Highly bright perovskite light-emitting diodes enabled by retarded Auger recombination

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Article https://doi.org/10.1038/s41467-025-56001-x
# Highly bright perovskite light-emitting diodes enabled by retarded Auger recombination

Received: 10 June 2024

Accepted: 6 January 2025

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Zhiqi Li 1,2,3,8, Qi Wei 4,8, Yu Wang 1, Cong Tao 5, Yatao Zou 5, Xiaowang Liu 5,
Ziwei Li 6, Zhongbin Wu 5, Mingjie Li 4, Wenbin Guo 3, Gang Li 2,
Weidong Xu 5,7 & Feng Gao 1

One of the key advantages of perovskite light-emitting diodes (PeLEDs) is their
potential to achieve high performance at much higher current densities
compared to conventional solution-processed emitters. However, state-of-theart PeLEDs have not yet reached this potential, often suffering from severe
current-efficiency roll-off under intensive electrical excitations. Here, we
demonstrate bright PeLEDs, with a peak radiance of 2409 W sr −1 m −2 and negligible current-efficiency roll-off, maintaining high external quantum efficiency
over 20% even at current densities as high as 2270 mA cm −2 . This significant
improvement is achieved through the incorporation of electron-withdrawing
trifluoroacetate anions into three-dimensional perovskite emitters, resulting in
retarded Auger recombination due to a decoupled electron-hole wavefunction. Trifluoroacetate anions can additionally alter the crystallization dynamics
and inhibit halide migration, facilitating charge injection balance and
improving the tolerance of perovskites under high voltages. Our findings shed
light on a promising future for perovskite emitters in high-power light-emitting applications, including laser diodes.

Metal halide perovskite emitters are emerging as promising candidates for next-generation light-emitting diodes (LEDs) due to their
low cost, high color purity, tunable emission wavelength, compatibility with low-temperature processing, and flexible module
fabrication 1–7 . Over recent years, numerous strategies have been
developed to improve the external quantum efficiencies (EQE) of
perovskite LEDs (PeLEDs), including defect passivation 8, light outcoupling management 9, interface engineering 10, compositional and
dimensional tailoring 11,12, as well as optimization of crystallization

processes 13, boosting the EQE to more than 20% 14,15 and gradually
approaching state-of-the-art organic LEDs (OLEDs) 16,17 . Beyond
achieving high peak quantum efficiencies, it has been generally
believed that the perovskite emitters can potentially achieve better
performance at high current density compared to OLEDs, making
them desirable for high-power applications, such as outdoor display,
lighting and potentially laser diodes 18,19 .

Despite the fact that the peak brightness of PeLEDs is usually
higher than that of OLEDs, their performance remains significantly

1Department of Physics, Chemistry and Biology (IFM), Linköping University, Linköping, Sweden. 2Department of Electrical and Electronic Engineering,
Photonic Research Institute (PRI), Research Institute of Smart Energy (RISE), The Hong Kong Polytechnic University, Hong Kong, China. 3 State Key Laboratory
of Integrated Optoelectronics, College of Electronic Science and Engineering, Jilin University, Changchun, China. 4 Department of Applied Physics, The Hong
Kong Polytechnic University, Hong Kong, China. 5 Frontiers Science Center for Flexible Electronics, Institute of Flexible Electronics (IFE), Northwestern
Polytechnical University, Xi’an, China. 6 Hunan Institute of Optoelectronic Integration, College of Materials Science and Engineering, Hunan University,
Changsha, China. 7 Henan Institute of Flexible Electronics (HIFE) and School of Flexible Electronics (SoFE), Henan University, 379 Mingli Road, 450046
Zhengzhou, China. 8 These authors contributed equally: Zhiqi Li, Qi Wei. [e-mail: ming-jie.li@polyu.edu.hk; guowb@jlu.edu.cn; gang.w.li@polyu.edu.hk;](mailto:ming-jie.li@polyu.edu.hk)
[ifewdxu@nwpu.edu.cn; feng.gao@liu.se](mailto:ifewdxu@nwpu.edu.cn)

below expectations due to a pronounced decrease in quantum efficiencies over increasing current densities 20,21 . This current-efficiency
roll-off can be mainly attributed to Auger recombination—the primary
nonradiative paths in conventional III-V quantum wells and chalcogenide quantum dot LEDs—also playing a critical role in
perovskites 18,22,23 . For typical three-dimensional (3D) perovskites, the
Auger recombination constant is ~10 −27 —10 −28 cm 6 s −1, markedly faster
by two to three orders of magnitude than III-V quantum well emitters
(~10 −30 cm 6 s −1 ) 24,25 . This effect is even more pronounced in prevailing
mixed-dimensional and quantum dot perovskite emitters due to
enhanced carrier confinement and localization 26,27 . More seriously,
unbalanced charge injection and relevant leakages in working LEDs
exacerbates Auger recombination even at low current densities, due to
enhancement in many-body interactions 28 . In addition, fast degradation of PeLEDs at high driving voltages is another critical issue limiting
their performance at high current densities as a result of facilitated ion
movement and increased Joule heating 29–32 . As such, to achieve high
performance of PeLEDs under high current densities, so that highpower applications and electrically pumped lasers can be realized, it is
necessary to simultaneously mitigate Auger recombination and device
degradation at high excitations.

In this report, we demonstrate bright and highly stable PeLEDs by
simultaneously slowing down Auger recombination, balancing charge
injection and suppressing ion migration. Specifically, we have managed to decrease the Auger recombination constant by an order of
magnitude in 3D perovskite emitters, achieved by the reduced exciton

Fig. 1 | Device structure and performance characteristics of our PeLEDs.
a Device configuration and cross-sectional SEM image of the PeLED. The scale bar
represents 100 nm. b Electroluminescence (EL) spectra of the device at different
voltages. c Dependence of EQE on the current density (EQE–J). d Dependence of

binding energies in perovskites and the alleviation of charge accumulation at interfaces. Additionally, ion migration is effectively suppressed by introducing additional ionic interactions. All these effects
are realized by incorporating trifluoroacetate anions in well-passivated
3D perovskites with dense thin-film coverage. Consequently, the
champion PeLED shows a negligible current-efficiency roll-off up to
2000 mA cm −2, yielding a high radiance of 2409 W sr −1 m −2 . The high
radiance is also coupled with excellent operational stability with a halflifetime of 142 h under a large current density of 100 mA cm −2, representing highly bright and highly stable direct-current (DC) driven
PeLEDs.

Results
Device performance
We use formamidinium lead tri-iodide (FAPbI3) as the parental perovskites, with 5-ammonium valeric acid iodide (5AVAI) 33,34 and trifluoroacetate cesium (CsTFA) 35,36 as the additives to enhance the film
qualities. The perovskite films were prepared by spin-coating a mixture
of FAI, PbI2, 5AVAI, CsTFA with a ratio of 1.28:1:0.1: 0.12 in a DMF:DMSO
solution (details in methods). The addition of 5AVAI and an extra
amount of FAI is crucial for the formation of the α-phase of FAPbI333,34.

Unless otherwise stated, we refer to the samples without and with
CsTFA as F- and FCT-films/devices, respectively. We also compared the
samples enriched with CsI in the same molar ratio as CsTFA (0.12
equivalent to Pb 2+ ) to assess the contribution of the TFA − anions to the
film properties (Supplementary Fig. 1).

EQE on the radiance (EQE-R). e Dependence of current density and radiance on the
voltage (J–V–R). f A comparison of device efficiency and radiance with the state-ofthe-art NIR PeLEDs reported in the literatures (Supplementary Table 1). g Stability
of the device measured at a high current density of 100 mA cm −2 .

Fig. 2 | Carrier recombination dynamics of perovskite films. TA spectra for (a) Fand (b) FCT-films, respectively. c Power dependent IPL0 for F- and FCT-films. The
IPL0 values were extracted from TRPL spectra at time zero. d PLQY as a function of
excitation power density for F-and FCT-films. Proportion of the recombination ratio
as a function of carrier density for the F- (e) and FCT-films (f), respectively.

PeLEDs were fabricated with the multilayered device structure,
consisting of indium tin oxide (ITO)/polyethylenimine ethoxylated
(PEIE)-modified zinc oxide (ZnO)/perovskite/poly(9,9-dioctyl-fluorene-co-N-(4-butylphenyl) diphenylamine) (TFB)/molybdenum oxide
(MoO3)/Au (Fig. 1a). The cross-sectional scanning electron microscope
(SEM) image (Fig. 1a) shows the formation of dense perovskites films in
the emitting layer with a thickness of 80 nm. Our PeLEDs have been
engineered to ensure exceptional electrical excitation tolerance and
maintain stable electroluminescence (EL) spectra even under high
driving voltages of up to 6.0 V, as shown in the characteristics of EL and
EQE versus voltage curves (Fig. 1b and Supplementary Fig. 2, respectively). The FCT-devices show a notably higher peak EQE of 21.4% than
that of F-ones (~18.9%). With increasing current density, the EQE values
of FCT-devices remain above 20% until reaching a current density as
high as 2270 mA cm −2 (Fig. 1c). The devices exhibit minimal currentefficiency roll-off with EQE reductions of 5% at a high radiance of
around 2000 W sr −1 m −2, a stark contrast to the significant drop in
control devices (Fig. 1d). The negligible current-efficiency roll-off gives
rise to a high peak radiance of 2409 W sr −1 m −2 at a current density of
3368 mA cm −2 (Fig. 1e). We believe that such a decent performance
caused by CsTFA addition is mainly achieved by their TFA − anions
instead of Cs + cations, as the addition of only CsI leads to a mild effect
(Supplementary Fig. 1). We compare our device efficiency/radiance
with the state-of-the-art NIR PeLEDs, which is shown in Fig. 1f and
further summarized in Supplementary Table 1. In short, the best
reported high-performance PeLEDs with EQE over 20% commonly
show a moderate radiance 2,7,8,37, and our results represent a significant
improvement for DC-driven PeLEDs.

We next measured the operational lifetime of F- and FCT-PeLEDs
with a high constant current density of 100 mA cm −2 . State-of-the-art
NIR PeLEDs with the best stability, which showed a half operational
lifetime (T50) of 11,539 h at 5 mA cm −2, demonstrated 120.3 h at such a
harsh working condition of 100 mA cm −2 for accelerating
degradation 37 . As shown in Fig. 1g, while the F-devices exhibit a short
T50 of 11 h, our FCT-counterparts deliver a drastically enhanced

g Derived values of recombination rate constant for the F- and FCT-films, where k1
was extracted from low-fluence TRPL spectra, and k2 and k3 were extracted from
both TRPL and TA measurements. The error bars indicate the confidence intervals
of the fitted rates.

lifetime reaching 142 h. The excellent stability of our optimized devices
is also indicated by the evolution of peak EQE upon voltage scanning
cycles (0 ~ 6.0 V). As shown in Supplementary Fig. 3, the peak EQE of
F-devices drop significantly after several cycles. In contrast, the peak
EQE of FCT-devices can preserve half of the highest value until 52
cycles under the same scanning voltage span. The comparison of the
stability of our PeLEDs with those of state-of-the-art is also summarized
in Supplementary Table 1, where the much-improved lifetime at such a
high electrical excitation represents a notable improvement compared
with those of state-of-the-art PeLEDs.

Recombination kinetics in PeLEDs

To clarify the differences in device performance, particularly at high
current densities, we performed a series of optical measurements to
investigate the recombination kinetics of perovskites. We conducted
transient absorption (TA) and fluence-dependent time-resolvedphotoluminescence (TRPL) measurements. The characteristics and
relevant 2D pseudo-colour images of TA measurements are shown in
Supplementary Fig. 4 and Fig. 2a, b respectively. We observe two primary photobleaching bands situated at 450 nm and 780 nm in both
cases, which can be ascribed to the typical sub-band of iodideperovskites and radiative recombination respectively 38 . Notably, no
distinct charge carrier transfers in the different time scales and additional photobleaching features can be found, clearly suggestive of the
absence of low-dimensional phases. These findings are in line with the
X-ray diffraction (XRD) patterns and grazing-incidence wide-angle
X-ray scattering (GIXWAS) measurements, where only the signatures
of 3D perovskites were visible (Supplementary Fig. 5).

We then investigated the scaling of the initial PL intensity (at t = 0,
IPL0) with excitation density from TRPL measurements (Supplementary
Fig. 6) to uncover the nature of the carrier recombination 24 . As shown
in Fig. 2c, the IPL0 for both samples exhibit a quadratic dependence on
carrier density, suggesting that the photogenerated carriers predominantly undergo bimolecular recombination, typical of free carrier
processes. Additionally, we observe that the PL intensity gradually

deviated from the quadratic relationship with the increase of the carrier concentration, indicating a growing prominence of the Auger
process 39 . A similar phenomenon is also visible in the evolution of
relative PL quantum yields (PLQYs) as a function of carrier density for
both cases (Fig. 2d). Notably, the relative PLQYs of FCT-films are much
more tolerant to high carrier density compared to F-samples, suggesting a retarded Auger recombination 18,24 . These observations are
consistent with the device characteristics, where the peak EQE
appeared at high current densities reaching hundreds of mA cm −2 . In
addition, we note that PLQYs of FCT-films are also higher at low carrier
densities (<~10 15 cm −3 ), owing much to a mitigated trap-assisted nonradiative recombination. Such a difference eventually leads to a higher
absolute PLQY of 81% for FCT-films compared to 62% of F-films at the
best excitation conditions (Supplementary Fig. 7), in line with the
higher peak EQE of FCT-devices than F-ones.

Having confirmed that bimolecular recombination is the dominant recombination pathway and the absence of low-dimensional
phases, we are able to quantify the recombination kinetics with a
combined use of TA and fluence-dependent TRPL by the following
equation 40 :

dn

dt = � k1n � k2n2 � k3n3 ð1Þ

where k1 is assigned to the trap-assisted nonradiative recombination
constant, k2 is the bimolecular recombination (band-to-band recombination) rate constant, k3 is the trimolecular (Auger) recombination rate
constant. Here, k1 was experimentally determined by TRPL spectra at
low fluences, while k2 and k3 are deduced from both TRPL and TA tests
(Supplementary Figs. 4, 6, 8, and 9). We show the extracted ratios of the
various recombination pathways upon carrier densities in Fig. 2e, f and
summarized the average recombination constants in Fig. 2g and
Supplementary Table 2. The TRPL-derived rates are in good agreement
with that extracted from TA measurements. Compared to F-films, we
observe largely retarded mono-molecular and Aguer recombination
constants but close bimolecular recombination rates in FCT-samples.
These contrasts thus lead to a much more dominant role of radiative
recombination in FCT-films across a wide range of excitation fluence
(Fig. 2e, f). The average k3 of FCT-films extracted by TRPL and TA is
(1.15 ± 0.64) × 10 −28 cm 6 s −1 and (5.28 ± 0.43) × 10 −29 cm 6 s −1, respectively.
Notably, the best sample showed an ultra-slow Auger rate of
1.29 × 10 −29 cm 6 s −1 (obtained by TA). These results are among the best
so far for lead-halide perovskites and are only one order of magnitude
faster than that of III-V emitters (e.g., GaInN, ~2 × 10 -30 cm 6 s −1 ) 39,41,42 . The
notably suppressed Auger recombination in FCT-films is critical for the
mitigated current-efficiency roll-off in the devices. More detailed
discussions are provided in Supplementary Note 1.

Carrier injection kinetics in PeLEDs
Despite the slow-downed Auger recombination in FCT-films with
optical excitation, we understand that Augur losses could happen at a
low carrier density with electrical excitation, especially in the case of
unbalanced charge injections. Given that both devices of interest
employed identical charge transport layers, we focused our investigations on the perovskites. Ultraviolet photoelectron spectroscopy
(UPS) was used to examine changes in the energy levels of the F- and
FCT-films. The results reveal that the introduction of CsTFA causes a
shift in the secondary electron cutoff from 17.25 eV to 17.43 eV (Supplementary Fig. 10a). Further calculations indicate that the valence
band maximum (VBM) of the perovskite shifts from −5.46 eV to
−5.26 eV. The resulting flat-band energy level diagrams are illustrated
in Supplementary Fig. 10b, which display that the VBM of the perovskite becomes more closely aligned with the highest occupied
molecular orbital (HOMO) level of TFB. Given that the electron
mobility of the ZnO electron transport layer is typically much higher

than that of TFB, the improved hole injection facilitates more balanced
charge transport.

Another critical observation we noted is the significant difference
in morphology between the F- and FCT-films. As illustrated in the SEM
and atomic force microscopy (AFM) topographical images (Fig. 3a, b),
the dense FCT-films exhibit a small root-mean-square (rms) roughness
of 1.67 nm, in stark contrast to the F-films, which display prominent
pinholes and an enhanced rms roughness of 5.42 nm. Achieving such a
uniform thin film in the FCT-cases is very challenging for 3D perovskite
thin-film emitters, considering the ultra-thin-film thickness typically
required to enhance charge carrier confinement for efficient
recombination.

We realize that the notable changes on film morphology may
affect the electrical properties of devices 43 . We thus conducted simulations of the electric potential distribution across simplified device
architectures—one featuring densely packed emissive films and the
other with discrete perovskite grains (Fig. 3c, d) 44 . The simulations
reveal that in both configurations, the electric potential gradient is
concentrated near the hole transport layer (TFB)/anode interface,
diminishing once moving away from the interfaces (region I and II in
Fig. 3c and region III in Fig. 3d). However, it is worthy to mention that
excessive holes accumulate at the HTL–ETL interface once increasing
the bias, leading to additional charge transport pathways and hence
potential electrical shunts (Fig. 3e). Figure 3f displays a further analysis
regarding the quantification of charge carrier losses dependent on the
surface coverage and applied bias. We find that the electrical shunt can
be negligible at low bias even for those with low surface coverage.
However, the loss of charge carriers significantly increases once the
bias exceeds 4.0 V, suggesting an increasing current-efficiency roll-off
over enhancement in bias.

To link the simulation and experimental results, we fabricated
devices with a structure of ITO/PEIE-modified ZnO/TFB/MoO3/Au and
compared their J–V characteristics with those with perovskite emissive
layers (Supplementary Fig. 11). Consistent with our simulation results,
we note that the current densities of the devices without perovskites are
more than one order of magnitude lower at a small bias (<2.0 V), suggesting that the charge injection barrier between TFB and ZnO is large
enough to prevent the shunts at the TFB/ZnO interfaces for PeLED
devices. In addition, the current density of the ZnO/TFB-only device
quickly rises with increasing the bias and eventually reaches the same
order of magnitude as the perovskite device at high voltages. These
results align with our simulation results that the electrical shunts can be
very severe at the high bias necessary to achieve high radiance. These
findings demonstrate that dense film morphology enables minimized
detrimental effects such as charge accumulation and electrical shunts,
thereby mitigating the issues of current-efficiency roll-off 45,46 .

To confirm and generalize our findings, we investigated two
additional types of devices with controlled surface coverages of perovskites. One approach was to modify the morphology by varying the
quantity of excess FAI into FAPbI3, with the molar ratios of FAI to PbI2
ranging from 2.0:1 to 2.8:1. The other case is to introduce DMF vapour
atmosphere during thermal annealing to control the grain growth,
using a fixed FAI to PbI2 ratio of 2.0:1 47 . In brief, we find that the films
with high surface coverage commonly give rise to enhanced performance at high current densities, thereby supporting above analysis
that the dense film morphology is crucial. The corresponding device
characteristics and film morphologies are detailed in Supplementary
Figs. 12 and 13, and a more thorough description is available in Supplementary Note 2.

Roles of TFA − addition

Having concluded that the decent performance of FCT-devices at high
current densities mainly stems from the slowed-down Auger recombination rates of perovskites as well as balanced charge injection, the
next question that arises is why TFA − addition can achieve these

Fig. 3 | Perovskite film characteristics and device electric simulations. a SEM
images of F- and FCT-films. The scale bar represents 1 μm. b, AFM and 3D AFM
images of F- and FCT-films. The scale bar represents 1 μm. The electric simulation
and simulated electric field profile of devices with low (c) and full coverage (d)

effects. In addition, the observed mitigated trap-assisted recombination caused by TFA − is possible to be correlated with the prolonged
operational lifetime. To understand the underlying mechanisms,
probing the TFA − interaction with perovskites is crucial. We conducted
X-ray photoelectron spectroscopy (XPS) and time-of-flight secondary
ion mass spectrometry (ToF-SIMS) measurements, aiming to verify the
presence of TFA − within the perovskite films and investigate their
distribution perpendicular to the substrate. Different from the F-films,
the core-level spectra of C 1s and F 1s, as depicted in XPS images
(Fig. 4a), clearly reveal the distinct features corresponding to −O−C=O
and −CF3 functional groups, confirming the presence of TFA − on the
surface of the perovskite films. In addition, ToF-SIMS measurements
show that the TFA − ions are evenly distributed throughout the perovskite films (Supplementary Fig. 14). These results not only confirm
the incorporation of TFA − into the perovskite films but also help to
establish the connection to the modified recombination dynamics and
film morphology.

Previous reports in II-IV semiconductor nanocrystals suggest that
a decouple of wave-function overlap between electrons and holes
usually gives rise to a decreased Auger process. Such an effect can be
achieved by either constructing a type II core-shell heterojunction or a
careful design of organic ligands on the surface 48–50 . Given that TFA − is
a strong and well-known electron-withdrawing moiety, we used density functional theory (DFT) calculation to unveil the changes in electron cloud density of crystal surface with and without TFA adsorption.
We examined the surface lattice structures of FAPbI3 and its variant,
where one Cs + cation substitutes FA + (denoted as FAxCs1-xPbI3) for
comparison (Fig. 4b). As expected, the strong electron-withdrawing
capability of −CF3 moiety leads to an electron polarization and thus

perovskite films (the color represents the electric potential gradient distribution
inside of the devices). e Simulated holes distribution after carrier injection in
PeLEDs at the bias of 6 V. f Dependence of loss ratio on bias and surface coverage.

results in the reduction of electron cloud density in the lattice. This
redistribution results in a decrease in wave-function overlap, reducing
Coulomb electron-hole interaction and thereby decreasing Auger
recombination. Our analysis on temperature-dependent PL enabled us
to determine the exciton binding energies (Supplementary Fig. 15),
which are 44 meV for F-films and about 28 meV for FCT-films, respectively. The decrease in exciton binding energy often leads to a retarded
Auger process due to the changes in Coulomb interactions within the
materials. Additionally, the electron localization function (ELF) and the
Bader charge analysis (Supplementary Fig. 16) of the samples suggests
that the atoms of the FAPbI3 loss electrons in the presence of TFA. This
obviously changed electron distribution of lattices is attributed to
large electronegativity of the dangling F atoms in the TFA.

Having revealed the underlying reasons behind the decreased
Auger recombination rates, we proceeded to investigate the origins of
morphological variation, which is commonly associated with a change
of crystallization process. We then performed 1 H nuclear magnetic
resonance ( 1 H NMR) and Fourier-transform infrared spectroscopy
(FT-IR) measurements, aiming to gather more information about the
chemical interactions between perovskite precursors and TFA − anions,
which could be signs of the changes in the formation of intermediates
and thus varied crystallization process. In Fig. 4c, we present the
changes in the chemical shifts of the active protons in FA + upon
the addition of CsTFA (highlighted by the red square). For both
cases, the chemical shifts moved toward the low field after TFA −

addition in the respective systems (FAI/CsTFA and PbI2/FAI/CsTFA),
confirming the interactions between FA + cations and TFA − anions. This
can be assigned to the ionic interactions between positively charged
FA + and negatively charged TFA −, or additional hydrogen bonding

Fig. 4 | Mechanisms of TFA − anions on improving the PeLED performance.
a Core-level spectra of F 1 s and C 1 s obtained from high-resolution XPS of F- and
FCT-films. b Lattice structures and corresponding electron cloud density for
FAPbI3, FA1-xCsxPbI3, and FA1-xCsxPbI3 with TFA − adsorption, respectively. The
isosurface is 0.1 eV/Å 3 . c 1 H NMR spectra of different materials dissolved in DMSOd6. The red and grey squares highlight the characteristic peaks of FA + . d ATR-FTIR
spectroscopy data for CsTFA and CsTFA−PbI2 samples (with molar ratio of 1:1).

interactions. In addition, the FT-IR results showed that the C=O
vibration of TFA − shifts to a lower wavenumber by mixing with PbI2,
accompanied by the shift of the C–F vibration frequency to a higher
wavenumber (Fig. 4d). All these observations demonstrated the
interactions of TFA − anions with both FA + and Pb 2+ cations 35 .

We found that the additional interactions of TFA − with perovskite
precursors not only facilitate the perovskite formation but also
homogenize the grain growth. The former isevident from the presence
of 3D perovskites within the precursor films (before thermal annealing), as their diffraction (Fig. 4e) and absorption features can be clearly
identified by XRD and UV-Vis absorption measurements (Supplementary Fig. 17). In comparison, films without TFA − displayed no prominent
perovskite signals, despite the presence of PbI2 signal. In the insets of
Fig. 4e, the digital images visually capture the differences between
F- and FCT-precursor films, highlighting the effect of TFA − addition on
facilitating perovskite nucleation. We attribute these discrepancies to
the reduced formation of lead-Lewis base adducts with TFA − addition,
that is, Pb• DMSO in the current case; Pb• DMSO can slow down perovskite crystallization as they require extra energy to release the

e XRD of the F- and FCT-films before thermal annealing. The insert are the photographs of the perovskite films before and after thermal annealing.
f Determination of desorption energy (Ed) of the I − on different perovskite surface:
With iodide vacancies (I), prefect structures (II), and passivated with TFA − anions.
g ToF-SIMS characteristics of F- and FCT-devices before and after electrical aging.
Here, the red circle highlights the increased accumulation of I − ions close to the
anode for F- devices.

coordinating solvents 35 . This speculation can be further evident from
FT-IR spectra that the S=O stretching vibration of FAI•PbI2•DMSO
complexes shifts to a higher wavenumber once incorporating TFA −

anions, confirming a reduced formation of intermediate
FAI•PbI2•DMSO adducts (Supplementary Fig. 18). Additionally, our
dynamic light scattering (DLS) analysis revealed that TFA − addition
reduces the size of the halide plumbate colloids, from an average size
of 164 nm down to 105 nm (Supplementary Fig. 19) in precursor solutions. This suggests a more uniform nucleation process and a lower
total free energy during perovskite phase formation. Taken together,
these observations consistently point to a faster and more uniform
crystallization process with TFA addition, which is usually beneficial
for formation of dense and smooth perovskite films.

Interestingly, although fast crystallization of perovskites commonly leads to high densities of defects 35, we still observed reduced
trap-assisted non-radiative recombination rates in FCT-films (Fig. 2g).
This contrast suggests an effective capability of TFA − in defect
passivation 35,36 . We thus investigated the adsorption geometry of TFA −

on vacancy-involved crystal lattice by DFT calculations. Our

calculations indicate that TFA − anions tend to adsorb on the surface of
lattice instead of being in the crystals (Supplementary Fig. 20), as a
result of their large ionic radius (2.38 Å). Notably, the adsorption
energy (Eads) of TFA − anion on perovskite crystals (exposed lead
cations) is as high as 4.55 eV, much larger than that of the iodide anions
(~3.11 eV). This difference indicates a strong tendency of TFA − to work
with defect involved perovskite surfaces, and hence the passivation
effect can be rationalized.

One consensus is that the reduced formation of vacancy-type
defects contributes to the stabilization of ions within perovskite
structures. To establish the correlation between ion movement and
device stability, we calculated the desorption energy (Ed) of surface
iodides (I − ) in different local environment (Fig. 4f). A lower Ed indicates
facilitated ion migration and thus accelerated device degradation, and
vice versa. For pristine crystals and those with VI, the Ed is 1.01 eV and
1.35 eV, respectively, suggesting that iodide anions adjacent to the
defects are more likely to migrate out of the lattice. In contrast, the Ed
values significantly increase up to 2.28 eV for the crystal lattice healed
by TFA −, indicative of a greatly enhanced difficulty for neighboring
iodides to become mobile 29 .

To experimentally confirm the role of TFA − in suppressing ion
migration, we investigated the ion distribution in fresh and aged
devices by TOF-SIMS measurements. The aged devices were measured
at a fixed current density of 100 mA cm −2 for 10 h ahead of tests. The
comprehensive TOF-SIMS results for the devices of interest are shown
in Supplementary Fig. 21. Figure 4g displays the evolution of iodide
distribution in F- (upper panel) and FCT- devices (bottom panel). The
former exhibits a distinct accumulation of iodides at the interface
between HTL and the contacts, whereas the FCT-devices don’t show a
substantial variation in iodide distribution after aging. These results
are consistent with the previous report that halide migration toward
the anode is the predominant factor limiting the operational lifetime of
PeLEDs 51 .

Discussion
In summary, we achieve highly bright and one of most stable perovskite light-emitting diodes reported to date, overcoming significant
limitations of current-efficiency roll-off at high current densities. This
breakthrough is achieved by introducing trifluoroacetate anions
(TFA − ) into the three-dimensional structure of perovskites. The most
critical effect observed is a substantial reduction in Auger recombination, primarily attributed to decreased electron-hole wave-function
overlap. Moreover, we demonstrate that dense and smooth perovskite
films are crucial for balancing charge injection at high electrical excitation and thus high brightness. Further studies reveal that the use of
TFA − anions in perovskites suppresses halide migration and hence
contributes to decent device stability. Our demonstration of highperformance PeLEDs under intense electrical excitation opens possibilities for their use in high-power applications including the development of perovskite laser diodes.

Methods
Materials

Formamidinium iodide (FAI) and 5-ammonium valeric acid iodide
(5AVAI) were bought from Greatcell solar. Cesium trifluoroacetate
(CsTFA) was purchased from Alfa Aesar. Other chemicals were
obtained from Sigma–Aldrich. All materials were used without
extra treatments. Zinc oxide (ZnO) was synthesized based on
previous reports 2,7–9 . 2.8 mmol tetramethylammonium hydroxide
pentahydrate (TMAH·5H2O) in 5 mL ethanol was added into
1.5 mmol zinc acetate hydrate (Zn(Ac)2·2H2O) in 15 mL dimethyl
sulfoxide (DMSO). Then, the solution was stirred at 40 °C for 3 h
to form ZnO nanoparticles. The obtained ZnO was washed with
ethanol and ethyl acetate for 2 times, and finally dispersed in
ethanol (8 mL).

Perovskite precursors
0.15 mmol CsTFA, FAI, lead iodide (PbI2), and 5AVAI
(n(CsTFA):n(FAI):n(PbI2):n(5AVAI) = 0.12:1.28:1:0.1) were dissolved in
N, N-dimethylformamide:dimethyl sulfoxide (DMF:DMSO = 9:1) mixture to form FCT-precursors. 0.15 mmol FAI, PbI2, and 5AVAI were
dissolved in DMF:DMSO (9:1) mixture at a molar ratio of 1.4:1:0.3 to
form F-precursors. 0.13 mmol FAI, PbI2, and 5AVAI are dissolved in
DMF at a molar ratio of (2.8, 2.4, or 2):1:0.15 to form FAI-2.8, FAI-2.4,
and FAI-2.0 based precursors. The precursor solution was stirred
overnight and filtered with polytetrafluoroethylene filters (0.22 μm)
before use.

Device fabrication

ZnO was spin-coated onto the precleaned indium tin oxide (ITO) substrates at 5000 r.p.m for 30 s, and annealed at 120 °C for 10 min in air.
Polyethylenimine ethoxylated (PEIE) in isopropanol (1.5 mg mL −1 ) was
spun on ZnO at 5000 r.p.m. for 30 s, and annealed at 100 °C for 10 min.
Then, the perovskite precursor was deposited onto the ZnO/PEIE layers
at 500 r.p.m. for 3 s and 6000 r.p.m. for 30 s, respectively. The FCT- and
F-precursor films were annealed for 15 min at 90 °C and 100 °C,
respectively. After cooling, phenethylammonium iodide (PEAI,
0.25 mg mL −1 in isopropanol) was dynamically spun on the perovskite
layer at 4000 r.p.m for 30 s. Next, poly(9,9-dioctyl-fluorene-co-N-(4butylphenyl)diphenylamine) (TFB, 12 mg mL −1 in chlorobenzene) was
spun at 3000 r.p.m. for 45 s. Finally, 6 nm of molybdenum oxide (MoO3)
and 100 nm of gold (Au) were sequentially deposited by thermal evaporation. For optimized FAI-2.8, FAI-2.4, and FAI-2.0 based devices: ZnO
was spin-coated onto ITO at 4000 r.p.m for 30 s without annealing.
Then, PEIE (1.2 mg mL −1 in IPA) was deposited on ZnO at 5000 r.p.m. for
30 s, and annealed at 100 °C for 10 min. The perovskite precursor was
spun onto ZnO/PEIE at 500 r.p.m. for 3 s and 4000 r.p.m. for 30 s,
respectively. The precursor film was annealed at 100 °C for 10 min.

Characterizations

PeLEDs measurements were conducted using a Keithley 2400 source
meter and an integration sphere equipped with a QE Pro spectrometer.
The stability test was taken on the same test platform in a nitrogenfilled glovebox with a constant room temperature of 20 °C. SEM was
measured using a JEOL JSM-7500F microscope. AFM was performed
using a Bruker Dimension Icon microscope. FTIR was tested using a
commercial FTIR testing equipment. XRD was tested using a X-ray
diffractometer (Panalytical X’Pert Pro) with an X-ray tube (Cu Kα,
λ = 1.5406 Å). XPS was measured using a Scienta ESCA 200 spectrometer with a monochromatic Al (Kɑ) X-ray source. ToF-SIMS tests were
conducted using a ToF-SIMS.5 instrument from IONTOF, operated in
the spectral mode using a 25 keV Bi 3+ primary ion beam with an ion
current of 0.78 pA. 1 H NMR was collected using a Bruker Ultra Shield
Plus 400 MHz NMR system. PLQYs was recorded using a Quanta-Phi
integrating sphere with a Fluorolog system under the excitation
wavelength of 400 nm. TA measurements were performed using a
Helios setup. The transient dynamics in fs-ns time region (50 fs - 7 ns)
was acquired by Helios that works in a nondegenerate pump–probe
configuration. The pump pulses were generated from an optical
parametric amplifier (OPerA Solo) that was pumped by a 1-kHz
regenerative amplifier (Coherent Libra, 800 nm, 50 fs, 4 mJ). A mode
lock Ti-sapphire oscillator (Coherent Vitesse, 100 fs, 80 MHz) was used
to seed the amplifier. The probe pulse was a white light continuum
generated by passing the 800 nm fs pulses through a 2 mm sapphire
plate for visible part (420–780 nm).

Device simulation

Perovskite devices were modelled with the software of COMSOL
Multiphysics 6.0 and the corresponding Semiconductor Module, the
main method of which is based on solving the electrostatic equation
and the drift–diffusion equations for electrons and holes. 2D

axisymmetric model was used to evaluate the influence of surface
coverage in 3D situation by reducing the calculation cost in the 3D
model. Electrical potentials were calculated through Poisson equation
by taking both carrier dynamics (recombination and transport) and
electrostatics into consideration. The detailed model is based on the
as-mentioned device consisting of ITO/ZnO/PEIE/perovskite/TFB/
MoO3/Au, where TFB and PEIE-modified ZnO are defined as HTLs and
ETLs, respectively. The coverage ratio is defined as the area ratio of
Region I (perovskite area) and Region I + Region II (perovskite area and
transport contact) while the loss ratio is taken as the area-averaged
product ratio of electrons and holes in the area without perovskite
(Region II in Fig. 3c) to that in the region with perovskite coverage
(Region I in Fig. 3c) to estimate the possible loss at the interface
through the charge accumulation. The loss ratio can also be affected
by the specific interfacial recombination coefficients, and it is not
included in the discussion because of our focus on the impact of
perovskite coverage on the device performance.

~~R~~

region I n × pÞ=arearegion I

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Loss Ratio =

ð region II n × pÞ=arearegion II

~~R~~

ð region I n × pÞ=arearegion I

× 100% ð2Þ

First-principles calculation

DFT calculations were conducted in the framework of the density
functional theory using CP2K package (version 7.1) with a plane-wave
density cutoff of 500 Ry. The DFT-D3 approach was employed to
account for the electronic exchange-correlation interaction in the
presence of the organic molecules. CP2K with BFGS scheme was
employed to fully relax the crystals structures, and the force convergence criterion was settled to 4.5 × 10 −4 hartree/bhor. GoedeckerTeter-Hutter (GTH) pseudopotentials were used to describe atomic
core electrons, and the valence electron orbitals were expanded into
DZVP-MOLOPT-SR-GTH basis sets. For the geometry optimization, the
system multiplicity was relaxed until converged.

Data availability
The published article includes all data analyzed and necessary to draw
the conclusions of this study in the figures and tables of the main text
and Supplemental Information. Further information and requests
should be directed to the corresponding authors.

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Acknowledgements
This work was supported by the National Natural Science Foundation of
China (Grants 52250060, 62274135, 62288102, 52302167, 62175048),
Key project of Ningbo Natural Science Foundation (Grants
20221JCGY01049). F.G. is a Wallenberg Scholar and acknowledges the
financial support from the Swedish Strategic Research Foundation
(SIP21-0151), the European Research Council Consolidator Grant (LEAP,
101045098), the Olle Engkvists Stiftelse, and the Swedish Government

Strategic Research Area in Materials Science on Functional Materials at
Linköping University (faculty grant SFO-Mat-LiU no. 2009-00971). G. Li
acknowledges the financial support from the Research Grants Council
of Hong Kong (GRF Grant Nos. 15221320, CRF C7018-20G, C4005-22Y),
the Shenzhen Science and Technology Innovation Commission (Project
No. JCYJ 20200109105003940), the Hong Kong Innovation and Technology Commission (GHP/205/20SZ), the Hong Kong Polytechnic University (the Sir Sze-yuen Chung Endowed Professorship Fund (8-8480),
PRI strategic Grant (1-CD7X), RISE Strategic Grant (Q-CDBK)). M. Li
acknowledges the financial support from the Research Grants Council
of Hong Kong (Project No. 25301522, 15301323, 15300824, C5003-24E),
National Natural Science Foundation of China (22373081).

Author contributions
F. G. and W. X. conceived the idea and supervised the project; Z. L., Q.
W., and W. G. performed the experiments and analyzed the data; Y. W.
performed device simulation and analyzed the data; Q. W. performed
transient absorption under the supervision of M. L.; C. T., Y.Z. and Z. W. L.
contributed to the data analysis. G. L. supervised Z. L. and contributed to
the interpretation of results; W. X., F. G., and Z. L. wrote the manuscript;
X. L. and Z. W. provided revisions to the manuscript. All authors discussed the results and commented on the manuscript.

Funding
Open access funding provided by Linköping University.

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/s41467-025-56001-x.](https://doi.org/10.1038/s41467-025-56001-x)

Correspondence and requests for materials should be addressed to
Mingjie Li, Wenbin Guo, Gang Li, Weidong Xu or Feng Gao.

Peer review information Nature Communications thanks Lianfeng Zhao
and the other anonymous reviewer(s) for their contribution to the peer
review of this work. A peer review file is available.

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