Ultrahigh energy-storage in lead-free ceramic capacitors via local structure design

Domain: biology
Published: 2026-03-31
Editorial outcome: major_revision

Abstract
Dielectric capacitors are crucial energy-storage modules in pulsed- and high-power devices.
However, the simultaneous enhancement in recoverable energy density (W ) and efficiency (η)
rec
still remains challenge owing to the restrictive relationship between the maximum polarization
(P ), remanent polarization (P) and electric breakdown strength (E ). To address this, we
max r b
propose a strategy of local polar structure design in BiFeO -based ceramics. By incorporating
3
NaNbO to create embedded, persistent polar nanoregions within a weakly polar matrix, we
3
achieve a giant polarization difference ΔP (56.4 μC cm-2), a low P (3.6 μC cm-2) and a large E
Sr b
(66 kV mm-1), endowing an ultrahigh W
rec
of 14.5 J cm-3 and aS high η of 88%. The local structure,
E
directly visualized via the 2D/3D atomic displacemen
R
t mapping, enables a high P
max
under an
applied field and a low P at zero field, a mechanism ePxplicitly validated by phase-field simulation.
r
Meanwhile, the optimized microstructure and eNnhanced insulating property contribute to a giant
I
E
b
. This work provides insights to overcom e the paradox between multiple parameters and paves
E
a feasible route for the cutting-edge enLergy-storage applications.

Introduction
As one of the important functional materials, dielectric ceramics have provoked ever-
increasing interest in capacitive energy-storage in recent years1-3. Unlike chemical batteries and
electrochemical capacitors, the electric energy stored and released by dielectric capacitors is based
on the polarization response to applied electric field (E), allowing for ultrafast charge/discharge
speed, giant power density, high working voltage and excellent reliability. Hence, the dielectric
capacitors hold promising applications in pulsed-power and high-power fields, such as inverters
and converters in hybrid vehicles4-6. Generally, the energy-storage performance (ESP) of dielectric
S
ceramics, characterized by recoverable energy density (WS) and efficiency (η), is closely
rec
E
associated to the polarization electric-field (P-E) curve. Specifically, a high maximum polarization
R
(P ), a small remanent polarization (P) (or a high Ppolarization difference ΔP = P -P) and a
max r max r
large electric breakdown strength (E ) are desirNable for the high W , while the low polarization
b rec
I
hysteresis characterized by a small P
r
is re sponsible for the high η4. Relaxor ferroelectric (RFE)
E
ceramics are considered as potential aLlternative for energy-storage, resulting from their slim P-E
C
loop and unique polar nanoregions (PNRs). However, the prevailing W < 10 J cm-3 poses
rec
I
T
challenges for their usage in advanced electronic devices requiring compact size and light weight4-
R
6.
A
Numerous strategies have been implemented to improve W based on the regulation of
rec
electric field durability and polarization behavior. On the one hand, the E of RFE ceramics not
b
only restricts their operating electric field, but also possesses a positive correlation with W 6,7.
rec
Compositional optimizations are widely employed to reduce grain size, improve bandgap and
insulating properties3,8. Novel processing techniques, including tape-casting and hot-press
sintering, are generally adopted to enhance the density and reduce the porosity, thereby improving
E 9,10. Additionally, microstructural designs like core-shell structure, 0-3 type composite and
b
multilayer structure, help prevent the development of local electric branch, further improving E 11-
b
13. On the other hand, the element selection, polymorphic PNRs, chemical framework and entropy
engineering are the popular and feasible approaches to regulate the polarization behavior1,8,14-18.

For example, Bi-based matrix or end-member, including BiFeO , Bi(Mg Zr )O , etc., is
3 0.5 0.5 3
beneficial to retain a relatively high P owing to the strong hybridization between Bi3+ 6s2 lone
max
pair electrons and O2- 2p orbitals4,8,15. High-entropy assisted Bi Na TiO -based RFE ceramics
0.5 0.5 3
exhibited ultra-small polymorphic domains, which weakened the polarization anisotropy and
reduced the switching energy barriers. This polarization configuration finally resulted in a small
P, delayed polarization saturation while maintaining high P , thereby endowing superior
r max
ESP14,17.

Despite great advancements in dielectric capacitors, the intrinsic constraint between P , P
S max r
and E has not yet been addressed, limiting the further enhanceSment of ESP. Typically, P and P
b max r
E
typically increase or decrease synchronously, creating
R
a challenge for obtaining a low P
r
while
maintaining a high P , which finally yields a trade-Poff between W and η6. Additionally, a high
max rec
P generally indicates that more induced charNges emerge at two surfaces of a dielectric, leading
max
I
to a large electrostatic force and prematur e electro-mechanical breakdown (i.e., a low E
b
)4,19.
E
Therefore, it still remains a great challeLnge to concurrently realize high W
rec
and η through breaking
C
the interdependent limitations by P , P and E .
max r b
I
T
To address this issue, we propose the local structure design to overcome the restraint between
R
different parameters inA (0.6-x)BiFeO
3
-0.4SrTiO
3
-xNaNbO
3
+0.1wt% MnO (denote as BF-ST-xNN-
Mn) ceramics, thereby achieving high W and η simultaneously. 0.6BF-0.4ST ceramic with a
rec
morphotropic phase boundary demonstrates significant potential for energy-storage capacitors,
due to its high intrinsic polarization ~40 μC cm-2 20,21. Yet, the strong ferroelectricity and high
leakage of 0.6BF-0.4ST ceramics enables lossy square P-E loops and inferior ESP (Fig. S1a).
Herein, the perovskite end-member NN was incorporated to disrupt the long-range ferroelectric
orders and promote the formation of PNRs while retaining highly polarized local clusters3,18,
leading to a slim P-E loop with a relatively high P and a significantly reduced P.
max r
Meanwhile,
NN possess a large band gap (E ~ 3.4 eV), which is favorable for improving resistivity and
g
enhance E . Additionally, the synthesis of phase-pure BF-based ceramics is challenging due to
b
volatilization of Bi O and the reduction of Fe3+ (Fe3+ + e- → Fe2+) during sintering, which
2 3
typically leads to high dielectric loss and leakage. In this study, MnO was added as a sintering aid
to mitigate these issues. By helping to maintain the charge balance, MnO suppresses the formation
of oxygen vacancies and secondary phase, thereby reducing leakage current and enhancing E 4,8.
b
Resultantly, the restrictive relationships between P , P and E were significantly mitigated,
max r b
endowing an ultrahigh W of 14.5 J/cm3 and a high η of 88% under 65 kV/mm in 0.46BF-0.4ST-
rec
0.14NN-Mn ceramics.

Results
2.1 Energy-storage properties

Fig. 1 | Energy-storage properties of x = 0.14. E-dependent a unipolar P-E loops and b
corresponding W , η. Comparison in c ΔP and E , d W /E and η between BF-based bulk ceramics.
rec b rec
e W as a function of time under increasing E. f Comparison in W between BF-based bulk ceramics.
d d
A series of BF-ST-xNN-Mn (x = 0-0.16) ceramics were synthesized by solid-state reaction,
and their P-E loops under 10 kV mm-1 were depicted in Figs. S1, S2a. With increasing NN content,
the P-E loops become gradually slimmer, accompanied by decreased P and P from 39.7 and
max r
30.4 μC cm-2 (x = 0) to 8.5 and 0.2 μC cm-2 (x = 0.16), owing to the enhanced relaxor behavior
(Fig. S3). Accordingly, the η progressively increases from 13.2% at x = 0 to 91.3% at x = 0.16 (Fig.
S2b). Fig. S2c exhibits the P-E loops of BF-ST-xNN-Mn ceramics under the respective maximum
applied E. Samples were prepared with a thickness of ~80 μm and an electrode area of ~ 0.79 mm2.
This sample geometry, combined with the high relative permittivity (Fig. S3), suggests a relatively
lower susceptibility to fringing-field effect and parasitic capacitance (Fig. S4)7. A high P
max
exceeding 40 μC cm-2 under applied E is maintained across compositions x = 0.05-0.16, due to the
strong hybridization between Bi-O orbitals8,15. Particularly, the composition x = 0.14 achieves a
giant P of 60 μC cm-2, a low P of 3.6 μC cm-2, and the highest ΔP of 56.4 μC cm-2 under 65
max r
kV mm-1, thereby delivering an optimal ESP with W = 14.5 J cm-3 and η = 88% (Fig. S2d). These
rec
results suggest that x = 0.14 achieves a good balance between P and P, ΔP and E, and W and
maSx r rec
η. It was therefore selected as the representative composition fSor further study.

Fig. 1a displays the E-dependent P-E loops for x = 0.14. The P-E loop always retains a slender
shape across all fields, with both P and P increasiPng steadily as the field rises (Fig. S5). This is
max r
attributed to the E-induced long-rang ferroeNlectric orders, which rapidly return the initial
I
disorganized state after removing E. Ultima tely, the W
rec
increases monotonically from 0.6 J cm-3
E
at 10 kV mm-1 to 14.5 J cm-3 at 65 kLV mm-1, while the η remains largely constant (Fig. 1b). To
C
appreciate the ESP of x = 0.14, the comparison in the key parameters, including ΔP, E , W , and
b rec
I
T
η, between BF-based bulk ceramics is presented in Figs. 1c, 1d, and S6. A high W is dominated
R
by the collective effecAt of high ΔP and large E
b
. However, these two parameters are typically in
direct competition, fundamentally limiting the enhancement of W in dielectric ceramics4,6.
rec
Remarkably, the x = 0.14 composition mitigates this trade-off, achieving a giant ΔP of 56.4 μC
cm-2 simultaneously with a high E , resulting in a record-high W of 14.5 J cm-3 among BF-based
b rec
ceramics (Fig. 1c). Considering the inequality for W obtained under different E, the factor of
rec
merit W /E is adopted as another indicator to further evaluate the ESP. As shown in Fig. 1d, the
rec
simultaneous optimization of a high W /E and a high η also remains a considerable challenge.
rec
For example, BF-BaHf Ti O -Sr(Al Nb )O ceramic exhibits an ultrahigh W of 12.2 J cm-
0.05 0.95 3 0.5 0.5 3 rec
3 and a high η of 90.1% under 83 kV mm-1, but its W /E is only 0.1516. And BF-BaTiO -
rec 3
(Na Sm )NbO ceramic achieves a W of 13.1 J cm-3 under 65 kV mm-1 with a high W /E of
0.7 0.1 3 rec rec
0.20, but the η is only 78.3%22. However, x = 0.14 in this work overcomes the conventional
restriction between P , P and E , thereby exhibiting a high W /E of 0.22 and a high η (88%)
max r b rec
simultaneously, which demonstrates a superior comprehensive ESP among BF-based ceramic
capacitors.

The E-dependent overdamped charge-discharge curves of x = 0.14 are depicted in Fig. S7,
which is a direct route to obtain the discharge energy density (W ). A sharp current peak under
d
each E reveals the fast charge-discharge speed. W is determined by the time-dependent current
d
R∫I2(t)dt
curves according to the equation: W = , where R, I, V are the load resistance (13 kΩ),
d
V
current and active volume of sample4-6, respectively, as shoSwn in Fig. 1e. W increases
d
monotonically from 0.4 J cm-3 at 10 kV mm-1 to 10.3 J cm-3 at 65 kV mm-1, far beyond currently
reported BF-based bulk ceramics (Fig. 1f). Notably, W Rof x = 0.14 is smaller than W (14.5 J cm-
d rec
3) under the same E, which is ascribed to the clamping effect of domains and the viscous force in
the direct charge-discharge process4. Additionally, the discharge time t is 3.3 μs, which is
0.9
normally utilized to evaluate the discharge time and extracted from the time to reach 90% of W ,
d
as marked in Fig. 1e.

Fig. S8a presents the cumuIlative cycle dependent P-E loops of x = 0.14 ceramic measured at
25 kV mm-1. No significanRt variation in P-E loop can be observed with increasing cycling numbers,
demonstrating excellent fatigue resistant. Accordingly, the W of 3.25 ± 0.02 J cm-3 and the η of
rec
88 ± 1% are achieved over 1-106 cycles (Fig. S8b). Fig. S8c displays the unipolar P-E loops
measured at various temperatures under 25 kV mm-1. As the temperature elevates from 30 to 120C,
the P and P slightly increase owing to the enhanced leakage loss and/or ionic conduction at
max r
elevated temperature. Consequently, W slightly decreases from 3.3 to 3.1 J cm-3 while η show a
rec
more pronounced reduction from 88.5 to 81.5% (Fig. S8d), revealing a satisfactory thermal
stability.

2.2 Characterization of local structure
To probe the underlying mechanism for the outstanding ESP, transmission electron
microscopy (TEM) was employed to directly explore the 2D local microstructure of x = 0.14.
Conventional ferroelectric materials typically exhibit lamellar domains visible in dark-field TEM
imaging8,23. However, such long-range domains are absent in the x = 0.14 sample (Fig. 2a). Instead,
x = 0.14 displays a slim P-E loop (Fig. S2), characteristic of relaxor behavior. This combination
suggests the presence of PNRs or fine-scale nanodomains. Moreover, x = 0.14 exhibits a long-
range average cubic structure (Fig. 2b, 2c), consistent with X-ray diffraction (XRD) patterns (Fig.
S9). The high-angle annular dark-field (HAADF) image recorded by atomic-resolution STEM
along the [001] is shown in Fig. 2d. The larger and brighter dots are A-site atoms while the smaller
and dimmer atoms are B-site atoms, due to their different ionic radius and atomic masses11. The
diffused c/a ratio ranging from 0.97 to 1.06 indicates the fluctuating local lattice distortion (Fig.
2e). The intrinsic ferroelectric polarization stems from the strucStural distortion. So, this fluctuating
lattice distortion yields various local polarizations24,25. The atomic displacement vectors,
determined by calculating the offset of A/B-site atoPms from the center of nearest four B/A-site
atoms via 2D Gaussian function fitting1,24, are ilNlustrated in Fig. 2f. The length and color of arrows
represent the magnitude and direction of these vectors, respectively. The atomic displacement
vectors reveal regions with polar orieLntations corresponding to rhombohedral (R), orthorhombic
(O), tetragonal (T), and cubic (C)-like symmetries (Fig. 2g). This local symmetry frustration and
the coexistence of polymorphic polar configurations are characteristic of relaxor systems with
compositional disorderA and are consistent with the formation of PNRs. A wide distribution of polar
direction suggests that these polar clusters are connected by the monoclinic (M) symmetry low
symmetry in the transitional regions (Fig. 2h), which facilitates reducing the polarization
anisotropy and polarization switching barrier, thereby resulting in a decreased polarization
hysteresis, a low P and an enhanced η. Moreover, 1-4 nm sized polar clusters with polar magnitude
r
up to 18 pm are embedded into the matrix (Fig. 2i), demonstrating the strong local polar clusters,
which does favor to retain a high P /ΔP under applied E, and thus enhance W .
max rec

To elucidate the role of distinct ions in the formation of polar clusters, the atomic pair-
distribution function (PDF) based on the neutron total scattering is performed, which reveals the
local structure information from subnanometre to nanoscale26. Fig. 3a depicts the refinement
results of real-space data G(r) and reciprocal-space S(Q) data of x = 0.14 derived from the reverse
Monte Carlo (RMC) algorithm using a 20 × 20 × 20 perovskite supercell (~80 × 80 × 80 Å)27,28.
The perfect matching between experimental and calculated data suggests high reliability. Fig. 3b
exhibits the PDF G(r) pattern in the low-r range, manifesting the first coordination environment
for different atomic bands. The extracted partial A/B-O PDFs further indicates a large difference
in the local distribution of different atomic pairs (Fig. 3c). The obvious separated peaks of Bi-O,
Sr-O and Na-O atomic pairs suggest the strong local polar displacement of Bi, Sr and Na29. The
specific atomic polar displacement vector can be obtained from the deviation between A/B-site ion
and the center of AO and BO polyhedron. The 3D polar projections of Bi3+ along [001] are
shown in Fig. 3d. The random distribution of polar vectors indicates the coexisting of R, T and O
symmetry. In addition, the Bi3+ achieves the largest average polar displacement of 0.38 Å among
all ions (Fig. 3e), ascribing to its unique electronic structure3,30. And the average polar
displacement of Nb5+ is larger than that of Fe3+ and Ti4+, indicating the improvement of local polar
displacement at the B-site. The reconstructed 3D atomic model and polar vector distributions (Fig.
3f-i) reveal significant A-site positional disorder and local chemical clustering. These features,
commonly observed in multi-component perovskites, contribute to the nanoscale polar
heterogeneity that underpins the relaxor behavior. The inStroduced Na+ ions are randomly
distributed in the matrix consisted by Bi3+ and Sr2+, which tend to gather together to form nanoscale
clustering. Similar phenomenon is also found in the Pdistribution of different oxygen octahedron.
Accordingly, a wide polar direction distributioNn, as well as several 1-4 nm polar clusters with a
high polarization as high as 50 μC cm-3, is observed, consistent with STEM results. This result
suggests that the engineered local struLcture is induced by the heterogeneous distribution of atoms
with differing ionic radii, electronic configurations, and polarizabilities3,18, that collectively
facilitate a high P /ΔP, a low P, and superior ESP.
max r

2.3 Phase-field simulation for the role of polar structure
To further elucidate the role of local polar structure, phase-field simulations were performed
to visualize polarization vectors and their evolution under external E. The simulated configuration,
shown in Fig. 4a, exhibits polarization vectors with diverse orientations, representing polymorphic
nanodomains. Fig. 4b shows polar fluctuations with spontaneous polarization ~0.3 C m-2 randomly
embedded in the weak-polar matrix, closely resembling the polarization configurations of x = 0.14
composition. With increasing E, the polarization of matrix rapidly aligns with E. In contrast, the
embedded local polar fluctuations retain their initial orientation. As E is further increased, these
local polar fluctuations gradually align along the E, promoting the development of enhanced long-
range ferroelectric orders and a further increase in overall polarization. This behaviour indicates
that the existence of strong local fluctuations delays the polarization saturation and facilitates a
higher P under high E31, a trend consistent with the E-dependent P-E loops observed for x =
max
0.14 (Fig. 1a). Upon removal of the external E, the polarization vectors largely revert to their initial
randomized and low magnitude state, suggesting a low energy barrier for domain switching and
minimal energy dissipation during charging/discharging procesSs, thereby leading to a low P.
r
This
mechanism illustrates how relaxors with pronounced local polar structure can circumvent the
conventional trade-off between P and P under anP applied field, which concurrently enhancing
max r
the W and η.
rec

2.4 Mechanism for electric breakdown
The large E also plays an indispensable role in achieving excellent ESP of x = 0.14. As shown
b
in Fig. 5a, the distribution of actual E for each sample matches well with the Weibull distribution
b
analysis1,16. The average E elevates from 13 kV mm-1 for x = 0 to 66 kV mm-1 for x = 0.14 (Fig.
b
5b), resulting from complicated factors, including but not limited to microstructure, insulating
properties4. The dense microstructure with homogeneous element distribution is observed (Fig.
S10, S11). The average grain size decreases significantly from micron- to submicron-scale with
increasing NN content (Fig. 5b and S12). The smaller grain size corresponds to the more grain
boundaries, acting as barriers for the electric breakdown. To demonstrate this, the finite element
simulation on the distribution of the local E, electric potential distribution, and electric tree is
implemented (Fig. 5c, 5d, S13, and S14). It is well-documented that the ε of grain is much higher
r
than that of grain boundary23, leading to the higher E distribution in the grain boundary.
Additionally, the reduced grain size enables the more uniform distribution of local E and electric
potential. The starting electric breakdown stems from the weaSk areas with concentrated E. With
increasing time duration, the electric branches gradually propagate through the grains and grain
boundary. Particularly, the grain boundary is more easPily to attract the electric branches, dispersing
and dissipating the externally loaded electrostaNtic energy. Accordingly, the propagation speed of
electric tree slows down in x = 0.14 with h igher grain boundary density, leading to an enhanced
E
b .

The insulating properties can be analyzed from two aspects. On the one hand, Fig. 5e, S15
depicts the impedance spectra of BF-ST-xNN-Mn ceramics, where Zʹ/Zʺ is the real/imaginary part
of complex impedancAe. The composition-dependent total resistivity, extracted from the low
frequency intercept of the arc on the Zʹ axis, steadily elevates from x = 0 to x = 0.14 (Fig. S16a).
The activation energy E , calculated from the Arrhenius law (Fig. S16b), achieves a peak value at
a
x = 0.14 (Fig. 5e), suggests the optimized insulating performance. On the other hand, the band gap
E represents the energy to excite electrons jumping from the top of valence band to the bottom of
g
conduction band4. As plotted in Fig. 5f, S17, the improved E also suggests the enhanced high
g
insulating properties, leading to an improved E of x = 0.14.
b

In summary, BF-ST-xNN-Mn ceLramics with an ultrahigh W
rec
of 14.5 J cm-3 and a high η of
88% were simultaneously obtained via local structure design. Experimentally, the atomic
displacement vectors mapping derived from STEM images and 3D atomic model from neutron
total scattering usingA RMC method were performed to identify the local polar structure.
Theoretically, the positive effect of local structure on the polarization behavior was justified by
phase-field simulation. Moreover, the large E originated from the synergistic effect of multiple
b
factors, including reduced grain size, enhanced insulating properties, and so on. Accordingly, a
giant ΔP of 56.4 μC cm-2 under a giant E of 65 kV mm-1 was achieved, indicating the delayed
restrictive relationship between P , P and E . These results establish a viable approach to
max r b
designing high-performance dielectric ceramics for the cutting-edge capacitive energy-storage.

Methods
The detail in experimental procedure is provided in the Supplementary Information.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This manuscript reports ultrahigh recoverable energy density in the BFO-STO-NN system and attributes the high performance to a combination of defect-dipole engineering and local-structure design. However, the manuscript suffers from major conceptual inconsistencies, unresolved contradictions, and insufficient experimental validation of key mechanisms. As it stands, I cannot recommend publication.\n1. A central issue lies in the phase-field simulations that are used to justify the role of defect dipoles. The simulations are presented without essential information such as the meaning of the color scale, the magnitude of polarization represented, or the physical interpretation of the 1% and 4% defect-dipole concentrations. Without clear definitions or scale bars, it is impossible to assess the physical relevance of the simulated domain patterns. More critically, the simulated structures still show large and continuous ferroelectric domains even at 4% defect level, which contradicts the manuscript’s repeated claim that defect dipoles destroy long-range order to form polar nanoregions. Despite only modest changes in the simulated domain structures, the authors report dramatic differences in the P–E loops, yet no mechanistic explanation is provided to reconcile this inconsistency. This disconnect between the simulation, and the experimental interpretation is a major flaw that undermines the proposed mechanism.\n2. The manuscript also presents contradictory arguments regarding the origin of the observed nanodomain or relaxor-like behavior. In the simulation, the authors attribute the formation of “nanodomains” directly to defect dipoles generated by acceptor doping. Experimentally, however, they state that defect concentrations must remain low in order to preserve insulating properties, implying that defect levels cannot be high enough to generate significant structural disruption. At the same time, the authors attribute the emergence of polar nanoregions to NaNbO3 substitution, rather than to defects. These conflicting explanations are never resolved. If PNRs primarily arise from compositional complexity rather than defect dipoles, then the simulation, which assumes defect-driven nanodomain formation, does not reflect the actual material system. The mechanistic framework therefore appears internally inconsistent.\n3. Furthermore, the roles of NaNbO3 and MnO are described in a confusing and contradictory manner. NaNbO3 is variously referred to as a sintering aid, a relaxor inducer, a source of strong local polar clusters, and a substituent entering the BiFeO3 lattice. Mn is alternately treated as an acceptor dopant and as a processing additive, but the manuscript does not provide chemical or spectroscopic evidence to establish whether Mn truly incorporates into the lattice, what oxidation state it adopts, or whether it indeed produces oxygen vacancies. The electron paramagnetic resonance data shows nearly identical vacancy concentrations for x = 0 and x = 0.14, directly contradicting the assertion that defect dipoles are responsible for the pronounced evolution of the P–E behavior. Thus, the central chemical mechanism, the formation and effect of defect dipoles, is not convincingly demonstrated.\n4. The manuscript also overinterprets the STEM and neutron PDF results. While these techniques do reveal local structural heterogeneity, such heterogeneity is common in BFO-based and multi-component perovskites even without intentional defect engineering. The manuscript attributes all observed polar distortions to the combined effects of NN incorporation and defect dipoles, but alternative and well-known causes, chemical clustering, A-site size disorder, local symmetry frustration, are not considered. As a result, the structural data are used to reinforce a specific mechanistic narrative that is not uniquely supported by the evidence.\n5. The exceptionally high energy-storage performance requires more rigorous validation. The measurements use thin (~80 μm) samples, which can overrate breakdown strengths, and no discussion is provided on fringing-field corrections or parasitic capacitance, effects that have recently been shown to cause major overestimation of Wrec in dielectric materials. Given the very high values reported, it is important that the authors demonstrate that the measurements are free from such artifacts.", "reviewer_2": "The manuscript investigates the improvement of energy-storage properties in BiFeO3-based ceramics through the incorporation of NaNbO3 and MnO. The topic is relevant to the development of high-performance lead-free dielectric capacitors, and the manuscript presents comprehensive characterizations of polarization behavior and microstructure. However, several critical issues must be addressed before the manuscript can be considered for publication.\n1. The parameter x—used throughout the manuscript to describe compositional variation—is not explicitly defined when first introduced. Since the composition is central to the manuscript, a clear definition of how x relates to NaNbO3 and MnO contents is necessary for reproducibility and readability.\n2. While the manuscript suggests that NaNbO₃ contributes to relaxor behavior and MnO introduces defect-dipole polarization, the mechanistic interpretation remains largely qualitative. It would significantly strengthen the paper to provide a more systematic analysis of how incremental changes in composition influence microstructure, polarization, defect chemistry, and energy-storage performance. In particular, the roles of NaNbO₃, MnO, BiFeO₃, and SrTiO₃ should be quantitatively correlated with the observed changes in physical properties, rather than described only conceptually.\n3. The manuscript frequently highlights the excellent performance of the x = 0.14 sample, yet does not sufficiently justify why this particular composition is optimal or how the selection was made. It would be helpful to include quantitative trends across all compositions and to clarify the criteria for determining the optimal point.\n4. The authors highlight the potential of dielectric capacitors for pulsed-power and high-power applications, such as inverters and converters in hybrid vehicles. However, for such applications, fatigue stability and temperature-dependent energy storage performance are essential requirements, but the manuscript has not provided the corresponding experimental validation.\n5. The manuscript provides comparative plots of ΔP, Eb, Wrec, and η versus previously reported BF-based ceramics. However, the figure caption does not specify the literature sources of the data being compared. For transparency and reproducibility, each data point being compared should be clearly referenced in the figure caption or Supporting Information.\n6. The authors report impressive energy-storage metrics for their BiFeO₃-based ceramics; however, a very recent study (DOI: 10.1038/s41467-025-63584-y) has already demonstrated higher Wrec and η values, together with a more comprehensive performance evaluation (including temperature stability and fatigue cycling), which the present manuscript has not addressed. To strengthen the scientific novelty, the authors are encouraged to explicitly clarify how their materials-design strategy or mechanistic insight differs from and advances beyond current state-of-the-art approaches.", "reviewer_3": "The manuscript titled “Ultrahigh energy-storage in lead-free ceramic capacitors via defect-dipole engineering and local structure design” is a work which focuses on creating sintered ceramics for use in capacitors as the dielectric medium, enhancing the capacitor performance using defect-dipole engineering. The paper builds on the existing (1-x)BiFeO3-xSrTiO3 system by adding MnO as a sintering aid in the matrix before sintering and NaNbO3 as the inhibitor of long range ferroelectric order. This serves to increase the recoverable energy density in the ceramic capacitor and increases the breakdown voltage and the overall efficiency of it. The work serves to showcase an advance in lead free ferroelectric capacitors with enhanced performance and in my opinion is worthy of publication given the comments below are addressed thoroughly:\n1. The novelty of this approach is not completely clearly as written in the paper. Is this the first time this multi-component matrix has been used to achieve such an effect? For example, a recent paper (Nat Commun 16, 8580 (2025)) shows a similar enhancement in capacitor behaviour through an increased entropy effect. Is the local structure effect similar to the high entropy effect and if not, how would one differentiate the origin of such effects. This is a key point in future design of such systems.\n2. Is the NaNbO3 spread uniformly across the sintered ceramic, or does it tend to form clusters? The theoretical analysis in Fig 1 and the images in Fig 2 suggest a uniform distribution, and so does Fig 3, but the discussion from line 180-205 is hard to parse for anyone who is not a PDF expert. This needs to be made more clear, especially what each figure subpanel represents and how finally fig 3f-I are obtained/calculated. This discussion is very dense with jargon and hard to follow in its current form.\n3. Can the full composition of the BF-ST be discussed somewhere in the main text, and also of the MnO. Is the BF-ST the 0.6BF-0.4ST? How much MnO is added for incorporation, and why is it now listed as something that could sit in the B site?\n4. The overall presentation of the paper will benefit with major improvements in clarity and a coherent presentation of the work and comparisons, especially in the discussions. All the separate sections make sense on their own, but it is very hard to read the paper from intro to conclusion and put a thread through all the data and results presented. Does Figure 1 serve as the motivation as a whole, because it is not addressed anywhere else again? Could the authors please connect all the different results better as to how they all come together? The data looks great, but it is difficult to parse all of it together and seems very disjointed.\n5. Related to point 4, there are many typos and also sentences which seem like incomplete clauses. I list some of them here:\nSentence starting on line 127, it's not clear what defect content has to do with nanodomains\nSentence on line 134, what are after mentioned displacement vectors\nSentence on line 150, the logic of the suggestion of existence of PNRs and nanodomains is unclear. Discuss further please. All that can be seen is a lack of macroscopic ferroelectric order\nSentence on line 159 starting with “as knows...” seems like a throwaway statement\nThe sentence on line 203-204 claims a polarization value but does not account for any oxygen vacancy effects which might introduce electronic perturbations to the total polarization, so is it correct to state such a value based on just the cationic displacements?\nThe sentence starting on 224 is too long and mentions weak coupling which has only been mentioned before on line 172 but it is not clear how that conclusion is drawn, about the weak coupling? Is the weak coupling between different polar clusters or between polar clusters and defect-dipoles?\nThese are only to name a few. Please correct such sentences and if a logical conclusion is drawn, please explain how such a conclusion was reached.\n6. Please also redefine the acronyms and the variables not just in the abstract but also in the main text. This would improve the readability of the paper a lot."}, "author_response": "We sincerely thank all reviewers for spending their time reviewing our manuscript, NCOMMS-25-88618, entitled \"Ultrahigh energy-storage in lead-free ceramic capacitors via local structure design\", and for providing insightful comments. In the following, we present a point-by-point response to the reviewers’ comments. We hope that these changes adequately address the concerns raised. All changes have been highlighted within the manuscript.\nReviewer #1:\nThis manuscript reports ultrahigh recoverable energy density in the BFO-STO-NN system and attributes the high performance to a combination of defect-dipole engineering and local-structure design. However, the manuscript suffers from major conceptual inconsistencies, unresolved contradictions, and insufficient experimental validation of key mechanisms. As it stands, I cannot recommend publication.\nResponse: We sincerely thank the reviewer for the time and careful consideration given to our manuscript, as well as for the insightful suggestions, which have been invaluable in strengthening our work. We have incorporated these comments into the revised manuscript, and provide a point-by-point response to each suggestion below.\nComment-1: A central issue lies in the phase-field simulations that are used to justify the role of defect dipoles. The simulations are presented without essential information such as the meaning of the color scale, the magnitude of polarization represented, or the physical interpretation of the 1% and 4% defect-dipole concentrations. Without clear definitions or scale bars, it is impossible to assess the physical relevance of the simulated domain patterns. More critically, the simulated structures still show large and continuous ferroelectric domains even at 4% defect level, which contradicts the manuscript’s repeated claim that defect dipoles destroy long-range order to form polar nanoregions. Despite only modest changes in the simulated domain structures, the authors report dramatic differences in the P-E loops, yet no mechanistic explanation is provided to reconcile this inconsistency. This disconnect between the simulation, and the experimental interpretation is a major flaw that undermines the proposed mechanism.\nResponse: We sincerely thank the reviewer for this critical and insightful comment, which identified a major conceptual flaw in our original manuscript. We agree that the phase-field simulations presented to justify the role of defect dipoles were insufficient and created a significant disconnect with the proposed mechanism.\nIn light of this, we have completely removed the discussion on defect dipoles from the manuscript, as it was not central to our core experimental findings and its justification was inadequate. Instead, we have refocused our interpretation on the role of local polar structure (polymorphic nanodomains and fluctuations), which is strongly supported by our experimental data.\nTo address the specific concerns regarding simulation clarity and provide a robust justification for this refined mechanism, we have taken the following actions:\n1. Rewritten the simulation section: We have performed new phase-field simulations to visualize the evolution of local polar structure under an electric field. All figures now include clear color bars to indicate polarization direction and magnitude, and scale bars to define physical dimensions.\n2. Provided full simulation details: All essential parameters—including mesh size (128 × 128), grid spacing (Δx = Δy = 4 nm), system scale (512 × 512 nm), boundary conditions (periodic), Landau coefficients, and electrostrictive/elastic constants—are now explicitly provided in the revised Supplementary Materials.\n3. Reconciled simulation with experiment: The revised simulation directly models the coexistence of strong local polar fluctuations within a weak-polar matrix. It demonstrates how this configuration delays polarization saturation (enhancing Pmax) and allows for easy re-randomization upon field removal (reducing Pr), thereby directly and consistently explaining the characteristic slim P-E loops and excellent energy storage performance observed in our relaxor sample (x = 0.14).\nFig. R1 (Fig. 4 in the revised manuscript) Phase-field simulation for the role of local structure. Evolution of a polarization angle and b magnitude of relaxors with strong local polar structure during loading and unloading E.\nSpecific revisions in the manuscript:\n1. The discussion on the role of local structure by phase-field simulation has been provided in revised main text. Please see page 10-11: “To further elucidate the role of local polar structure, phase-field simulations were performed to visualize polarization vectors and their evolution under external E. The simulated configuration, shown in Fig. 4a, exhibits polarization vectors with diverse orientations, representing polymorphic nanodomains. Fig. 4b shows polar fluctuations with spontaneous polarization ~0.3 C m-2 randomly embedded in the weak-polar matrix, closely resembling the polarization configurations of x = 0.14 composition. With increasing E, the polarization of matrix rapidly aligns with E. In contrast, the embedded local polar fluctuations retain their initial orientation. As E is further increased, these local polar fluctuations gradually align along the E, promoting the development of enhanced long-range ferroelectric orders and a further increase in overall polarization. This behaviour indicates that the existence of strong local fluctuations delays the polarization saturation and facilitates a higher Pmax under high E, a trend consistent with the E-dependent P-E loops observed for x = 0.14 (Fig. 1a). Upon removal of the external E, the polarization vectors largely revert to their initial randomized and low magnitude state, suggesting a low energy barrier for domain switching and minimal energy dissipation during charging/discharging process, thereby leading to a low Pr. This mechanism illustrates how relaxors with pronounced local polar structure can circumvent the conventional trade-off between Pmax and Pr under an applied field, which concurrently enhancing the Wrec and η.”\n2. We have added a dedicated subsection detailing all phase-field simulation parameters, as quoted in the original response, to ensure full reproducibility and transparency. Please see page S4 in revised Supplementary Materials: “In the phase field simulation, a 128 × 128 mesh with periodic boundary conditions was used, and the grid space in real space was chosen to be Δx = Δy =4 nm. So, the scale for domain morphology is 512 × 512 nm. The normal ferroelectrics are randomly distributed in the weak-polar matrix, of which difference are reflected in fL. Based on the theoretical parameters of BaTiO3, the electrostrictive coefficients and elastic constants listed below: Q11 = 0.055 m4 C-2, Q12 = -0.023 m4 C-2, Q44 = 0.0315 m4 C-2, C11 = 17.8×1010 J m-3, C12 = 9.6×1010 J m-3, C44 = 12.2×1010 J m-3. For the normal ferroelectrics, the Landau free energy parameters are: α1 = -6.0×107 J m C-2, α11 = -1.1×109 J m5 C-4, α111 = 1.33×1010 J m9 C-6, α12 = 3.23×108 J m5 C-4, α112 = 4.47×109 J m9 C-6, α123 = 4.91×109 J m9 C-6. For weak-polar matrix, the Landau free energy parameters are: α1 = -1.79×106 J m C-2, α11 = -3.07×107 J m5 C-4, α111 = 2.57×109 J m9 C-6, α12 = 6.08×108 J m5 C-4, α112 = 6.95×109 J m9 C-6, α123 = 1.31×109 J m10 C-6.”\nComment-2: The manuscript also presents contradictory arguments regarding the origin of the observed nanodomain or relaxor-like behavior. In the simulation, the authors attribute the formation of “nanodomains” directly to defect dipoles generated by acceptor doping. Experimentally, however, they state that defect concentrations must remain low in order to preserve insulating properties, implying that defect levels cannot be high enough to generate significant structural disruption. At the same time, the authors attribute the emergence of polar nanoregions to NaNbO3 substitution, rather than to defects. These conflicting explanations are never resolved. If PNRs primarily arise from compositional complexity rather than defect dipoles, then the simulation, which assumes defect-driven nanodomain formation, does not reflect the actual material system. The mechanistic framework therefore appears internally inconsistent.\nResponse: Thanks for raising this important point regarding the mechanistic inconsistency in our original manuscript. We agree that the attribution of polar nanoregions (PNRs) to defect dipoles in the simulations created a fundamental disconnect with the experimental system, where PNRs arise primarily from compositional heterogeneity.\nWe have fully addressed this by removing all discussion of defect dipoles and refocusing the entire narrative on a unified, composition-driven mechanism. The revised manuscript now consistently attributes the formation of PNRs and local polar structure to the substitution of NaNbO3 (NN). The role of local structure is now justified by phase-field simulations that model compositional disorder directly, eliminating the previous contradiction.\nSpecific revisions in the manuscript:\n1. Clarification of the origin of PNRs is revised in the main text. Please see page 4: “Herein, the perovskite end-member NN was incorporated to disrupt the long-range ferroelectric orders and promote the formation of PNRs while retaining highly polarized local clusters3,18, leading to a slim P-E loop with a relatively high Pmax and a significantly reduced Pr.” This establishes compositional substitution as the sole primary cause from the outset.\n2. Unified explanation for local structure is added in revised main text. Please see page 10: “This result suggests that the engineered local structure is induced by the heterogeneous distribution of atoms with differing ionic radii, electronic configurations, and polarizabilities3,18, that collectively facilitate a high Pmax/ΔP, a low Pr, and superior ESP.”\n3. Updated simulation justification (please see above response to comment-2).\nCorrespondingly, the phase-field simulations in the revised manuscript (Fig. 4) no longer reference defect dipoles. Instead, they are framed as modeling the local polar fluctuations and polymorphic nanodomains inherent to a relaxor system with compositional disorder, thereby providing direct and consistent support for the proposed mechanism.\nWe believe these revisions have resolved the internal inconsistency by presenting a single, coherent mechanism where compositional complexity (NN substitution) drives the formation of PNRs and local polar structure, which in turn govern the macroscopic energy storage properties. Thanks again for the guidance in strengthening this core argument.\nComment-3: Furthermore, the roles of NaNbO3 and MnO are described in a confusing and contradictory manner. NaNbO3 is variously referred to as a sintering aid, a relaxor inducer, a source of strong local polar clusters, and a substituent entering the BiFeO3 lattice. Mn is alternately treated as an acceptor dopant and as a processing additive, but the manuscript does not provide chemical or spectroscopic evidence to establish whether Mn truly incorporates into the lattice, what oxidation state it adopts, or whether it indeed produces oxygen vacancies. The electron paramagnetic resonance data shows nearly identical vacancy concentrations for x = 0 and x = 0.14, directly contradicting the assertion that defect dipoles are responsible for the pronounced evolution of the P-E behavior. Thus, the central chemical mechanism, the formation and effect of defect dipoles, is not convincingly demonstrated.\nResponse: Thank you for this important comment, which has helped us clarify a key source of confusion in our original manuscript. We acknowledge the description of the roles of NN and MnO was inconsistent and that the evidence for defect dipole formation through Mn doping was not conclusively demonstrated.\nIn the revised manuscript, we have taken two primary actions to resolve these issues:\n1. We have removed all claims regarding defect dipoles from the manuscript, as they were not central to our experimental findings and lacked sufficient supporting evidence. This directly addresses the core critique that the chemical mechanism was not convincingly demonstrated.\n2. We have refined and clarified the distinct roles of NN and MnO to eliminate contradiction and present a unified, evidence-based narrative.\nThe revised manuscript now presents the following consistent framework:\n• Role of NN: NN is treated consistently as a substituent entering the perovskite lattice. Its primary role is to induce relaxor behavior by disrupting long-range ferroelectric order and promoting polar nanoregions (PNRs), as supported by our structural and electrical data. A secondary, beneficial role is its contribution to a larger effective band gap, which helps improve resistivity and breakdown strength (Eb).\n• Role of MnO: Based on the literature for BiFeO3-based systems and our processing conditions, MnO is described as a sintering aid and processing additive. Its accepted function is to suppress the formation of oxygen vacancies and secondary phases by maintaining charge balance during high-temperature processing, thereby reducing leakage current and enhancing Eb. We no longer claim it acts as an acceptor dopant that forms defect dipoles, as our EPR data showed similar vacancy concentrations across compositions, which does not support a doping-driven defect evolution.\nSpecific revisions in the manuscript:\nThe role of NN and MnO is revised in the main text. Please see page 4-5: “Herein, the perovskite end-member NN was incorporated to disrupt the long-range ferroelectric orders and promote the formation of PNRs while retaining highly polarized local clusters3,18, leading to a slim P-E loop with a relatively high Pmax and a significantly reduced Pr. Meanwhile, NN possess a large band gap (Eg ~ 3.4 eV), which is favorable for improving resistivity and enhance Eb. Additionally, the synthesis of phase-pure BF-based ceramics is challenging due to volatilization of Bi2O3 and the reduction of Fe3+ (Fe3+ + e- → Fe2+) during sintering, which typically leads to high dielectric loss and leakage. In this study, MnO was added as a sintering aid to mitigate these issues. By helping to maintain the charge balance, MnO suppresses the formation of oxygen vacancies and secondary phase, thereby reducing leakage current and enhancing Eb4,8”\nWe believe these revisions resolve the previous contradictions, align the narrative with the experimental evidence, and present a clear, consistent mechanism for the observed performance. We thank the reviewer for prompting this essential clarification.\nComment-4: The manuscript also overinterprets the STEM and neutron PDF results. While these techniques do reveal local structural heterogeneity, such heterogeneity is common in BFO-based and multi-component perovskites even without intentional defect engineering. The manuscript attributes all observed polar distortions to the combined effects of NN incorporation and defect dipoles, but alternative and well-known causes, chemical clustering, A-site size disorder, local symmetry frustration, are not considered. As a result, the structural data are used to reinforce a specific mechanistic narrative that is not uniquely supported by the evidence.\nResponse: Thanks for this insightful comment. We agree that our original interpretation of the STEM and neutron PDF data was overly narrow, attributing structural heterogeneity exclusively to a specific mechanism without adequately considering other common and contributing factors in complex perovskites. In the revised manuscript, we have broadened our discussion to present a more balanced and evidence-based interpretation. We maintain that the combination of STEM and neutron PDF is powerful for characterizing local structure. However, we now frame the observed heterogeneity—including atomic displacements, symmetry variations, and chemical clustering—as characteristic features of a compositionally complex relaxor system, rather than as direct evidence for a singular origin.\nSpecific revisions in the manuscript:\nThe chemical clustering, A-site size disorder and local symmetry frustration have been added in revised main text. Please see page 8 (STEM analysis): “The atomic displacement vectors reveal regions with polar orientations corresponding to rhombohedral (R), orthorhombic (O), tetragonal (T), and cubic (C)-like symmetries (Fig. 2g). This local symmetry frustration and the coexistence of polymorphic polar configurations are characteristic of relaxor systems with compositional disorder and are consistent with the formation of PNRs.”, and page 10 (Neutron PDF analysis): “The reconstructed 3D atomic model and polar vector distributions (Fig. 3f-i) reveal significant A-site positional disorder and local chemical clustering. These features, commonly observed in multi-component perovskites, contribute to the nanoscale polar heterogeneity that underpins the relaxor behavior.”\nFurthermore, throughout the manuscript, we have refined our overall mechanistic narrative to be more precise. We now consistently attribute the primary origin of the relaxor behavior and PNRs to the compositional complexity introduced by NN substitution, which inherently leads to A-site disorder, local symmetry frustration, and chemical clustering. The advanced local probes (STEM and PDF) are thus presented as tools that quantitatively characterize this inherent disorder, rather than as evidence for an externally imposed defect-dipole mechanism.\nWe believe these revisions correct the overinterpretation, align the discussion with well-established concepts in relaxor ferroelectrics, and provide a more robust and credible foundation for our conclusions. Thanks again for guiding us toward this more rigorous presentation.\nComment-5: The exceptionally high energy-storage performance requires more rigorous validation. The measurements use thin (~80 μm) samples, which can overrate breakdown strengths, and no discussion is provided on fringing-field corrections or parasitic capacitance, effects that have recently been shown to cause major overestimation of Wrec in dielectric materials. Given the very high values reported, it is important that the authors demonstrate that the measurements are free from such artifacts.\nResponse: Thanks for raising this important methodological point regarding the rigorous validation of energy storage performance. We acknowledge the necessity of addressing potential overestimation due to geometric factors, fringing fields, and parasitic capacitance, especially given the high values reported. We have taken this concern seriously and conducted additional control experiments to validate our measurements. In the revised manuscript, we have added a dedicated section in the Supplementary Materials to address this issue directly.\nWe have added a discussion citing the recent critical work by Li et al. (Nat. Commun. 2025, 16, 608, doi: 10.1038/s41467-025-55855-5) on measurement artifacts in thin dielectric samples, framing it as an essential consideration for reporting reliable energy storage properties. For our primary sample (x = 0.14, thickness t = 0.08 mm, electrode diameter d = 1 mm), the d/t ratio is 12.5. According to the scaling principles highlighted in the literature, this geometry, combined with the material's high relative permittivity (~450 at room temperature, Fig. S3), suggests a relatively lower susceptibility to fringing-field overestimation. To conclusively rule out significant artifact, we prepared and tested a control sample with a more conventional geometry (t = 0.2 mm, d = 9.0 mm, d/t = 45). The measured recoverable energy density (Wrec) for this control sample showed a deviation of only ~9% compared to the primary thin sample, as illustrated in Fig. R2. This minor deviation confirms that the exceptionally high Wrec values reported in our study are not a major artifact of sample geometry and robustly represent the material's intrinsic performance.\nAccording to the reviewer’s comment, we have revised the manuscript and added Fig. S4 for the validation of energy storage measurement in the Supplementary Materials.\nFig. R2 (Fig. S4 in revised Supplementary Materials) P-E loops of x = 0.14 composition under 19 kV mm-1. S1: 0.08 mm in thickness, 1 mm in electrode diameter. S2: 0.2 mm in thickness, 9.0 mm in electrode diameter.\nReviewer #2:\nThe manuscript investigates the improvement of energy-storage properties in BiFeO3-based ceramics through the incorporation of NaNbO3 and MnO. The topic is relevant to the development of high-performance lead-free dielectric capacitors, and the manuscript presents comprehensive characterizations of polarization behavior and microstructure. However, several critical issues must be addressed before the manuscript can be considered for publication.\nResponse: We greatly appreciate the reviewer for their time to review our manuscript and give us the enlightening suggestions, which are very helpful for improving this work. We have revised our manuscript accordingly, and the point-by-point responses to the reviewer’s suggestions are enclosed as follows.\nComment-1: The parameter x-used throughout the manuscript to describe compositional variation-is not explicitly defined when first introduced. Since the composition is central to the manuscript, a clear definition of how x relates to NaNbO3 and MnO contents is necessary for reproducibility and readability.\nResponse: Thanks for the valuable comment. The parameter x is the NaNbO3 content in the chemical formula (0.6-x)BiFeO3-0.4SrTiO3-xNaNbO3+0.1wt% MnO. We have defined this parameter in the revised manuscript at its first usage.\nSpecific revisions in the manuscript:\nPlease see page 4: “…we propose the local structure design to overcome the restraint between different parameters in (0.6-x)BiFeO3-0.4SrTiO3-xNaNbO3+0.1wt% MnO (denote as BF-ST-xNN-Mn) ceramics…”\nComment-2: While the manuscript suggests that NaNbO3 contributes to relaxor behavior and MnO introduces defect-dipole polarization, the mechanistic interpretation remains largely qualitative. It would significantly strengthen the paper to provide a more systematic analysis of how incremental changes in composition influence microstructure, polarization, defect chemistry, and energy-storage performance. In particular, the roles of NaNbO3, MnO, BiFeO3, and SrTiO3 should be quantitatively correlated with the observed changes in physical properties, rather than described only conceptually.\nResponse: We appreciate the valuable comments. We agree that a more systematic and quantitative analysis would significantly strengthen the paper. In the revised manuscript, we have addressed this by removing the ambiguous defect-dipole narrative and instead providing a clear, stepwise correlation between compositional changes, microstructure evolution, and the resultant electrical properties. We have structured this analysis into three key aspects:\n1. Microstructure\nWe have added a detailed analysis showing that increasing NN content systematically refines the grain size from 1.6 μm (x = 0) to 0.4 μm (x = 0.16) (Fig. S10). This quantitative trend is directly linked to the dramatic improvement in Eb, which increases from 12.8 kV mm-1 (x = 0) to 66.0 kV mm-1 (x = 0.14). The analysis explicitly states that the increased grain boundary density acts as barriers to charge transport, a relationship vividly supported by electric field simulation (Figs. 5c, d). This provides a quantitative microstructure-property link.\n2. Polarization and energy storage performance\nFigs. R3 and R4a depict the P-E loops of BF-ST-xNN-Mn ceramics under the same E of 10 kV mm-1. With increasing NN content, the P-E loops become gradually slimmer, accompanied by decreased Pmax and Pr from 39.7 and 30.4 μC cm-2 (x = 0) to 8.5 and 0.2 μC cm-2 (x = 0.16), consistent with enhanced relaxor behavior. Accordingly, the η progressively increases from 13.2% at x = 0 to 91.3% at x = 0.16 (Fig. R4b). Fig. R4c exhibits the P-E loops of BF-ST-xNN-Mn ceramics under the respective maximum applied E. Samples were prepared with a thickness of ~80 μm and an electrode area of ~ 0.79 mm2. A high Pmax exceeding 40 μC cm-2 under applied E is maintained across compositions x = 0.05-0.16, yielding a high ΔP. Particularly, the composition x = 0.14 achieves a giant Pmax of 60 μC cm-2, a low Pr of 3.6 μC cm-2, and the highest ΔP of 56.4 μC cm-2 under 65 kV mm-1, thereby delivering an optimal ESP with Wrec = 14.5 J cm-3 and η = 88% (Fig. R4d). These results suggest that x = 0.14 achieves a good balance between Pmax and Pr, ΔP and E, and Wrec and η.\nFig. R3 (Fig. S1 in revised Supplementary Materials) P-E loop of 0.6BF-0.4ST ceramics under 10 kV mm-1.\nFig. R4 (Fig. S2 in revised Supplementary Materials) a P-E loops of BF-ST-xNN-Mn ceramics under 10 kV mm-1, and corresponding b Wrec, Wtotal, η. c P-E loops of BF-ST-xNN-Mn ceramics under the maximum applied E, and corresponding d Wrec, Wtotal, η.\n3. Clarified and distinct roles of each component\nThe qualitative description of roles has been refined and grounded in the observed quantitative trends:\nNaNbO3 was incorporated to disrupt the long-range ferroelectric orders to form PNRs while retaining highly polarized clusters in local regions. Meanwhile, NaNbO3 possess a large band gap (Eg ~ 3.4 eV), being conducive to the improved resistivity and enhanced Eb.\nThe additive MnO, described consistently as a sintering aid, can also suppress oxygen vacancies and leakage, thereby supporting the achievement of high Eb. Its role is now separate from the polarization mechanism.\nBiFeO3 provides the foundational high polarization potential via Bi-O hybridization, explaining the retained high Pmax even in the relaxor state (Fig. 3). Additionally, SrTiO3’s role is forming a morphotropic phase boundary (MPB) with BiFeO3, cited as the basis for achieving high intrinsic polarization in the host matrix ~40 μC cm-2 (Energy Environ. Sci. 2020, 13, 2938-2948, doi: 10.1039/d0ee02104k).\nSpecific revisions in the manuscript:\n1. The discussion for the microstructure has been added in revised Supplementary Materials. Please see page S11: “Fig. S10 | Typical microstructure morphologies of a x = 0, b x = 0.05, c x = 0.08, d x = 0.10, e x = 0.12, f x = 0.14, and g x = 0.16. Please note that the scale ruler of x = 0 is different with others. All samples exhibit dense, pore-free microstructure. With increasing NN content, the average grain size decreases significantly from 1.6 μm (x = 0) to 0.4 μm (x = 0.16). The grain refinement suggests higher density of grain boundaries, which act as space-charge depletion layers that impede charge carrier transport, resulting in a dramatically improved Eb from 12.8 kV mm-1 (x = 0) to 66.0 kV mm-1 (x = 0.14).”\n2. The evolution of polarization and energy storage performance has been revised in the main text. Please see page 5-6: “A series of BF-ST-xNN-Mn (x = 0-0.16) ceramics were synthesized by solid-state reaction, and their P-E loops under 10 kV mm-1 were depicted in Figs. S1, S2a. With increasing NN content, the P-E loops become gradually slimmer, accompanied by decreased Pmax and Pr from 39.7 and 30.4 μC cm-2 (x = 0) to 8.5 and 0.2 μC cm-2 (x = 0.16), consistent with enhanced relaxor behavior (Fig. S3). Accordingly, the η progressively increases from 13.2% at x = 0 to 91.3% at x = 0.16 (Fig. S2b). Fig. S2c exhibits the P-E loops of BF-ST-xNN-Mn ceramics under the respective maximum applied E. Samples were prepared with a thickness of ~80 μm and an electrode area of ~ 0.79 mm2. This sample geometry, combined with the high relative permittivity (Fig. S3), suggests a relatively lower susceptibility to fringing-field effect and parasitic capacitance (Fig. S4)7. A high Pmax exceeding 40 μC cm-2 under applied E is maintained across compositions x = 0.05-0.16, yielding a high ΔP. Particularly, the composition x = 0.14 achieves a giant Pmax of 60 μC cm-2, a low Pr of 3.6 μC cm-2, and the highest ΔP of 56.4 μC cm-2 under 65 kV mm-1, thereby delivering an optimal ESP with Wrec = 14.5 J cm-3 and η = 88% (Fig. S2d). These results suggest that x = 0.14 achieves a good balance between Pmax and Pr, ΔP and E, and Wrec and η. It was therefore selected as the representative composition for further study.”\n3. The roles of BF-ST, NN and MnO are revised in the main text. Please see page 4-5: “0.6BF-0.4ST ceramic with a morphotropic phase boundary demonstrates significant potential for energy-storage capacitors, due to its high intrinsic polarization ~40 μC cm-2 20,21.”, and “Herein, the perovskite end-member NN was incorporated to disrupt the long-range ferroelectric orders and promote the formation of PNRs while retaining highly polarized local clusters3,18, leading to a slim P-E loop with a relatively high Pmax and a significantly reduced Pr. Meanwhile, NN possess a large band gap (Eg ~ 3.4 eV), which is favorable for improving resistivity and enhance Eb. Additionally, the synthesis of phase-pure BF-based ceramics is challenging due to volatilization of Bi2O3 and the reduction of Fe3+ (Fe3+ + e- → Fe2+) during sintering, which typically leads to high dielectric loss and leakage. In this study, MnO was added as a sintering aid to mitigate these issues. By helping to maintain the charge balance, MnO suppresses the formation of oxygen vacancies and secondary phase, thereby reducing leakage current and enhancing Eb4,8.”\nComment-3: The manuscript frequently highlights the excellent performance of the x = 0.14 sample, yet does not sufficiently justify why this particular composition is optimal or how the selection was made. It would be helpful to include quantitative trends across all compositions and to clarify the criteria for determining the optimal point.\nResponse: Thanks for the valuable comment. Fig. R3 and R4 depict the P-E loops of BF-ST-xNN-Mn ceramics under the same E of 10 kV mm-1 and breakdown strength. Thought the Pmax at the same E decreases with increasing NN, but the increasing Eb benefits achieving the higher Pmax, yielding a high Wrec = 14.5 J cm-3 and a high η = 88%. These results suggest that x = 0.14 shows a good balance between Pmax and Pr, ΔP and E, and Wrec and η. Therefore, x = 0.14 was selected as representative for the further study.\nSpecific revisions in the manuscript:\nThe reason for selecting x = 0.14 as the optimal composition has been provided in revised main text. Please see page 5-6: “A series of BF-ST-xNN-Mn (x = 0-0.16) ceramics were synthesized by solid-state reaction, and their P-E loops under 10 kV mm-1 were depicted in Figs. S1, S2a. With increasing NN content, the P-E loops become gradually slimmer, accompanied by decreased Pmax and Pr from 39.7 and 30.4 μC cm-2 (x = 0) to 8.5 and 0.2 μC cm-2 (x = 0.16), consistent with enhanced relaxor behavior (Fig. S3). Accordingly, the η progressively increases from 13.2% at x = 0 to 91.3% at x = 0.16 (Fig. S2b). Fig. S2c exhibits the P-E loops of BF-ST-xNN-Mn ceramics under the respective maximum applied E. Samples were prepared with a thickness of ~80 μm and an electrode area of ~ 0.79 mm2. This sample geometry, combined with the high relative permittivity (Fig. S3), suggests a relatively lower susceptibility to fringing-field effect and parasitic capacitance (Fig. S4)7. A high Pmax exceeding 40 μC cm-2 under applied E is maintained across compositions x = 0.05-0.16, yielding a high ΔP. Particularly, the composition x = 0.14 achieves a giant Pmax of 60 μC cm-2, a low Pr of 3.6 μC cm-2, and the highest ΔP of 56.4 μC cm-2 under 65 kV mm-1, thereby delivering an optimal ESP with Wrec = 14.5 J cm-3 and η = 88% (Fig. S2d). These results suggest that x = 0.14 achieves a good balance between Pmax and Pr, ΔP and E, and Wrec and η. It was therefore selected as the representative composition for further study.”\nComment-4: The authors highlight the potential of dielectric capacitors for pulsed-power and high-power applications, such as inverters and converters in hybrid vehicles. However, for such applications, fatigue stability and temperature-dependent energy storage performance are essential requirements, but the manuscript has not provided the corresponding experimental validation.\nResponse: Thanks for the valuable comment. According to the reviewer’s suggestion, we have performed additional experiments on fatigue stability and temperature-dependent energy storage.\nFig. R5a exhibits the cumulative cycle dependent P-E loops of x = 0.14 ceramic measured at 25 kV mm-1. No significant variation in P-E loop can be observed with increasing cycling numbers. So, the Wrec of 3.25±0.02 J cm-3 and the η of 88±1% are achieved in 1-106 cycles (Fig. R5b), suggesting the excellent anti-fatigue performance. Fig. R5c plots the unipolar P-E loops of x = 0.14 measured at various temperatures under 25 kV mm-1. As the temperature elevates from 30 to 120C, the Pmax and Pr of ceramics slightly increase gradually owing to the high temperature induced leakage loss and/or ionic conduction. As a result, a decreased Wrec from 3.3 J cm-3 to 3.1 J cm-3 and a reduced η from 88.5% to 81.5% are observed (Fig. R5d), revealing satisfactory thermal stability.\nFig. R5 (Fig. S8 in the revised Supporting Information) Unipolar P-E curves of x = 0.14 at different (a) cumulative cycle numbers, and (c) temperature, and corresponding (b), (d) Wrec and η.\nSpecific revisions in the manuscript:\nThe fatigue stability and temperature-dependent energy storage performance are provided in Fig. S8 (Supplementary Material), and related discussion is added in revised main text. Please see page 7: “Fig. S8a presents the cumulative cycle dependent P-E loops of x = 0.14 ceramic measured at 25 kV mm-1. No significant variation in P-E loop can be observed with increasing cycling numbers, demonstrating excellent fatigue resistant. Accordingly, the Wrec of 3.25±0.02 J cm-3 and the η of 88±1% are achieved over 1-106 cycles (Fig. S8b). Fig. S8c displays the unipolar P-E loops measured at various temperatures under 25 kV mm-1. As the temperature elevates from 30 to 120C, the Pmax and Pr slightly increase owing to the enhanced leakage loss and/or ionic conduction at elevated temperature. Consequently, Wrec slightly decreases from 3.3 to 3.1 J cm-3 while η show a more pronounced reduction from 88.5 to 81.5% (Fig. S8d), revealing a satisfactory thermal stability.”\nComment-5: The manuscript provides comparative plots of ΔP, Eb, Wrec, and η versus previously reported BF-based ceramics. However, the figure caption does not specify the literature sources of the data being compared. For transparency and reproducibility, each data point being compared should be clearly referenced in the figure caption or Supporting Information.\nResponse: Thanks for the valuable comment. Following the reviewer’s suggestion, the detailed information of BF-based ceramics from literatures has been provided in the Table S1 in revised Supplementary Material.\nTable S1 Comparison of energy storage performance of BF-based ceramics\nE (kV mm-1) ΔP (μC cm-2) Wrec (J cm-3) η (%) Wd (J cm-3) Composition Ref.\n0.35BiFeO3-0.65SrTiO3 75 38 8.4 90 3.3 12\n0.88(0.67BiFeO3-0.33BaTiO3)-0.12Na0.73Bi0.09NbO3 41 33.9 5.57 83.8 -- 13\n0.61BiFeO3-0.33BaTiO3-0.06Ba(Mg1/3Nb2/3)O3 12.5 32.3 1.56 75 -- 14\n0.61BiFeO3-0.33BaTiO3-0.06La(Mg1/2Ti1/2)O3 13 33.3 1.66 82 -- 15\n0.60BiFeO3-0.34BaTiO3-0.06Sr(Al0.5Nb0.5)O3 15.5 37.1 1.75 81 -- 16\n0.56BiFeO3-0.30BaTiO3-0.14AgNbO3-5‰ mol CuO 19.5 25.4 2.11 84 -- 17\n0.97(0.67BiFeO3-0.34BaTiO3)-0.3(Sr0.7Bi0.2)TiO3 18 34.9 2.4 90.4 0.18 18\n0.60BiFeO3-0.34BaTiO3-0.06Ba(Zn1/3Ta2/3)O3 16 43.41 2.56 72 -- 19\n0.59BiFeO3-0.33BaTiO3-0.08BaBi2Nb2O9 23 30.1 3.08 85.6 0.12 20\n0.7(Bi0.97Nd0.03)[Fe0.95(Li0.5Nb0.5)0.05]O3-0.3BaTiO3-0.1 wt% MnO2 29 30.8 3.2 73.7 -- 21\n0.46Bi1.02FeO3-0.29BaTiO3-0.25Bi0.5Na0.5TiO3-1 wt% Nb2O5 25 34.9 3.9 80 3.3 22\n0.53BiFeO3-0.3Ba(Hf0.05Ti0.95)O3-0.17NaNbO3-0.1 wt% MnO2 34.5 37.7 5.2 88 3.7 23\n0.53BiFeO3-0.3Ba(Hf0.05Ti0.95)O3-0.17NaTaO3-0.1 wt% MnO2 42.5 35.7 6.3 86.6 3.31 24\n0.7(0.67BiFeO3-0.33BaTiO3)-0.3Ca0.85Bi0.05Sm0.05TiO3 30 38.9 5.26 82.4 1.21 25\n0.85(0.67BiFeO3-0.33BaTiO3)-0.15Sr(Nb0.5Al0.5)O3 30 32.7 3.95 85.9 1.15 26\n0.7Bi0.8Sm0.2FeO3-0.3BaTiO3 22 41.1 4.1 78.1 -- 27\n0.6(0.67BiFeO3-0.33BaTiO3)-0.4CaTiO3 38 29.3 5.03 89.7 0.68 28\n0.5BiFeO3-0.5(0.88(Ba0.8Sr0.2)TiO3-0.12Bi(Zn2/3Nb1/3)O3) 68 42.01 7.4 81 3 29\n(Bi0.7Ba0.3)(Fe0.8Ti0.2)O3-0.7(Bi0.8Ta0.2)O3 55 36.0 10.3 68 -- 30\n0.55BiFeO3-0.3BaHf0.05Ti0.95O3-0.15Sr(Al0.5Nb0.5)O3 83 38.1 12.2 90.1 6.6 31\n0.97(0.67BiFeO3-0.33BaTiO3)-0.03SmNbO4 35.5 44.76 6.12 80 0.14 32\n0.59BiFeO3-0.33BaTiO3-0.08Na(Nb0.85Ta0.15)O3 33 45.9 6.1 81.4 3.8 33\n0.5BiFeO3-0.3BaTiO3-0.2(Na0.7Sm0.1)NbO3 65 51.9 13.1 78.3 -- 34\n0.4BiFeO3-0.6Ba(Ti0.8Zr0.2)O3 21 44.8 4.1 86.2 -- 35\n0.4BiFeO3-0.6(0.9(Ba0.75Sr0.25)TiO3-0.1Bi(Zn2/3Ta1/3)O3) 60 42.7 8.3 85.8 5.8 36\n0.5BiFeO3-0.3BaTiO3-0.2Ca(Cr0.5Nb0.5)O3 61 41.6 9.72 72 0.63 37\n0.7(0.6BiFeO3-0.4Sr0.7Bi0.2TiO3)-0.3BaTiO3 13.5 20.8 1.2 88.6 -- 38\n0.38BiFeO3-0.5Bi0.5Na0.5TiO3-0.12BaTiO3 14 39.9 2.1 67 0.012 39\n0.57BiFeO3-0.33BaTiO3-0.1Ag(Nb0.85Ta0.15)O3 42.5 43.4 7.04 80.5 -- 40\n(Bi0.5Ba0.1Sr0.1Ca0.2Na0.1)(Fe0.5Ti0.3Zr0.1Nb0.1)O3 66.4 43.2 13.3 78 2.23 41\n0.50BiFeO3-0.40BaTiO3-0.10CaHfO3 41 32.1 4.7 79 0.31 42\n0.52BiFeO3-0.28BaTiO3-0.2La(Mg1/2Zr1/2)O3 28 34.3 4.2 75.2 0.41 43\nBi0.83Sm0.17FeO3-0.95(0.85BaTiO3-0.15Bi(Mg0.5Zr0.5)O3)+2.1 wt%MnO2 23.9 38.3 4.1 88 -- 44\n0.59BiFeO3-0.33BaTiO3-0.08Sr(Mg1/6Zn1/6Ta1/3Nb1/3)O3 19 31.4 2.4 75 -- 45\n0.5(Bi0.7Dy0.2)FeO3-0.5BaTiO3 18.5 36.9 2.7 67 -- 46\n0.52BiFeO3-0.4SrTiO3-0.08NaNbO3 30 38.1 4.4 72 0.15 47\n((Bi0.85Nd0.1Sm0.05)0.5Ba0.5)(Fe0.5Ti0.5)O3 25 53.2 4.9 81 -- 48\n0.86(0.65BiFeO3-0.35BaTiO3)-0.14La(Mg2/3Ta1/3)O3 39 28.96 4.92 85 -- 49\n0.57BiFeO3-0.1Bi0.1Na0.7NbO3-0.33BaTiO3-0.15 wt%MnCO3 26 27.4 3.06 81.4 0.14 50\n0.57BiFeO3-0.4SrTiO3-0.03Bi(Zn0.5Zr0.5)O3 25.5 25.4 2.31 58 0.17 51\n0.59BiFeO3-0.33BaTiO3-0.08Bi2LaTiNbO9 25 29.3 3.26 80.8 0.9 52\n0.48BiFeO3-0.4SrTiO3-0.12La(Mg2/3Nb1/3)O3+0.1 wt%MnO2 45 32.6 6.3 74.3 4.5 53\n0.53BiFeO3-0.33BaTiO3-0.14Sr(Sc1/2Nb1/2)O3 47 51 8.5 87 5.0 8\n0.43BiFeO3-0.45SrTiO3-0.12BaTiO3+0.1 wt%MnO2 50.4 44.6 7.3 86.3 -- 54\n0.6BiFeO3-0.3BaTiO3-0.1(La0.5Na0.5)(Mg1/3Nb2/3)O3 50 42.4 8.3 84.6 -- 55\nComment-6: The authors report impressive energy-storage metrics for their BiFeO3-based ceramics; however, a very recent study (DOI: 10.1038/s41467-025-63584-y) has already demonstrated higher Wrec and η values, together with a more comprehensive performance evaluation (including temperature stability and fatigue cycling), which the present manuscript has not addressed. To strengthen the scientific novelty, the authors are encouraged to explicitly clarify how their materials-design strategy or mechanistic insight differs from and advances beyond current state-of-the-art approaches.\nResponse: Thanks for this good comment regarding novelty and state-of-the-art context.\nWe have carefully considered the recent benchmark study (DOI: 10.1038/s41467-025-63584-y) and agree that a clearer distinction of our work’s unique contribution is necessary. The key distinctions are as follows:\n1. Material system: While the cited excellent work utilizes Bi0.5Na0.5TiO3 (BNT)-based systems, our study focuses on BiFeO3 (BF)-based ceramics. This platform offers a fundamentally higher theoretical polarization ceiling (~100 µC cm-2) due to the strong hybridization between Bi 6s and O 2p orbitals. This provides a greater intrinsic \"design space\" for optimizing polarization characteristics, which is a distinct advantage of our chosen material family.\n2. Design strategy: The state-of-the-art strategy exemplified by Wang et al. effectively employs entropy-assisted polymorphic domain engineering to prevent premature polarization saturation while maintaining high Pmax, thereby achieving excellent energy storage performance. In contrast, our work demonstrates the local structure design to achieve a high Pmax under applied E and a low Pr after zero E, leading to a high ΔP of 56.4 μC cm-2, an ultrahigh Wrec of 14.5 J cm-3 and a high η of 88% under 65 kV mm-1. This result validates local structure engineering as a distinct and effective pathway to decouple Pmax and Pr, which directly translates to our high performance.\nSpecific revisions in the manuscript:\n1. The novelty is revised in Abstract part. Please see page 2: “However, the simultaneous enhancement in recoverable energy density (Wrec) and efficiency (η) still remains challenge owing to the restrictive relationship between the maximum polarization (Pmax), remanent polarization (Pr) and electric breakdown strength (Eb). To address this, we propose a strategy of local polar structure design in BiFeO3-based ceramics. By incorporating NaNbO3 to create embedded, persistent polar nanoregions within a weakly polar matrix, we achieve a giant polarization difference ΔP (56.4 μC cm-2), a low Pr (3.6 μC cm-2) and a large Eb (66 kV mm-1), endowing an ultrahigh Wrec of 14.5 J cm-3 and a high η of 88%. The local structure, directly visualized via the 2D/3D atomic displacement mapping, enables a high Pmax under an applied field and a low Pr at zero field, a mechanism explicitly validated by phase-field simulation.”\n2. This state-of-the-art strategy (Ref. 17 in revised manuscript) has been added in Introduction part to strengthen the progress in energy storage. Please see page 4: “High-entropy assisted Bi0.5Na0.5TiO3-based RFE ceramics exhibited ultra-small polymorphic domains, which weakened the polarization anisotropy and reduced the switching energy barriers. This polarization configuration finally resulted in a small Pr, delayed polarization saturation while maintaining high Pmax, thereby endowing superior ESP14,17.”\n3. The merit of BF-ST systems is revised in the manuscript. Please see page 4: “0.6BF-0.4ST ceramic with a morphotropic phase boundary demonstrates significant potential for energy-storage capacitors, due to its high intrinsic polarization ~40 μC cm-2 20,21.”\nWe believe this explicit clarification strengthens the scientific novelty and contribution of our manuscript.\nReviewer #3:\nThe manuscript titled “Ultrahigh energy-storage in lead-free ceramic capacitors via defect-dipole engineering and local structure design” is a work which focuses on creating sintered ceramics for use in capacitors as the dielectric medium, enhancing the capacitor performance using defect-dipole engineering. The paper builds on the existing (1-x)BiFeO3-xSrTiO3 system by adding MnO as a sintering aid in the matrix before sintering and NaNbO3 as the inhibitor of long range ferroelectric order. This serves to increase the recoverable energy density in the ceramic capacitor and increases the breakdown voltage and the overall efficiency of it. The work serves to showcase an advance in lead free ferroelectric capacitors with enhanced performance and in my opinion is worthy of publication given the comments below are addressed thoroughly:\nResponse: We greatly appreciate the reviewer for their time to review our manuscript and give us the enlightening suggestions, which are very helpful for improving this work. We have revised our manuscript accordingly, and the point-by-point responses to the reviewer’s suggestions are enclosed as follows.\nComment-1: The novelty of this approach is not completely clearly as written in the paper. Is this the first time this multi-component matrix has been used to achieve such an effect? For example, a recent paper (Nat Commun 16, 8580 (2025)) shows a similar enhancement in capacitor behaviour through an increased entropy effect. Is the local structure effect similar to the high entropy effect and if not, how would one differentiate the origin of such effects. This is a key point in future design of such systems.\nResponse: Thanks for raising this critical point regarding novelty and the differentiation of mechanisms. In the revised manuscript, we have clarified the unique contribution of our work by explicitly contrasting our \"local structure design\" with the \"high-entropy effect.\" The key distinction lies not in the final goal—which is indeed similar—but in the specific design principle and atomic-scale origin of the polar heterogeneity.\n1. Similarity: We agree that both our work and the recent high-entropy study (Nat Commun 16, 8580 (2025)) share a common objective: to optimize nanoscale polarization configurations to overcome the intrinsic trade-off between high Pmax and low Pr, thereby achieving superior energy storage performance.\n2. Differences: The fundamental distinction lies in the source and nature of the nanoscale disorder:\n• High-Entropy Effect: Achieves polarization control primarily through configurational entropy by incorporating a large number of cations on equivalent lattice sites. This creates a chemical landscape that statistically suppresses long-range order and promotes fine, polymorphic domains.\n• Local Structure Design (This Work): We employ a targeted compositional modification using a specific perovskite end-member NN to intentionally create a heterogeneous structure with distinct local environments. This is not a statistical cocktail but a designed system where the local polar clusters are embedded in a weakly polar matrix, as directly evidenced by our atomic-scale probes. To our knowledge, this is the first demonstration that such a designed local polar structure—featuring embedded, strong polar clusters—can be engineered in BiFeO3-based ceramics to directly govern energy storage performance. While entropy stabilization is a powerful top-down strategy, our approach offers a complementary, bottom-up design principle based on creating and leveraging specific local structural motifs. This provides a new and differentiated pathway for future material design, where targeted local chemistry, rather than maximized configurational entropy, is used to tailor polar landscapes.\nSpecific revisions in the manuscript:\n1. The novelty is revised in Abstract part. Please see page 2: “However, the simultaneous enhancement in recoverable energy density (Wrec) and efficiency (η) still remains challenge owing to the restrictive relationship between the maximum polarization (Pmax), remanent polarization (Pr) and electric breakdown strength (Eb). To address this, we propose a strategy of local polar structure design in BiFeO3-based ceramics. By incorporating NaNbO3 to create embedded, persistent polar nanoregions within a weakly polar matrix, we achieve a giant polarization difference ΔP (56.4 μC cm-2), a low Pr (3.6 μC cm-2) and a large Eb (66 kV mm-1), endowing an ultrahigh Wrec of 14.5 J cm-3 and a high η of 88%. The local structure, directly visualized via the 2D/3D atomic displacement mapping, enables a high Pmax under an applied field and a low Pr at zero field, a mechanism explicitly validated by phase-field simulation.”\n2. This state-of-the-art strategy (Ref. 17 in revised manuscript) has been added in Introduction part to strengthen the progress in energy storage. Please see page 4: “High-entropy assisted Bi0.5Na0.5TiO3-based RFE ceramics exhibited ultra-small polymorphic domains, which weakened the polarization anisotropy and reduced the switching energy barriers. This polarization configuration finally resulted in a small Pr, delayed polarization saturation while maintaining high Pmax, thereby endowing superior ESP14,17.”\n3. Origin of local structure has been revised in the manuscript. Please see page 4: “Herein, the perovskite end-member NN was incorporated to disrupt the long-range ferroelectric orders and promote the formation of PNRs while retaining highly polarized local clusters3,18, leading to a slim P-E loop with a relatively high Pmax and a significantly reduced Pr.” and page 10: “This result suggests that the engineered local structure is induced by the heterogeneous distribution of atoms with differing ionic radii, electronic configurations, and polarizabilities3,18, that collectively facilitate a high Pmax/ΔP, a low Pr, and superior ESP.”\nComment-2: Is the NaNbO3 spread uniformly across the sintered ceramic, or does it tend to form clusters? The theoretical analysis in Fig 1 and the images in Fig 2 suggest a uniform distribution, and so does Fig 3, but the discussion from line 180-205 is hard to parse for anyone who is not a PDF expert. This needs to be made more clear, especially what each figure subpanel represents and how finally fig 3f-I are obtained/calculated. This discussion is very dense with jargon and hard to follow in its current form.\nResponse: Thanks for the valuable comment.\n1. Composition homogeneity and local structure heterogeneity\nAt the micro-scale, EDS mapping confirms that the constituent elements of BF, ST and NN are uniformly distributed in the sintered ceramics (Fig. R6). At the atomic scale, however, local clustering of like atoms occurs, leading to the formation of distinct polar clusters with dimensions on the order of 1-4 nm (as illustrated in Fig. 2 and Fig. 3).\nFig. R6 (Fig. S11 in the revised Supplementary Materials) EDS mapping for x = 0.14.\n2. Reverse Monte Carlo refinement\nThe Bragg Rietveld refinement was done with GSAS-II software, using the neutron data from four banks (bank 2~5) to get the long-range average structure parameters. Reverse Monte Carlo (RMC) simulation was performed based on the RMCprofile software. The 20 × 20 × 20 supercells of pseudocubic (Pm-3m) structure (about 80 × 80 × 80 Å3) containing 40000 atoms and randomly distributed A-site atoms were established. In the RMC simulations, the Bragg intensities, G(r) and S(Q) data were fitted simultaneously under the bond valence sum and coordination constraints. A-site atoms were allowed to swap, and each RMC modelling was run to generate more than 103 move per atom to get converged. Detailed analysis of 3D polar displacements, unit-cell polar vectors, and polar projection can be referred to our previous works. To get the statistical analysis results, 20 refined 3D atom configurations were merged together.\n3. Calculation of polar displacement (D⃗) and unit-cell polar vector (P⃗)\nFrom the refined atom configurations, for each A-site atom (r⃗A), the coordinates of the neighboring oxygen atoms (r⃗Oi) can be extracted. The geometric center of the AO12 polyhedra can be calculated as 1/12 ∑ r⃗Oi. Similarly, the surrounding six oxygen atoms of each B-site atom (r⃗B) can be found. The A-site atom polar displacement vector D⃗A and the B-site atom polar displacement vector D⃗B can be calculated by the equation (1) and equation (2), respectively. Notably, the A-site atom without 12 physically reasonable oxygen bonds, or these B-site atom without 6 physically reasonable oxygen bonds were discarded.\nD⃗A = r⃗A - 1/12 ∑ r⃗Oi (1)\nD⃗B = r⃗B - 1/6 ∑ r⃗Oi (2)\nFor unit-cell polar vector calculation, one perovskite unit cell contains 8 A-site atoms, 1 B-site atom, and 6 O atoms. Firstly, the coordinates of 6 neighboring oxygen atoms r⃗Oi, and the coordinates of 8 neighboring A-site atoms of each B-site atom r⃗Ai can be found. In the perovskite unit-cell, the center of the oxygen polyhedral is 1/6 ∑ r⃗Oi, and the center of the A-site atom is 1/8 ∑ r⃗Ai. Subsequently, the unit-cell polar vector (P⃗) can be calculated by equation (3):\nP⃗ = [qA × (1/8 ∑ r⃗Ai - 1/6 ∑ r⃗Oi) + qB × (r⃗Bi - 1/6 ∑ r⃗Oi)] / V (3)\nWhere the q indicates the ion charge, V is the perovskite unit-cell volume.\nSpecific revisions in the manuscript:\n1. The EDS mapping have been provided in Fig. S11, Supplementary Materials.\n2. The methods for RMC refinement and polar vector calculation have been provided in revised Supplementary Materials, please see pages S4-S5.\nComment-3: Can the full composition of the BF-ST be discussed somewhere in the main text, and also of the MnO. Is the BF-ST the 0.6BF-0.4ST? How much MnO is added for incorporation, and why is it now listed as something that could sit in the B site?\nResponse: Thanks for the valuable comment. The full composition for this work is (0.6-x)BiFeO3-0.4SrTiO3-xNaNbO3+0.1wt% MnO (denoted as BF-ST-xNN-Mn). The BF-ST is 0.6BF-0.4ST, and the content of MnO is 0.1wt%.\nIn the revised manuscript, the role of MnO is updated. The additive MnO, acting as sintering aid, can maintain the charge balance and suppress the oxygen vacancies.\nSpecific revisions in the manuscript:\n1. The full composition has been provided in revised manuscript. Please see page 4: “…we propose the local structure design to overcome the restraint between different parameters in (0.6-x)BiFeO3-0.4SrTiO3-xNaNbO3+0.1wt% MnO (denote as BF-ST-xNN-Mn) ceramics…”, and “0.6BF-0.4ST ceramic with a morphotropic phase boundary demonstrates significant potential for energy-storage capacitors, due to its high intrinsic polarization ~40 μC cm-2 20,21.”\n2. The role of MnO is revised in the main text. Please see page 5: “Additionally, the synthesis of phase-pure BF-based ceramics is challenging due to volatilization of Bi2O3 and the reduction of Fe3+ (Fe3+ + e- → Fe2+) during sintering, which typically leads to high dielectric loss and leakage. In this study, MnO was added as a sintering aid to mitigate these issues. By helping to maintain the charge balance, MnO suppresses the formation of oxygen vacancies and secondary phase, thereby reducing leakage current and enhancing Eb4,8.”\nComment-4: The overall presentation of the paper will benefit with major improvements in clarity and a coherent presentation of the work and comparisons, especially in the discussions. All the separate sections make sense on their own, but it is very hard to read the paper from intro to conclusion and put a thread through all the data and results presented. Does Figure 1 serve as the motivation as a whole, because it is not addressed anywhere else again? Could the authors please connect all the different results better as to how they all come together? The data looks great, but it is difficult to parse all of it together and seems very disjointed.\nResponse: Thanks for the great suggestion.\nWe have thoroughly revised the presentation and discussion of the manuscript. The revised narrative is structured as follows:\nFirst, we present the energy storage performance of the BF-ST-xNN-Mn ceramics, highlighting that the composition with x = 0.14 achieves a superior comprehensive performance among BF-based ceramic capacitors, with a high recoverable energy density (Wrec = 14.5 J cm-3) and efficiency (η = 88%) at 65 kV mm-1.\nSubsequently, to uncover the physical origin of this excellent performance, we employ advanced local probes—atomic-resolution STEM and neutron total scattering—to characterize the unique local structure of the x = 0.14 composition. The positive role of this local structure in governing polarization behavior is further validated by phase-field simulations, which align with the experimental observations.\nFinally, given that a high breakdown strength (Eb) is indispensable for achieving such performance, we also investigate the mechanism responsible for the enhanced Eb in the x = 0.14 ceramic.\nComment-5: Related to point 4, there are many typos and also sentences which seem like incomplete clauses. I list some of them here:\nResponse: Thanks for the careful reading. We have revised the typos and sentences in the manuscript.\nSentence starting on line 127, it's not clear what defect content has to do with nanodomains.\nResponse: Sorry for the confusion. We have removed all content regarding the defect-dipoles in the revised manuscript.\nSentence on line 134, what are after mentioned displacement vectors.\nResponse: Thanks for the careful reading. This unclear sentence has been removed in the revised manuscript.\nSentence on line 150, the logic of the suggestion of existence of PNRs and nanodomains is unclear. Discuss further please. All that can be seen is a lack of macroscopic ferroelectric order.\nResponse: Thanks for the valuable comment. Fig. 2(a) exhibits the TEM micrography of x = 0.14, which shows no evidence of long-range ferroelectric order at macroscopic scale. Yet, x = 0.14 shows slim P-E loop typical of relaxor behavior (Fig. S1(f)). This contrast suggests the presence of polar nanoregions (PNRs) or nanodomains at a finer scale.\nSpecific revisions in the manuscript:\nThe discussion on the existence of PNRs has been revised in the main text. Please see page 8: “Conventional ferroelectric materials typically exhibit lamellar domains visible in dark-field TEM imaging8,23. However, such long-range domains are absent in the x = 0.14 sample (Fig. 2a). Instead, x = 0.14 displays a slim P-E loop (Fig. S2), characteristic of relaxor behavior. This combination suggests the presence of PNRs or fine-scale nanodomains.”\nSentence on line 159 starting with “as knows...” seems like a throwaway statement\nResponse: Thanks for the good comment. These words have been removed in the revised manuscript.\nThe sentence on line 203-204 claims a polarization value but does not account for any oxygen vacancy effects which might introduce electronic perturbations to the total polarization, so is it correct to state such a value based on just the cationic displacements?\nResponse: Thanks for the good comment. In our initial model, the 3D polar structure was constructed based solely on cationic displacements, without incorporating the effect of defect dipoles. As we have now removed the defect-dipole narrative from the revised manuscript, this does not affect the current discussion or conclusions\nThe sentence starting on 224 is too long and mentions weak coupling which has only been mentioned before on line 172 but it is not clear how that conclusion is drawn, about the weak coupling? Is the weak coupling between different polar clusters or between polar clusters and defect-dipoles?\nResponse: Thanks for the good comment. The conclusion on the weak coupling is removed in the revised manuscript, and the long sentence is also revised. Please see page 6: “A high Wrec is dominated by the collective effect of high ΔP and large Eb. However, these two parameters are typically in direct competition, fundamentally limiting the enhancement of Wrec in dielectric ceramics4,6. Remarkably, the x = 0.14 composition mitigates this trade-off, achieving a giant ΔP of 56.4 μC cm-2 simultaneously with a high Eb, resulting in a record-high Wrec of 14.5 J cm-3 among BF-based ceramics (Fig. 1c).”.\nThese are only to name a few. Please correct such sentences and if a logical conclusion is drawn, please explain how such a conclusion was reached.\nResponse: Thanks for the good comment. We have revised the whole manuscript carefully.\nComment-6: Please also redefine the acronyms and the variables not just in the abstract but also in the main text. This would improve the readability of the paper a lot.\nResponse: Thanks for the good comment. The acronyms and the variables are defined in both the abstract and main text."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors have satisfactorily addressed my major concerns from the previous round. In particular, the mechanistic narrative is now consistent and supported by local structure characterization. I therefore support publication in Nature Communications.", "reviewer_2": "The authors have carefully revised the manuscript and provided detailed responses to the reviewers’ comments. The additional experiments, particularly the fatigue cycling and temperature-dependent energy-storage measurements, significantly improve the completeness of the study. Overall, the revised manuscript is substantially improved. However, several issues still require clarification before the manuscript can be considered for publication.\n1. The mechanistic interpretation remains largely qualitative. While the authors have removed the previously proposed defect-dipole explanation and instead attributed the enhanced energy-storage performance to compositional disorder and local structural heterogeneity, the current discussion still relies primarily on phenomenological correlations (e.g., grain size evolution and polarization behavior) rather than direct experimental evidence. The mechanistic claims should therefore be moderated and presented as plausible interpretations rather than definitive conclusions.\n2. Compared with previously reported studies (eg. DOI: 10.1038/s41467-025-63584-y), the improvement in energy-storage performance appears to be relatively modest. In addition, the novelty of the present work relative to recent state-of-the-art reports remains somewhat unclear. The authors are therefore encouraged to clarify the specific novelty and the scientific advance of the present work.", "reviewer_3": "I thank the authors for revising the manuscript and recommend publication."}, "author_response": "Response Letter\nWe sincerely thank all reviewers for spending their time reviewing our manuscript, NCOMMS-25-88618A, entitled \"Ultrahigh energy-storage in lead-free ceramic capacitors via local structure design\", and for providing insightful comments. In the following, we present a point-by-point response to the reviewers’ comments. We hope that these changes adequately address the concerns raised. All changes have been highlighted within the manuscript.\nReviewer #1:\nThe authors have satisfactorily addressed my major concerns from the previous round. In particular, the mechanistic narrative is now consistent and supported by local structure characterization. I therefore support publication in Nature Communications.\nResponse: We sincerely thank the reviewers for their time and effort in reviewing our work, as well as the recognition of our work.\nReviewer #2:\nThe authors have carefully revised the manuscript and provided detailed responses to the reviewers’ comments. The additional experiments, particularly the fatigue cycling and temperature-dependent energy-storage measurements, significantly improve the completeness of the study. Overall, the revised manuscript is substantially improved. However, several issues still require clarification before the manuscript can be considered for publication.\nResponse: We greatly appreciate the reviewers for their time to review our manuscript and give us the enlightening suggestions, which are very helpful for improving this work. We have revised our manuscript accordingly, and the point-by-point responses to the reviewer’s suggestions are enclosed as follows.\nResponse:Thanks for your enlightening comment. More discussion and interpretations on the evolution of grain size and polarization behavior have been added in the revised manuscript.\nSpecific revisions in the manuscript:\n More discussion on the grain size evolution. Please see page S12, Supplementary Materials: “Fig. S10 | Typical microstructure morphologies of a x = 0, b x = 0.05, c x = 0.08, d x = 0.10, e x = 0.12, f x = 0.14, and g x = 0.16. Please note that the scale ruler of x = 0 is different with others. All samples exhibit dense, pore-free microstructure. With increasing NN content, the average grain size decreases significantly from 1.6 μm (x = 0) to 0.4 μm (x = 0.16). On the one hand, the oxygen vacancies facilitate the mobility of grain boundaries. However, the addition of MnO retains the charge balance of BF-BT-NN ceramics and suppress the oxygen vacancies, thereby hampering the grain boundary mobility and delaying the grain growth11. On the other hand, the incorporation of Na+ and Nb5+ with different ionic radius increases the lattice strain energy, also disturbing the grain boundary mobility and delaying the grain growth12. The grain refinement suggests higher density of grain boundaries, which act as space-charge depletion layers that impede charge carrier transport, resulting in a dramatically improved Eb from 12.8 kV mm-1 (x = 0) to 66.0 kV mm-1 (x = 0.14).”\n More discussion on the general polarization behaviour related to P-E loops has been added in the main text. Please see page 4-5: “With increasing NN content, the P-E loops become gradually slimmer, accompanied by decreased Pmax and Pr from 39.7 and 30.4 μC cm-2 (x = 0) to 8.5 and 0.2 μC cm-2 (x = 0.16), owing to the enhanced relaxor behavior (Fig. S3). Accordingly, the η progressively increases from 13.2% at x = 0 to 91.3% at x = 0.16 (Fig. S2b)” ... “A high Pmax exceeding 40 μC cm-2 under applied E is maintained across compositions x = 0.05-0.16, due to the strong hybridization between Bi-O orbitals. Particularly, the composition x = 0.14 achieves a giant Pmax of 60 μC cm-2, a low Pr of 3.6 μC cm-2, and the highest ΔP of 56.4 μC cm-2 under 65 kV mm-1, thereby delivering an optimal ESP with Wrec = 14.5 J cm-3 and η = 88% (Fig. S2d).”, and “Fig. 1a displays the E-dependent P-E loops for x = 0.14. The P-E loop always retains a slender shape across all fields, with both Pmax and Pr increasing steadily as the field rises (Fig. S4). This is attributed to the E-induced long-rang ferroelectric orders, which rapidly return the initial disorganized state after removing E. Ultimately, the Wrec increases monotonically from 0.6 J cm-3 at 10 kV mm-1 to 14.5 J cm-3 at 65 kV mm-1, while the η remains largely constant (Fig. 1b).”\n Additionally, the further discussion on polarization behavior is presented in the characterization of local structures, because the polarization is highly related to the structure and specific elements. Please see page 8: “The atomic displacement vectors reveal regions with polar orientations corresponding to rhombohedral (R), orthorhombic (O), tetragonal (T), and cubic (C)-like symmetries (Fig. 2g). This local symmetry frustration and the coexistence of polymorphic polar configurations are characteristic of relaxor systems with compositional disorder and are consistent with the formation of PNRs. A wide distribution of polar direction suggests that these polar clusters are connected by the monoclinic (M) symmetry low symmetry in the transitional regions (Fig. 2h), which facilitates reducing the polarization anisotropy and polarization switching barrier, thereby resulting in a decreased polarization hysteresis, a low Pr and an enhanced η. Moreover, 1-4 nm sized polar clusters with polar magnitude up to 18 pm are embedded into the matrix (Fig. 2i), demonstrating the strong local polar clusters, which does favor to retain a high Pmax/ΔP under applied E, and thus enhance Wrec.”, and page 10: “Accordingly, a wide polar direction distribution, as well as several 1-4 nm polar clusters with a high polarization as high as 50 μC cm-3, is observed, consistent with STEM results. This result suggests that the engineered local structure is induced by the heterogeneous distribution of atoms with differing ionic radii, electronic configurations, and polarizabilities3,18, that collectively facilitate a high Pmax/ΔP, a low Pr, and superior ESP.”\nComment-2: Compared with previously reported studies (eg. DOI: 10.1038/s41467-025-63584-y), the improvement in energy-storage performance appears to be relatively modest. In addition, the novelty of the present work relative to recent state-of-the-art reports remains somewhat unclear. The authors are therefore encouraged to clarify the specific novelty and the scientific advance of the present work.\nResponse:Thanks for this critical point regarding the novelty and scientific advance of our work.\n Novelty\nWe employ a targeted compositional modification using a specific perovskite end-member NN to intentionally create a heterogeneous structure with distinct local environments. This is not a statistical cocktail but a designed system where the local polar clusters are embedded in a weakly polar matrix, as directly evidenced by our atomic-scale probes (Figs. 2-3). This unique polar configuration facilitates achieving a low Pr under zero E and a high Pmax/ΔP under applied E. The phase-field simulation illustrates how relaxors with pronounced local polar structure can circumvent the conventional trade-off between Pmax and Pr under an applied field (Fig. 4), which concurrently enhancing the Wrec and η.\n Scientific Advance\nBiFeO3-based ceramics generally exhibit a strong ferroelectricity and a high leakage, leading to a large Pr, a low Eb and a poor energy storage performance (ESP). Our work demonstrates the local structure design to achieve a high Pmax under applied E and a low Pr after zero E, leading to a high ΔP of 56.4 μC cm-2, an ultrahigh Wrec of 14.5 J cm-3 and a high η of 88% under 65 kV mm-1. Although the Wrec value is a litter lower than previously report (DOI: 10.1038/s41467-025-63584-y), the comprehensive ESP, achieved in our work, outperforms the currently reported BiFeO3-based ceramic capacitors (Figs. 1c, 1d and S6), suggesting that the primary obstacle for the enhancement in ESP of BiFeO3-based systems have been addressed. Moreover, this result also provides insights to overcome the paradox between multiple parameters and paves a feasible route for the exploitation of high ESP in other material systems.\nReviewer #3:\nI thank the authors for revising the manuscript and recommend publication.\nResponse: We sincerely thank the reviewers for their time and effort in reviewing our work, as well as the recognition of our work.\n5"}]

Decision letter