs, which is ultimately important for controlling particle density and distribution.

Crystals 2021, 11, 878 38 of 58
Figure 33. SEM and HAADF-STEM images of NCA after 1500 cycles. (a) Cross-section image of a secondary particle. (b)
a representative GB at the microcrack line indicated in (a). Fast Fourier transform (FFT) pattens from (c) the GB and (d)
the bulk. (e) a representative GB beyond the microcrack line indicated in (a). (f) the EELS spectra of Ni L-edge, (g) Co L-
edge, and (h) O K-edge acquired at the points indicated in (b). (i) Normalized Ni concentration (relative to Co) from the
GB to the bulk derived from EELS spectra [245]. Figure 33 reprinted with permission from the Royal Society of Chemistry
Copyright 2021.
4. Corelating GB Characterization with Properties
4.1. Electrical Properties
In polycrystalline oxides, GBs and HIs play an important role in functional properties
such as ionic conductivity [247,248]. To measure these properties, techniques such as elec-
trochemical impedance spectroscopy (EIS) are used that can measure conductivity of ions
and distinguish the contribution of each component of the system to the measured con-
ductivities [249,250]. It is common for GBs to behave different than the bulk and affect the
conductivity values, whether reduce or increase them. For example, in many oxide sys-
tems GBs tend to hinder oxygen vacancy movement thereby drastically reducing the con-
ductivity values. To better understand the mechanisms responsible for these observations,
a comprehensive analysis of charge carriers responsible for conductivity or in other words
defect chemistry at the GB is required and has been studied to good extent [71]. Previ-
ously, an intrinsic space charge-layer was discovered that is localized at the GB which
induces local redistribution of charge carriers in the vicinity of the GB core. This results in
the depletion of charge carriers, reducing the conductivity at the GB [251]. In the following
example, STEM- EELS was used to link nanoscale GB defect chemistry to the oxygen ion
conductivity across the GBs Ca-doped ceria [21]. Figure 34a presents an ADF image of a
GB where both grains are oriented near a zone axis orientation (Figure 34) [21]. The con-
trast is sensitive to the Z or atomic number of the elements. The decrease in ADF intensity
at the GB is an indication of a decrease in high angle scattering of the electrons. This agrees
with the EELS results shown on the side, presenting an increase in Ca2+ concentration at
and near GB core compared to the bulk, as Ca2+ is lighter than Ce4+ therefore we expect
lower signals where it segregates. The amount of Ca segregated at the GB was measured
using EELS and is shown in Figure 34b. The segregation is maximized in the case of 5
mol% Ca added to the system, which is also followed by a 2 times higher conductivity
value (Figure 34c).
f
h
i

Crystals 2021, 11, 878 39 of 58
Figure 34. (a) ADF AC-STEM image of a GB in 5CCO, the associated elemental EELS map was acquired from the area in
the dashed box and is on the right. (b) All EELS measurements of GB Ca2+ concentration, (c) Influence of nominal Ca2+
concentration, x, on GB conductivity measured at 300 °C [21]. Reprinted from Nanoscale, Enhanced ionic conductivity in
electroceramics by nanoscale enrichment of grain boundaries with high solute concentration, William J. Bowman, Made-
leine N. Kelly, Gregory S. Rohrer, Cruz A. Hernandez and Peter A. Crozier, 2017, 9, 17293–17302, with permission from
the Royal Society of Chemistry (2017).
Another major application of S/TEM characterization technique is the field of all
solid-state lithium batteries, where ceramic oxides are employed as solid electrolytes
[252]. These batteries are of interest for electrochemical energy storage applications due
to their high energy density, safety, long cycle life and wide range of operation tempera-
ture. The candidate oxides include cubic garnet-type Li7La3Zr2O12 (LLZO), perovskite-
type Li3xLa2/3−xTiO3 (LLTO) and NASICON NLiM2(PO4)3 (M = Ti, Ge, Zr, Hf) [253–258].
These structures are stable in air and with respect to Li metal and other high voltage cath-
odes. However, there are a few downsides including low ionic conductivity at various
interfaces such as the anode/electrolyte/cathode, and GBs in the bulk. For example, the
perovskite Li0.34 La0.51TiO2.94 has shown bulk ionic conductivity of 1 × 10−3 S cm−1 and total
ionic conductivity of higher than 2 × 10−5 S cm−1 at room temperature, meaning GBs tend
to decrease conductivity [257]. Therefore, researchers began investigating this phenome-
non to design new interfaces with higher efficiencies. Doping has been effective in im-
proving total conductivity values in some oxides [259]. In general, doping of elements
such as Al3+/Ge4+at the Li sites and Ta5+/Nb5+ at the Zr sites has stabilized the cubic garnet
and reduced Li ion migration activation energy thereby increasing the Li ion conductivity

Crystals 2021, 11, 878 40 of 58
[260,261]. The substitution of the Zr4+ sites by Ge4+ improved ionic conductivity to be 4.78
× 10−4 S cm−1 at 20 °C with an activation energy of 0.29 eV which are among the best in the
field [262,263]. This dopant has also improved the sintering process by reducing the sin-
tering temperature and increasing the density.
Another example is the poor GB conductivity in the Li-ion conducting solid electro-
lyte LLTO. Figure 35 shows HAADF-STEM image of GBs with dark contrast denoted as
GB type I (upper portion) and normal contrast denoted as type II (lower portion). The
difference between atomic structure of these types is evident in Figure 35b, c. While the
GB type I has significantly deviated from the bulk and perovskite atomic structure disap-
peared, the GB type II experienced a much smaller lattice mismatch compared to that in
type I and the perovskite structure remained. It is discovered using TEM that most GBs
have severe structural and chemical changes from normal perovskite to compensate for
the random orientation of the adjacent grains including such as reduction of Ti4+ to Ti3+,
depletion of Li, decrease in La content and deformation of Ti-O polyhedral (Figure 36). In
addition, the structure of type I GB was schematically depicted in Figure 37 based on the
HAADF-STEM and EELS analysis. The intensities of each element were defined by the
atomic column intensity in the bulk. As seen in EELS results, the GB was filled with O and
Ti and possessed a non-perovskite crystal structure. It was concluded that the existing GB
is not energetically preferred for either Li accommodation or transport, giving rise to the
poor GB conductivity. Papers including one by Ma et al. [20] perfectly demonstrates the
correlation between atomic-scale GB structure and ionic conductivity in Li superionic con-
ductors [27,261,264].
Figure 35. (a) HADDF-STEM of a GB exhibiting both dark and normal contrast regions, labelled as
type I and type II, respectively. The {001} planes of the alternating L-rich/La-poor layers of differ-
ent regions in the grain were marked to highlight the existence of nanodomains (b) magnified type
I GB feature. (c) magnified type II GB feature [20]. Reproduced from [20] with permission from the
Royal Society of Chemistry, Copyright 2014.

Crystals 2021, 11, 878 41 of 58
Figure 36. EELS data of (a) La-M4,5, (b) Li-K, (c) Ti-L2,3, and (d) O-K edges for the type I GB and the
bulk. The spectra were normalized to the integrated intensity of the Ti-L2,3 edge [20]. Reproduced
from [20] with permission from the Royal Society of Chemistry, Copyright 2014.
Figure 37. Schematic of the atomic configuration of the type I GB based on the HAADF-STEM
images and EELS analysis, along with an illustration of the Li site distribution across type I GB
[20]. Reproduced from [20] with permission from the Royal Society of Chemistry, Copyright 2014.
The atomic structures and the band gaps of four typical GBs in α-Al2O3 by stem and
valence electron energy-loss spectroscopy. It was found that the band gaps of the GBs are
narrowed by 0.5–2.1 eV compared with that of 8.8 eV in the bulk [161].
4.2. Thermal Properties
Another functional property that can be affected by the segregations at the GBs are
thermal conductivity and thermoelectric characteristics. While most of the literature in

Crystals 2021, 11, 878 42 of 58
this area is related to metals and non-oxide ceramics [265–267], there still exist a few stud-
ies that correlate thermal conductivity to segregation at the GB. In a work done by Boyle,
et al. [268], it was proven that the addition of bismuth dopant to the calcium cobaltite
structure, has improved Seebeck factor (a measure of the magnitude of an induced ther-
moelectric voltage in response to a temperature difference across the material [269]) and
thermal conductivity, increasing the power factor in this material that can potentially be
used as thermoelectric generators, coolers, etc. Segregation of bismuth at the GB plane is
evident in Figure 38 captured using TEM and HAADF imaging [268].
Figure 38. Nanostructure examination of Ca2.7Bi0.3Co4O9. (a) Low magnification TEM diffraction contrast image shows
typical GB and the adjacent two grains. The pink arrow C1 and the green arrows C2 indicate the different c-axis orientations
of such two neighboring grains and the misorientation angle is about 6°. (b), Low magnification STEM-Z contrast image
from the same GB as indicated in (a). (c) High magnification STEM Z-contrast image of the same GB indicated in b, show-
ing two segregation sites. (d)Temperature dependence of the thermal conductivity of Ca3-xBix Co4O9 with x = 0, 0.1, 0.2, 0.3
and 0.4 [268]. Figure 38 reprinted from Journal of the European Ceramic Society, 36, Cullen Boyle, Paulo Carvillo, Yun
Chen, Ever J. Barbero, DustinMcintyre, Xueyan Song, Grain boundary segregation and thermoelectric performance en-
hancement of bismuth doped calcium cobaltite, 601–607, Copyright (2015), with permission from Elsevier.
Furthermore, it was shown previously that the combination of lattice dopants substi-
tution and secondary phase segregation at the GB has dramatically increased the energy
conversion efficiency in Bi-doped CaMnO3-. In this case CuO segregates at the GB, lead-
ing to a significant increase of both the Seebeck coefficient and electrical conductivity (Fig-
ure 39) [268].

Crystals 2021, 11, 878 43 of 58
Figure 39. (a) HRTEM image taken from the Cu-added Bi-doped Ca MnO3−δ samples. Fourier transformation from the
corresponding nanodomains with different orientation. (b)Temperature dependence of thermal conductivity in Ca1–
xBixMnCuyO3−δ [270]. Reprinted from ACS Applied Materials & Interfaces, Grain Boundary Phase Segregation for Dramatic
Improvement of the Thermoelectric Performance of Oxide Ceramics, Xueyan Song, Sergio A. Paredes Navia, Liang Liang,
Cullen Boyle, Cesar-Octavio Romo-De-La-Cruz, Bryan Jackson, Alec Hinerman, Megan Wilt, Jacky Prucz, and Yun Chen-
with, 2018, 10, 45, 39018–39024, with permission from American Chemical Society (2018).
Using S/TEM techniques, it is possible to learn about grain texture and orientations
to correlate that to the properties, as done in a work by Shi et al. [271]. It has been demon-
strated in the literature that GB has a huge impact on thermal conductivities of polycrys-
talline ceramics, including oxides [270,272,273]. GB engineering has shown to be a better
way for improving thermal resistance of materials such as nanocrystalline yttria-stabilized
zirconia and alumina ceramics and refractories [273]. While the impact of GBs is evident,
a more comprehensive study of the nature of the structure and chemistry of the bounda-
ries and their interaction with phonons is required to optimize such properties.
4.3. Mechanical Properties
As a general type of planar defect, GBs in ceramics play an important role in their
deformation mechanisms, such as GB sliding and intergranular fracture. Changes in the
structure and chemistry of GBs could strongly influence the mechanical properties of ce-
ramics. Tremendous research efforts have been devoted to enhancing the mechanical
properties via GB engineering approaches that include controlling the grain size [90,274–
277]; doping with dilute amount of rare earth elements [90,278,279]; forming glassy
nanolayer films at GBs [200,280]; and optimizing the GB character distribution by intro-
ducing a high frequency of low-angle GBs [281]. Despite these important improvements,
fundamental breakthroughs in the understanding of the correlations between GB struc-
ture and mechanical response is needed for the development of ceramic components with
superior mechanical reliability.
Accurately determining the structure, chemical composition and bonding state at
GBs is critical to unveil the structure-mechanical property relationships. GB non-stoichi-
ometry often strongly influences the mechanical properties. As one of the most widely
used structural ceramics, Al2O3 has a weak resistance to creep at high temperatures via GB
sliding, while its creep resistance could be effectively improved by the addition of rare
earth elements. Ikuhara, et al. [278,279] observed a segregation layer rich in Ln cations at
GBs using HRTEM, EDXS and EELS. Such non-stoichiometry at GBs contributes to the
dopant-dependent creep resistance in Ln oxide-doped Al2O3, although questions remain

Crystals 2021, 11, 878 44 of 58
regarding its atomistic mechanisms (Figure 40). Later studies identified the atomic-scale
location of rare earth atoms at the expanded sites along GBs, which leads to an increase in
the number of bonds and bond strength at GBs [200]. This finding suggests the enhance-
ment in creep resistance could be affected by the ionic radii of rare earth dopants. Towards
three-dimensional characterization of the GB segregation, a buried GB in Y-doped Al2O3
was studied under off-axis illumination condition using HAADF STEM, which revealed
an ordered two-dimensional structure of Y dopants on the GB with local disordering. In
addition to dopant segregation, locations of oxygen vacancies on GBs were determined
by HAADF and ADF STEM and EELS, underlining the importance of oxygen non-stoichi-
ometry on the mechanical properties in oxides [7,282].
Figure 40 (a) High temperature creep properties of pure Al2O3 and different Ln oxides-doped Al2O3; (b) ELNES Al-L1 edge
peak for Lu2O3-doped Al2O3 in the grain interior and at a GB. (c) HRTEM micrograph of a GB in Lu2O3-doped Al2O3. (d)
EDXS point spectra from the grain interior and the GB, corresponding to the two points in (c). (e) and (g) are HAADF
STEM images of a Σ31 [0001] tilt GB in pure Al2O3, Y-doped Al2O3, respectively. Brighter columns in (g) are the locations
of segregated Y atoms. (f) and (h) are the same images as in (e) and (g), respectively, and with schematic overlay to show
the arrangement of atomic columns. It is observed that Y atoms are at the center of the seven-member ring units along the
GB [200]. Figures 40a–d reprinted from Materials Science and Engineering: A, Impurity effects on grain boundary strength
in structural ceramics, Yuichi Ikuhara, Hidehiro Yoshida, Taketo Sakuma, 2001, 24–30, with permission from Elsevier.
Figures 40e–h reprinted from Grain Boundary Strengthening in Alumina by Rare Earth Impurities, J. P. Buban, K. Matsu-
naga, J. Chen, N. Shibata, W. Y. Ching, T. Yamamoto, Y. Ikuhara, Science 2006, 311, 212–215. Reprinted with permission
from AAAS.
Furthermore, microstructural characterization of mechanically tested ceramics could
shed light on their deformation mechanisms. Using HR-TEM, the atomic structure of
Al2O3 recovered from shock loading was characterized, and a transition in deformation
mechanisms from intergranular fracture and dislocation activities near GBs at low impact
pressure to deformation twinning at high impact pressure has been observed [283]. In B4C,
which is an important super-hard lightweight ceramic, the deformation and failure mech-
anism has long been controversial. Acquiring experimental evidence on atomic structure

Crystals 2021, 11, 878 45 of 58
of B4C is limited by its complex crystal structure with a C-B-C chain and B11C icosahedra
and low Z-contrast of boron and carbon. Recent developments in aberration-corrected
scanning transmission electron microscopy enables microstructural characterization with
sub angstrom spatial resolution, particularly for light elements. Using ABF STEM imag-
ing, detailed structure within the amorphous shear bands in nanoindentation tested B4C
was investigated. Distorted icosahedra structures were found randomly distributed in the
amorphous bands, offering insights into the deformation-induced amorphization mecha-
nisms of B4C [284]. Guo et al. characterized the GB structure in nanocrystalline B4C after
quasi-static bending tests of cantilever nanobeams. They reported 2 nm wide amorphous
phase at GBs, which is likely induced by GB sliding and corresponds to the breakdown of
Hall–Petch relationship in ceramics [181].
4.4. Magnetic Properties
GBs and HIs can affect magnetic behavior of materials such as oxides. For example,
lower magnetoresistance at T⪡ TC in a polycrystalline La0.67Ba0.33MnO2.99 compared to it is
single crystal form is related to high volume of GBs present in the former [285]. In another
example, ferromagnetic properties of nanocrystalline Mn-doped ZnO thin films were im-
proved due the existence of greater GB specific area compared to the critical values
[286].BF-TEM studies show that the wurtzite lattice of ZnO is separated by amorphous
layers whose thickness increases with the Mn concentration which is controllable (Figure
41). Amorphous phases at GBs and HIs can potentially affect other magnetic properties
such as coercivity (resistance of a magnetic material to demagnetization). This is evident
in a work by Sepehri-amin et al. [287], where B-enriched amorphous intergranular phases
cause the domain wall pinning that is correlated to high coercivity observed in B-doped
Sm (Fe0.8Co0.2)12. Lattice distortion and strain in HIs can affect magnetic flux pinning cen-
ters leading to high critical density in SrZrO3 [288].
Figure 41 represents the BF-TEM image of Mn-doped ZnO films with (a) 10 and (b)
15 at% Mn concentration. The direct resolution of the lattice allows the observation of the
crystalline ZnO grains in wurtzite structure, and it is clearly visible that amorphous
phases lie at the areas between the grains or in other words, GBs [289].
Figure 41. BF-TEM micrographs of doped ZnO thin films with (a) 10 at% Mn ZnO grains and (b) 15 at% Mn ZnO grains,
surrounded by amorphous layers. The insets in (b) show the fast Fourier transform patterns for crystalline and amorphous
regions, A, B and C, respectively [289]. Figure 41 reprinted by permission from Springer Nature Customer Service Centre
GmbH: Springer Nature, JETP Letters, Grain boundary layers in nanocrystalline ferromagnetic zinc oxide, Straumal, B.B.;
Myatiev, A.A.; Straumal, P.B.; Mazilkin, A.A.; Protasova, S.G.; Goering, E.; Baretzky, B., M.H., Copyright © 2021.

Crystals 2021, 11, 878 46 of 58
The examples mentioned above and more in literature [290,291] prove the im-
portance of GBs and HIs in magnetic behavior of oxides, and therefore highlight the im-
portance of their characterization.
4.5. Optical Properties
Polycrystalline ceramics can be transparent, translucent, or opaque [40]. Transparent
polycrystalline ceramics have become of interest since Neodymium-doped yttrium alu-
minum garnet (Nd:YAG) was introduced as a lasing medium for solid state lasers in 1995
[201] and have applications in armors, opto-electrical devices such as solar cells, displays,
and circuitries. Compared to single crystals, polycrystalline ceramics are cost-effective,
easier to produce, exhibit good mechanical, thermal and chemical stability, and shape-
control availability [292,293]. However, certain defects such as impurities and porosity
inside their microstructure scatter and refract light leading to optical losses, thus limiting
wide employment of transparent polycrystalline ceramics [203,294]. GBs are one main
area accommodating such defects and can impact properties of these materials. The GB
structure can also control grain size and phase composition, thus playing a significant role
in transmittance of polycrystalline ceramics. Some boundaries with higher doping con-
centrations show a low reflectance, whereas other boundaries with secondary phases
show a bold difference in refractive index with respect to the bulk, becoming the major
scattering sources [295]. In addition to impurities, grain size related birefringence also af-
fects the transparency of polycrystalline ceramics based on Rayleigh approximation
[296,297]. Therefore, it is important to characterize such defects and interfaces to achieve
optimal properties in polycrystals [12–14] For example, S/TEM techniques can help detect
the origin of light scattering at interface to improve optical quality. Segregation of ele-
ments at the GBs in polycrystalline oxides affect optical properties such as transmittance
[148]. In a work by Trunec, et al. [298] polycrystalline alumina was doped with Zr and
spinel nanoparticles to improve in-line transmittance. According to STEM EDXS results,
Zr segregated at the GBs shown in Figure 42. According to this study, the segregation has
impeded GB motion to achieve grain size refinement in dense alumina structure, thereby
lowering light loss by birefringence.
Figure 42. Morphology and element mapping of grain boundaries in Zr-doped Al2O3. (a) STEM
image of a triple junction; (b) DF-STEM image and STEM-EDXS elemental mapping of the selected
GB area [298]. Figure 42 reprinted from Journal of the European Ceramic Society, 35, Martin
Trunec, Karel Maca, Radim Chmelik, Polycrystalline alumina ceramics doped with nanoparticles
forincreased transparency, 1001–1009, Copyright (2014), with permission from Elsevier.
In addition to intentional doping, small amounts of impurities in raw material can
impact the microstructure and properties of transparent magnesium aluminate spinel
(MgAl2O4) [299]. Impurities such as carbon, sulfur and silicon tend to segregate and form

Crystals 2021, 11, 878 47 of 58
amorphous phases at GBs and triple junctions (Figure 43a–c) and restrict grain growth by
GB pinning and solute dragging mechanisms. Hence, sub-micrometer grains increase GB
volume fraction and impurity segregation at GBs discontinues light transmittance by in-
crease in light scattering, leading to opacity (Figure 43f). LiF reacts with impurities, form-
ing volatile species. These products can be easily removed in high temperature to clean
GB and facilitate grain growth. Clean GBs greatly improve transparency of the MgAl2O4
that can be used as transparent armor and IR-seeking missile domes (Figure 43d–g).
Figure 43. Microstructure of MgAl2O4 spinel by low-purity powder hot-pressed at 1200 °C without LiF (a,b,c,f) and with
1 wt% LiF (d,e,g,a) TEM image of sub-micrometer grains; (b) TEM image of amorphous phase at interface between ma-
trix grains and sub-micrometer grains; (c) TEM image of amorphous phase at triple junctions; (d) SEM image of matrix
grains (e) SEM image of micrometer grains; (f) digital image of a MgAl2O4 pellet without LiF; (g) digital image of a
MgAl2O4 pellet with 1 wt% LiF [299]. Figure 43 reprinted from International Journal of the American Ceramic Society,
10, Marc Rubat du Merac and Ivar E. Reimanis, Effect of Impurities and LiF Additive in Hot-Pressed Transparent Mag-
nesium Aluminate Spinel, E33–E48, Copyright (2012), with permission from John Wiley and Sons.
5. Conclusion and Future Perspective
Recent and continual advancements in electron microscopy hardware, data acquisi-
tion and processing algorithms, three-dimensional analyses, and innovative workflows
provide enormous potential to the field of interfacial characterization. An overarching
goal of high-resolution analysis is to such fine-scale data to develop models capable of
predicting macroscopic structure-property relationships. However, it is indeed challeng-
ing to collect a statistically significant amount of observations/data using S/TEM, simply
because one needs to prepare many samples and analyze tens, hundreds or even thou-
sands of interfaces, all of which might vary slightly in their atomic scale nature (e.g.,
atomic structure and chemistry). Therefore, there is significant potential in automation of
such measurements using machine learning and artificial intelligence.
It is imperative to consider the limitations of S/TEM while using and planning to use
it. Direct observations made with the help of powerful electron microscopes give us val-
uable insights to the underlying science of interfaces. But these techniques focus on a small
volume of material, likely subject to multiple specimen preparation steps as mentioned
earlier, and by the time the specimen is observed, the material may be somewhat different
than the material in bulk form. Thus, the authors strongly suggest that (potential) users
keep this in mind and be aware of possible impacts of preparation steps.
There is a clear lack in literature on heterointerface characterization in multiphase
polycrystalline ceramics. The added complexity of such systems due to existence of

Crystals 2021, 11, 878 48 of 58
several phases and defects makes such studies more challenging and resource intensive.
While the majority ceramics community focuses on single phase oxides due to their prom-
ising properties, mixed phase counterparts have proven to be of importance in various
fields such as structural materials, thermoelectric materials, and mixed ion electron con-
ductors (where one or more phases dominate the electron conduction while the other/oth-
ers maintain ion conduction). It is expected that advancements in high-throughput auto-
mated EM, including the ones mentioned in this paper, will enable more investigation of
GBs/HIs in multiphase polycrystalline systems.
Lastly, in-situ and in-operando studies of interfaces in functional materials, and those
sensitive to our atmosphere and electron beam irradiation are very promising, even
though there are limited studies at present. This will likely change with recent and future
technological and workflow advancements in EM, such as sample preparation and cryo-
EM methods for sensitive materials, such as alkali ion conductors. Such dynamic experi-
ments on interfaces (and nanomaterials more generally) are particularly exciting due to
the possibility of observing reactions, degradation, and mechanisms underlying interface
functionality.
In closing, while there are many experimental and practical considerations to make
when undertaking such a study, the authors hope this article paves the way for more sci-
entists and engineers interested in using S/TEM techniques to study and develop interfa-
cial sciences, particularly in ceramic oxides.
Author Contributions: H.V. and W.J.B. conceived the idea of the review. H.V. wrote most of the
paper and H.V. and W.J.B. did critical revision. K.S. contributed to useful discussions and wrote
part of the introduction and the EDXS technique section. X.W. wrote the mechanical properties and
in-situ S/TEM sections. J.L.W. wrote the introduction to TEM and STEM with inputs from H.G. and
J.M. H.G. wrote parts of the optical and magnetic properties sections. All authors have read and
agreed to the published version of the manuscript.
Funding: This research was funded by the National Science Foundation Materials Research Science
and Engineering Center program through the UC Irvine Center for Complex and Active Materials
(NSF DMR-2011967).
Acknowledgments: H.V. and X.W. acknowledge funding by the National Science Foundation Ma-
terials Research Science and Engineering Center program through the UC Irvine Center for Com-
plex and Active Materials (NSF DMR-2011967). K.S. acknowledges NSF DMR Grant 442660-21130
and support from US Department of Education Graduate Assistance in Areas of National Need
(GAANN) Fellowship. J.L.W. and W.J.B. acknowledge funding under the award NSF DMR-2042638.
J.M. acknowledges funding under the award NSF Graduate Research Fellowship. H.G. and W.J.B.
acknowledge funding from the American Chemical Society’s Petroleum Research Fund Doctoral
New Investigator Grant. W.J.B. acknowledges funding from the UC Irvine School of Engineering
new faculty set-up funding.
Conflicts of Interest: The authors declare no conflict of interest.
References
1. Reynolds, W.T. Interfaces in Crystalline Materials By, A.P. Sutton (University of Oxford) and R. W. Balluffi (MIT). Oxford
University Press: New York. 1995. xxvii + 819 pp. $165. ISBN 0-19-851385-2. J. Am. Chem. Soc. 1997, 119, 2343–2343.
2. Pryds, N.; Esposito, V.; Dk, N.; Dk, V. When two become one: An insight into 2D conductive oxide interfaces. J. Electroceramics
2017, 38, 1–23.
3. Klie, R.F.; Ito, Y.; Stemmer, S.; Browning, N.D. Observation of oxygen vacancy ordering and segregation in Perovskite oxides.
Ultramicroscopy 2001, 86, 289–302.
4. Carter, C.B.; Norton, M.G. Ceramic Materials: Science and Engineering; Springer: Berlin/Heidelberg, Germany, 2007; Volume 716.
5. Lin, Y.; Fang, S.; Su, D.; Brinkman, K.S.; Chen, F. Enhancing grain boundary ionic conductivity in mixed ionic-electronic
conductors. Nat. Commun. 2015, 6, 1–9.
6. Feng, B.; Lugg, N.R.; Kumamoto, A.; Ikuhara, Y.; Shibata, N. Direct Observation of Oxygen Vacancy Distribution across Yttria-
Stabilized Zirconia Grain Boundaries. ACS Nano 2017, 11, 11376–11382.
7. Bowman, W.J.; Darbal, A.; Crozier, P.A. Linking Macroscopic and Nanoscopic Ionic Conductivity: A Semiempirical Framework
for Characterizing Grain Boundary Conductivity in Polycrystalline Ceramics. ACS Appl. Mater. Interfaces 2020, 12, 507–517.

Crystals 2021, 11, 878 49 of 58
8. Bowman, W.J.; Zhu, J.; Sharma, R.; Crozier, P.A. Electrical conductivity and grain boundary composition of Gd-doped and
Gd/Pr co-doped ceria. Solid State Ion. 2015, 272, 9–17.
9. Orlovskaya, N.; Browning, N. (Eds.). Mixed Ionic Electronic Conducting Perovskites for Advanced Energy Systems; Springer Science
& Business Media: Secaucus, NJ, USA, 2012; Volume 173.
10. Dong, Y.; Zhang, Z.; Alvarez, A.; Chen, I.-W. Potential jumps at transport bottlenecks cause instability of nominally ionic solid
electrolytes in electrochemical cells. Acta Mater. 2020, 199, 264–277.
11. Giuntini, D.; Torresani, E.; Chan, K.T.; Blankenburg, M.; Saviot, L.; Bor, B.; Domènech, B.; Shachar, M.; Müller, M.; Olevsky,
E.A.; et al. Iron oxide-based nanostructured ceramics with tailored magnetic and mechanical properties: Development of
mechanically robust, bulk superparamagnetic materials. Nanoscale advances, 1(8), pp.3139–3150. Nanoscale Adv. 2019, 1, 3139–
3150.
12. Gilde, G.; Patel, P.; Sands, J.; Patterson, P.; Blodgett, D.; Duncan, D.; Hahn, D. Evaluation of hot isostatic pressing parameters
on the optical and ballistic properties of spinel for transparent armor. Army Res. Lab. Aberd. Proving Ground Md 2006, 88, 2747–
2751.
13. Sutorik, A.C.; Gilde, G.; Kilczewski, S.; Patel, P.; Sands, J.M. Development of transparent ceramic spinel (MgAl2O4) for armor
applications. In Proceedings of the Optics InfoBase Conference Papers, Optical Society of America, Jackson Hole, WY, USA, 13–
17 June 2010; p. OWA6.
14. Lange, F.F.; Clarke, D.R. Morphological Changes of an Intergranular Thin Film in a Poly crystalline Spinel. J. Am. Ceram. Soc.
1982, 65, 502–506.
15. Smith, D.; Auvray, S.; Absi, J.; Kadiebu, S.; Fayette, S. Grain-boundary thermal resistance in polycrystalline oxides: Alumina,
tin oxide, and magnesia. High. Temp. High. Press. 2003, 35–36, 93–99.
16. Szot, K.; Rodenbücher, C.; Bihlmayer, G.; Speier, W.; Ishikawa, R.; Shibata, N.; Ikuhara, Y. Influence of dislocations in transition
metal oxides on selected physical and chemical properties. Crystals 2018, 8, 241.
17. Kingery, W.D.; Bowen, H.K.; Uhlmann, D.R. 1976. Introduction to Ceramics; John Wiley & Sons: Hoboken, NJ, USA, 1976; Volume
17.
18. Mebane, D.S.; De Souza, R.A. A generalised space-charge theory for extended defects in oxygen-ion conducting electrolytes:
From dilute to concentrated solid solutions. Energy Environ. Sci. 2015, 8, 2935–2940.
19. Xu, X.; Liu, Y.; Wang, J.; Isheim, D.; Dravid, V.P.; Phatak, C.; Haile, S.M. Variability and origins of grain boundary electric
potential detected by electron holography and atom-probe tomography. Nat. Mater. 2020, 19, 887–893.
20. Ma, C.; Chen, K.; Liang, C.; Nan, C.W.; Ishikawa, R.; More, K.; Chi, M. Atomic-scale origin of the large grain-boundary resistance
in perovskite Li-ion-conducting solid electrolytes. Energy Environ. Sci. 2014, 7, 1638–1642.
21. Bowman, W.J.; Kelly, M.N.; Rohrer, G.S.; Hernandez, C.A.; Crozier, P.A. Enhanced ionic conductivity in electroceramics by
nanoscale enrichment of grain boundaries with high solute concentration. Nanoscale 2017, 9, 17293–17302.
22. Tong, X.; Bowman, W.J.; Mejia-Giraldo, A.; Crozier, P.A.; Mebane, D.S. New Data-Driven Interacting-Defect Model Describing
Nanoscopic Grain Boundary Compositions in Ceramics. J. Phys. Chem. C 2020, 124, 23619–23625.
23. Alekseeva, I.; Dymshits, O.; Tsenter, M.; Zhilin, A.; Golubkov, V.; Denisov, I.; Skoptsov, N.; Malyarevich, A.; Yumashev, K.
Optical applications of glass-ceramics. J. Non Cryst. Solids 2010, 356, 3042–3058.
24. Cao, X.Q.; Vassen, R.; Stoever, D. Ceramic materials for thermal barrier coatings. J. Eur. Ceram. Soc. 2004, 24, 1–10.
25. Handwerker, C.A.; Morris, P.A.; Coble, R.L. Effects of Chemical Inhomogeneities on Grain Growth and Microstructure in Al2O3.
J. Am. Ceram. Soc. 1989, 72, 130–136.
26. Shibata, N.; Findlay, S.D.; Azuma, S.; Mizoguchi, T.; Yamamoto, T.; Ikuhara, Y. Atomic-scale imaging of individual dopant
atoms in a buried interface. Nat. Mater. 2009, 8, 654–658.
27. An, J.; Park, J.S.; Koh, A.L.; Lee, H.B.; Jung, H.J.; Schoonman, J.; Sinclair, R.; Gür, T.M.; Prinz, F.B. Atomic scale verification of
oxide-ion vacancy distribution near a single grain boundary in YSZ. Sci. Rep. 2013, 3, 1–6.
28. Lu, X.; Xia, G.; Lemmon, J.P.; Yang, Z. Advanced materials for sodium-beta alumina batteries: Status, challenges and
perspectives. J. Power Sources 2010, 195, 2431–2442.
29. Yoon, H.; Choi, M.; Lim, T.W.; Kwon, H.; Ihm, K.; Kim, J.K.; Choi, S.Y.; Son, J. Reversible phase modulation and hydrogen
storage in multivalent VO 2 epitaxial thin films. Nat. Mater. 2016, 15, 1113–1119.
30. Oshima, Y.; Lee, S.; Takayanagi, K. Visualization of lithium ions by annular bright field imaging. Microscopy 2017, 66, 15–24.
31. Bowman, W.J.; Hernandez, C.A.; McGuinness, K.; Crozier, P.A. Quantifying and Correlating the Composition and Conductivity
of Grain Boundaries in Ca-doped CeO2 Electrolytes, an AC-STEM EELS Study. Microsc. Microanal. 2015, 21, 1727–1728.
32. Garvie, L.A.J.; Buseck, P.R. Determination of Ce4+/Ce3+ in electron-beam-damaged CeO2 by electron energy-loss spectroscopy.
J. Phys. Chem. Solids 1999, 60, 1943–1947.
33. Lei, Y.; Ito, Y.; Browning, N.D.; Mazanec, T.J. Segregation Effects at Grain Boundaries in Fluorite-Structured Ceramics. J. Am.
Ceram. Soc. 2002, 85, 2359–2363.
34. Hojo, H.; Mizoguchi, T.; Ohta, H.; Findlay, S.D.; Shibata, N.; Yamamoto, T.; Ikuhara, Y. Atomic structure of a CeO2 grain
boundary: The role of oxygen vacancies. Nano Lett. 2010, 10, 4668–4672.
35. Dufour, L.; Monty, C. Surfaces and Interfaces of Ceramic Materials; Springer Science & Business Media: Berlin/Heidelberg,
Germany, 2012.
36. McGee, T.D. Grain boundaries in ceramic materials. In Materials Science Research; Springer: Berlin/Heidelberg, Germany, 1965;
pp. 3–32.

Crystals 2021, 11, 878 50 of 58
37. Ernst, F.; Kienzle, O.; Rühle, M. Structure and Composition of Grain Boundaries in Ceramics. J. Eur. Ceram. Soc. 1999, 19, 665–
673.
38. Clarke, D. Grain Boundaries In Polycrystalline Ceramics. Annu. Rev. Mater. Sci. 2003, 17, 57–74.
39. Kingery, W.D. The chemistry of ceramic grain boundaries. Pure Appl. Chem. 1984, 56, 1703–1714.
40. Burgers, J.M. Geometrical considerations concerning the structural irregularities to be assumed in a crystal. Proc. Phys. Soc. 1940,
52, 23–33.
41. Syed, K.; Xu, M.; Ohtaki, K.K.; Kok, D.; Karandikar, K.K.; Graeve, O.A.; Bowman, W.J.; Mecartney, M.L. Correlations of grain
boundary segregation to sintering techniques in a three-phase ceramic. Materialia 2020, 14, 100890.
42. Smallman, R.E.; Ngan, A.H.W. Chapter 10—Surfaces, grain boundaries and interfaces. In Modern Physical Metallurgy, 8th ed.;
Elsevier: Oxford, UK, 2014; pp. 415–442, ISBN 978-0-08-098204-5.
43. Porter, D.A.; Easterling, K.E. Phase Transformations in Metals and Alloys (Revised Reprint); CRC Press: Boca Raton, FL, USA, 2009.
44. Olmsted, D.L.; Foiles, S.M.; Holm, E.A. Survey of computed grain boundary properties in face-centered cubic metals: I. Grain
boundary energy. Acta Mater. 2009, 57, 3694–3703.
45. Randle, V. ‘Special’ boundaries and grain boundary plane engineering. Scr. Mater. 2006, 54, 1011–1015.
46. Ghamarian, I.; Samimi, P.; Rohrer, G.S.; Collins, P.C. Determination of the five parameter grain boundary character distribution
of nanocrystalline alpha-zirconium thin films using transmission electron microscopy. Acta Mater. 2017, 130, 164–176.
47. Guziewski, M.; Banadaki, A.D.; Patala, S.; Coleman, S.P. Application of Monte Carlo techniques to grain boundary structure
optimization in silicon and silicon-carbide. Comput. Mater. Sci. 2020, 182, 109771–109771.
48. Echeverri Restrepo, S.; Tamayo Giraldo, S.; Thijsse, B.J. Using artificial neural networks to predict grain boundary energies.
Comput. Mater. Sci. 2014, 86, 170–173.
49. Kim, H.K.; Ko, W.S.; Lee, H.J.; Kim, S.G.; Lee, B.J. An identification scheme of grain boundaries and construction of a grain
boundary energy database. Scr. Mater. 2011, 64, 1152–1155.
50. Soifer, Y.M.; Verdyan, A.; Kazakevich, M.; Rabkin, E. Nanohardness of copper in the vicinity of grain boundaries. Scr. Mater.
2002, 47, 799–804.
51. Han, J.; Thomas, S.L.; Srolovitz, D.J. Grain-Boundary Kinetics: A Unified Approach. Prog. Mater. Sci. 2018, 98, 386–476.
52. Pérez-Pérez, F.J.; Smith, R. Structural changes at grain boundaries in bcc iron induced by atomic collisions. Nucl. Instrum.
Methods Phys. Res. Sect. b Beam Interact. Mater. At. 2020, 164, 487–494.
53. Dickey, E.C.; Fan, X.; Pennycook, S.J. Structure and chemistry of yttria-stabilized cubic-zirconia symmetric tilt grain boundaries.
J. Am. Ceram. Soc. 2001, 84, 1361–1368.
54. Tochigi, E.; Nakamura, A.; Shibata, N.; Ikuhara, Y. Dislocation Structures in Low-Angle Grain Boundaries of α-Al2O3. Crystals
2018, 8, 133.
55. Gottstein, G. Physical Foundations of Materials Science; Springer-Verlag: Berlin Heidelberg, Germany, 2004; ISBN 978-3-540-40139-
1.
56. Dillon, S.J.; Tang, M.; Carter, W.C.; Harmer, M.P. Complexion: A new concept for kinetic engineering in materials science. Acta
Mater. 2007, 55, 6208–6218.
57. Wang, Y.; Wang, C.; Yuan, L.; Cai, R.; Liu, X.; Li, C.; Zhou, G. Coincidence-site-lattice twist boundaries in bicrystalline α-Fe2O3
nanoblades. J. Phys. Chem. C 2014, 118, 5796–5801.
58. Sato, Y.; Buban, J.P.; Mizoguchi, T.; Shibata, N.; Yodogawa, M.; Yamamoto, T.; Ikuhara, Y. Role of Pr segregation in acceptor-
state formation at ZnO grain boundaries. Phys. Rev. Lett. 2006, 97, 106802.
59. Shibata, N.; Painter, G.S.; Satet, R.L.; Hoffmann, M.J.; Pennycook, S.J.; Becher, P.F. Rare-earth adsorption at intergranular
interfaces in silicon nitride ceramics: Subnanometer observations and theory. Phys. Rev. B 2005, 72, 140101.
60. Imai, H. Mesostructured crystals: Growth processes and features. Prog. Cryst. Growth Charact. Mater. 2016, 62, 212–226.
61. Schuler, J.D.; Rupert, T.J. Materials selection rules for amorphous complexion formation in binary metallic alloys. Acta Mater.
2017, 140, 196–205.
62. Cantwell, P.R.; Frolov, T.; Rupert, T.J.; Krause, A.R.; Marvel, C.J.; Rohrer, G.S.; Rickman, J.M.; Harmer, M.P. Grain boundary
complexion transitions. In Annual Review of Materials Research; Annual Reviews Inc: Palo Alto, CA, USA, 2020; Volume 50, pp.
465–492.
63. Mwema, F.; Akinlabi, E.; Oladijo, O. Sputtered Thin Films: Theory and Fractal Descriptions; CRC Press: Boca Raton, FL, USA, 2021.
64. Liu, H. Growth kinetics of thin film epitaxy. In 21st Century Surface Science—A Handbook; IntechOpen: London, UK, 2020.
65. Hubler, G.K. Pulsed Laser Deposition. MRS Bull. 1992, 17, 26–29.
66. Sediva, E.; Bowman, W.J.; Gonzalez-Rosillo, J.C.; Rupp, J.L. Investigation of the eightwise switching mechanism and its
suppression in SrTiO3 modulated by humidity and interchanged top and bottom platinum and LaNiO3 electrode contacts. Adv.
Electron. Mater. 2019, 5, 1800566.
67. Schweiger, S.; Pfenninger, R.; Bowman, W.J.; Aschauer, U.; Rupp, J.L. Designing strained interface heterostructures for
memristive devices. Adv. Mater. 2017, 29, 1605049.
68. Sheth, J.; Chen, D.; Kim, J.J.; Bowman, W.J.; Crozier, P.A.; Tuller, H.L.; Misture, S.T.; Zdzieszynski, S.; Sheldon, B.W.; Bishop,
S.R. Coupling of strain, stress, and oxygen non-stoichiometry in thin film Pr 0.1 Ce 0.9 O 2−δ. Nanoscale 2016, 8, 16499–16510.
69. Garbayo, I.; Struzik, M.; Bowman, W.J.; Pfenninger, R.; Stilp, E.; Rupp, J.L. Glass-Type Polyamorphism in Li-Garnet Thin Film
Solid State Battery Conductors. Adv. Energy Mater. 2018, 8, 1702265.
70. Taylor, R.I.; Coad, J.P.; Brook, R.J. Grain Boundary Segregation in Al2O3. J. Am. Ceram. Soc. 1974, 57, 539–540.

Crystals 2021, 11, 878 51 of 58
71. Gregori, G.; Merkle, R.; Maier, J. Ion conduction and redistribution at grain boundaries in oxide systems. Prog. Mater. Sci. 2017,
89, 252–305.
72. Rohrer, G.S. Grain boundary energy anisotropy: A review. J. Mater. Sci. 2011, 46, 5881–5895.
73. Krivanek, O.L.; Isoda, S.; Kobayashi, K. Lattice imaging of a grain boundary in crystalline germanium. Philos. Mag. 1977, 36,
931–940.
74. Gust, M.; Goo, G.; Wolfenstine, J.; Mecartney, M.L. Influence of Amorphous Grain Boundary Phases on the Superplastic
Behavior of 3-mol%-Yttria-Stabilized Tetragonal Zirconia Polycrystals (3Y-TZP). J. Am. Ceram. Soc. 1993, 76, 1681–1690.
75. Steele, B.C.H. Mass transport in materials incorporated in electrochemical energy conversion systems. Solid State Ion. 1984, 12,
391–406.
76. Guo, X.; Sigle, W.; Maier, J. Blocking Grain Boundaries in Yttria-Doped and Undoped Ceria Ceramics of High Purity. J. Am.
Ceram. Soc. 2003, 86, 77–87.
77. Kelly, T.F.; Miller, M.K. 2007. Atom probe tomography. Rev. Sci. Instrum. 2007, 78, 031101.
78. von Harrach, H.S.; Klenov, D.; Freitag, B.; Schlossmacher, P.; Collins, P.C.; Fraser, H.L. 2010. Comparison of the detection limits
of EDS and EELS in S/TEM. Microsc. Microanal. 2010, 16, 1312–1313.
79. Diercks, D.R.; Tong, J.; Zhu, H.; Kee, R.; Baure, G.; Nino, J.C.; O’Hayre, R.; Gorman, B.P. Three-dimensional quantification of
composition and electrostatic potential at individual grain boundaries in doped ceria. J. Mater. Chem. A 2016, 4, 5167–5175.
80. Ross, F.M.; Minor, A.M. In situ transmission electron microscopy. In Springer Handbooks; Springer: Berlin/Heidelberg, Germany,
2019; pp. 101–187.
81. Vurpillot, F.; Lefebvre, W.; Cairney, J.M.; Oberdorfer, C.; Geiser, B.P.; Rajan, K. Advanced volume reconstruction and data
mining methods in atom probe tomography. MRS Bull. 2016, 41, 46–51.
82. Wu, Q.; Yu, Z.; Wang, Y.; Diercks, D.; Gorman, B.P.; Rickman, J.M.; Harmer, M.P.; Chan, H.M. Influence of codoping with Hf
and La on grain-boundary transport in alumina. J. Am. Ceram. Soc. 2021, 104, 514–523.
83. Burton, G.L.; Ricote, S.; Foran, B.J.; Diercks, D.R.; Gorman, B.P. Quantification of grain boundary defect chemistry in a mixed
proton-electron conducting oxide composite. J. Am. Ceram. Soc. 2020, 103, 3217–3230.
84. Mutas, S.; Klein, C. Importance of the Protective Layers and the Specimen Preparation for Reproducible APT Results. Microsc.
Microanal. 2011, 17, 730–731.
85. Gupta, D. Diffusion, solute segregations and interfacial energies in some material: An overview. Interface Sci. 2003, 11, 7–20.
86. Foiles, S.M. Temperature dependence of grain boundary free energy and elastic constants. Scr. Mater. 2010, 62, 231–234.
87. Munoz, N.E.; Gilliss, S.R.; Carter, C.B. Remnant grooves on alumina surfaces. Surf. Sci. 2004, 573, 391–402.
88. Saylor, D.M.; Rohrer, G.S. Measuring the influence of grain-boundary misorientation on thermal groove geometry in ceramic
polycrystals. J. Am. Ceram. Soc. 1999, 82, 1529–1536.
89. Kelly, M.N.; Bojarski, S.A.; Rohrer, G.S. The temperature dependence of the relative grain-boundary energy of yttria-doped
alumina. J. Am. Ceram. Soc. 2017, 100, 783–791.
90. Yoshida, H.; Yokoyama, K.; Shibata, N.; Ikuhara, Y.; Sakuma, T. High-temperature grain boundary sliding behavior and grain
boundary energy in cubic zirconia bicrystals. Acta Mater. 2004, 52, 2349–2357.
91. Dillon, S.J.; Harmer, M.P.; Rohrer, G.S. The relative energies of normally and abnormally growing grain boundaries in alumina
displaying different complexions. J. Am. Ceram. Soc. 2010, 93, 1796–1802.
92. Williams, D.B.; Carter, C.B. Transmission Electron. Microscopy: A Textbook for Materials Science; Springer: Berlin/Heidelberg,
Germany, 2009; p. 760, ISBN 978-0-387-76500-6.
93. Thomas, J.; Gemming, T. Analytical Transmission Electron. Microscopy: An. Introduction for Operators; Springer Science & Business
Media: Secaucus, NJ, USA, 2014.
94. Ikuhara, Y. Towards new transmission electron microscopy in advanced ceramics. J. Ceram. Soc. Jpn. 2002, 110, 139–145.
95. Rukari, T.; Babita, A. Review Article Transmission Electron Microscopy—An Overview. Int. Res. J. Invent. Pharm. Sci. 2013, 1, 1–
7.
96. Winey, M.; Meehl, J.B.; O’Toole, E.T.; Giddings, T.H. Conventional transmission electron microscopy. Mol. Biol. Cell 2014, 25,
319–323.
97. Hawkes, P.W.; Spence, J.C. Springer Handbook of Microscopy; Springer Nature: Berlin/Heidelberg, Germany, 2019.
98. Pennycook, S.J. A Scan Through the History of STEM. In Scanning Transmission Electron Microscopy; Springer: New York, NY,
USA, 2011; pp. 1–90.
99. Pennycook, S.J.; Nellist, P.D. (Eds.) 2011. Scanning Transmission Electron. Microscopy: Imaging and Analysis; Springer Science &
Business Media: Secaucus, NJ, USA, 2011; ISBN 978-1-4419-7200-2.
100. Bacon, N.J.; Corbin, G.J.; Dellby, N.; Hrncirik, P.; Krivanek, O.L.; McManama-Smith, A.; Murfitt, M.F.; Szilagyi, Z.S.Nion
UltraSTEM: An aberration-corrected STEM for imaging and analysis. Microsc. Microanal. 2005, 11, 1422–1423.
101. MacLaren, I.; Ramasse, Q.M. Aberration-corrected scanning transmission electron microscopy for atomic-resolution studies of
functional oxides. Int. Mater. Rev. 2014, 59, 115–131.
102. Botton, G.; Prabhudev, S. Analytical electron microscopy. In Springer Handbook of Microscopy; Springer: Cham, Switzerland, 2019;
pp. 345–453.
103. Hren, J. Introduction to Analytical Electron. Microscopy; Springer Science & Business Media: Secaucus, NJ, USA, 2013; ISBN 978-
1-4757-5581-7.

Crystals 2021, 11, 878 52 of 58
104. Shindo, D.; Oikawa, T. Analytical Electron. Microscopy for Materials Science; Springer Science & Business Media: Secaucus, NJ,
USA, 2013; ISBN 978-4-431-66988-3.
105. Joy, D.C.; Roming, A.D., Jr.; Goldstein, J.; Goldstein, J.I. Principles of Analytical Electron. Microscopy; Springer Science & Business
Media: Secaucus, NJ, USA, 1986; ISBN 978-0-306-42387-1.
106. Scott, M.C.; Chen, C.C.; Mecklenburg, M.; Zhu, C.; Xu, R.; Ercius, P.; Dahmen, U.; Regan, B.C.; Miao, J. Electron tomography at
2.4-ångström resolution. Nature 2012, 483, 444–447.
107. De, L.O.; Roglie, B.; Lecture, N. The wave nature of the electron. Nobel Lect. 1929, 12, 244–256.
108. Ruska, E. The Development of the Electron Microscope and of Electron Microscopy (Nobel Lecture). Angew. Chem. Int. Ed. Engl.
1987, 26, 595–605.
109. Haider, M.; Rose, H.; Uhlemann, S.; Schwan, E.; Kabius, B.; Urban, K. A spherical-aberration-corrected 200 kV transmission
electron microscope. Ultramicroscopy 1998, 75, 53–60.
110. Batson, P.E.; Dellby, N.; Krivanek, O.L. Sub-ångstrom resolution using aberration corrected electron optics. Nature 2002, 418,
617–620.
111. Yamashita, S.; Kikkawa, J.; Yanagisawa, K.; Nagai, T.; Ishizuka, K.; Kimoto, K. Atomic number dependence of Z contrast in
scanning transmission electron microscopy. Sci. Rep. 2018, 8, 12325–12325.
112. Hawkes, P.W. The correction of electron lens aberrations. Ultramicroscopy 2015, 156, A1–A64.
113. Hart, J.L.; Lang, A.C.; Leff, A.C.; Longo, P.; Trevor, C.; Twesten, R.D.; Taheri, M.L. Direct detection electron energy-loss
spectroscopy: A method to push the limits of resolution and sensitivity. Sci. Rep. 2017, 7, 1–14.
114. Zou, X.; Hovmöller, S.; Oleynikov, P. Electron Crystallography: Electron Microscopy and Electron Diffraction; Oxford University
Press: Oxford, UK, 2012; Volume 9780199580, p. 344, ISBN 978-0-19-173121-1.
115. Abelson, A.; Qian, C.; Salk, T.; Luan, Z.; Fu, K.; Zheng, J.G.; Wardini, J.L.; Law, M. Collective topo-epitaxy in the self-assembly
of a 3D quantum dot superlattice. Nat. Mater. 2020, 19, 49–55.
116. Plapcianu, C.; Valsangiacom, C.; Schaffer, J.E.; Wieg, A.; Garay, J.; Stanciu, L. Spark plasma sintering studies of nanosize
lanthanide-doped ceria obtained by sol-gel method. J. Optoelectron. Adv. Mater. 2011, 13, 1101.
117. Fultz, B.; Howe, J.M. Transmission Electron. Microscopy and Diffractometry of Materials; Springer Science & Business Media:
Secaucus, NJ, USA, 2012.
118. Kotaka, Y. Direct visualization method of the atomic structure of light and heavy atoms with double-detector Cs-corrected
scanning transmission electron microscopy. Appl. Phys. Lett. 2012, 101, 133107–133107.
119. Krivanek, O.L.; Chisholm, M.F.; Nicolosi, V.; Pennycook, T.J.; Corbin, G.J.; Dellby, N.; Murfitt, M.F.; Own, C.S.; Szilagyi, Z.S.;
Oxley, M.P.; et al. Atom-by-atom structural and chemical analysis by annular dark-field electron microscopy. Nature 2010, 464,
571–574.
120. Findlay, S.D.; Shibata, N.; Sawada, H.; Okunishi, E.; Kondo, Y.; Ikuhara, Y. Dynamics of annular bright field imaging in scanning
transmission electron microscopy. Ultramicroscopy 2010, 110, 903–923.
121. Findlay, S.D.; Saito, T.; Shibata, N.; Sato, Y.; Matsuda, J.; Asano, K.; Akiba, E.; Hirayama, T.; Ikuhara, Y. Direct imaging of
hydrogen within a crystalline environment. Appl. Phys. Express 2010, 3, 116603–116603.
122. Hirsch, S.G.; Walck, S.D.; LaSalvia, J.C.; Swab, J.J. Transmission Electron Microscopy Characterization of Knoop Indentation Inelastic
Deformation Regions in Three Commercial Silicon Carbides; US Army Research Laboratory: Adelphi, MD, USA, 2018.
123. Giannuzzi, L.A.; Stevie, F.A. A review of focused ion beam milling techniques for TEM specimen preparation. Micron 1999, 30,
197–204.
124. Ayache, J.; Beaunier, L.; Boumendil, J.; Ehret, G.; Laub, D. Sample Preparation Handbook for Transmission Electron. Microscopy:
Techniques; Springer Science & Business Media: Secaucus, NJ, USA, 2020; Volume 2.
125. Rigort, A.; Plitzko, J.M. Cryo-focused-ion-beam applications in structural biology. Arch. Biochem. Biophys. 2015, 581, 122–130.
126. Zachman, M.J.; Asenath-Smith, E.; Estroff, L.A.; Kourkoutis, L.F. Site-specific preparation of intact solid-liquid interfaces by
label-free in situ localization and Cryo-focused ion beam lift-out. Microsc. Microanal. 2016, 22, 1338–1349.
127. Barna, Á.; Pécz, B.; Menyhard, M. TEM sample preparation by ion milling/amorphization. Micron 1999, 30, 267–276.
128. Ünlü, N. Preparation of high quality Al TEM specimens via a double-jet electropolishing technique. Mater. Charact. 2008, 59,
547–553.
129. Pourbabak, S.; Orekhov, A.; Schryvers, D. Twin-jet electropolishing for damage-free transmission electron microscopy specimen
preparation of metallic microwires. Microsc. Res. Tech. 2021, 84, 298–304.
130. Baena, V.; Schalek, R.L.; Lichtman, J.W.; Terasaki, M. Serial-section electron microscopy using automated tape-collecting
ultramicrotome (ATUM). In Methods in Cell Biology; Academic Press Inc.: London, UK, 2019; Volume 152, pp. 41–67, ISBN 978-
0-12-817018-2.
131. Schrand, A.M.; Schlager, J.J.; Dai, L.; Hussain, S.M. Preparation of cells for assessing ultrastructural localization of nanoparticles
with transmission electron microscopy. Nat. Protoc. 2010, 5, 744–757.
132. Sugar, J.D.; El Gabaly, F.; Chueh, W.C.; Fenton, K.R.; Tyliszczak, T.; Kotula, P.G.; Bartelt, N.C. High-resolution chemical analysis
on cycled LiFePO4 battery electrodes using energy-filtered transmission electron microscopy. J. Power Sour. 2014, 246, 512–521.
133. Fischione, P.E.; Williams, R.E.; Genç, A.; Fraser, H.L.; Dunin-Borkowski, R.E.; Luysberg, M.; Bonifacio, C.S.; Kovács, A. A small
spot, inert gas, ion milling process as a complementary technique to focused ion beam specimen preparation. Microsc. Microanal.
2017, 23, 782–793.

Crystals 2021, 11, 878 53 of 58
134. Van Mierlo, W.; Geiger, D.; Robins, A.; Stumpf, M.; Ray, M.L.; Fischione, P.; Kaiser, U. Practical aspects of the use of the X2
holder for HRTEM-quality TEM sample preparation by FIB. Ultramicroscopy 2014, 147, 149–155.
135. Wang, H.; Srot, V.; Fenk, B.; Laskin, G.; Mannhart, J.; van Aken, P.A. An optimized TEM specimen preparation method of
quantum nanostructures. Micron 2021, 140, 102979–102979.
136. Liu, S.-S.; Toh, S.; Daio, T.; Koyama, M.; Matsumura, S. Microstructure Observation of Ni/YSZ Boundary by TEM and STEM.
ECS Trans. 2013, 57, 1401–1405.
137. Srot, V.; Wang, Y.; Salzberger, U.; Fenk, B.; Kelsch, M.; Minola, M.; Salluzzo, M.; De Luca, G.M.; Keimer, B.; van Aken, P.A.
Improved sample preparation of beam-sensitive ultra-thin cuprate films. Microsc. Microanal. 2019, 25, 686–687.
138. Aitkaliyeva, A.; Madden, J.W.; Miller, B.D.; Cole, J.I.; Gan, J. Comparison of preparation techniques for nuclear materials for
transmission electron microscopy (TEM). J. Nucl. Mater. 2015, 459, 241–246.
139. Ennos, A.E. The origin of specimen contamination in the electron microscope. Br. J. Appl. Phys. 1953, 101–106.
140. Isabell, T.C.; Fischione, P.E.; O’Keefe, C.; Guruz, M.U.; Dravid, V.P. Plasma cleaning and its applications for electron microscopy.
Microsc. Microanal. 1999, 5, 126–135.
141. Bennett, J.C.; Egerton, R.F. NiO Test Specimens for Analytical Electron Microscopy: Round-Robin Results. Microsc. Microanal.
1995, 1, 143–149.
142. Egerton, R.F. An introduction to EELS. In Electron Energy-Loss Spectroscopy in the Electron Microscope; Springer: Berlin/Heidelberg,
Germany, 2011; pp. 1–28.
143. Krivanek, O.L.; Ahn, C.C.; Keeney, R.B. Parallel detection electron spectrometer using quadrupole lenses. Ultramicroscopy 1987,
22, 103–115.
144. Menon, N.K.; Krivanek, O.L. Synthesis of Electron Energy Loss Spectra for the Quantification of Detection Limits. Microsc.
Microanal. 2002, 8, 203–215.
145. Crozier, P.A. Quantitative elemental mapping of materials by energy-filtered imaging. Ultramicroscopy 1995, 58, 157–174.
146. Egerton, R.F.; Cheng, S.C. Measurement of local thickness by electron energy-loss spectroscopy. Ultramicroscopy 1987, 21, 231–
244.
147. Malis, T.; Cheng, S.C.; Egerton, R.F. EELS log-ratio technique for specimen-thickness measurement in the TEM. J. Electron.
Microsc. Tech. 1988, 8, 193–200.
148. Zhang, F.; Vanmeensel, K.; Batuk, M.; Hadermann, J.; Inokoshi, M.; van Meerbeek, B.; Naert, I.; Vleugels, J. Highly-translucent,
strong and aging-resistant 3Y-TZP ceramics for dental restoration by grain boundary segregation. Acta Biomater. 2015, 16, 215–
222.
149. Zhang, Z.; Rauch, E.F.; Véron, M. Twinning analyses in a magnesium alloy with tilting series scanning method using a TEM
based orientation mapping system. Mater. Lett. 2013, 111, 192–196.
150. Overwijk, M.H.F.; Reefman, D. Maximum-entropy deconvolution applied to electron energy-loss spectroscopy. Micron 2000,
31, 325–331.
151. Egerton, R.F. Limits to the spatial, energy and momentum resolution of electron energy-loss spectroscopy. Ultramicroscopy 2007,
107, 575–586.
152. Nelayah, J.; Kociak, M.; Stéphan, O.; de Abajo, F.J.G.; Tencé, M.; Henrard, L.; Taverna, D.; Pastoriza-Santos, I.; Liz-Marzán, L.M.;
Colliex, C. Mapping surface plasmons on a single metallic nanoparticle. Nat. Phys. 2007, 3, 348–353.
153. Kimoto, K.; Kothleitner, G.; Grogger, W.; Matsui, Y.; Hofer, F. Advantages of a monochromator for bandgap measurements
using electron energy-loss spectroscopy. Micron 2005, 36, 185–189.
154. Eccles, J.W.L.; Bangert, U.; Bromfield, M.; Christian, P.; Harvey, A.J.; Thomas, P. UV-Vis plasmon studies of metal nanoparticles.
In Journal of Physics: Conference Series (Vol. 241, No. 1, p. 012090). IOP Publishing. J. Phys. Conf. Ser. 2010, 241, 012090.
155. Aguiar, J.A.; Reed, B.W.; Ramasse, Q.M.; Erni, R.; Browning, N.D. Quantifying the low-energy limit and spectral resolution in
valence electron energy loss spectroscopy. Ultramicroscopy 2013, 124, 130–138.
156. Jeangros, Q.; Faes, A.; Wagner, J.B.; Hansen, T.W.; Aschauer, U.; Hessler-Wyser, A.; Dunin-Borkowski, R.E. In situ redox cycle
of a nickel–YSZ fuel cell anode in an environmental transmission electron microscope. Acta Mater. 2010, 58, 4578–4589.
157. Rouviere, J.L.; Béché, A.; Martin, Y.; Denneulin, T.; Cooper, D. Improved strain precision with high spatial resolution using
nanobeam precession electron diffraction. Appl. Phys. Lett. 2013, 103, 241913–241913.
158. Cizek, P.; Sankaran, A.; Rauch, E.F.; Barnett, M.R. Observation of (sub)grain clusters in the as-deposited and in situ annealed
nanocrystalline nickel using automated crystal orientation mapping. Scr. Mater. 2012, 67, 685–688.
159. Krivanek, O.L.; Lovejoy, T.C.; Dellby, N.; Aoki, T.; Carpenter, R.W.; Rez, P.; Soignard, E.; Zhu, J.; Batson, P.E.; Lagos, M.J.; et al.
Vibrational spectroscopy in the electron microscope. Nature 2014, 514, 209–212.
160. Bowman, W.J.; March, K.; Hernandez, C.A.; Crozier, P.A. Measuring bandgap states in individual non-stoichiometric oxide
nanoparticles using monochromated STEM EELS: The Praseodymium–ceria case. Ultramicroscopy 2016, 167, 5–10.
161. Wei, J.; Ogawa, T.; Feng, B.; Yokoi, T.; Ishikawa, R.; Kuwabara, A.; Matsunaga, K.; Shibata, N.; Ikuhara, Y. Direct measurement
of electronic band structures at oxide grain boundaries. Nano Lett. 2020, 20, 2530–2536.
162. Krivanek, O.L.; Paterson, J.H. Elnes of 3d transition-metal oxides. I. Variations across the periodic table. Ultramicroscopy 1990,
32, 313–318.
163. Manoubi, T.; Colliex, C.; Rez, P. Quantitative electron energy loss spectroscopy on M45 edges in rare earth oxides. J. Electron.
Spectrosc. Relat. Phenom. 1990, 50, 1–18.

Crystals 2021, 11, 878 54 of 58
164. Colliex, C.; Cosslett, V.E.; Leapman, R.D.; Trebbia, P. Contribution of electron energy loss spectroscopy to the development of
analytical electron microscopy. Ultramicroscopy 1976, 1, 301–315.
165. Vincent, R.; Midgley, P.A. Double conical beam-rocking system for measurement of integrated electron diffraction intensities.
Ultramicroscopy 1994, 53, 271–282.
166. Frechero, M.A.; Rocci, M.; Sánchez-Santolino, G.; Kumar, A.; Salafranca, J.; Schmidt, R.; Díaz-Guillén, M.R.; Durá, O.J.; Rivera-
Calzada, A.; Mishra, R.; et al. Paving the way to nanoionics: Atomic origin of barriers for ionic transport through interfaces. Sci.
Rep. 2015, 5, 17229–17229.
167. Stöger-Pollach, M.; Bukvišová, K.; Schwarz, S.; Kvapil, M.; Šamořil, T.; Horák, M. Fundamentals of cathodoluminescence in a
STEM: The impact of sample geometry and electron beam energy on light emission of semiconductors. Ultramicroscopy 2019,
200, 111–124.
168. Beck, G. Contribution to the Theory of the Cherenkov Effect. Phys. Rev. 1948, 74, 795–802.
169. Chen, J.; Sekiguchi, T.; Yang, D.; Yin, F.; Kido, K.; Tsurekawa, S. Electron-beam-induced current study of grain boundaries in
multicrystalline silicon. J. Appl. Phys. 2004, 96, 5490–5495.
170. Chernyak, L.; Osinsky, A.; Temkin, H.; Yang, J.W.; Chen, Q.; Asif Khan, M. Electron beam induced current measurements of
minority carrier diffusion length in gallium nitride. Appl. Phys. Lett. 1996, 69, 2531–2533.
171. Leach, C. Grain boundary structures in zinc oxide varistors. Acta Mater. 2005, 53, 237–245.
172. Leach, C. Crystal plane influence of the EBIC contrast in zinc oxide varistors. J. Eur. Ceram. Soc. 2001, 21, 2127–2130.
173. Bhattacharyya, A.; Reimer, J.D.; Ritz, K.N. Breakdown voltage characteristics of thin oxides and their correlation to defects in
the oxide as observed by the EBIC technique. IEEE Electron. Device Lett. 1986, 7, 58–60.
174. Midgley, P.A.; Eggeman, A.S. Precession electron diffraction—A topical review. IUCrJ 2015, 2, 126–136.
175. Ophus, C. Four-Dimensional Scanning Transmission Electron Microscopy (4D-STEM): From Scanning Nanodiffraction to
Ptychography and Beyond. Microsc. Microanal. 2019, 25, 563–582.
176. Watanabe, M.; Williams, D. Development of Diffraction Imaging for Orientation Analysis of Grains in Scanning Transmission
Electron Microscopy. Microsc. Microanal. 2007, 13, 962–963.
177. Bober, D.B.; Kumar, M.; Rupert, T.J. Nanocrystalline grain boundary engineering: Increasing Σ3 boundary fraction in pure Ni
with thermomechanical treatments. Acta Mater. 2015, 86, 43–54.
178. Izadi, E.; Darbal, A.; Sarkar, R.; Rajagopalan, J. Grain rotations in ultrafine-grained aluminum films studied using in situ TEM
straining with automated crystal orientation mapping. Mater. Des. 2017, 113, 186–194.
179. Idrissi, H.; Kobler, A.; Amin-Ahmadi, B.; Coulombier, M.; Galceran, M.; Raskin, J.P.; Godet, S.; Kübel, C.; Pardoen, T.; Schryvers,
D. Plasticity mechanisms in ultrafine grained freestanding aluminum thin films revealed by in-situ transmission electron
microscopy nanomechanical testing. Appl. Phys. Lett. 2014, 104, 101903.
180. Garner, A.; Gholinia, A.; Frankel, P.; Gass, M.; MacLaren, I.; Preuss, M. The microstructure and microtexture of zirconium oxide
films studied by transmission electron backscatter diffraction and automated crystal orientation mapping with transmission
electron microscopy. Acta Mater. 2014, 80, 159–171.
181. Guo, D.; Song, S.; Luo, R.; Goddard, W.A. III.; Chen, M.; Reddy, K.M.; An, Q. Grain boundary sliding and amorphization are
responsible for the reverse Hall-Petch relation in superhard nanocrystalline boron carbide. Phys. Rev. Lett. 2018, 121, 145504–
145504.
182. Harks, P.P.R.M.L.; Mulder, F.M.; Notten, P.H.L. In situ methods for Li-ion battery research: A review of recent developments.
J. Power Sources 2015, 288, 92–105.
183. Wu, F.; Yao, N. Advances in windowed gas cells for in-situ TEM studies. Nano Energy 2015, 13, 735–756.
184. Miller, B.K. Development and Application of Operando TEM to a Ruthenium Catalyst for CO Oxidation; Arizona State University:
Phoenix, AR, USA, 2016.
185. Jiang, N. Electron beam damage in oxides: A review. Rep. Prog. Phys. 2015, 79, 016501–016501.
186. Egerton, R.F. Radiation damage to organic and inorganic specimens in the TEM. Micron 2019, 119, 72–87.
187. Jiang, N.; Spence, J.C.H. Electronic ionization induced atom migration in spinel MgAl 2O4. J. Nucl. Mater. 2010, 403, 147–151.
188. Smith, R.; Bacorisen, D.; Uberuaga, B.P.; Sickafus, K.E.; Ball, J.A.; Grimes, R.W. Dynamical simulations of radiation damage in
magnesium aluminate spinel, MgAl2O4. J. Phys. Condens. Matter 2005, 17, 875–891.
189. Ishimaru, M.; Afanasyev-Charkin, I.V.; Sickafus, K.E. Ion-beam-induced spinel-to-rocksalt structural phase transformation in
MgAl2O4. Appl. Phys. Lett. 2000, 76, 2556–2558.
190. Bouchet, D.; Colliex, C. Experimental study of ELNES at grain boundaries in alumina: Intergranular radiation damage effects
on Al-L23 and O-K edges. Ultramicroscopy 2003, 96, 139–152.
191. Nakamura, R.; Ishimaru, M.; Yasuda, H.; Nakajima, H. Atomic rearrangements in amorphous Al2O3 under electron-beam
irradiation. J. Appl. Phys. 2013, 113, 064312–064312.
192. Yao, L.; Majumdar, S.; Äkäslompolo, L.; Inkinen, S.; Qin, Q.H.; van Dijken, S. Electron-Beam-Induced Perovskite–
Brownmillerite–Perovskite Structural Phase Transitions in Epitaxial La2/3Sr1/3MnO3 Films. Adv. Mater. 2014, 26, 2789–2793.
193. Nord, M.; Vullum, P.E.; Hallsteinsen, I.; Tybell, T.; Holmestad, R. Assessing electron beam sensitivity for SrTiO3 and
La0.7Sr0.3MnO3 using electron energy loss spectroscopy. Ultramicroscopy 2016, 169, 98–106.
194. Lin, Y.; Wen, J.; Hu, L.; McCarthy, J.A.; Wang, S.; Poeppelmeier, K.R.; Marks, L.D. Electron-induced Ti-rich surface segregation
on SrTiO3 nanoparticles. Micron 2015, 68, 152–157.
195. West, G.D.; Perkins, J.M.; Lewis, M.H. Characterisation of fine-grained oxide ceramics. J. Mater. Sci. 2004, 39, 6687–6704.

Crystals 2021, 11, 878 55 of 58
196. Winterstein, J.P.; Carter, C.B. Electron-beam damage and point defects near grain boundaries in cerium oxide. J. Eur. Ceram. Soc.
2014, 34, 3007–3018.
197. Yoshida, H.; Kuwabara, A.; Yamamoto, T.; Ikuhara, Y.; Sakuma, T. High temperature plastic flow and grain boundary chemistry
in oxide ceramics. J. Mater. Sci. 2005, 40, 3129–3135.
198. Dillon, S.J.; Harmer, M.P. Multiple grain boundary transitions in ceramics: A case study of alumina. Acta Mater. 2007, 55, 5247–
5254.
199. Höche, T.; Kenway, P.R.; Kleebe, H.-J.; Rühle, M.; Morris, P.A. High-Resolution Transmission Electron Microscopy Studies of a
Near Σ11 Grain Boundary in α-Alumina. J. Am. Ceram. Soc. 1994, 77, 339–348.
200. Matsunaga, K.; Nishimura, H.; Hanyu, S.; Muto, H.; Yamamoto, T.; Ikuhara, Y. HRTEM study on grain boundary atomic
structures related to the sliding behavior in alumina bicrystals. Appl. Surf. Sci. 2005, 241, 75–79.
201. Matsui, K.; Yoshida, H.; Ikuhara, Y. Nanocrystalline, Ultra-Degradation-Resistant Zirconia: Its Grain Boundary Nanostructure
and Nanochemistry. Sci. Rep. 2014, 4, 4758.
202. Stemmer, S.; Vleugels, J.; Van Der Biest, O. Grain boundary segregation in high-purity, yttria-stabilized tetragonal zirconia
polycrystals (Y-TZP). J. Eur. Ceram. Soc. 1998, 18, 1565–1570.
203. Sokol, M.; Ratzker, B.; Kalabukhov, S.; Dariel, M.P.; Galun, E.; Frage, N. Transparent polycrystalline magnesium aluminate
spinel fabricated by spark plasma sintering. Adv. Mater. 2018, 30, 1706283.
204. Nuns, N.; Béclin, F.; Crampon, J. Grain-Boundary Characterization in a Nonstoichiometric Fine-Grained Magnesium Aluminate
Spinel: Effects of Defect Segregation at the Space-Charge Layers. J. Am. Ceram. Soc. 2009, 92, 870–875.
205. Avila-Paredes, H.J.; Kim, S. The effect of segregated transition metal ions on the grain boundary resistivity of gadolinium doped
ceria: Alteration of the space charge potential. Solid State Ion. 2006, 177, 3075–3080.
206. Bueno, P.R.; Leite, E.R.; Oliveira, M.M.; Orlandi, M.O.; Longo, E. Role of oxygen at the grain boundary of metal oxide varistors:
A potential barrier formation mechanism. Appl. Phys. Lett. 2001, 79, 48–50.
207. Zhang, Z.; Sigle, W.; Rühle, M.; Jud, E.; Gauckler, L.J. Microstructure characterization of a cobalt-oxide-doped cerium-
gadolinium-oxide by analytical and high-resolution TEM. Acta Mater. 2007, 55, 2907–2917.
208. Merkle, K.L.; Smith, D.J. Atomic Structure of Symmetric Tilt Grain Boundaries in NiO. Phys. Rev. Lett. 1987, 59, 2887–2890.
209. Sternlicht, H.; Bojarski, S.A.; Rohrer, G.S.; Kaplan, W.D. Quantitative differences in the Y grain boundary excess at boundaries
delimiting large and small grains in Y doped Al2O3. J. Eur. Ceram. Soc. 2018, 38, 1829–1835.
210. Sternlicht, H.; Rheinheimer, W.; Mehlmann, A.; Rothschild, A.; Hoffmann, M.J.; Kaplan, W.D. The mechanism of grain growth
at general grain boundaries in SrTiO3. Scr. Mater. 2020, 188, 206–211.
211. Jia, C.L.; Urban, K. Atomic-Resolution Measurement of Oxygen Concentration in Oxide Materials. Science 2004, 303, 2001–2004.
212. Feng, B.; Yokoi, T.; Kumamoto, A.; Yoshiya, M.; Ikuhara, Y.; Shibata, N. Atomically ordered solute segregation behaviour in an
oxide grain boundary. Nat. Commun. 2016, 7, 1–6.
213. Chiang, Y.; Wang, H.; Lee, J. HREM and STEM of intergranular films at zinc oxide varistor grain boundaries. J. Microsc. 1998,
191, 275–285.
214. Winnubst, A.J.A.; Kroot, P.J.M.; Burggraaf, A.J. AES/STEM grain boundary analysis of stabilized zirconia ceramics. J. Phys. Chem.
Solids 1983, 44, 955–960.
215. Pyrz, W.D.; Blom, D.A.; Sadakane, M.; Kodato, K.; Ueda, W.; Vogt, T.; Buttrey, D.J. Atomic-level imaging of Mo-VO complex
oxide phase intergrowth, grain boundaries, and defects using HAADF-STEM. Proc. Natl. Acad. Sci. USA 2010, 107, 6152–6157.
216. Schusteritsch, G.; Ishikawa, R.; Elmaslmane, A.R.; Inoue, K.; McKenna, K.P.; Ikuhara, Y.; Pickard, C.J. Anataselike Grain
Boundary Structure in Rutile Titanium Dioxide. Nano Lett. 2021, 21, 2745–2751.
217. Ikuhara, Y. Grain boundary atomic structures and light-element visualization in ceramics: Combination of Cs-corrected
scanning transmission electron microscopy and first-principles calculations. Electron. Microsc. 2011, 60, S173–S188.
218. Gale, J.D. GULP: A computer program for the symmetry-adapted simulation of solids. J. Chem. Soc. Faraday Trans. 1997, 93, 629–
637.
219. Tong, W.; Yang, H.; Moeck, P.; Nandasiri, M.I.; Browning, N.D. General schema for [0 0 1] tilt grain boundaries in dense packing
cubic crystals. Acta Mater. 2013, 61, 3392–3398.
220. Yuan, R.; Zhang, J.; He, L.; Zuo, J.-M. Training Artificial Neural Networks for Precision Orientation and Strain Mapping Using 4D
Electron. Diffraction Datasets; Elsevier: Amesterdam, Netherlands, 2021.
221. Boland, T.M.; Rez, P.; Crozier, P.A.; Singh, A.K. Impact of Aliovalent Alkaline-Earth metal solutes on Ceria Grain Boundaries:
A density functional theory study. Acta Mater. 2021, 205, 116481.
222. Ishihara, S.; Tochigi, E.; Ishikawa, R.; Shibata, N.; Ikuhara, Y. Atomic structures of Ti-doped α-Al2O3 Σ13 grain boundary with
a small amount of Si impurity. J. Am. Ceram. Soc. 2020, 103, 6659–6665.
223. Saito, M.; Wang, Z.; Ikuhara, Y. Selective impurity segregation at a near-Σ5 grain boundary in MgO. J. Mater. Sci. 2014, 49, 3956–
3961.
224. Yang, H.; Kotula, P.G.; Sato, Y.; Chi, M.; Ikuhara, Y.; Browning, N.D. Segregation of Mn2+ dopants as interstitials in SrTiO3
grain boundaries. Mater. Res. Lett. 2014, 2, 16–22.
225. Song, Z.; Xie, Z.H. A literature review of in situ transmission electron microscopy technique in corrosion studies. Micron 2018,
112, 69–83.
226. Cho, J.; Phuah, X.L.; Li, J.; Shang, Z.; Wang, H.; Charalambous, H.; Tsakalakos, T.; Mukherjee, A.K.; Wang, H.; Zhang, X.
Temperature effect on mechanical response of flash-sintered ZnO by in-situ compression tests. Acta Mater. 2020, 200, 699–709.

Crystals 2021, 11, 878 56 of 58
227. Kondo, S.; Shibata, N.; Mitsuma, T.; Tochigi, E.; Ikuhara, Y. Dynamic observations of dislocation behavior in SrTiO 3 by in situ
nanoindentation in a transmission electron microscope. Appl. Phys. Lett. 2012, 100, 1–5.
228. Kondo, S.; Mitsuma, T.; Shibata, N.; Ikuhara, Y. Direct observation of individual dislocation interaction processes with grain
boundaries. Sci. Adv. 2016, 2, 1–8.
229. Grosso, R.L.; Muccillo, E.N.; Muche, D.N.; Jawaharram, G.S.; Barr, C.M.; Monterrosa, A.M.; Castro, R.H.; Hattar, K.; Dillon, S.J.
In situ transmission electron microscopy for ultrahigh temperature mechanical testing of ZrO2. Nano Lett. 2020, 20, 1041–1046.
230. Vikrant, K.S.N.; Grosso, R.L.; Feng, L.; Muccillo, E.N.; Muche, D.N.; Jawaharram, G.S.; Barr, C.M.; Monterrosa, A.M.; Castro,
R.H.; García, R.E.; et al. Ultrahigh temperature in situ transmission electron microscopy based bicrystal coble creep in zirconia
I: Nanowire growth and interfacial diffusivity. Acta Mater. 2020, 199, 530–541.
231. Gao, P.; Wang, Z.; Fu, W.; Liao, Z.; Liu, K.; Wang, W.; Bai, X.; Wang, E. In situ TEM studies of oxygen vacancy migration for
electrically induced resistance change effect in cerium oxides. Micron 2010, 41, 301–305.
232. Zhang, Z. Surface effects in the energy loss near edge structure of different cobalt oxides. Ultramicroscopy 2007, 107, 598–603.
233. Stemmer, S.; Sane, A.; Browning, N.D.; Mazanec, T.J. Characterization of oxygen-deficient SrCoO3-δ by electron energy-loss
spectroscopy and Z-contrast imaging. Solid State Ion. 2000, 130, 71–80.
234. Klie, R.F.; Zheng, J.C.; Zhu, Y.; Varela, M.; Wu, J.; Leighton, C. Direct measurement of the low-temperature spin-state transition
in LaCoO 3. Phys. Rev. Lett. 2007, 99, 047203–047203.
235. Klie, R.F.; Buban, J.P.; Varela, M.; Franceschetti, A.; Jooss, C.; Zhu, Y.; Browning, N.D.; Pantelides, S.T.; Pennycook, S.J. Enhanced
current transport at grain boundaries in high-T c superconductors. Nature 2005, 435, 475–478.
236. Browning, N.D.; Buban, J.P.; Nellist, P.D.; Norton, D.P.; Chisholm, M.F.; Pennycook, S.J. The atomic origins of reduced critical
currents at [001] tilt grain boundaries in YBa2Cu3O7− δ thin films. Phys. C Supercond. Its Appl. 1998, 294, 183–193.
237. Zhao, Y.; Feltes, T.E.; Regalbuto, J.R.; Meyer, R.J.; Klie, R.F. In Situ Electron Energy Loss Spectroscopy Study of Metallic Co and Co
Oxides; American Institute of Physics AIP: College Park, MD, USA, 2010; Volume 108, p. 063704.
238. Song, X.; Daniels, G.; Feldmann, D.M.; Gurevich, A.; Larbalestier, D. Electromagnetic, atomic structure and chemistry changes
induced by Ca-doping of low-angle YBa2Cu3O7-δ grain boundaries. Nat. Mater. 2005, 4, 470–475.
239. Bischoff, J.; Motta, A.T. EFTEM and EELS analysis of the oxide layer formed on HCM12A exposed to SCW. J. Nucl. Mater. 2012,
430, 171–180.
240. Jiang, N.; Spence, J.C.H. Interpretation of Oxygen K pre-edge peak in complex oxides. Ultramicroscopy 2006, 106, 215–219.
241. Wang, Z.L.; Yin, J.S.; Jiang, Y.D. EELS analysis of cation valence states and oxygen vacancies in magnetic oxides Micron 2000,
31, 571–580.
242. Yuan, Y.; Amine, K.; Lu, J.; Shahbazian-Yassar, R. Understanding materials challenges for rechargeable ion batteries with in situ
transmission electron microscopy. Nat. Commun. 2017, 8, 1–14.
243. Sun, Y.Y.; Hou, P.Y.; Zhang, L.C. Mitigating the Microcracks of High-Ni Oxides by in Situ Formation of Binder between
Anisotropic Grains for Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2020, 12, 13923–13930.
244. Park, K.J.; Hwang, J.Y.; Ryu, H.H.; Maglia, F.; Kim, S.J.; Lamp, P.; Yoon, C.S.; Sun, Y.K. Degradation mechanism of Ni-enriched
NCA cathode for lithium batteries: Are microcracks really critical? ACS Energy Lett. 2019, 4, 1394–1400.
245. Liu, X.; Zhan, X.; Hood, Z.D.; Li, W.; Leonard, D.N.; Manthiram, A.; Chi, M. Essential effect of the electrolyte on the mechanical
and chemical degradation of LiNi 0.8 Co 0.15 Al 0.05 O 2 cathodes upon long-term cycling. J. Mater. Chem. A 2021, 9, 2111–2119.
246. Jo, Y.R.; Koo, B.; Seo, M.J.; Kim, J.K.; Lee, S.; Kim, K.; Han, J.W.; Jung, W.; Kim, B.J. Growth kinetics of individual Co particles
Ex-solved on SrTi0. 75Co0. 25O3-δ polycrystalline perovskite thin films. J. Am. Chem. Soc. 2019, 141, 6690–6697.
247. Zajac, W.; Molenda, J. Electrical conductivity of doubly doped ceria. Solid State Ion. 2008, 179, 154–158.
248. Guo, X.; Waser, R. Electrical properties of the grain boundaries of oxygen ion conductors: Acceptor-doped zirconia and ceria.
Prog. Mater. Sci. 2006, 51, 151–210.
249. Steele, B.C.H.; Heinzel, A. Materials for fuel-cell technologies. Nature 2001, 414, 345–352.
250. Mogensen, M.; Sammes, N.M.; Tompsett, G.A. Physical, chemical and electrochemical properties of pure and doped ceria. Solid
State Ion. 2000, 129, 63–94.
251. Avila-Paredes, H.J.; Choi, K.; Chen, C.T.; Kim, S. Dopant-concentration dependence of grain-boundary conductivity in ceria: A
space-charge analysis. J. Mater. Chem. 2009, 19, 4837–4842.
252. Luo, J. Interfacial engineering of solid electrolytes. J. Mater. 2015, 1, 22–32.
253. Murugan, R.; Thangadurai, V.; Weppner, W. Fast Lithium Ion Conduction in Garnet-Type Li7La3Zr2O12. Angew. Chem. Int. Ed.
2007, 46, 7778–7781.
254. Li, Y.; Han, J.T.; Wang, C.A.; Xie, H.; Goodenough, J.B. Optimizing Li+ conductivity in a garnet framework. J. Mater. Chem. 2012,
22, 15357–15361.
255. Liu, Y.; Chen, J.; Gao, J. Preparation and chemical compatibility of lithium aluminum germanium phosphate solid electrolyte.
Solid State Ion. 2018, 318, 27–34.
256. Inaguma, Y.; Liquan, C.; Itoh, M.; Nakamura, T.; Uchida, T.; Ikuta, H.; Wakihara, M. High ionic conductivity in lithium
lanthanum titanate. Solid State Commun. 1993, 86, 689–693.
257. Kimura, K.; Wagatsuma, K.; Tojo, T.; Inada, R.; Sakurai, Y. Effect of composition on lithium-ion conductivity for perovskite-
type lithium-strontium-tantalum-zirconium-oxide solid electrolytes. Ceram. Int. 2016, 42, 5546–5552.
258. Cui, Y.; Mahmoud, M.M.; Rohde, M.; Ziebert, C.; Seifert, H.J. Thermal and ionic conductivity studies of lithium aluminum
germanium phosphate solid-state electrolyte. Solid State Ion. 2016, 289, 125–132.

Crystals 2021, 11, 878 57 of 58
259. Sharafi, A. Microstructural and Interface Engineering of Garnet-Type Fast Ion-Conductor for Use in Solid-State Batteries
(Doctoral dissertation). 2007. Available online: https://deepblue.lib.umich.edu/handle/2027.42/140865 (accessed on 19 May
2021).
260. Allen, J.L.; Wolfenstine, J.; Rangasamy, E.; Sakamoto, J. Effect of substitution (Ta, Al, Ga) on the conductivity of Li 7La 3Zr 2O
12. J. Power Sources 2012, 206, 315–319.
261. Hu, S.; Li, Y.F.; Yang, R.; Yang, Z.; Wang, L. Structure and ionic conductivity of Li7La3Zr2−xGexO12 garnet-like solid electrolyte
for all solid state lithium ion batteries. Ceram. Int. 2018, 44, 6614–6618.
262. Thompson, T.; Yu, S.; Williams, L.; Schmidt, R.D.; Garcia-Mendez, R.; Wolfenstine, J.; Allen, J.L.; Kioupakis, E.; Siegel, D.J.;
Sakamoto, J. Electrochemical window of the Li-ion solid electrolyte Li7La3Zr2O12. ACS Energy Lett. 2017, 2, 462–468.
263. Thangadurai, V.; Narayanan, S.; Pinzaru, D. Garnet-type solid-state fast Li ion conductors for Li batteries: Critical review. Chem.
Soc. Rev. 2014, 43, 4714–4727.
264. Ahmad, M.M. Estimation of the concentration and mobility of mobile Li+ in the cubic garnet-type Li7La3Zr2O12. RSC Adv.
2015, 5, 25824–25829.
265. Pickett, W.E.; Feldman, J.L.; Deppe, J. Thermal transport across boundaries in diamond structure materials. Model. Simul. Mater.
Sci. Eng. 1996, 4, 409–419.
266. Sood, A.; Cheaito, R.; Bai, T.; Kwon, H.; Wang, Y.; Li, C.; Yates, L.; Bougher, T.; Graham, S.; Asheghi, M.; et al. Direct
visualization of thermal conductivity suppression due to enhanced phonon scattering near individual grain boundaries. Nano
Lett. 2018, 18, 3466–3472.
267. Kinoshita, T.; Munekawa, S. Effect of grain boundary segregation on thermal conductivity of hot-pressed silicon carbide. Acta
Mater. 1997, 45, 2001–2012.
268. Boyle, C.; Carvillo, P.; Chen, Y.; Barbero, E.J.; Mcintyre, D.; Song, X. Grain boundary segregation and thermoelectric
performance enhancement of bismuth doped calcium cobaltite. J. Eur. Ceram. Soc. 2016, 36, 601–607.
269. Blundell, S.J.; Blundell, K.M. Concepts in Thermal Physics; Oxford University Press, Oxford, UK: 2010; Volume 9780199562091, p.
512, ISBN 978-0-19-171823-6.
270. Song, X.; Paredes Navia, S.A.; Liang, L.; Boyle, C.; Romo-De-La-Cruz, C.O.; Jackson, B.; Hinerman, A.; Wilt, M.; Prucz, J.; Chen,
Y. Grain boundary phase segregation for dramatic improvement of the thermoelectric performance of oxide ceramics. ACS Appl.
Mater. Interfaces 2018, 10, 39018–39024.
271. Shi, Z.; Su, T.; Zhang, P.; Lou, Z.; Qin, M.; Gao, T.; Xu, J.; Zhu, J.; Gao, F. Enhanced thermoelectric performance of Ca 3 Co 4 O
9 ceramics through grain orientation and interface modulation. J. Mater. Chem. A 2020, 8, 19561–19572.
272. Yang, H.S.; Bai, G.R.; Thompson, L.J.; Eastman, J.A. Interfacial thermal resistance in nanocrystalline yttria-stabilized zirconia.
Acta Mater. 2002, 50, 2309–2317.
273. Smith, D.S.; Fayette, S.; Grandjean, S.; Martin, C.; Telle, R.; Tonnessen, T. Thermal resistance of grain boundaries in alumina
ceramics and refractories. J. Am. Ceram. Soc. 2003, 86, 105–111.
274. Watanabe, T. Grain boundary engineering: Historical perspective and future prospects. J. Mater. Sci. 2011, 46, 4095–4115.
275. Krause, A.R.; Cantwell, P.R.; Marvel, C.J.; Compson, C.; Rickman, J.M.; Harmer, M.P. Review of grain boundary complexion
engineering: Know your boundaries. J. Am. Ceram. Soc. 2019, 102, 778–800.
276. Maehara, Y.; Langdon, T.G. Superplasticity in ceramics. J. Mater. Sci. 1990, 25, 2275–2286.
277. Ryou, H.; Drazin, J.W.; Wahl, K.J.; Qadri, S.B.; Gorzkowski, E.P.; Feigelson, B.N.; Wollmershauser, J.A. Below the hall–petch
limit in nanocrystalline ceramics. ACS Nano 2018, 12, 3083–3094.
278. Ikuhara, Y.; Yoshida, H.; Sakuma, T. Impurity effects on grain boundary strength in structural ceramics. Mater. Sci. Eng. A 2001,
319–321, 24–30.
279. Matsunaga, K.; Nishimura, H.; Muto, H.; Yamamoto, T.; Ikuhara, Y. Direct measurements of grain boundary sliding in yttrium-
doped alumina bicrystals. Appl. Phys. Lett. 2003, 82, 1179–1181.
280. Kondo, S.; Ishihara, A.; Tochigi, E.; Shibata, N.; Ikuhara, Y. Direct observation of atomic-scale fracture path within ceramic grain
boundary core. Nat. Commun. 2019, 10, 1–6.
281. Coleman, S.P.; Hernandez-Rivera, E.; Behler, K.D.; Synowczynski-Dunn, J.; Tschopp, M.A. Challenges of Engineering Grain
Boundaries in Boron-Based Armor Ceramics. JOM 2016, 68, 1605–1615.
282. Watanabe, T.; Tsurekawa, S. Control of brittleness and development of desirable mechanical properties in polycrystalline
systems by grain boundary engineering. Acta Mater. 1999, 47, 4171–4185.
283. Chen, M.W.; McCauley, J.W.; Dandekar, D.P.; Bourne, N.K. Dynamic plasticity and failure of high-purity alumina under shock
loading. Nat. Mater. 2006, 5, 614–618.
284. Reddy, K.M.; Liu, P.; Hirata, A.; Fujita, T.; Chen, M.W. Atomic structure of amorphous shear bands in boron carbide. Nat.
Commun. 2013, 4, 1–5.
285. Ju, H.L.; Sohn, H. Role of grain boundaries in double exchange manganite oxides La1-xAxMnO3 (A = Ba, Ca). Solid State Commun.
1997, 102, 463–466.
286. Straumal, B.B.; Mazilkin, A.A.; Protasova, S.G.; Myatiev, A.A.; Straumal, P.B.; Schütz, G.; Van Aken, P.A.; Goering, E.; Baretzky,
B. Magnetization study of nanograined pure and Mn-doped ZnO films: Formation of a ferromagnetic grain-boundary foam.
Phys. Rev. B Condens. Matter Mater. Phys. 2009, 79, 205206–205206.

Crystals 2021, 11, 878 58 of 58
287. Sepehri-Amin, H.; Tamazawa, Y.; Kambayashi, M.; Saito, G.; Takahashi, Y.K.; Ogawa, D.; Ohkubo, T.; Hirosawa, S.; Doi, M.;
Shima, T.; et al. Achievement of high coercivity in Sm (Fe0. 8Co0. 2) 12 anisotropic magnetic thin film by boron doping. Acta
Mater. 2020, 194, 337–342.
288. Zhang, J.; Wang, W.; Wang, T.; Jiang, L.; Wang, N.; Sun, D.; Zhao, X.; Wang, M.; Qi, Y. Nanoscale characterization of the doped
SrZrO3 nanoparticles distribution and its influence on the microstructure of Bi2Sr2CaCu2O8+ δ film. J. Alloy. Compd. 2021, 858,
157650.
289. Straumal, B.B.; Myatiev, A.A.; Straumal, P.B.; Mazilkin, A.A.; Protasova, S.G.; Goering, E.; Baretzky, B. Grain boundary layers
in nanocrystalline ferromagnetic zinc oxide. JETP Lett. 2010, 92, 396–400.
290. Park, Y.J.; Kang, J.H.; Kim, A.H.; Kim, T.H.; Lim, T.Y.; Kim, D.H. Structural, electrical, and magnetic properties of BiFeO3-
Y3Fe5O12 bulk ceramics and sputtered thin films. J. Magn. Magn. Mater. 2021, 535, 168058.
291. Li, S.; Pan, J.; Gao, F.; Zeng, D.; Qin, F.; He, C.; Dodbiba, G.; Wei, Y.; Fujita, T. Structure and magnetic properties of coprecipitated
nickel-zinc ferrite-doped rare earth elements of Sc, Dy, and Gd. J. Mater. Sci. Mater. Electron. 2021, 1–16, doi:10.21203/rs.3.rs-
277352/v1.
292. Chen, C.F.; Brennecka, G.L.; Synowicki, R.A.; Tegtmeier, E.L.; Brand, M.J.; Montalvo, J.D.; Ivy, J.; Cherepy, N.J.; Seeley, Z.; Payne,
S.A. Transparent polycrystalline Gd2Hf2O7 ceramics. J. Am. Ceram. Soc. 2018, 101, 3797–3807.
293. Zhao, W.; Anghel, S.; Mancini, C.; Amans, D.; Boulon, G.; Epicier, T.; Shi, Y.; Feng, X.Q.; Pan, Y.B.; Chani, V.; et al. Ce3+ dopant
segregation in Y3Al5O12 optical ceramics. Opt. Mater. 2011, 33, 684–687.
294. Palmero, P.; Bonelli, B.; Fantozzi, G.; Spina, G.; Bonnefont, G.; Montanaro, L.; Chevalier, J. Surface and mechanical properties
of transparent polycrystalline YAG fabricated by SPS. Mater. Res. Bull. 2013, 48, 2589–2597.
295. Lyberis, A.; Patriarche, G.; Gredin, P.; Vivien, D.; Mortier, M. Origin of light scattering in ytterbium doped calcium fluoride
transparent ceramic for high power lasers. J. Eur. Ceram. Soc. 2011, 31, 1619–1630.
296. Klimke, J.; Trunec, M.; Krell, A. Transparent tetragonal yttria-stabilized zirconia ceramics: Influence of scattering caused by
birefringence. J. Am. Ceram. Soc. 2011, 94, 1850–1858.
297. Zhang, Y. Making yttria-stabilized tetragonal zirconia translucent. Dent. Mater. 2014, 30, 1195–1203.
298. Trunec, M.; Maca, K.; Chmelik, R. Polycrystalline alumina ceramics doped with nanoparticles for increased transparency. J. Eur.
Ceram. Soc. 2015, 35, 1001–1009.
299. Merac, M.R.; Reimanis, I.E.; Smith, C.; Kleebe, H.-J.; Müller, M.M. Effect of Impurities and LiF Additive in Hot-Pressed
Transparent Magnesium Aluminate Spinel. Int. J. Appl. Ceram. Technol. 2013, 10, E33–E48.