# Investigating the orientation dependence on functional properties in Bi0.5Na0.5TiO3-BaTiO3 films

Ilham Hamdi Alaoui, Abdelilah Lahmar, Oumayna Mlida, Freddy Ponchel, Antonio da Costa, Françoise Le Marrec, Jamal Belhadi, Alexandre Ysebaert, Anna Cantaluppi, Marie-Hélène Chambrier, et al.

► To cite this version:

Ilham Hamdi Alaoui, Abdelilah Lahmar, Oumayna Mlida, Freddy Ponchel, Antonio da Costa, et al. Investigating the orientation dependence on functional properties in Bi0.5Na0.5TiO3-BaTiO3 films. Materials Research Bulletin, 2025, 189, pp.113475. 10.1016/j.materresbull.2025.113475. hal-05029568

HAL Id: hal-05029568

https://u-picardie.hal.science/hal-05029568v1

Submitted on 13 May 2025

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Research Papers

# Investigating the orientation dependence on functional properties in Bi<sub>0.5</sub>Na<sub>0.5</sub>TiO<sub>3</sub>-BaTiO<sub>3</sub> films

<figure><img src="image_2.png" /></figure>

Ilham Hamdi Alaoui<sup>a</sup>, Abdelilah Lahmar<sup>a</sup>, Oumayma Mlida<sup>b</sup>, Freddy Ponchel<sup>c</sup>, Antonio Da Costa<sup>b</sup>, Françoise Le Marrec<sup>a</sup>, Jamal Belhadi<sup>a</sup>, Alexandre Ysebaert<sup>a</sup>, Anna Cantaluppi<sup>a, ID</sup>, Marie-Hélène Chambrier<sup>b</sup>, Rachel Desfeux<sup>b, ID</sup>, Denis Remiens<sup>c</sup>, Anthony Ferri<sup>b</sup>, Nathalie Lemée<sup>a, *, ID</sup>

ARTICLE INFO

Keywords:

Epitaxy

Sodium bismuth titanate

Relaxor

Lead free material

Piezoelectricity

Energy storage

ABSTRACT

The morphotropic phase boundary (MPB) composition in lead free (1-x) Bi<sub>0.5</sub>Na<sub>0.5</sub>TiO<sub>3</sub> - x BaTiO<sub>3</sub> (BNTBT) solid solution has attracted extensive research due to its significant potential for piezoelectric and high-power energy storage applications. Here, epitaxial (001) and (111) BNTBT films with composition around the MPB are investigated. A complex domain pattern is evidenced for both film orientation, due to the coexistence of a weak polar phase and a strong polar ferroelectric phase. An electric field induced phase switching is shown in both (001) and (111) oriented film, as well as a weakening of the polar state in the (111) BNTBT film. The enhanced ergodic relaxor state in the (111) BNTBT film gives rise to a reduced piezoelectric response and improved energy storage performances. The epitaxial symmetry engineering is shown to provide a complementary approach to the composition strategy to improve the functional properties in BNTBT films.

# 1. Introduction

Perovskite ferroelectric materials have diverse functionalities and are widely used in various kinds of modern electronic devices, such as sensors and actuators for instance, but have also great significance for high-power dielectric energy storage. Lead-based ferroelectric materials display high performance but the toxicity of lead will restrict their future use in devices. Environmental concerns and RoHS directives have given rise to extensive research in the past two decades into lead-free ferroelectric materials, among them the lead free relaxor (Bi,Na)TiO<sub>3</sub> [1].

The lead-free relaxor (1-x) Bi<sub>0.5</sub>Na<sub>0.5</sub>TiO<sub>3</sub> - x BaTiO<sub>3</sub> (BNTBT) solid solution has been thoroughly investigated because of its significant potential as piezoelectric material [2] and also since it is a competitive ferroelectric material for energy storage [3,4]. This system exhibits a morphotropic phase boundary (MPB) around x = 0.06, which separates ferroelectric tetragonal phase (P4mm) from a ferroelectric rhombohedral phase (R3c) associated to a large mechanical response [5]. As established in Pb(Zr<sub>x</sub>Ti<sub>1-x</sub>)O<sub>3</sub> [6], the instability of the polarization state

at the MPB is known to facilitate the rotation polarization and domain wall motion, both of which induce an enhanced piezoelectric response. However the MPB of BNTBT presents distinct specificities due in particular to its relaxor ferroelectric behavior in relation to structural disorder. In the vicinity of the MPB composition, there is the coexistence of the R3c ferroelectric phase with the P4bm phase which is characterized as a weakly polar ferrielectric phase resulting from an uncompensated antiferroelectric order [7]. This P4bm phase acts as a depolarization field and restrain the domain boundary, giving rise to a complex polar nanodomain structure whose morphology is related to the R3c/P4bm ratio. The coexisting R3c/P4bm symmetries induce a disorder which is going to influence the structural as well as the functional properties. For instance, due to these specificities, multiple polar topologies were achieved in BNT based relaxors [8]. The relaxor characteristics of BNTBT make this system attractive for energy storage performances due to the slim P-E loops and corresponding high maximum polarization and low remnant polarization.

There are still some drawbacks with BNTBT. Regarding the energy

<footnote><sup>a</sup> Laboratoire de Physique de la Matière Condensée, Université de Picardie Jules Verne, Amiens UR 2081, France</footnote>

<footnote><sup>b</sup> Univ. Artois, CNRS, Centrale Lille, Univ. Lille, UMR 8181, Unité de Catalyse et Chimie du Solide (UCCS), F-62300 Lens, France</footnote>

<footnote><sup>c</sup> Institut d'Electronique, de Microélectronique et de Nanotechnologie - Site de Valenciennes, UMR 8520 CNRS, Université Polytechnique Hauts-de-France (UPHF), Valenciennes, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: nathalie.lemee@u-picardie.fr (N. Lemée).</footnote>
storage performances, the coercive field and the remnant polarization are too high. Concerning the piezoelectric properties, the d33 is relatively low and the piezoelectric response drops when crossing a relatively low depolarization temperature. In order to improve the functional properties, different substitutions were investigated giving rise to a complex composition design [3,9].

Most of studies are reported on ceramics, single crystals and polycrystalline films. The epitaxial growth of thin film strongly influences the domain structure and domain wall motion. Strain engineering was widely used to tune the functional performances of the ferroelectric perovskite materials [10]. This approach was used to shift the ferroelectric phase transition at higher temperature under compressive strain for instance. Ferroelectric domain engineering is essential to enhance the functional properties, and can be achieved in epitaxial thin films by tuning the strain induced by the epitaxial growth but also through epitaxial symmetry engineering [11,12]. It was thus reported in PbTiO3 [11] and BaTiO3 [12] thin films that the threefold rotational symmetry induced by a (111) orientation of a SrTiO3 substrate led to structure instabilities and complex domain structures. The strain effect was studied in epitaxial BNT thin films [13–17] and BNTBT superlattices [18]. The tensile strain was reported to contribute to an improvement of the dielectric properties. The investigation of the orientation dependence of functional properties in epitaxial BNT based thin film is scarce. Macroscopic and nanoscale electrical properties of (100) and (110) epitaxial BNT based film were studied [13–16]. A strong orientation dependence of the piezoelectric properties in thin films was reported [15,16]. A giant energy storage effect was also evidenced in (100) and (111) (Bi<sub>0.5</sub>Na<sub>0.5</sub>)<sub>0.9118</sub>La<sub>0.02</sub>Ba<sub>0.0582</sub>(Ti<sub>0.97</sub>Zr<sub>0.03</sub>)O<sub>3</sub> epitaxial thin film [17].

In this report, we investigate (001) and (111) BNTBT films epitaxially grown on LaAlO3 (001) or (111) substrates. Our objective is to compare the functional properties, piezoelectric performances as well as ability for energy storage, for these two orientations, with a composition around the morphotropic phase boundary in the BNTBT system. The strain state is determined by X-ray diffraction (XRD), the temperature dependence of dielectric and ferroelectric properties is studied, and the results are interpreted in relation with the characteristics of the domain structure evidenced by piezoresponse force microscopy (PFM). We show here that the (111) epitaxial orientation strengthens the ergodic relaxor state in BNTBT film, giving rise to a reduced piezoelectric response and improved energy storage performances.

## 2. Experimental details

Films were grown by pulsed laser deposition from a 1 mol.% Manganese-enriched (0.94 (Bi<sub>0.55</sub>Na<sub>0.55</sub>) TiO<sub>3</sub>–0.06 BaTiO<sub>3</sub> (BNTBT) ceramic. The excess in bismuth and sodium aims to compensate the losses during the deposition process. Doping with manganese should enable to reduce leakage current [19]. They were deposited on (001) and (111) LaAlO<sub>3</sub> (LAO) single crystal substrates, buffered with a conductive layer of La<sub>0.5</sub>Sr<sub>0.5</sub>CoO<sub>3</sub> (LSCO), applying the same deposition conditions. A KrF excimer laser (248 nm, pulse duration = 20 ns) with a laser energy density of 2 J.cm<sup>-2</sup> and a laser frequency of 2 Hz was used. The LSCO buffer layer as well as the BNTBT film were deposited at 740 °C in 0.3 mbar of O<sub>2</sub>.

In order to investigate room temperature structural properties, ω – 2θ and reciprocal space map measurements were performed using a four-circle high resolution D8 Discover Bruker diffractometer. A double reflection channel-cut Ge monochromator was used to select the Cu Ka<sub>1</sub> radiation (1.5406 Å).

An FEI Quanta 200 FEG Environmental scanning electron microscope (SEM) was used to analyze the microstructure of the films. The surface morphology of the BNTBT films was investigated by Atomic Force microscopy (AFM) in PinPoint mode using a NX10 microscope (Park Systems) operating at room temperature and under ambient conditions. Local ferroelectric and piezoelectric properties were studied

by using the Piezoresponse Force Microscopy (PFM) technique available on the microscope. The ferroelectric domains pattern was revealed by PFM imaging. Both out-of-plane (OP) and in-plane (IP) PFM images were simultaneously recorded. Additionally, both local switching properties and local electromechanical activities of the films were measured by recording phase and amplitude piezoelectric hysteresis loops using the spectroscopic tool of the PFM. For such measurements, a driving voltage of 2 V was applied to the local probe while the PFM loops were recorded in the remnant mode to promote electromechanical activity against electrostatic interaction. To enhance the detected electromechanical signal, PFM analyses were carried out at the contact resonance frequency of the cantilever. The nanoscale electric performances of the samples were given by both measuring the coercive voltage $V_c$ on the PFM loops and by evaluating the piezoelectric activity from the displacement measured at saturation on the amplitude PFM butterfly-shape loops [20]. For these experiments, Pt/Ir-coated conductive tips (nanosensors PPP-EFM probes with a nominal ~ 2.8 N.m⁻¹ spring constants) were used.

The dielectric measurements were carried out between 100 Hz and 1 MHz using a Solartron Impedance analyzer SI-12,060 with a probing AC electric field amplitude of 100 mV. The polarization-electric field (P-E) hysteresis loops were measured at 20 kHz using a TF Analyzer 3000, aixACCT system. All electrical measurements were performed in metal–dielectric–metal geometry using sputtered platinum circular top electrodes with 250 µm diameters deposited through a shadow mask.

The piezoelectric d33eff coefficient was measured using the experimental setup shown in Fig. 1. The experimental procedure is based on the converse piezoelectric effect. Using a HP 4192 impedancemeter, a superposition of bias-voltage (VDC) and sinusoidal voltage (VAC) is applied between the top and bottom electrodes of the sample. Therefore, a strain is induced into the film under test along the out-of-plane direction. The sample surface displacement is recorded by a single-beam Laser head interferometer (Polytec OFV-512). The OFV-5000rometer controller (coupled with the VD-08 digital velocity decoder) measures a voltage (VΔI in Fig. 1) proportional to the phase shift between the incident laser beam and the reflected laser beam by one of the top electrodes. This setup provides a maximum velocity resolution of 5 nm. s⁻¹ for a maximum linearity error of 0.1 % within a frequency range of 0–25 kHz. Finally, using the Polytech E-220 data acquisition units, an FFT analysis is performed on all the measured voltages (VAC, VDC and VΔI) and the resulting data are transmitted to a PC for processing. A software specially designed to control the system, collect and synchro-nize the data to calculate the piezoelectric coefficient. Then, for a particular bias-voltage VDC the piezoelectric coefficient d33eff is deduced from VΔI and VAC FFT data. This procedure can be repeated for several selected values of bias-voltages to obtain the butterfly-like hysteresis loop commonly found in piezoelectric materials. In this work, an AC voltage VAC of fixed amplitude (1 V) and frequency (10 kHz) is applied and the VDC bias-voltage loop range from –8 V to 8 V corresponding to a static electric field of approximately ± 250 kV/cm.

## 3. Results and discussion

The SEM top views of the (001) and (111) BNTBT samples are shown in Fig. 2a and b. The morphology is irregular for both samples, and is characteristic of an island growth [21]. The size of the islands is obviously reduced in the (111) BNTBT. The thicknesses of the conductive LSCO layer, of the BNTBT films, are estimated to 85 nm and 200 nm respectively, from SEM investigations carried out on the cross section of the film (Fig. 2c). These thickness values are in agreement with the deposition rate about 0.03 nm / laser shot, determined by X-ray reflectivity measurements on thinner films.

The X-ray diffraction (XRD) diagrams for both film (Fig. 3a and b) show a preferential out-of-plane orientation. No peak splitting is
<figure><img src="image_3.png" /><figcaption>Fig. 1. Experimental setup for $d_{33}$ piezoelectric coefficient measurement.</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>Fig. 2. SEM images showing the microstructures of the (001) oriented BNTBT film grown on the (001) LAO (a) and on the (111) oriented LAO (b) substrate. A cross section of the (001) BNTBT sample on LSCO on (001) LAO is presented (c).</figcaption></figure>

observed therefore the average structure of the films is determined here as pseudocubic. The coexistence of different phases is not excluded since no observable lattice distortion has already been reported in unpoled BNTBT ceramics [22]. A (001) or (111) preferential out of plane orientation induced by the single crystal substrate is clearly evidenced for the BNTBT film and the LSCO conductive layer. Reciprocal space maps (RSMs) were measured around the (103) reflection of LAO for the film grown on (001) LAO (Fig. 3c) and around the (132) reflection of LAO for the sample deposited on the (111) LAO substrate (Fig. 3d). Concerning the LSCO layer, it is partially relaxed whatever the substrate, however it can be seen that when grown on (001) LAO, a part of the layer is coherently strained (Fig. 3c). At room temperature, LSCO is cubic (Fm-3m space group, $a = 3.805 Å$) and LAO is cubic (Pm3m, $a = 3.787 Å$). It has to be pointed out that unlike the (001) LAO, the (111) LAO substrates are characterized by a twinned structure, which gives rise to two reflections along the [001] direction. This difference in the twinning characteristics, can certainly contribute to explain the lack of coherent growth of LSCO layer in the case of (111) LAO substrate. In the coherently strained LSCO layer grown on (001) LAO substrate, the tetragonality c/a is 1.02. In the partially relaxed LSCO layer, this tetragonality decreases to 1.00 in relationship with a relaxation of the compressive strain. The in-plane lattice constants in both part of the LSCO layer are higher than in the bulk (3.812 Å in the strained layer, 3.849 Å in the relaxed one) that can be probably explained by the presence of vacancies related to oxygen for example.

The (001) BNTBT film is epitaxially grown but does not share the same in-plane lattice constant with the LAO substrate. At room tem-

perature, bulk BNTBT around $x \sim 0.06$ is pseudo-cubic (space group R3c, $a = 3.8991 Å$, $\alpha = 89.7^\circ$). The BNTBT film is partially relaxed and the out-of-plane and in-plane lattice constants are $d_{003} = 1.3073 Å$ and $d_{100} = 3.9210 Å$, respectively. The epitaxial strain along the [hkl] direction is defined as $\epsilon_{hkl} = \frac{d_{hkl, film} - d_{hkl, bulk}}{d_{hkl, bulk}}$ with $d_{hkl, film}$ and $d_{hkl, bulk}$ being the lattice constant of the film and bulk BNTBT, respectively. The out-of-plane and in-plane distortion can be calculated as $\epsilon_{003, out} = 0.59 \%$ and $\epsilon_{100, in} = 0.56 \%$, respectively. The values reported here for the (001) BNTBT film are consistent with an average pseudo-cubic symmetry.

Concerning the film grown on (111) LAO, the lattice constants are calculated using the (132) RSM (Fig. 3d). In the BNTBT film, the out-of-plane and in-plane lattice constants are $d_{222} = 1.1233 Å$ and $d_{110} = 2.7550 Å$, respectively. In the BNTBT film, the in-plane and out of plane epitaxial strain can be estimated as $\epsilon_{110, in} = 0.18 \%$ and $\epsilon_{222, out} = -0.72 \%$, respectively. The (111) BNTBT film is thus under in-plane tensile strain and out-of-plane compressive strain. Hence, the lattice distortions are different in the (001) and (111) BNTBT films.

The characteristics of ferroelectric domains patterns in the as-grown samples were investigated using OP-PFM and IP-PFM imaging (Fig. 4). The topographic images (Fig. 4a, b) show irregular grain shapes for both orientations, but donut-shaped aggregates are observed for the (111)-oriented BNTBT film. The average grain size is larger for the (001) -oriented BNTBT film compared to the (111) -oriented one. Both samples display mosaic-like ferroelectric domain structures as reported in BNTBT crystals [23]. The amplitude OP-PFM and IP-PFM images reveal the existence of domain patterns with a strong amplitude response
<figure><img src="image_5.png" /><figcaption>**Fig. 3.** ω-20 scans measured on the BNTBT film grown on (001) LAO (a) and on (111) (b). Reciprocal space maps (RSM) around the (103) peak of the (001) LAO for the (001) BNTBT film (c) and around the (132) peak of LAO for the (111) BNTBT film (d). All the Miller indices used are based on the pseudo-cubic lattice.</figcaption></figure>

(highlighted by yellow circles in Fig. 4c–f) surrounded by regions with a very weak signal (highlighted by pink circles in Fig. 4c–f). Within these areas, the phase signal is complex with various contrasts. The result is a pattern of irregular domains, hundreds of nanometers in size. It clearly emphasizes the presence of two phases in the samples. As previously explained, several phases coexist in the vicinity of the MPB composition. The phase diagrams of BNTBT reported in the literature [7,24] evidence the coexistence of the P4bm and R3c phases at room temperature for x ~ 0.06. The P4bm symmetry exhibits ferrielectric features with a weakly polar response and has a weak spontaneous polarization, while R3c symmetry shows ferroelectric features with a strong polarization [8]. The regions with low deformation could be attributed to this P4bm phase, which has already been evidenced in the films [25–27]. The domains with a strong out-of-plane amplitude signal, could be characteristic of a polar rhombohedral R3c phase since an in-plane amplitude signal is also detected for these domains. The domain structure in rhombohedral ferroelectrics can be rather complicated due to the various possibilities of equivalent polarization directions [13,14, 28–30]. In BNT based materials, the polar domain structure is rather complex, labyrinthine, stripe, bubble and polar-skyrmion domains were evidenced [31]. This chemically driven domain evolution is related to the nonergodic relaxor (NR) state to the ergodic relaxor (ER) state, and the P4bm symmetry is reported to limit the size of the R3c domains. Here the domain patterns present some similarities with those reported

for BNTBT piezoceramic or thin film [25,27]. It is noticed that the size of the domains is smaller in the (111) oriented BNTBT film. Even if the P4bm and R3c phases are present in both samples, the crystallographic orientation of the film has an obvious influence on the domain pattern.

The polarization and corresponding switching current-electric field hysteresis loops recorded at different temperatures are shown in Fig. 5 and in Fig. 6. At room temperature, the (001) BNTBT exhibits typical ferroelectrical loop with remnant polarization (2 P<sub>r</sub>) of 17.5 µC.cm<sup>-2</sup> and a coercive field (2 E<sub>c</sub>) of 276 kV.cm<sup>-1</sup>. Two current switching peaks produced by ferroelectric switching of the polar domains occurs at the coercive field (E<sub>c</sub>) and are shifted along the positive electric field axis, as expected when there is an imprint field E<sub>imp</sub> (12 kV/cm) induced by an asymmetric metal-dielectric interface [32]. In BNT based materials, a nonergodic relaxor state consisting of short-range ordered polar nano regions (PNRs) is reported [33]. A ferroelectric state is induced from the nonergodic relaxor state under an applied electric field and the long range order can be induced and stable upon electric field removal. This is what is observed at room temperature for the (001) sample. On heating, the PE loops become more slanted and more pinched. The emergence of additional current peaks in I-E loops is progressively evidenced. From 140 °C, 4 current peaks labelled I<sub>1</sub>(I'<sub>1</sub>) and I<sub>2</sub>(I'<sub>2</sub>) are clearly observed. The occurrence of four peaks was already reported in BNT based ceramics, and TEM investigations as well as synchrotron studies have shown that it was the signature of a combination of domain
<figure><img src="image_6.png" /><figcaption>Fig. 4. AFM ([a, b] Topography) and PFM ([c, d] OP amplitude, [e, f] IP amplitude, [g, h] OP phase, [i, j] IP phase) images recorded on the (001) BNTBT (left side) and the (111) BNTBT (right side). The pink (yellow) circles point out the absence (presence) of PFM amplitude for some ferroelectric domains. In the images the scale bar is 100 nm.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>Fig. 5. Room temperature PE loops (a) and current electric field (b) measured for the (001) BNTBT and the (111) BNTBT. The data were recorded during the second full cycle of electric field application.</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>Fig. 6. PE loops (a) and current electric field (b) measured at several temperatures for the (001) BNTBT (a, b) and the (111) BNTBT (c, d).</figcaption></figure>

switching and electric field induced phase switching [34–38]. Consid-
ering the data reported in the litterature, the peaks I₁ and I'₁ can be
attributed to the ferroelectric domain switching and to the electric field
induced transition from weakly polar to strongly polar ferroelectric
phase symmetry. Regarding the peaks I₂ and I'₂, they are the signature of
a ferroelectric to ergodic relaxor phase transition, with the decomposi-
tion of ferroelectric domains into polar nanodomains. During the
loading of a positive field, above the I₁ peak, a strong polar structure is

established and is stable upon unloading. When a small negative field is
applied, the strong polar state turns into a weak polar phase (I₂ peak).
With the maximum negative applied field, the weakly polar phase
changes again to a polar structure (I'₁ peak). On reversal of the applied
electric field, the phase switching from a polar structure to the non polar
phase occurs (I'₂ peak). As evidenced by the PFM investigations, both
samples exhibit weakly polar and strongly polar ferroelectric phases, so
phase switching is coherent with the characteristics of the sample.
The shapes of the PE and IE loops are very different for the (111) BNTBT (Fig. 5). A slimmer PE loop is evidenced, characterized by a significantly lower remnant polarization (2 *P*<sub>r</sub> ~ 9.7 µC.cm<sup>-2</sup>) than in the (001) BNTBT, and a slightly lower coercive field (2 *E*<sub>c</sub> ~ 226 kV.cm<sup>-1</sup>) is determined. The pinched PE loop and the smaller remnant polarization in the (111) BNTBT indicates a strong weakening of the polar state. As the temperature increases, the four broad peaks in I-E loops are still observed with a fairly low intensity contrast. The polar nanoregions turn into the long range order (I<sub>1</sub> peak), but the transformation back to polar nanoregions (I<sub>2</sub> peak) occurs before the positive field returns to zero and the current value remains almost constant, meaning that the field induced transitions are reversible. This behavior is a signature of an enhanced ergodic relaxor state in the (111) BNTBT.

The dependences of the dielectric permittivity (ε') and the dissipa-
tion factor (tan δ) from room temperature to 300 °C at several fre-
quencies are shown in Fig. 7 for both samples. The temperature of T<sub>m</sub>
corresponding to the dielectric maximum is not observed in the inves-
tigated temperature range. The dielectric permittivity spectra result
from the coexisting weakly polar and strongly polar ferroelectric PNRS.
The dielectric constants of the materials decreased with increasing fre-
quency, which can be attributed to the fact that at higher frequencies
fewer dipoles can follow the applied alternating electric field.

A weaker temperature dependence of the dielectric constant is
observed in the (111) BNTBT sample. Compared with the (001) BNTBT
sample, a reduction of the permittivity and of the loss tangent is evi-
denced in the (111) BNTBT film. Due to the enhancement of the relaxor
feature in the (111) oriented BNTBT sample, the PNRs are more easily
switchable resulting in reduced dielectric losses [39,40]. This decrease
in permittivity when comparing the dielectric properties of the (001)
and (111) BNTBT films is in agreement with the results reported for
BNLBTZ films [17]. It could be related to a higher ratio of the weakly
polar phase in the (111) BNTBT film, which remains to be investigated.
The evolution of the domain pattern as a function of film orientation,
can also influence the mobility of the domain walls and consequently the
dielectric permittivity. The temperature dependent dielectric permit-
tivity accounts from the R3c to P4bm transition and the thermal evo-
lution of P4bm PNRs. The R3c PNRs are expected to transform into
P4bm PNRs with increased temperatures [41]. An anomaly is detected
in dielectric constant as well as in dielectric losses for both samples. It
occurs at around 170 °C for the (111) BNTBT and is shifted around 190
°C for (001) BNTBT. This anomaly is not characterized by a frequency
dispersion and we surmise that it can be attributed to the strongly polar
ferroelectric to the weakly polar phase transition [41]. The deepening of
the ergodic degree in the (111) BNTBT film could explain the decrease in
the temperature transition.

For BNT-based thin films, typical d<sub>33</sub> values reported in the literature range from tens to hundreds of pm.V<sup>-1</sup> [16,19,30,42–45]. The d<sub>33</sub> coefficient is strongly dependent on the composition [43], deposition conditions [30], thickness [45] and structural characteristics

(polycrystalline, textured or epitaxial growth) of the film [16]. The
values can also vary depending on whether local or macroscopic
piezoelectric properties are probed using techniques, such as piezores-
ponse force microscopy (PFM) or laser interferometry methods,
respectively. Moreover, it is worth noting that values obtained using
PFM, a complex technique, should be interpreted with caution. The
macroscopic effective electromechanical response (d<sub>33eff</sub>) was measured
at room temperature for both samples (Fig. 8). The (001) BNTBT film
shows a higher d<sub>33eff</sub> (around 40 pm/V) than the (111) BNTBT sample
(around 24 pm/V), whose d<sub>33eff</sub> increases more slowly with voltage. It
could be explained by the strengthened ergodicity in the (111) BNTBT
sample, which imposes a higher voltage to induce the long-range
ferroelectric order [46]. The (111) orientation of the BNTBT film re-
sults in a significant decrease in the electromechanical response. Further
investigations, including a transmission electron microscopy, are
needed to determine if a change in the ratio between the weakly polar
and ferroelectric phases could contribute to this result.

The typical remnant phase and amplitude PFM loops recorded over
the polar domains observed at the free surface are displayed in Fig. 9.
The characteristic square-shaped in-phase PFM hysteresis loops, with a
phase difference of nearly 180° between the two opposite polarization
states, along with well-saturated amplitude PFM hysteresis loops
confirm the distinct local switching properties of the ferroelectric do-
mains for both samples. The average coercive voltage is slightly higher
for the (001) oriented BNTBT film but remains very close to the value
measured for the (111) oriented film, i.e. 3.9 V and 3.5 V, respectively.
Regarding the amplitude PFM butterfly loops, a slightly higher ampli-
tude of about 30 % is observed for the (001) oriented BNTBT film
compared to the (111) oriented one. At the local scale, the largest local

<figure><img src="image_9.png" /><figcaption>Fig. 8. Comparison of d<sub>33eff</sub> piezoelectric coefficient for (001) and (111) BNTBT films.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>Fig. 7. Temperature dependence of the real part of the dielectric constant ε' and of the dissipation factor tan δ for the (001) (a) and (111) (b) BNTBT films.</figcaption></figure>

<figure><img src="image_11.png" /></figure>
<figure><img src="image_12.png" /><figcaption>Fig. 9. Amplitude (a, b) and remnant phase (c, d) PFM loops recorded on the (001) BNTBT (a, c) and (111) BNTBT (b, d).</figcaption></figure>

piezoresponse is confirmed in the (001) BNTBT film.

It was recently reported in the literature that in the MPB composition of the BNTBT system, the maximum d33 is achieved in the mixed, non ergodic and ergodic relaxor state [47]. The reduction in the value of the d33eff in the (111) BNTBT sample could be explained by the enhanced ergodic state evidenced in this film. This value is also controlled by the crystallographic orientation which is in relation with the engineered domain configuration. In the case of rhombohedral 0.95 BNT – 0.05 BT single crystals, a lower d33 value was determined for the (111) oriented single crystal than for the (001) oriented single crystal [48]. It was also reported in BNTBT-ST thin film [16]. Under the applied electric field along the [001] direction, four possible polar directions ([111], [111], [111], [111]) are expected to align along the electric field direction, giving rise to a high rhombohedral lattice distortion. Along the [111] direction, the single domain state is expected to be less stable and associated to a lower piezoelectric response. In this way, the crystallographic orientation can also contribute to the lower d33eff in the (111) BNTBT film. However it cannot be considered as the dominant contribution since it should lead to a remnant polarization in the (001) BNTBT sample that should be 1/√3 of that in the (111) BNTBT film. Here the ferroelectric investigations have shown a larger remnant polarization value for the (001) BNTBT sample.

The enhanced ergodic relaxor state evidenced in the (111) BNTBT film by ferroelectric measurements, may contribute to improved energy storage properties, which are now determined from the PE hysteresis loops presented in Fig. 6. The interesting parameters were calculated using the following equations:

$$W_{\text{rec}} = \int_{\text{Pr}}^{P_{\text{max}}} \text{EdP} \qquad (1)$$

$$W_{\text{total}} = \int_{0}^{P_{\text{max}}} \text{EdP} \qquad (2)$$

$$\eta = \frac{W_{\text{rec}}}{W_{\text{total}}} \times 100 \qquad (3)$$

where, Wrec represents the density of the recoverable energy, Wtotal is the total energy and η is the associated efficiency.

The evolutions of the recoverable energy density (Wrec) and the related efficiency (η) as a function of temperature from 25 °C to 300 °C, are depicted in Fig. 10. A weak evolution of the energy storage density as a function of temperature was observed for both films. The Wrec was determined to be between 1.3 and 1.7 J.cm⁻³ for (001) BNTBT and between 1.4 and 1.9 J.cm⁻³ for (111) BNTBT. However, even though Wrec is comparable in both films, a significant difference is detected in terms of efficiency, which is <50 % for the (001) oriented film but reaches 80 % at high temperature for the film oriented along the [111] direction. The lower energy losses in the (111) BNBT film explain this result.

In order to benchmark the ability of the investigated films for energy storage application and give an objective comparison with some promising equivalent materials, we used the performance coefficient defined by Hao Pan et al. [49] as:

$$F = \frac{W_{\text{rec}}}{1 - \eta} \qquad (4)$$

In Fig. 11, the temperature dependence of F is presented for the samples studied in this work and is compared with those determined under a similar applied electric field for BNTBT polycrystalline films, grown either by PLD or by chemical solution (CSD) on platinum-buffered silicon substrate [26]. From room temperature to 150 °C, the BNTBT processed by sol gel on Pt-substrate exhibits better performance comparing to the films grown by PLD. However, the (001) and (111) BNTBT films studied in this work show a better thermal stability of the performance coefficient. Above 150 °C, the performance coefficient of the (111) BNTBT film increases significantly due to an improved related efficiency in this temperature range. This result makes the (111)
<figure><img src="image_13.png" /><figcaption>**Fig. 10.** Temperature dependence of W<sub>rec</sub> and η for: (a) (001) BNTBT film; (b) (111) BNTBT film.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>**Fig. 11.** Comparison of the temperature dependence of the performance coefficient F of BNTBT films studied in this work with those reported in [26].</figcaption></figure>

orientation in BNTBT based thin films very interesting for energy storage capacitors operating above 150 °C.

We note here that only a low electric field of 600 kV.cm⁻¹ was used to calculate the energy storage parameters, which are therefore not comparable with data reported in the literature measured at very high fields. The energy storage performances can be further enhanced by increasing the applied electric field but also by adjusting the composition in (111) oriented BNT-based thin films through doping or cation substitution promoting the ergodic state.

### 3. Conclusion

In summary, the orientation dependence on functional properties was investigated in epitaxial (001) and (111) BNTBT films grown on LAO substrates. The film orientation has a strong influence on the domain pattern. The morphology of the domain in the vicinity of the MPB composition is tuned by the coexistence between a weak polar ferrielectric phase and a strong polar ferroelectric phase. The [111] orientation of the film is shown here to promote smaller domain size for both phases and leads to an enhanced ergodicity of the relaxor. It gives rise to a slimmer PE hysteresis loop with a significant reduction of the remnant polarization, which contributes to an improvement of the energy storage efficiency. The [001] orientation is more effective for enhancing the piezoelectric response.

In BNT based system, the disruption of the long-range ferroelectric order is usually achieved by adjusting the composition through an appropriate doping leading to a more and more complex composition design. In films, our work evidences that the epitaxial symmetry engineering provides a complementary approach to engineer complex polar nanodomain structure in BNT based system. Depending on the intended applications, in the field of piezoelectricity or energy storage, the orientation of the film in addition to its composition will have to be appropriately selected.

### CRediT authorship contribution statement

**Ilham Hamdi Alaoui:** Writing – original draft, Visualization, Investigation, Data curation.
**Abdelilah Lahmar:** Writing – review & editing, Visualization, Validation, Formal analysis, Conceptualization.

**Oumayma Mlida:** Visualization, Investigation. **Freddy Ponchel:** Writing – review & editing, Visualization, Investigation, Data curation.

**Antonio Da Costa:** Visualization, Investigation, Data curation.

**Françoise Le Marrec:** Visualization, Investigation, Formal analysis.

**Jamal Belhadi:** Writing – review & editing, Validation, Methodology.

**Alexandre Ysebaert:** Visualization, Investigation. **Anna Cantaluppi:** Visualization, Resources.
**Marie-Hélène Chambrier:** Validation.

**Rachel Desfeux:** Writing – review & editing, Validation.
**Denis Remiens:** Validation.
**Anthony Ferri:** Writing – review & editing,

Validation.
**Nathalie Lemée:** Writing – review & editing, Writing – original draft, Supervision, Methodology, Funding acquisition,

Conceptualization.

### Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

### Acknowledgments

This work was supported by the Region of Hauts de France (OPPEN project, Grant No 20002828).

I. H. A. gratefully acknowledges financial support from the Region of Hauts de France and Amiens Metropole, and the platform of electron microscopy at university of Picardie Jules Verne for their support and assistance of Dr Arash Jamali in this work. A. F., A. D. C., O. M, M. H. C. and R. D. thank Chevreul Institute (FR 2638), Ministère de l'Enseignement Supérieur, de la Recherche et de l'Innovation, Hauts-de-France Region, Fonds Européen de Développement Régional (FEDER) and Major Domain of Interest (DIM) "Eco-Energy Efficiency" of Artois University, for the financial support of this work, particularly for funding the NX10 microscope.

### Data availability

Data will be made available on request.
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