Synergistic proton conduction via Ca-vacancy coupled with Li+-bridge in Ca5(PO4)3OH

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https://doi.org/10.1038/s43246-024-00719-6
# Synergistic proton conduction via Ca-vacancy coupled with Li + -bridge in Ca5(PO4)3OH

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Nabeela Akbar 1,6, Qing Pang 2,6, Yaokai Lu 1,6, Yifu Jing 3, Manish Singh 4, Jun Wang1, Bin Zhu 1,
Faze Wang 1 & Sining Yun 5

Proton conductivity plays a crucial role in the advancement of materials for proton ceramic fuel cells

(PCFCs) and a variety of electrochemical devices. Traditional approaches to enhancing proton

conductivity in perovskites have largely relied on doping strategies to induce structural oxygen

vacancies. However, these methods have yet to overcome the challenges associated with achieving
desired proton conductivity. Here, we introduce an approach wherein intermediate Li + ions act as a

bridge linked to Ca vacancies, fostering a mechanism for accelerated proton transport. Utilizing
protonated Ca5(PO4)3OH-H(Li) as an electrolyte, we achieve a proton conductivity of 0.1 S cm −1 and a
fuel cell performance of 661 mW cm −2 at an operational temperature of 550 °C for realizing low

temperature PCFCs. This proton transport synergy overcomes traditional doping limitations, enabling
the advancement of proton-conducting electrolytes and enhancing the efficiency of proton

conducting electrolyte fuel cells, with implications in energy conversion and storage technologies.

PCEFCs emerge as trailblazers in the dynamic landscape of electrochemical
energy conversion, navigating the complexities of both high and lowtemperature domains (HT and LT) 1–7 . Discerning the nuances between HT
and LT proton fuel cells proves pivotal in comprehending the tailored
characteristics of these fuel cells designed for diverse operational ranges 8–11 .
The proton ceramic fuel cells (PCFCs) are robust challenger because of their
operational temperature is typically above 300 °C 12–16 .

PCFCs do not utilize noble catalyst metals in electrodes and electrochemical reactions; instead, they heavily rely on proton conductivity in
ceramic electrolyte materials. Notably, commonly employed PCFC electrolytes,groundedinyttrium-dopedBaZrO3 andBaCeO3 perovskiteoxides,
exhibit a constrained proton conductivity ranging from 10 −2 to 10 −3 S cm −1

at approximately 600 °C 17–19 . Therefore, there is a growing demand for
developing materials with high proton conductivity, around 0.1 S cm −1 at
temperatures between 300 and 500 °C 20,21 . This temperature range is ideal
for combining the advantages of high-temperature and low-temperature
fuel cells while minimizing their drawbacks, thus facilitating in the commercialization of fuel cells 22 . Conversely, LT-PCEFCs, such as PEMFCs that
operate at lower temperatures, require noble catalyst electrodes like Pt and

depend on specific high humidity levels to maintain sufficient proton
conductivity for optimal power output 23,24 .

Both HT and LT fuel cells present distinctive advantages and disadvantages but LT PCFCs hold even more attractive potential in certain
aspects especially in low cost for future commercialization. The global
landscape of research and development evinces heightened interest in both
types of fuel cells. Notably, PEMFCs have already achieved notable commercialization milestones, finding application especially in the automotive
and some interest in power generation sectors 24,25 .

Notably, Qian et al. recently spotlighted the potential of proton conducting materials, demonstrating exceptional proton conductivity through
Cd vacancies in a CdPS3-based membrane 26 . However, their achievement
primarily residesin the LT domain, evident in the material’s dependence on
high relative humidity levels and operational temperatures not surpassing 100 °C.

In clarifying this distinction, this work propels the mechanism of metal
vacancy-induced proton transport into PCFCs, elucidating the synergistic
interplay between Ca vacancies and Li + bridging ions, exemplified in
Ca5(PO4)3OH. This breakthrough not only bridges the gap betweenLT and

1School of Energy and Environment, Southeast University, No.2 Si Pai Lou, Nanjing, 210096, China. 2College of Science, Xi’an University of Architecture and
Technology, Xi’an, 710055 Shaanxi, PR China. 3 Department of Material Science, Shenzhen MSU-BIT University, Shenzhen, 518000, China. 4 Department of
Metallurgical and Materials Engineering, Indian Institute of Technology Patna, Bihta, Bihar, 801106, India. 5 School of Materials Science and Engineering, Xi’an
University of Architecture and Technology, Xi’an, Shaanxi, 710055, China. 6 These authors contributed equally: Nabeela Akbar, Qing Pang, Yaokai Lu.

[e-mail: zhu-bin@seu.edu.cn; fazewang@seu.edu.cn; yunsining@xauat.edu.cn](mailto:zhu-bin@seu.edu.cn)

HT PCEFCs but also underscores the pivotal role of PCFCs in the broader
context of energy conversion technologies.

Ca5(PO4)3OH, characterized by the hydroxyapatite structure, emerges
as a focal point of our investigation 27–29 . While extensively studied for its
crystal structure and properties, our research unveils an intriguing phenomenon of significant enhancement in proton conductivity facilitated by
calcium vacancies within Ca5(PO4)3OH.

In addition, Furuseki, Horiuchi, and Bystrov’s previous study provides
vital insights into the proton-conducting properties of hydroxyapatite
(HAp) composites, displaying the variety and possible applications of this
material 30–32 . Our study uniquely focuses on characterizing calcium vacancies and protonation within Ca5(PO4)3OH-H(Li), incorporating insights
derived from nuclear magnetic resonance (NMR) spectroscopy.

The apex of our investigation reveals an exceptional proton conductivity approaching 0.1 S cm −1 at 500 °C for Ca5(PO4)3OH-H(Li), presenting a compelling avenue for augmenting fuel cell performance. In the
subsequent sections, we delve into experimental details, results, and the farreaching implications of our findings, offering a comprehensive exploration
of the remarkable properties of Ca5(PO4)3OH and its potential transformative impact on future energy technologies. This discovery represents
advancement towards developing efficient and cost-effective proton-conductive materials. It also introduces new possibilities in energy conversion
and storage applications, highlighting the importance of high-temperature
PCFCs in sustainable energy solutions.

Results and discussions
Structure and morphology
Figure 1A presents the XRD patterns of the four samples which well match
with the XRD diffraction pattern of Ca5(PO4)3OH (JCPDS No.73-0293)
with a space group structure of P63/m. The diffraction angle of about 29° in
K- Ca5(PO4)3OH and Li- Ca5(PO4)3OH belongs to CaCO3 (JCPDS No. 860147), and this diffraction peak disappeared in acetic acid treated
H-Ca5(PO4)3OH, which furtherindicatesthat the peak belongs to CaCO333.

Since Ca 2+ is partially replaced by other ions K +, Li + and H + which have
different ion radii, the unit cell dimensions of the material will be affected,
resulting in different degrees of deviation of the diffraction peak in the XRD
patterns 34 .

In order to further analyze changes in crystal structure the XRD patterns were refined using GSAS II software. The refinement results are presented in Fig. 1B. As shown in Supplementary Fig. S1, due to H + (8 × The
ionicradiusof10 −6 Å)andLi + (0.69 Å)issmallerthanthatofCa 2+ (1 Å),and

thelatticeparametersofLi-Ca5(PO4)3OHandH-Ca5(PO4)3OHaresmaller
compared to Ca5(PO4)3OH, indicating that Li + and H + enter the interior of
the lattice. Although the ionic radius of K + (1.38 Å) is larger than that of
Ca 2+, the decrease in crystal cell parameters has also been mentioned in
other literature. This abnormal phenomenon is mainly due to the charge
imbalancecaused bythe substitutionof K + forCa 2+,whichcreatesvacancies
and leads to abnormal changes in crystal cell parameters 35 .

Supplementary Fig. S2 presents Scanning Electron Microscopy (SEM)
measurements which show that both Ca5(PO4)3OH and Ca5(PO4)3OHH(Li)areplatestructure,andthemorphologyremainsconsistentbeforeand
after ions substitution. However, due to the acid treatment the particle size
has reduced from 120 nm (Ca5(PO4)3OH-H(Li)) to 70 nm (Ca5(PO4)3OH)
as acid treatment can introduce functional groups or alter the surface
roughness of a material, which may result in a decrease in apparent size.

Supplementary Fig. S3 shows the high-resolution transmission electron microscopy (HR-TEM) image of Ca5(PO4)3OH-H(Li). The particle
displays well-defined crystalline fringes with a lattice spacing of 0.27 nm
correspond to the (112) plane of Ca5(PO4)3OH, indicating that the crystal
structure was not affected by the protonation. Through EDS observation of
element distribution, it can be seen that Ca, P and O are uniformly distributed in H-Ca5(PO4)3OH, and the content of K is lower compared to
other elements. This is mainly ascribed to multiple washing and ion
exchange processes during the preparation of H- Ca5(PO4)3OH. Due to the
limitations of quantitative EDS analysis, the distribution of H and Li light
elements cannot be observed.

Characterization of calcium vacancies and protonation in Ca5
(PO4)3OH-H and Ca5(PO4)3OH-Li samples
As above mentioned, because of the EDS quantitative testing limitation, the
light elements cannot be routinely analyzed. Hence, ICP-AES was performed to determine the concentration of various elements in the
Ca5(PO4)3OH-H and Ca5(PO4)3OH-Li samples, and the results are listed in
Supplementary Table S1. According to the relative atomic mass of each
element, the corresponding molar concentration of each element can be
calculated. The chemical formula of calcium hydroxy phosphate is Ca5
(PO4)3OH, and the P content is constant during protonation. Therefore, we
assume that the molar ratio of P in Ca5(PO4)3OH-H(Li) is 3.1159312842.
Based on the calculated molar concentration, it can be determined that the
molar ratio between the various elements is approximately K: Li: Ca:
P = 0.0075:0.122:5.19:3.11. It can be seen that there is an order of magnitude
difference between the content of K and other elements, which can be

Fig. 1 | X-ray Diffraction analysis. (A) XRD patterns of Ca5(PO4)3OH, Ca5(PO4)3OH-K, Ca5(PO4)3OH-Li and Ca5(PO4)3OH-H(Li). (B) Refined XRD patterns of
Ca5(PO4)3OH, Ca5(PO4)3OH-K, Ca5(PO4)3OH-Li and Ca5(PO4)3OH-H(Li).

Fig. 2 | Structual characterizations. A FTIR spectra of Ca5 (PO4)3OH-H and Ca5(PO4)3OH (B) (a) 31 P NMR spectra of H-Ca5 (C) XPS Spectra of Ca5(PO4)3OH and
H-Ca5(PO4)3OH (D) HR-TEM image of H- Ca5(PO4)3OH powder.

ignored, and this is also consistent with EDS results. Based on the charge
balance and valence states of other elements, the molar ratio of H in
Ca5(PO4)3OH-H(Li) can be calculated, thereby determined the chemical
formula of Ca5(PO4)3OH-H as Ca4.89Li0.03H0.19 (PO4) 3OH. Ca-vacancy
promotes proton transport as it creates a defect site with a low activation
energy for proton hopping. In oxide materials the Ca-vacancy can act as a
proton acceptor, which can facilitate the dissociation of protonic species like
H2O into H + and OH- ions. The generated H + ions then migrate through
the Ca-vacancy site via proton hopping, which has a lower activation energy
compared to other pathways. Therefore, the presence of Ca-vacancy defects
can significantly enhance the proton conductivity of the materials 36 . In
addition, the presence of Ca vacancies influences the hydration dynamics
within the material, creating favorable hydration shells around the protonic
species. These hydration effects promote the mobility of protons and contributeto theirfacilitated transportthroughthelattice,ultimatelyenhancing
the overall proton conductivity of the material. The Ca vacancies are not
merelyemptyspacesbutactasactivesitesthatsignificantlyinfluenceproton
dynamics within the material.

Figure 2A shows the FTIR results of Ca5 (PO4) 3OH-H. The functional
groups contained in Ca5 (PO4)3OH-H can be clearly identified by infrared
spectroscopy. The absorption peak at about 3408 cm −1 is attributed to the
stretching vibration of O-H, and 1026 cm −1 belongs to the antisymmetric
stretching vibration of PO43-. The absorption peaks of 602 cm−1 and

stretching vibration of PO43-. The absorption peaks of 602 cm−1 and

563 cm −1 are due to the asymmetric angular vibration of PO43−. The
absorption peak located around 877 cm −1 belongs to the bending vibration
of CO32 36,37. The content of CaCO3 in Ca5 (PO4)3OH is higher than that in

of CO32 36,37. The content of CaCO3 in Ca5 (PO4)3OH is higher than that in

Ca5(PO4)3OH-H, which is consistent with the change in peak intensity in
FTIR. The absorption peak located at 1456 cm −1 belongs to a variable angle
vibration of H2O. As the preparation process does not undergo high

temperature sintering, this may be due to residual H2O in the material
during the preparation process. More interestingly, the absorption peak at
2361 cm −1 could be attributed to the tensile vibration of P-H, while the
absorption peak of Ca5(PO4)3OH-H is much stronger than that of
Ca5(PO4)3OH, which is the result of the presence of Ca vacancies and
proton insertion.

The 31 P and 1H NMR spectra of Li-Ca5(PO4)3OH and
H-Ca5(PO4)3OH are shown in Fig. 2B. Whena calcium vacancy is occupied
by protons, the extranuclear environment around the atom changes, which
can be detected by nuclear magnetic resonancespectroscopy (NMR). When
the chemical environment of the atomic nucleus changes, the conditions
that can trigger resonance also change 38 . The same functional group has the
same ppmvalue,butdifferent chemicalenvironmentscanproducedifferent
chemical shifts 39 . The phosphate group has a tetrahedral structure.
According to the number of oxygen atoms bridged by each tetrahedron, the
phosphate bond is divided into four categories, represented by Qn
(n = 0,1,2,3). In the NMR spectra of 31 P of H- Ca5(PO4)3OH and LiCa5(PO4)3OH, there is only one single chemical shift peak, 3.353 ppm and
3.323 ppm, respectively. This indicates that P exists in the form of Q0, i.e.,
free PO43-, in H-Ca5(PO4)3OH and Li-Ca5(PO4)3OH. The difference in

chemical shifts between the two may be due to changes in the chemical
environmentaroundPcausedbyprotonsenteringtheCa5(PO4)3OHlattice,
and due to the relatively low content of embedded H, the difference is
relatively small, which is also consistent with the ICP-AES results. There are
significant differences in the 1 H spectra of H- Ca5(PO4)3OH and LiCa5(PO4)3OH. The chemical shifts at 5.793 ppm and 5.122 ppm, respectively, belong to the adsorbed water on the material surface, while the
chemical shifts at 0.302 ppm and 0.216 ppm reflect the O-H within the
crystal structure. It can be seen that the O-H intensity in H-Ca5(PO4)3OH is

Fig. 3 | Electrical and electrochemical performances. (A) EIS spectra obtained for
Ca5(PO4)3OH-H (Li) at various temperatures and (B) Arrhenius curves and calculated activation energies for Ca5 (PO4) 3OH and Ca5 (PO4) 3OH-H (C) Fuel cell

much higher than that in Li-Ca5(PO4)3OH, which indicates that after
protonation treatment, more H + enters the Ca5(PO4)3OH crystal, which is
also consistent with the FTIR results.

According to the X-ray photoelectron spectroscopy (XPS) spectra
shown in Fig. 2C, the valence states of Ca, P, and O ( + 2, +5, and –2,
respectively) remain same in both samples Ca5(PO4)3OH and
Ca5(PO4)3OH-H(Li), suggesting that all the samples keep charge balance
though protons are incorporated into the structure, the charge is balanced
by forming calcium vacancies. However, the intensity of the peaks of Ca 2p,
P 2p and O1s have been changed after protonation. Which can be attributed
tothealterationinthechemicalenvironmentoftheatomsafterprotonation,
leading to shifts in the binding energies and changes in the relative intensities of the spectral peaks. These changes provide valuable insights into the
chemical bonding and electronic structure variations induced by protonation. The presence of calcium vacancies (Ca-vacancies) induces alterations
in the local environment and the surrounding crystal structure, resulting in
the creation of distinctive proton-conductive pathways. These modified
pathways play a crucial role in facilitating the movement of protons through
the material. By reducing energy barriers, these pathways contribute to an
enhanced efficiency of proton transport within the material 40,41 . Further
analysis though HRTEM as shown in Fig. 2D we directly find that
H-Ca5(PO4)3OH contains a large number of tiny pores with diameters
ranging from 3 to 5 nm, which may be caused by the partial dissolution of
Ca 2+ during the preparation process, that directly proves the existence of Ca
vacancies in H- Ca5(PO4)3OH.

EIS measurements were conducted on Ca5 (PO4)3OH and Ca5
(PO4)3OH-Hunderopencircuitconditionsinafuelcell,andtheimpedance
spectra obtained are shown in Fig. 3A. The EIS data was fitted based on the
fitting circuit R0 (R1CPE1) (R2CPE2) for Ca5(PO4)3OH. In the equivalent
circuit, Ro represents the ohmic resistance which can be obtained by

performance comparison of Ca5 (PO4) 3OH-H and Ca5 (PO4) 3OH at 550 °C, and
(D) I-V and I-P curves of Ca5 (PO4) 3OH-H at different temperatures.

calculating the intercept on the real axis at high frequencies, R1 and R2
represent the grain-boundary resistances of the electrolyte and the polarization resistance of the electrode, respectively, whereas CPE is the constant
phaseelement.The fittingdataarelistedinSupplementaryTablesS2andS3.
As the temperature decreases for Ca5 (PO4)3OH, the bulk resistance of
Ca5(PO4)3OH increases from 0.606 Ω cm 2 to 2.134 Ω cm 2, and at different
temperatures, the bulk resistance is higher than the grain boundary resistance, indicating that ions in Ca5(PO4)3OH tend to transport through the
grain boundary. The conductivity of Ca5 (PO4)3OH at 550 °C is calculated
from the EIS to be 0.056 S cm −1 using formula σ ¼ R AL , and continues to
decrease as the temperature decreases, reaching 0.012 S cm −1 at 430 °C. The
EIS pattern of Ca5(PO4)3OH-H with temperature is consistent with that of
Ca5(PO4)3OH.

ComparedtoCa5(PO4)3OH,thebulkresistanceofCa5(PO4) 3OH-Hat
550 °C decreased from 0.606 Ω cm 2 to 0.277 Ω cm 2, the grain boundary
impedance decreased from 0.104 Ω cm 2 to 0.003 Ω cm 2, and the conductivity increased from 0.056 S cm −1 to 0.143 S cm −1, an increase of nearly
2.5times.Asthetemperature decreases,the impedance of Ca5 (PO4)3OH-H
gradually increases, with R0 and R1 increasing to 0.386 and 0.116 Ω cm 2 at
430 °C, respectively. The corresponding conductivity is 0.080 S cm −1, which
islowerthan0.1 Scm −1,butstillhigherthanthatofCa5(PO4)3OHat550 °C.
This indicates that the conductivity of Ca5(PO4)3OH can be effectively
improved by constructing a Ca vacancy. According to the above formula,
the activation energy corresponding to Ca5 (PO4)3OH and Ca5(PO4)3OHH can be obtained from the slope of the Arrhenius curve 42 . The activation
energies of Ca5(PO4)3OH and Ca5(PO4)3OH-H are 0.68 ± 0.064 eV and
0.274 ± 0.038 eV, respectively as shown in Fig. 3B.
Figure 3C shows the I-V (current voltage) and I-P (current power)
curves of a fuel cell with Ca5(PO4)3OH and Ca5(PO4)3OH-H as electrolytes
at 550 °C. It can be seen that at 550 °C, the OCV of Ca5 (PO4) 3OH-H is

Table 1 | Comparing our work with other fuel cell
performances at 550 °C

Composition OCV(V) Power
density(mW cm −2 )

Ref.

Ni-NCAL/3C-SiC/ZnO/
NCAL-Ni

Ni-NCAL/BCFZY-ZnO/
NCAL-Ni

1.1 270 45

1.01 643 1

Ni/LCN–NSDC/Ag 0.8 500 46

Ni-NCAL/SrTiO3 (STO)/
NCAL–Ni

0.99 620 47

Ni-NCAL/ZnO/NCAL-Ni 1.06 482 48

Ni-NCAL/Film TiO2/NCAL 1.1 364 49

NCAL/BCZY/ Ca5(PO4)3OHH(Li) /BZCY/NCAL

1.04 661 Our work

1.04 V, and the maximum power density (MPD) is 661 mW cm −2 . This
power output is about 2.3 times the power density of Ca5 (PO4)3OH at
550 °C, similar to the change in conductivity. The I-V and I-P characteristic
curves of a Ca5(PO4)3OH-H electrolyte fuel cell at lower temperatures are
shown in Fig. 3D. Ca5(PO4)3OH-H electrolyte fuel cell performance comparison with other fuel cell performances at 550 °C is presented in
Table 1.The OCV and MPD are 1.015 V, and 573.6 mW cm −2, 1.046 V and
483.4 mW cm −2,1.053 Vand402.2 mW cm −2,1.064 Vand330.5 mW cm −2

at 520, 490, 460 and 430 °C, respectively. The power outputs of the fuel cell
using Ca5 (PO4)3OH-H as the electrolyte is much higher than that of using
the Ca5(PO4)3OH electrolyte, further demonstrating that constructing
metal cation vacancies can effectively improve the conductivity of the
material. The electrostatic interactions between the Ca vacancies and protonic species exert a significant influence on the transport behavior of
protons within the lattice. The vacancies induce changes in the local electric
field, thereby facilitating the directional movement of protons 41 . This
alteration in the electric field, coupled with the pathways created by the
vacancies, promotes the efficient transport of protons through the material.
Consequently, this phenomenon leads to an enhanced power output in
applications such as fuel cells, underscoring the pivotal role of Ca vacancies
in optimizing the electrochemical performance of the material. NCAL was
utilized as an electrode material due to its excellent conductivity and catalytic properties. Oxides with a similar structure, like LiAl0.5Co0.5O2, have
demonstrated a conductivity of 0.1 S cm −1 at 500 °C and strong catalytic
activity for the oxygen reduction reaction (ORR). At specific temperatures,
an amorphous layer may develop on the surface of NCAL, which enhances
the formation of multiple oxygen vacancies, thereby facilitating improved
oxygen ion conduction. Moreover, the introduction of Li + ions into the
material can positively influence its catalytic characteristics and the performance of fuel cells. Liu et al. have highlighted the significant potential of
NCAL as an electrode material 43 . Additionally, Chen et al. have reported
promising results for NCAL electrodes in both the oxygen reduction reaction (ORR) and hydrogen oxidation reaction (HOR) 1 .

To further investigate proton conducting feature, the following
experimentswerecarriedout:i) Conductivityvshydrogenconcentration;ii)
Hydrogen concentration cell (HCC); iii) Conductivity isotopic effect; iv)
Conductivity measured using the proton conducting filter cell.

First,wedetermined theconductivityofCa5 (PO4)3OH-Hinhydrogen
at a different partial pressures or concentrations. According to relevant
literature, protons dissolved in oxides are considered to be interstitial protons Hi44. The diffusion of protons in oxides is positively proportional to the

partial pressure of hydrogen, as indicated by

Equation (1).

1 ��
2 H2 g ¼ H � i

-
H i � nPH � 2 1=2

� i þ e0K 1 ¼ H i �

H � 2 1=2 ð1Þ

The proton conductivity should be proportional to the concentration
ofhydrogen.AsshowninFig. 4Aastheconcentrationofhydrogenperoxide
increases, the arc of the impedance spectrum gradually decreases. The
specific impedance data are summarized in Supplementary Table S4. It can
be seen that with the increase in hydrogen concentration, R0 decreases from
0.287 Ω cm 2 to 0.194 Ω cm 2, and R1 also decreases from 0.101 Ω cm 2 to
0.051 Ω cm 2 . As shown in Fig. 4B, the conductivity of Ca5(PO4)3OH-H(Li)
increases with increasing hydrogen concentration, and at 20% and above
hydrogen concentration, it exhibits a conductivity of more than 0.1 S cm −1 .

To determine the proton transfer number (tH+) for Ca5(PO4)3OHH(Li), hydrogen concentration cells (pure and 5% H2) were applied using
silver as electrode and results are shown in Fig. 4C. The theoretical electromotive force is calculated according to the Nernst equation. The tH+ is
theratioofthemeasuredelectromotiveforcetothetheoreticalelectromotive
force. The measured proton transference number is in the range of 0.7 to
0.87 between 430 and 550 o C. Silver is not ideal electrode for hydrogen, the
measured proton transfer number may be lower than the real value.

Furthermore, we fabricated a cell with a configuration of NCAL/
BCZY/ Ca5(PO4)3OH-H(Li) /BZCY/NCAL, where BaZr0.3Ce0.6Y0.1O3-δ
(BCZY)servesasa filtertoblockotherionsandBCZYisaprotonconductor
at low-temperature. The I-V and I-P curves are shown in Fig. 4D where the
cell exhibits a maximum power output of 485 mW cm –2, which is 73.3 % of
the cell without BCZY. Possible causes of performance degradation include
i) the increased impedance of the BZCY layer itself, ii) two additional
interfaces between Ca5(PO4)3OH-H(Li) and BZCY layers, leading to more
polarization resistance 44 . The EIS in Fig. 4E could further prove the above
conclusionwiththe fittingdatalistedinSupplementaryTableS6.Theohmic
resistancesincreasedfrom0.43to0.60 Ω andthegrainboundaryresistances
increased from 0.0045 to 0.0106 Ω, for the cells with and without BZCY
layers. The results show that proton conduction is the dominant contribution to the conductivity. Finally, the isotopic effect measurements
further provide evidences for proton transport property as shown in Fig. 4F
where EIS measurements were obtained from H2 and D2 atmospheres,
respectively, a significant difference between them and larger resistance
displays that D + diffusion is slower than that of H + as a clear indication of
proton diffusion isotopic effect.

We further studied the proton transport isotope effect (KIE) of
Ca5(PO4)3OH-H(Li) and the results are presented in Fig. 4F and the EIS
fitting data is listed in Supplementary Table S5. For the Grotthuss
mechanism, the KIE is usually bigger than (mD/mH)1/2 ~ 1.4 because the
mass of D (mD) is twice that of H atom (mH). By contrast, the KIE for
vehicle mechanism is equal to the viscosity ratio of D2O to H2O, ~1.2. At
550 o C, the conductivity of Ca5(PO4)3OH-H(Li) in H2/air atmosphere is
2.04 times higher than that of Ca5(PO4)3OH-H(Li) in D2/air atmosphere,
confirming that the proton transport is governed by the Grotthuss
mechanism.ThedifferenceinisotopicmobilitybetweenH + andD + ionscan
be attributed to the influence of Ca vacancies. The presence of Ca vacancies
alters the local environment and energetics of proton transport within the
material. Consequently, any observed variations in the mobility of H + and
D + ions can be indirectly linked to the impact of these vacancies on the
mechanisms of proton transport. This disparity in isotopic mobility serves
as a valuable indicator, providing insights into how Ca vacancies affect and
modulate proton transport processes within the material.

The Ca vacancy plays a key role in improving proton ion conductivity.
Calcium vacancy generation is further induced by doping trivalent La ions
to form Ca5(PO4)3OH-La, by high-temperature solid-phase reaction. The
fuel cell performances of Ca5(PO4)3OH-La were tested which displayed a
maximum power density of 600 mW cm −2 at 550 °C, approximately 2 times
that of the original Ca5(PO4)3OH as shown in Supplementary Fig. S4 (A)),
and the electrochemical performance was measured, see Supplementary
Fig. S4(B). The power output of 392 mW cm −2 is obtained at 460 °C, which
is higher than the power density of original Ca5(PO4)3OH at 550 °C.
According to the EIS in Supplementary Fig. S4(C) the ionic conductivity of
Ca5(PO4)3OH-La reaches 0.118 S cm −1 at 550 o C (Supplementary Fig. S4
(D)) which is approximately 2 times higher than that of original

Fig. 4 | Conductivity analysis. (A) EIS spectra and (B) conductivity of Ca5 (PO4)3
OH-H varying with hydrogen concentration. C The measured electromotive force,
theoretical electromotive force, and H + transfer number from a hydrogen concentration cell based on Ca5(PO4)3OH-H(Li) electrolyte. D I-V and I-P characteristics of the Ca5(PO4)3OH-H(Li) electrolyte fuel cell with and without BZCY layer

Ca5(PO4)3OH. It can be seen that the cation vacancy produced by different
ways can effectively improve the conductivity of the materials.

We have further carried out theoretical modeling and calculations to
study proton transport mechanism. The schematic diffuse paths and calculated energy barriers for Li + -H + pair or H + are shown in Fig. 5A. The
corresponding relaxed structures of initial state and final state used for
energy barriers calculation are also shown in Fig. S5. We calculated and
compared various proton transport mechanisms and paths, e.g. i) The
calculated Li + -H + pair diffusion barrier along adjacent Ca vacancies is
2.73 eV,whichisveryhigh.TheLi-HbondsintheCavacanciesarerelatively
weak (~2.45 Å) and should not be able to move simultaneously, see Supplementary Fig. S6(a). This result is not quite as expected; ii) The diffusion
barrier of Ca vacancies filled with adjacent Li + associated bare H + diffusion
is approximately 0.35 eV; iii) while undoped by Li +, the initial result of H
along adjacent Ca vacancies is approximately 1.89 eV, which is also very
high being not a favorable path/mechanism. After the formation of the Ca 2+

vacancy in Ca5(PO4)3OH, single H + is prone to enter into the vacancy
region, bond with the surrounding O atom to form a stable structure
Supplementary Fig. S6(b). However, the imbalance of valence state still
exists, just like the presence of a negative charge in the vacancy region.
Therefore, when H + migrates between Ca 2+ vacancies, it will subject to
additional electrostatic attraction from the vacancy region, which leads to a
high ion diffusion barrier of 1.89 eV. Based on these theoretical calculations,
Fig. 5A illustrates the most favorable theoretical calculations for proton
transport paths/ mechanism, while is the most feasible in agreement with
experimental resultspresented in Fig. 2B, where 0.274 eV is obtained for the
proton transport activation energy. Based on both theoretical and experimental results, we proposed the proton transport mechanism via a H +

coupled Li + transport mechanism as presented by Fig. 5A.

WhentheCa 2+ vacancyis filledbyH + andLi + simultaneouslyasshown
in Supplementary Fig. S6(c), the valence state of the vacancy region is
balanced, and the electrostatic attraction on H + from vacancy region is

under H2/air atmosphere at 550 °C. E Corresponding electrochemical impedance
spectroscopy of the cells with and without BZCY layer and EIS fitting results are
listed in Supplementary Tables 3, 6. F Typical Nyquist plots for the Ca5(PO4)3OHH(Li) in the presence of H2 (red) and D2 (blue) at 550 o C and fitting results are listed
in Supplementary Tables 3, 5.

eliminated. In addition, the interaction between H + and Li + somewhat
weakens that between H + and O atom (The DFT results show that the H-O
bond length increases from 0.98 Å without Li to 1.01 Å with Li doped). The
above two factors lead to a large reduction of ion diffusion barrier of H +

when migrating between Ca 2+ vacancies. Therefore, in H + and Li + filled
Ca5(PO4)3OH, the diffusion barrier of H + along Ca 2+ vacancies is only
~0.35 eV (Fig. 5A).

Overall, Li + filling can effectively reduce the diffusion barrier of the H +

atom along the Ca vacancies, which may result in an enhanced proton
conduction. Figure 5B presents a detailed proton transport in
Ca5(PO4)3OH.TheexistenceofCavacanciesintroducesadditionalhopping
sites for protons within the lattice, facilitating easier and faster proton
migration through the material. These hopping sites create pathways that
enable protons to move more freely, ultimately enhancing proton conductivity overall. The study aims to elucidate the facilitated hopping
mechanism, considering factors such as electrostatic interactions, hydration
and protonation dynamics, and defect-mediated diffusion pathways.
Through this comprehensive analysis, the research contributes to a deeper
understanding of the proton transport mechanism facilitated by Ca
vacancies in Ca5(PO4)3OH. Introducing a H⁺-coupled Li⁺ mechanism to
occupy the Ca vacancies could indeed facilitate proton transport. Li⁺ ions
associated with the Ca vacancies, acting as intermediaries or bridges for
proton transport. They create sites or pathways for proton transfer. Protons
(H⁺) could then hop from one Li⁺ site to another, utilizing these occupied
vacancies as stepping stones for their movement. On the other hand, the
presence of Li⁺ ions within the vacancies facilitates a more controlled and
guided path for protons. This coupling mechanism allows for leveraging Li⁺
ions within the Ca vacancies introducing a potential strategy to enable
proton transport, addressing the size disparity between Ca vacancies and
protons. It creates a favorable scenario where the Li⁺ ions act as intermediaries, aiding in the movement of protons through the crystal lattice.
Further exploration and experimental validation of this proposed

Fig. 5 | DFT Calculations. (A) The energy barrier of Li-H pair or H + diffuses between adjacent Ca vacancies of CaPOH, and H + diffuses between Li + filled Ca vacancies of
CaPOH. The schematic diffuse pathways for Li-H pair and H + are also shown in the insets. B Proton conduction mechanism in Ca5(PO4)3OH.

mechanism could significantly contribute to understanding and enhancing
proton conductivity in Ca5(PO4)3OH-based materials. To enhance the
accuracy of the computational model, the formation energies of Ca
vacancies in Ca5(PO₄)₃OH were calculated using spin-polarized calculations. There exist two kinds of Ca vacancies in Ca5(PO₄)₃OH, see clearly
their structures in Fig. S7(b, c). The spin-polarized calculations reveal that a
7-coordinated Ca vacancy configuration (Ca-v2) exhibits a slightly lower
formation energy (Supplementary Table S7), indicating that it may be easier
to form. This finding supports the presence of both 9- and 7-coordinated Ca
vacanciesinthestructure,offeringadditionaldiffusionpathways.Thepath2
(Ca-v1→Ca-v2→Ca-v1) is also included [Fig. S7(d)], which takes advantage of these energetically favorable vacancies.

Regarding defined direction of bulk diffusion, there is a possibility of
proton diffusion along the c-axis, thereby introducing an additional pathway

[(Path 3: Ca-v1 → Ca-v1 along the c-axis, see in Fig. S7(e)]. This approach
aligns with our experimental observations, suggesting that the diffusion
mechanism is not confined to a single direction. These modifications provide
a more comprehensive understanding of the diffusion behavior in different
directions, ensuring alignment with experimental diffusion characteristics.
The insights gained from this detailed examination will significantly advance
the knowledge of proton-conductive materials, potentially impacting their
applications in various energy-related technologies.

Conclusions
This study represents a significant leap in materials science by unveiling a
pioneering method to boost proton conductivity via calcium vacancy and
lithium-ion bridging. The development of protonated Ca5(PO4)3OH-H(Li)
introduces a groundbreaking proton conduction mechanism, achieving
unprecedented material performance with exceptional proton conductivity
of 0.1 S cm −1 at 300–550 °C and superior PCFC performance, peaking at
661 mW cm −2 at 550 °C. This achievement not only advances the field of
ceramic proton conductors but also demonstrates the potential of
Ca5(PO4)3OH-H(Li) as a leading electrolyte for future PCFCs. The integration of calcium vacancies with lithium-ion bridging, fostering proton
transport pathways, enhances our understanding of proton conduction and
signifies the material’s broader impact on energy conversion and storage
technologies, marking a stride towards sustainable energy solutions.

Methods
Materials synthesis and fuel cell fabrication
The Ca5(PO4)3OH power was bought from Jiuding Chemical Co., Ltd.,
China. For the synthesis of Ca5(PO4)3OH containing Calcium (Ca)
vacancies, 10 g of Ca5(PO4)3OH power was firstly soaked in 400 mL of
0.25 M KCl solution and stirred for 2 h at room temperature. The resulting
suspensions were washed by deionized water to remove excess K + ions.

During the above treatments, the hydrated K ions inserted into
Ca5(PO4)3OH powder, and caused an equivalent loss of divalent Ca ions,
consequentlyleavingCa vacancies.400 mLof0.25 MLiClsolutionwasthen
added into the washed Ca5(PO4)3OH-K and stirred for 2 h at room temperature, in which K ions were replaced by Li ions. The resulting suspensions were washed by deionized water to remove excess Li + ions. To
synthesize Ca5(PO4)3OH-H(Li) power, the resulting Ca5(PO4)3OH-Li
power was redispersed in 0.5% Acetic acid solution for 3 h to exchange
partial Li ions with protons.

To verify further Ca-vacancy effect, we employed lanthanum-doped
Ca5(PO4)3OH (La0.25Ca4.75(PO4)3OH), which was prepared by hightemperature solid-phase reaction. Firstly, La(NO3)3 and Ca5(PO4)3OH
were fully ground according to stoichiometric ratios. The above powder was
thencalcined at 1000 °C with a heating rate of 5°/min in air for 4 h, followed
by adequately grounding of the powders to get the final sample. In this way,
two higher valency La 3+ replaced Ca 2+, to maintain neutrality, one extra Cavacancy must be created thus to enhance the Ca-vacancy concentration in
the sample.

To prepare BaZr0.3Ce0.6Y0.1O3-δ (BZCY), Ba, Zr, Ce, and Y metal
nitrates were dissolved in distilled water and citric acid was added as a
reacting agent. The solution was stirred and heated at 70 °Cfor 2 hours. The
compositionwas supposedto be BaZr0.3Ce0.6Y0.1O3-δ. The attained solution
was heated at 250 °C for 4 h, and a dark grey-black colored ash was found.
Finally, the resulting ash was successively calcined at 1000 °C for 3 hours to
yield BZCYpowder.ToauthenticatetheH + ionconductionofthefuelcell,a
distinct fuel cell device was fabricated using the similar assembly but adding
two O 2- ion blocking verythin layers of BaZr0.3Ce0.6Y0.1O3-δ on both sidesof
electrolyte. The fuel cell device was assembled in a design of NCAL/BCZY/
Ca5(PO4)3OH-H(Li) /BZCY/NCAL having same parameters regarding
active area and thickness, as used for the cell mentioned before i.e. NCALNi/ Ca5(PO4)3OH-H(Li)/NCAL-Ni.

Structural characterizations
The phase structure was measured by X-ray diffraction (XRD; Bruker D8
ADVANCE) with Cu Ka radiation source over the 2θ range of 10–80°,
operating at 45 kV and 40 mA. Morphological properties were characterized by scanning electron microscopy (SU8010, Hitachi Japan). The elemental composition was characterized by Inductively Coupled Plasma
Atomic Emission Spectrometry (ICP-AES, Agilent 5110). Transmission
electron microscopy (TEM, Talos F200S G2) was used to characterize the
materials. The functional groups were analyzed by Fourier Transform
Infrared Spectrometer (FTIR) conducted with PerkinElmer Frontier. X-ray
photoelectron spectra (XPS) was performed through Thermo Fisher,
ESCALAB 250Xi+ with the source of Al-Kα (hv=1486.6 eV).

To determine the chemical composition of each material, the sample
was measured by ICP-AES. We set the mole concentration of P as 1, and the
mole concentration of other elements, including Ca, O and Li, and the
corresponding chemical formula were determined relative to that of P. The
mole concentration of H was determined based on the principle of charge
balance and the valence states of other elements. Measurements of ion
conductivity 0.1 g powders were pressed under a hydraulic press pressure of
200–220 MPa to get electrolyte pellet with active area of 0.64 cm 2 . As prepared pellet was sandwiched between two NCAL (LiNi0.8Co0.15Al0.05O2)
symmetric electrodes pasted on Ni foam. The electrochemical impedance
spectra (EIS) were measured with a potentiostate (CHI) with the temperature from 550 o C to 400 o C. The applied frequency range was 1 MHz to
0.1 Hz with 10 mV bias voltage. The digital instrument (ITECH8511) was
used for measuring I–V (current-voltage) and I–P (current-power) characteristics. The conductivity was calculated according to the simulated
resistance in the EIS results.

Measurements of electronic conduction contribution
Response current testing under direct-current (DC) voltage was carried out
to identify the possible electronic conduction contribution to the conductivity of the Ca5(PO4)3OH-H(Li). Ag blocking electrodes were used to

measure the electronic conduction in the configuration of Ag/
Ca5(PO4)3OH-H(Li)/Ag under N2 atmosphere.

The electronic conduction for fuel cell condition is measured by quasi
in-situ measurement. After the end of the fuel cell test, N2 is introduced into
the anode and cathode immediately. Then the response current testing
under direct-current (DC) voltage was measured in the configuration of
NCAL/ Ca5(PO4)3OH-H(Li)/NCAL under N2 atmosphere.

Proton transference number measurements
The proton transfer number was test through the hydrogen concentration cell
intheconfigurationofAg/Ca5(PO4)3OH-H(Li)/Ag.5%H2 andpureH2 were
passed into the cathode and anode of the cell respectively with the same flow
rate. The electromotive force (EMF) of the cell was recorded on a computerized instrument (IT8511A +, ITECH Electrical Co., Ltd., China). The
theoreticalelectromotiveforcewastestaccordingtotheNernstequationEq. 2,

E ¼ RT 2F ln P P 1 2

ð2Þ

R is the gas constant; T is the absolute temperature; F is Faraday’s
constant, P1 is the hydrogen partial pressure on the anode side and P2 is the
hydrogen partial pressure on the cathode side. The proton transfer number
(tH+) is the ratio of the actual electromotive force to the theoretical electromotive force.

KIE measurements and analyses
The impedance spectra of the two fuel cells were tested in H2/O2 and D2/O2
atmospheres, respectively. If protons are the dominant charge carriersin the
material, there will be a significant difference for the H + /D + conductivity
due to the isotopic effect. The KIE is the ratio of ion conductivity of
membrane in H2/O2 atmosphere to the value in D2/O2 atmosphere. For the
Grotthuss mechanism, protons are transported by hydrogen bond breaking
and recombination. In comparison with the case of H2/O2, the proton
conductivity is decreased in D2/O2, whose mass (mD) is twice that of H
atom(mH).Theoretically,theKIEforGrotthussmechanismis(mD/mH)1/
2 ~ 1.4. The experimentally obtained KIE is usually in the range of 1.4 to 3.3
due to the influence of material’s structure and chemical condition. In
contrast,theKIEforvehiclemechanismis~1.2duetothedifferenttransport
groups.

Computational details and modeling
The density functional theory (DFT) calculations were implemented in the
Vienna ab initio simulationpackage(VASP) based onprojector-augmented
wave (PAW) pseudopotentials. The generalized gradient approximation
with the exchange-correction function of Perdew-Burke-Ernzerhof form
(GGA-PBE) was employed and the kinetic energy cutoff was set to be
400 eV for plane wave basis expansion. During structure optimization, the
convergence criteria of electronic self-consistent energy and atomic force
were respectively set to 10 −5 eV and 0.01 eV. To investigate the migration
property of H + along Ca vacancies in CaPOH material, a (2×1×1) supercell
was adopted with Brillouin zone (BZ) sampled by a Γ-centered (2×4×5)
Monkhorst-Pack grid, and the energy barrier was calculated by using the
climbing image nudged elastic band (CI-NEB) method.

Data availability
The data that support the findings of this study are available from the
authorsonreasonablerequest,seeauthorcontributionsforspecificdatasets.

Received: 19 June 2024; Accepted: 5 December 2024;

References
1. Xia, C. et al. Shaping triple-conducting semiconductor
BaCo0.4Fe0.4Zr0.1Y0.1O3-δ into an electrolyte for low-temperature
solid oxide fuel cells. Nat. Commun. 10, 1707 (2019).

2. Saini, D. S. et al. A Promising Proton Conducting Electrolyte
BaZr1-xHoxO3-δ (0.05 ≤ x ≤ 0.20) Ceramics for Intermediate
Temperature Solid Oxide Fuel Cells. Sci. Rep. 10, 3461 (2020).
3. Duan, C. et al. Readily processed protonic ceramic fuel cells with high
performanceatlowtemperatures. Science(1979) 349,1321–1326(2015).
4. Suzuki, T. et al. Impact of anode microstructure on solid oxide fuel
cells. Science (1979) 325, 852–855 (2009).
5. Rauf, S. et al. Highly active interfacial sites in SFT‐SnO2
heterojunction electrolyte for enhanced fuel cell performance via
engineered energy bands: envisioned theoretically and
experimentally. Energy Environ. Mater. 7, e12606 (2024).
6. Shah, M. Y. et al. ZnO/MgZnO heterostructure membrane with type II
band alignment for ceramic fuel cells. Energy Mater. 2, 200031 (2022).
7. Yu, Y. et al. Synergistic Proton and Oxygen Ion Transport in Fluorite
Oxide-Ion Conductor. Energy Mater. Adv. 5, 0081 (2024).
8. Sun, Y. et al. Advancements in cathode catalyst and cathode layer
design for proton exchange membrane fuel cells. Nat. Commun. 12,
5984 (2021).
9. Wachsman, E. D. & Lee, K. T. Lowering the temperature of solid oxide
fuel cells. Science (1979) 334, 935–939 (2011).
10. Xing, Y. et al. Proton shuttles in CeO2/CeO2-δ core–shell structure.
ACS Energy Lett. 4, 2601–2607 (2019).
11. Singh, M. et al. Recent advancement of solid oxide fuel cells towards
semiconductor membrane fuel cells. Energy Mater. 4, 400012 (2024).
12. Chen, G. et al. Advanced Fuel Cell Based on Perovskite La-SrTiO3
Semiconductor as the Electrolyte with Superoxide-Ion Conduction.
ACS Appl Mater. Interfaces 10, 33179–33186 (2018).
13. Kamran Yousaf Shah, M. A., Mushtaq, N., Rauf, S., Xia, C. & Zhu, B.
The semiconductor SrFe0.2Ti0.8O3-δ-ZnO heterostructure electrolyte
fuel cells. Int J. Hydrog. Energy 44, 30319–30327 (2019).
14. Qiao, Z. et al. Electrochemical and electrical properties of doped
CeO2-ZnO composite for low-temperature solid oxide fuel cell
applications. J. Power Sources 392, 33–40 (2018).
15. Wu, Y. et al. Proton conduction and fuel cell using the CuFe-oxide
mineral composite based on CuFeO2 structure. ACS Appl. Energy
Mater. 1, 580–588 (2018).
16. Akbar, N. et al. Space Charge Polarization Effect in Surface-Coated
BaTiO3 Electrolyte for Low-Temperature Ceramic Fuel Cell. ACS
Appl. Energy Mater. 7, 1128–1135 (2024).
17. Shim, J. H. Ceramics breakthrough. Nat. Energy 3, 168–169 (2018).
18. Wang, N. et al. Advanced cathode materials for protonic ceramic fuel
cells: recent progress and future perspectives. Adv. Energy Mater. 12,
2201882 (2022).
19. Dzul, I., Hernández, T., Hernández Carrillo, R., Peña, Y. & Rojo, T.
Synthesis and electric properties of perovskite Pr0.6 Ca0.4 Fe0.8 Co0.2
O3 for SOFC applications. Ion. (Kiel.) 20, 1031–1037 (2014).
20. Rauf, S. et al. Alternative Strategy for Development of Dielectric
Calcium Copper Titanate-Based Electrolytes for Low-Temperature
Solid Oxide Fuel Cells. Nanomicro Lett. 17, 1–23 (2025).
21. Bibi, B. et al. Emerging semiconductor ionic materials tailored by
mixed ionic-electronic conductors for advanced fuel cells. Adv.
Powder Mater. 3, 100231 (2024).
22. Goodenough, J. B. Oxide-ion conductors by design. Nature 404,
821–823 (2000).
23. Kraytsberg, A. & Ein-Eli, Y. Review of Advanced Materials for Proton
Exchange Membrane Fuel Cells. Energy Fuels 28, 7303–7330 (2014).
24. Wang, X. X., Swihart, M. T. & Wu, G. Achievements, challenges and
perspectives on cathode catalysts in proton exchange membrane fuel
cells for transportation. Nat. Catal. 2, 578–589 (2019).
25. Shen, G. et al. Multi-functional anodes boost the transient power and
durability of proton exchange membrane fuel cells. Nat. Commun. 11,
1191 (2020).
26. Qian, X. et al. CdPS3 nanosheets-based membrane with high proton
conductivity enabled by Cd vacancies. Science (1979) 370, 596–600
(2020).

27. Yashima, M. et al. Diffusion path and conduction mechanism of
protons in hydroxyapatite. J. Phys. Chem. C. 118, 5180–5187 (2014).
28. El Hammari, L., Laghzizil, A., Barboux, P., Saoiabi, A. & Lahlil, K.
Crystallinityand fluorinesubstitutioneffectsontheprotonconductivityof
porous hydroxyapatites. J. Solid State Chem. Fr. 177, 134–138 (2004).
29. Yamashita, K., Kitagaki, K. & Umegaki, T. Thermal instability and
proton conductivity of ceramic hydroxyapatite at high temperatures.
J. Am. Ceram. Soc. 78, 1191–1197 (1995).
30. Furuseki, T. & Matsuo, Y. Anhydrous Proton Conductivity in HApCollagen Composite. J. Compos. Sci. 6, 236 (2022).
31. Horiuchi, N. et al. Dielectric properties of fluorine substituted
hydroxyapatite: the effect of the substitution on configuration of
hydroxide ion chains. J. Mater. Chem. B 3, 6790–6797 (2015).
32. Bystrov, V., Paramonova, E., Bystrova, N., Sapronova, A. & Filippov,
S. Computational molecular nanostructures and mechanical/
adhesion properties of hydroxyapatite. Micro Nanostruct. Biol. Syst.
77–93 (2005).
33. Luo, X., Song, X., Cao, Y., Song, L. & Bu, X. Investigation of calcium
carbonate synthesized by steamed ammonia liquid waste without use
of additives. RSC Adv. 10, 7976–7986 (2020).
34. Jiang, X. et al. Structure and enhanced dielectric temperature stability
of BaTiO3-based ceramics by Ca ion B site-doping. J. Materiomics 7,
295–301 (2021).
35. Zhu, G. et al. Tuning the K+ concentration in the tunnels of α-MnO2 to
increase the content of oxygen vacancy for ozone elimination.
Environ. Sci. Technol. 52, 8684–8692 (2018).
36. Roumila, Y., Meziani, D., Belkhettab, I., Abdmeziem, K. & Trari, M.
Hydroxyapatite Fenton-like catalysis for methyl violet degradation.
Effect of operating parameters and kinetic study. J. Mater. Sci. Mater.
Electron. 34, 909 (2023).
37. Reig, F. B., Adelantado, J. V. G. & Moreno, M. C. M. M. FTIR
quantitative analysis of calcium carbonate (calcite) and silica (quartz)
mixtures using the constant ratio method. Application to geological
samples. Talanta 58, 811–821 (2002).
38. Liu, W. et al. Solar-induced direct biomass-to-electricity hybrid fuel
cell using polyoxometalates as photocatalyst and charge carrier. Nat.
Commun. 5, 3208 (2014).
39. Fulmer, G. R. et al. NMR chemical shifts of trace impurities: common
laboratorysolvents,organics,andgasesindeuteratedsolventsrelevantto
the organometallic chemist. Organometallics 29, 2176–2179 (2010).
40. Ji,C., He,B.,Yun, S.,Bai, X. &Lin, B.Thefracture mechanical behavior
simulation of calcium-deficient hydroxyapatite crystals by molecular
dynamics and first-principles calculation. J. Mech. Behav. Biomed.
Mater. 137, 105526 (2023).
41. Watras, A. et al. The role of the Ca vacancy in the determination of the
europium position in the energy gap, its valence state and
spectroscopic properties in KCa (PO3) 3. Phys. Chem. Chem. Phys.
16, 5581–5588 (2014).
42. Akbar, N. et al. Tunning tin-based perovskite as an electrolyte for
semiconductor protonic fuel cells. Int J. Hydrog. Energy 47,
5531–5540 (2022).
43. Liu, J. et al. High performance low-temperature solid oxide fuel cells
based on nanostructured ceria-based electrolyte. Nanomaterials 11,
2231 (2021).
44. Akbar, N., Paydar, S. & Wu, Y. Tuning an ionic-electronic mixed
conductor NdBa0.5Sr0. 5Co1. 5Fe0.5O5+δ for electrolyte functions of
advanced fuel cells. Int. J. Hydrog. Energy 46, 9847–9854 (2021).
45. Xing, Y. et al. Cubic silicon carbide/zinc oxide heterostructure fuel
cells. Appl. Phys. Lett. 117, 162105 (2020).
46. Zhu, B. et al. Schottky junction effect on high performance fuel cells
based on nanocomposite materials. Adv. Energy Mater. 5, 1614–6840
(2015).
47. Chen, G. et al. Electrochemical mechanisms of an advanced lowtemperature fuel cell with a SrTiO3 electrolyte. J. Mater. Chem. A
Mater. 7, 9638–9645 (2019).

48. Xia, C. et al. Study on zinc oxide-based electrolytes in lowtemperature solid oxide fuel cells. Materials 11, 40 (2017).
49. Dong, W. et al. Semiconductor TiO2 thin film as an electrolyte for fuel
cells. J. Mater. Chem. A Mater. 7, 16728–16734 (2019).

Acknowledgements
This work is supported by the funding from Science and Technology
Department of Jiangsu Province under Grant BE2022029.

Author contributions
Z.B.conceptualizethestudyandwrotetheinitialdraft ofthemanuscript.Y.L.
fabricated the fuel cells and conducted the electrochemical measurements.
N.A.further workedonthe manuscript toits final version. F.W.tookchargeof
designingandsupervisingtheelectrochemicalexperiments.Q.P.carriedout
DFT calculation under the supervision and instructions from S.Y., and Y.J.
evaluated and analyzed the data. M.S. and J.W. provided valuable
assistance in interpreting and analyzing the results. All authors engaged in
collaborative discussions to refine the findings and enhance the manuscript.

Competing interests
The authors declare no competing interests.

Additional information
Supplementary information The online version contains
supplementary material available at
[https://doi.org/10.1038/s43246-024-00719-6.](https://doi.org/10.1038/s43246-024-00719-6)

Correspondence and requests for materials should be addressed to
Bin Zhu, Faze Wang or Sining Yun.

Peerreviewinformation Communicationsmaterials thankstheanonymous
reviewers for their contribution to the peer review of this work. Primary
Handling Editors: Ruben Rizo and Jet-Sing Lee.

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