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Article
Carrier Mobility Enhancement in (121)-Oriented CsPbBr3 Perovskite
Films Induced by the Microstructure Tailoring of PbBr2 Precursor
Films
Xiaopeng Han, Xin Wang, Jianyong Feng, Huiting Huang, Zhi Zhu, Tao Yu,* Zhaosheng Li,
and Zhigang Zou
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ABSTRACT: It is important to prepare preferentially oriented films to enhance charge
carrier transport in the optoelectronic device applications. For the promising
optoelectronic material of CsPbBr3, obtaining its thin films with preferred crystal
orientation is highly desirable yet challenging. Herein, (121)-oriented CsPbBr3
perovskite films were successfully obtained by using HBr as the additive for PbBr2
precursor solution in the two-step solution method. Detailed investigations indicate that
microstructure tailoring of PbBr2 films via HBr additives plays a crucial role in achieving
(121)-oriented CsPbBr3 films. Theoretical calculations and experimental measurements
demonstrate high carrier mobility in (121)-oriented CsPbBr3 films, which accords well
with photovoltaic tests that the (121)-oriented CsPbBr3 film shows short-circuit
photocurrent density as much as 1.68 times than the (101)-oriented one. In comparison with the (101)-oriented CsPbBr3 solar cell,
the champion power conversion efficiency of the (121)-oriented CsPbBr3 solar cell increases from 2.56 to 6.91% owing possibly to
its higher coverage and carrier mobility. This work not only develops a pathway to prepare compact (121)-oriented CsPbBr3 films
but also highlights the importance of crystal orientation engineering in perovskite films for high-performance optoelectronic devices.
KEYWORDS: oriented CsPbBr3 films, PbBr2 precursor films, microstructure tailoring, carrier mobility enhancement, perovskite solar cells
1. INTRODUCTION
The crystal anisotropy originates from the different arrangements of atoms in the three-dimensional space, which makes
the physical and chemical properties of crystal materials vary
along different directions and across different surfaces.1 The
properties such as piezoelectricity,2 optoelectronics,3−5 thermal
conductivity,6 magnetics,7 catalysis, and so forth8,9 have been
proved to depend on the orientation of crystals. A well
understanding of crystal anisotropy and an ability to control
and predict crystal anisotropy are mostly subjects of interest
for researchers.10,11 In the field of optoelectronic device
applications, by affecting optical absorption, carrier mobility,
trap densities, and so forth, crystal anisotropy may contribute
significantly to device performance. Therefore, a comprehensive study of the relationship between crystal anisotropy and
optoelectronic properties is necessary to explore high-performance optoelectronic materials.
Metal halide perovskites are ideal candidates for a myriad of
applications as a result of their unique optoelectronic
properties, including large absorption coefficient, high
mobility, and long carrier diffusion length.12−24 Previous
studies have demonstrated that metal halide perovskite
materials, especially MAPbX3 (MA = CH3NH3+, X = Cl−,
Br− and I−),1,25,26 exhibit facet-dependent optoelectronic
properties. The influence of crystal anisotropy on optoelec© 2020 American Chemical Society
tronic and photovoltaic (PV) characteristics has recently raised
concerns in oriented MAPbX3 films and MAPbX3 single
crystals.5,27−30 For example, Cho et al. prepared MAPbI3
polycrystalline films with (112) and (220) orientations and
found that the charge transport was highly anisotropic.31 Bae et
al. fabricated MAPbI3 crystals with two different orientations to
the substrate; they demonstrated that crystal orientationinduced anisotropy of charge transfer led to significantly varied
PV performance.32 Zuo et al. reported that perovskite
photodetectors based on (110) plane-oriented MAPbBr3
single crystals showed a 153.33% enhancement of responsivity
compared to the device with the (100) orientation.33
Therefore, besides the crystallinity optimization and morphology tailoring strategies, crystal engineering via orientation
control is also an effective means to improve the performance
of optoelectronic devices.
CsPbBr3 can be applied in a variety of optoelectronic devices
because of its excellent optoelectronic properties and superior
Received: October 13, 2020
Accepted: December 8, 2020
Published: December 21, 2020
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moisture, oxygen, as well as thermal stability.34−38 Because
CsPbBr3 crystallizes in the orthorhombic system (Pnma space
group) at room temperature,39−45 a high dependence of
optoelectronic characteristics on crystal anisotropy in CsPbBr3
is expected because of its dissymmetry of the crystal structure.
However, until now, investigations about the anisotropy of
optoelectronic properties and their corresponding impacts on
PV behaviors for CsPbBr3 are still lacking.46,47 One possible
reason for such situation is the lack of suitable methods to
fabricate CsPbBr3 perovskite films with well-controlled crystal
orientations; the lack of a deep knowledge on the electronic
band structure and carrier mobility of CsPbBr3 could be the
other reason. Considering its importance in the optoelectronic
field and the urgent need for future performance improvement,
the relationship between crystal anisotropy and optoelectronic
properties in CsPbBr3 should not be neglected.
In this study, the (121)-oriented CsPbBr3 perovskite film
was successfully obtained by controlling the microstructure of
PbBr2 films in the two-step solution method. The amount of
HBr additives plays a crucial role in the microstructure
tailoring of PbBr2 precursor films. Then, based on morphology
and phase characterizations of PbBr2 and CsPbBr3 films, a
possible formation mechanism of (121)-oriented CsPbBr3 film
was proposed. The impact of crystal orientation on the PV
performances of CsPbBr3 perovskite solar cells (PSCs) was
probed by assembling (121)-oriented and (101)-oriented
CsPbBr3 devices, demonstrating the direct relation between
crystal orientation and PV performance in CsPbBr3 films.
Theoretical calculations and the space-charge limited current
(SCLC) measurements further reveal a high anisotropy of
carrier mobility in CsPbBr3, which explain the orientationdependent optoelectronic characteristics in CsPbBr3. The
present study highlights the importance of crystal orientation
engineering in perovskite films for high-performance optoelectronic devices and more importantly offers an effective means
to control and investigate the anisotropy of optoelectronic
properties in CsPbBr3.
Next, PbBr2 and x-HBr (x = 0, 10, 30, and 50 μL, respectively) in
DMF (1 M) solution was spin-coated on the FTO substrates at 2000
rpm for 30 s. After annealing at 75 °C for 30 min on a hot plate, the
as-obtained PbBr2 precursor film was immersed in a CsBr 2methoxyethanol solution of 15 mg mL−1 for 10 min at room
temperature. The resulting samples were rinsed with isopropanol,
dried in air, and annealed at 250 °C for 5 min on a hot plate to obtain
the CsPbBr3 perovskite films. For simplicity, the as-obtained CsPbBr3
films are abbreviated as HBr-0, HBr-10, HBr-30, and HBr-50,
respectively, according to the added amount of HBr. For example, the
HBr-30 sample was derived from the PbBr2 precursor solution (1 mL)
with 30 μL of added HBr. Finally, a carbon back-electrode with an
average area of 0.09 cm2 was deposited on the CsPbBr3 perovskite
film by a doctor-blade coating method, followed by drying at 120 °C
for 10 min.
2.3. Characterizations. The X-ray diffraction (XRD) patterns of
the as-prepared films were measured by the X-ray diffractometer
(Rigaku Ultima III) with Cu Kα radiation at a scan rate of 10° min−1.
The morphologies of the as-obtained films were observed using a
field-emission scanning electron microscope (FE-SEM, Nova NanoSEM230) at an operating voltage of 15 kV. The cross-section image
was captured by the Zeiss GeminiSEM 500 ultrahigh-resolution FESEM equipped with an Oxford EDS system. The absorption spectra of
the CsPbBr3 films on the c-TiO2/FTO substrates were obtained on an
UV−visible−near-infrared (UV−Vis−NIR) spectrophotometer (PerkinElmer, Lambda 950). An Oriel 92251A-1000 sunlight simulator
was used to provide the testing light AM 1.5G (100 mW cm−2). The
light intensity was calibrated at 100 mW cm−2 by the standard
reference of a Newport 91150V silicon cell before use. The electron
trap densities and mobility of the as-obtained CsPbBr3 films were
calculated by constructing the electron-only devices of FTO/c-TiO2/
perovskite/PCBM/C (PCBM, [6,6]-phenyl-C61-butyric acid methyl
ester) and measuring the current−potential curves from 0 to 3.5 V of
applied electric bias. In all of the current−voltage (J ≈ Vn) curves, the
Ohmic (n = 1), trap-filling (n > 3), and child (n = 2) regions were
observed with the increase of bias voltages. When operating in the
trap-free SCLC regime (n = 2), the dark current was well fitted by the
Mott−Gurney law: J = 9εε0μV2/8L3, where J is the current density, ε
is the relative dielectric constant, ε0 is vacuum permittivity, μ is the
charge mobility, V is the applied voltage, and L is the thickness of the
as-obtained CsPbBr3 films.
2.4. Computational Methods. The density functional theory
(DFT) calculations in this study are performed by VASP (5.2) with
the projected-augmented wave (PAW) method. The generalized
gradient approximation (GGA) in the scheme of Perdew−Burke−
Ernzerhof (PBE) is adopted for the exchange correlation functional.
The cut-off energy is 500 eV. Geometry relaxations are performed
until the residual forces on each ion converged to be smaller than 0.02
eV Å−1. All atomic models are visualized by the VESTA (3.3.1)
software.
2. EXPERIMENTAL SECTION
2.1. Materials. Lead bromide (PbBr2, ≥99.99%) and cesium
bromide (CsBr, ≥99.99%) were purchased from Xi’an Polymer Light
Technology Corp. Anhydrous N,N-dimethylformamide (DMF,
≥99.8%) was provided by Sigma-Aldrich. Anhydrous ethanol
(≥99.7%), isopropanol (≥98.5%), acetone (≥99.0%), and 2methoxyethanol (≥99.0%) were purchased from Sinopharm Chemical
Reagent Co., Ltd. Hydrobromic acid (HBr, 48% in water) was bought
from Aladdin Reagents (Shanghai) Co., Ltd. Conductive carbon paste
(ZF-G03-04-2) was supplied by Shanghai MaterWin New Materials
Co., Ltd. Fluorine-doped tin oxide (FTO) glass substrates (TEC A7,
8 Ω sq−1) was purchased from You Xuan Tech. As presented, all the
purchased precursor materials and solvents were used as received with
no further purification.
2.2. Fabrication of Solar Cells. FTO substrates (2 cm × 2 cm)
were cleaned sequentially by detergent, deionized water, acetone,
isopropanol, and anhydrous ethanol. Prior to use, dried FTO
substrates were treated by UV−ozone for 30 min. Then, a compact
TiO2 (c-TiO2) layer was deposited onto the FTO substrate by spincoating an ethanol solution of titanium butoxide (ca. 0.2 M) at 4500
rpm for 30 s. The annealing process was carried out in the muffle
furnace at 480 °C for 40 min.
As illustrated in Figure S1, the CsPbBr3 perovskite film was
fabricated by a two-step solution method. PbBr2 precursor solution
was prepared by dissolving 367 mg of PbBr2 in 1 mL of DMF while
stirring at 75 °C. Then, different amounts of HBr (0, 10, 30, and 50
μL, respectively) were introduced above PbBr2 precursor solution.
Article
3. RESULTS AND DISCUSSION
3.1. Morphology and Crystallographic Characterizations of CsPbBr3 Films. The morphological characteristics of obtained films were measured by SEM. As shown in
Figure 1, distinctive morphological features on CsPbBr3 films
were obtained when varying the amounts of the HBr additive.
For the HBr-0 film, considerable voids exist and seriously affect
the surface coverage of the CsPbBr3 layer. No obvious
morphological variation was observed on HBr-10 when
comparing with HBr-0, indicating almost unchanged growth
manner of the CsPbBr3 film with a small amount of the HBr
additive (10 μL). However, upon increasing the HBr additive
to 30 μL, the resulting HBr-30 film exhibits a smooth surface
with significantly improved surface coverage. This phenomenon suggests a possible change of the growth environment of
CsPbBr3 films by the introduction of the HBr additive. Further
increasing HBr additive to 50 μL, aggregates of cube-like
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XRD patterns (Figure 2a), the crystal structure of obtained
films can be assigned to orthorhombic CsPbBr3, with the space
group of Pnma. Specifically, XRD diffraction peaks at 15.2,
21.6, and 30.8° correspond to the (101), (121), and (202)
lattice planes of orthorhombic CsPbBr3, respectively.48
Interestingly, the relative intensities of these dominant
diffraction peaks vary significantly among different CsPbBr3
films. Using the (110) diffraction peak of FTO at 26.4° as the
internal standard, the peak intensity evolution of (101) and
(121) planes is analyzed. As presented in Figure 2b, HBr-0 and
HBr-10 films show high I(101)/IFTO values, while their I(121)/
IFTO values are greatly suppressed. In contrast, a decreased
I(101)/IFTO along with greatly enhanced I(121)/IFTO is observed
on HBr-30, indicating a possible variation of the crystal
orientation in it compared to HBr-0 and HBr-10 films. For the
HBr-50 film, the intensity ratios of its (101) and (121) peaks
relative to FTO are instead quite close to each other. Similarly,
the variation of characteristic peaks for CsPbBr3 films, as a
function of added volumes of HBr, can also be assessed by
comparing the integrated intensity ratio of the (101) and
(121) planes to the total dominant planes of CsPbBr3 (Figure
2c). This approach yields an identical trend of peak evolution
as Figure 2b. Therefore, judged from the areas and intensities
of diffraction peaks on as-derived CsPbBr3 films, the
introduction of the HBr additive is able to influence the
crystal growth and orientation in CsPbBr3 films.
To present a more quantitative analysis on the orientation
variation of CsPbBr3 films upon the use of the HBr additive,
the orientation index is introduced as defined by eq 1
Figure 1. SEM images of samples HBr-0, HBr-10, HBr-30, and HBr50, respectively. The as-obtained CsPbBr3 films are abbreviated as
HBr-0, HBr-10, HBr-30, and HBr-50, respectively, according to the
added amount of HBr. For example, the HBr-30 sample was derived
from the PbBr2 precursor solution (1 mL) with 30 μL of added HBr.
crystals dominate and lead to a quite rough surface on HBr-50.
The appearance of these cube-like crystals and their order
assembly imply that the crystal orientations in the CsPbBr3
films have been altered upon the application of the HBr
additive.
In order to get more insights into the impacts of the HBr
additive on the crystallographic behaviors of obtained CsPbBr3
films, XRD characterization was performed. On the basis of
Figure 2. (a) XRD patterns of samples HBr-0, HBr-10, HBr-30, and HBr-50, respectively. (b) Corresponding line chart of XRD peak intensity
ratios of the CsPbBr3 (101) and (121) to FTO substrate (110) planes. (c) Line chart of the XRD peak area ratio of (101) and (121) to the total
peak area [(101), (121), and (202) planes] of the CsPbBr3 sample. (d) Line chart of orientation indexes for CsPbBr3 samples. The as-obtained
CsPbBr3 films are abbreviated as HBr-0, HBr-10, HBr-30, and HBr-50, respectively, according to the added amount of HBr. For example, the HBr30 sample was derived from the PbBr2 precursor solution (1 mL) with 30 μL of added HBr.
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Figure 3. (a) Comparison of crystal structure and Pb−Br coordination number between the PbBr2 and CsPbBr3 crystal. (b) XRD patterns of PbBr2
films prepared with different addition amounts of HBr. (c) XRD patterns of 75 and 100 °C-annealed PbBr2 films without HBr, as well as 75 °Cannealed PbBr2 films with 30 μL of added HBr. (d) XRD patterns of CsPbBr3 films prepared with above obtained PbBr2 films (without HBr at 75
°C and 100 °C, respectively, and with 30 μL HBr at 75 °C). (e) Top-view SEM images of as-obtained PbBr2 films (without HBr at 75 and 100 °C,
respectively, and with 30 μL HBr at 75 °C).
IO(h′k′l′) =
Is(h′k′l′)
∑ Is(hkl)
orientation index values of the (101) diffraction peak, the
magnitude of (101) orientation in HBr-50 (1.70) is not as high
as HBr-0 (2.21) and HBr-10 (2.51). In sharp contrast, a
preferred (121) orientation is observed in HBr-30, as
manifested by its high IO(121) of 1.55 and suppressed
IO(101) of 0.68. Consequently, via controlling the amount of
added HBr in the PbBr2 precursor, the crystal orientation in
CsPbBr3 films can be well tuned, which provides a unique
platform for the investigation of the relationship between
crystal anisotropy and optoelectronic properties in CsPbBr3
films.
IJ(h′k′l′)
∑ IJ(hkl)
(1)
where IO(hḱ ĺ )́ is the orientation index of the (hḱ ĺ )́ plane, IS is
the peak intensity of the obtained CsPbBr3 sample, and IJ is the
peak intensity of the simulated XRD pattern for orthorhombic
CsPbBr3 (Figure S2).49−51 The orientation indexes of obtained
four CsPbBr3 films are summarized in Figure 2d. For HBr-0,
HBr-10, and HBr-50 films, the orientation indexes of (101)
diffraction peak are greater than 1, indicating that a preferred
(101) orientation exists in them; however, by comparing these
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3.2. Investigation on the Formation Mechanism of
Oriented CsPbBr3 Films. To reveal the formation mechanism of (121)-oriented CsPbBr3 perovskite films and understand the roles played by HBr, further investigations are
needed. First, it is well known that the growth of oriented films
is greatly dependent on the crystal structure/orientation of the
substrates,52,53 which implies the importance of selected
substrates, seeds, or reactants in growing crystallographically
oriented samples. By considering the conversion process from
PbBr2 to CsPbBr3, in brief, PbBr2 and Br− ions from CsBr form
the stable PbBr64− cubic frames, and Cs+ ions diffuse into the
interstice position of the cubic frames under the electrostatic
force to form ultimate perovskite CsPbBr3;54 there is a
possibility that the (121)-oriented CsPbBr3 perovskite films are
induced by the specific crystal orientation in the PbBr2
precursor film. Nevertheless, from the crystal structure point
view (Figure 3a), the Pb−Br coordination number in PbBr2
crystal is 7, as compared to 6 in CsPbBr3.55 This implies that
during the conversion from PbBr2 to CsPbBr3, breaking and
rearrangement of the Pb−Br bonds would occur; meanwhile,
no crystallographic similarity/relationship can be found
between PbBr2 and CsPbBr3. These analyses confirm that
the formation of (121)-oriented CsPbBr3 films from PbBr2
precursor films is a nonepitaxial growth process. Therefore, the
material property changes of PbBr2 films upon the addition of
HBr could be the underlying reasons for the formation of
(121)-oriented CsPbBr3 perovskite films.
The crystallographic behaviors of PbBr2 films with different
amounts of HBr were measured by XRD. As shown in Figure
3b, without the HBr additive, the as-derived PbBr2 film shows
weak PbBr2 characteristic peaks along with the intense signal
from FTO, which indicates poor crystallinity of the obtained
PbBr2 film. Similar results are observed on PbBr2 films with 10
and 50 μL of added HBr. However, when the added amount of
HBr is 30 μL, more diffraction peaks belonging to PbBr2
appear in the XRD pattern, and the intensities of these
diffraction peaks are also significantly enhanced. Combining
with the XRD results for CsPbBr3 (Figure 2), a possibility is
that the improved crystallinity of PbBr2 is responsible for the
growth of the (121)-oriented CsPbBr3 perovskite film. To
verify this possibility, a PbBr2 film with improved crystallinity
was prepared by increasing the annealing temperature from 75
to 100 °C. As expected, a higher crystallinity in PbBr2 is
realized, which compares well with the one obtained by 30 μL
of added HBr (Figure 3c). However, immersing the 100 °Cannealed PbBr2 film into CsBr 2-methoxyethanol solution
generates the CsPbBr3 film with poor crystallinity as well as no
sign of preferred orientation (Figure 3d). Therefore, improved
crystallinity of PbBr2 precursor films is not the key parameter
for producing (121)-oriented CsPbBr3 films.
To distinguish the difference between the crystalline PbBr2
obtained at 100 °C and the one achieved by adding 30 μL
HBr, their morphology difference was investigated. As shown
in Figure 3e, the PbBr2 film obtained by annealing at 100 °C
exhibits a smooth surface and limited voids, which is similar to
the morphology of PbBr2 film annealed at 75 °C. However, the
PbBr2 film prepared with 30 μL HBr shows a quite different
microstructure from the above two PbBr2 films annealed at 75
and 100 °C: there are a large number of nanochannels and
voids across the film surface; the film porosity of PbBr2
prepared with 30 μL HBr is thus significantly increased. It is
reasonable to consider that the introduction of HBr in PbBr2
precursor solutions may affect the solvent evaporation and
crystal growth processes of PbBr2 films, thereby leading to
highly porous PbBr2 films. Such a sharp morphology change in
PbBr2 films could contribute to the formation of (121)oriented CsPbBr3 films.
To prove the above hypothesis, the relationship between the
amount of added HBr and the morphologies of as-derived
PbBr2 films was studied. The corresponding top-view and
cross-sectional SEM images, as well as the calculated porosity
of PbBr2 films, are shown in Figures 4 and S3. The PbBr2 film
Article
Figure 4. Top-view and cross-sectional SEM images of PbBr2 films
prepared with (a) 0, (b) 10, (c) 30, and (d) 50 μL HBr, respectively.
fabricated without HBr presents a flat surface with a few voids
(porosity is about 3.0%); its film layer is composed of densely
packed amorphous small particles, as revealed by the crosssectional SEM image. A flat film layer along with more voids is
obtained for the PbBr2 film with 10 μL of added HBr; as
expected, the film porosity increases to 7.6%. Upon increasing
the HBr additive to 30 μL, a large number of nanochannels
appear on the surface of the PbBr2 film, which contribute to
high porosity (15.4%). The cross-sectional image displays that
the abundant nanochannels assembled by small PbBr2 particles
also exist throughout the entire film thickness. During the
conversion of PbBr2 to CsPbBr3 films, these abundant
nanochannels in PbBr2 films would provide sufficient space
for the infiltration of CsBr solutions and subsequent crystal
growth of CsPbBr3 films. When the HBr additive is further
increased to 50 μL, sharply increased film thickness and
roughness are observed in the corresponding PbBr2 film. These
SEM and porosity analyses confirm that via changing the
amounts of added HBr, the morphologies (porosity, particle
size, and particle assembly manner) of PbBr2 films can be well
tuned. The possible mechanism is discussed as follows. The
coordination number of Pb2+ changes from 7 in PbBr2 crystal
to 6 when dissolving in DMF, under which condition Pb2+ ions
are coordinated by 5 Pb2+ ions and 1 O atom of DMF ligands
to form the PbBr2-DMF adduct.55 When HBr is introduced
into PbBr2 DMF solutions, Pb−O bonds between PbBr2 and
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DMF are gradually replaced by Pb−Br bonds, owing to the
stronger coordination capability of Br− than DMF. Consequently, we speculate that the addition of HBr into PbBr2
DMF solutions would lead to the formation of the PbBr2−HBr
complex, which is probably responsible for the formation of
PbBr2 films with high porosity, as illustrated in Scheme 1.
Scheme 2. Illustration of the Growth Process of
Preferentially Oriented CsPbBr3 Filmsa
Article
Scheme 1. Illustration of the Stages in the Fabrication of
PbBr2 Films with Different Amounts of Added HBra
a
The different microstructures of PbBr2 films influence subsequent
growth direction of CsPbBr3, thereby resulting in different kinds of
CsPbBr3 crystal orientation.
growth of CsPbBr3 in all directions on the substrate. On the
basis of crystal growth theory, high energy crystal facets will
disappear rapidly during the crystal growth process, because
the growth rate perpendicular to a high-energy facet is much
faster than other low-energy facets; the preferred orientation of
the high-energy facet is then achieved.59,60 In this regard, for
the growth direction being perpendicular to the substrate
(without HBr), the high surface energy (101) facet will
disappear rapidly, and finally, the preferentially (101)-oriented
CsPbBr3 film forms. In comparison, for the unrestricted growth
process (with 30 μL of added HBr), a high-energy (101) facet
will disappear rapidly in all directions on the substrate, the
(121) facet that exhibits a relative large surface energy [but less
than (101)] then promotes the formation of the (121)oriented CsPbBr3 film. However, Scheme 2 is based on the
assumptions from the results of SEM and XRD measurements;
the exact mechanism for solution-phase growth of the (121)oriented CsPbBr3 film needs further investigation.
3.3. PV Performance and Characterization of Oriented CsPbBr3 Films. The control over the crystal
orientation in CsPbBr3 films enabled us to investigate its
effect on PV behaviors. To do so, the as-obtained CsPbBr3
films with different orientations were assembled with the
configuration of FTO/c-TiO2/CsPbBr3/carbon. The device
parameters for HBr-0, HBr-10, HBr-30, and HBr-50 were
extracted from the reverse J−V curves under simulated AM 1.5,
100 mW/cm2 sunlight illumination. The J−V curves of these
PSCs are illustrated in Figure 5a, and the detailed parameters
are summarized in Table 1. The device based on the HBr-0
film with (101) orientation achieves a power conversion
efficiency (PCE) of 2.56% with a short-circuit photocurrent
density (Jsc) of 4.44 mA cm−2, an open-circuit photovoltage
(Voc) of 1.11 V, and a fill factor (FF) of 0.52. Although there is
a similar morphology and orientation to the HBr-0, a slight
PCE improvement (3.02%) was achieved for the HBr-10-based
device, attributed to a higher degree of preferential orientation
than HBr-0. In contrast, the HBr-30 film with improved
surface coverage and (121) orientation-based device produces
a high PCE of 6.91%, along with a Jsc of 7.47 mA cm−2, a Voc of
1.36 V, and a FF of 0.68. The higher Voc and FF of HBr-30
compared to HBr-0 is not as surprising as it seems, as the latter
During the initial annealing stage of the wet film, solvent evaporation
occurs, but the PbBr2-DMF and/or PbBr2−HBr complex remains
within the film, the removal of which was further increased by
enthalpy. The PbBr2 films with different morphologies are obtained,
owing to the difference of the removal stage between PbBr2-DMF and
PbBr2−HBr complexes.
a
During the initial annealing stage of the wet film, solvent
evaporation occurs; however, the PbBr2-DMF and/or PbBr2−
HBr complexes remain predominantly in the precursor film.
The growth and crystallization of the precursor film to PbBr2
only proceeds as the complexes decompose and organic/Br
ligands are driven out of the precursor film with increased
enthalpy; however, the temperature and duration over which
this process occurs depend strongly on the stability of
complexes. With the annealing proceeding, relatively weakbonded PbBr2-DMF adducts decompose first. While for the
stable PbBr2−HBr complex, more heat is needed to trigger its
decomposition and release of HBr, during which coarsening of
the film may occur.56−58 Based on the above experimental
observations and related discussion, it can be concluded that
the addition of HBr in PbBr2 precursor solutions plays a vital
role in controlling the microstructure of resulting PbBr2 films.
The morphological modulation of PbBr2 films is correlated
to the crystal growth behaviors of CsPbBr3 films, and a possible
formation mechanism of the (121) preferential-oriented
CsPbBr3 film is proposed. As illustrated in Scheme 2, for the
dense PbBr2 film with less voids which well resembles the
PbBr2 film without HBr, during its conversion process to
CsPbBr3, only the surface region would react with solutionphase CsBr; the growth of CsPbBr3 in the lateral direction is
then constrained, resulting in its predominant growth
perpendicular to the substrate. While for the PbBr2 film with
abundant nanochannels (synthesized by the addition of 30 μL
HBr), CsBr solutions can penetrate and react sufficiently with
the PbBr2 layer; the voids within the PbBr2 film allow the
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Figure 5. PV characterization. (a) J−V characteristics of champion devices based on HBr-0, HBr-10, HBr-30, and HBr-50 films under simulated
AM 1.5G solar illumination of 100 mW cm−2 in reverse scans. (b) J−V curves of PSCs based on HBr-0 and HBr-30 films in the dark. (c)
Comparison of PV parameters (Jsc, Voc, FF, and PCE) for solar cells based on HBr-0, HBr-10, HBr-30, and HBr-50 films. (d) J−V curves with
forward scan and reverse scan, respectively, (e) steady-state maximum PCE outputs, (f) IPCE spectra for HBr-30 PSCs, respectively. The asobtained CsPbBr3 films are abbreviated as HBr-0, HBr-10, HBr-30, and HBr-50, respectively, according to the added amount of HBr. For example,
the HBr-30 sample was derived from the PbBr2 precursor solution (1 mL) with 30 μL of added HBr.
Table 1. PV Parameters of the All PSCs under the AM 1.5 G Solar Spectra with a Light Intensity of 100 mW cm−2a
devices
HBr-0
HBr-10
HBr-30
HBr-50
average
champion
average
champion
average
champion
average
champion
Jsc (mA cm−2)
Voc (V)
FF
PCE (%)
3.77 ± 0.97
4.44
4.07 ± 0.84
4.88
7.00 ± 0.55
7.47
5.57 ± 0.61
5.59
0.97 ± 0.14
1.11
1.05 ± 0.18
1.17
1.35 ± 0.09
1.36
1.25 ± 0.09
1.29
0.50 ± 0.07
0.52
0.53 ± 0.10
0.53
0.64 ± 0.05
0.68
0.59 ± 0.08
0.67
1.86 ± 0.69
2.56
2.18 ± 0.82
3.02
6.11 ± 0.80
6.91
4.11 ± 1.00
4.83
a
These data are extracted from 30 samples for each kind of devices.
and optical absorption of the HBr-0 and HBr-30 films should
be considered first. The cross-sectional SEM images of HBr-0
and HBr-30 films are shown in Figure S4. Similar thicknesses
are observed for HBr-0 (405 nm) and HBr-30 films (416 nm).
Furthermore, HBr-0 and HBr-30 films reveal nearly identical
absorption edge (Figure S5), indicating that the band gap of
them has not changed. Therefore, the film thickness and
sample has a poor surface coverage. The current leakage of the
HBr-0 based-device is larger than that of HBr-30 because of
the direct contact between c-TiO2 and carbon; this effect is
also in accordance with their dark current−potential characteristics (Figure 5b).40 Thus, the detrimental morphological
defects of abundant voids substantially reduce the Voc and
FF.61 With regard to the significant increase in Jsc, the thickness
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Figure 6. (a) Band structures of the orthorhombic CsPbBr3. (b) Dark J−V curves for the electron-only devices fitted with the Mott−Gurney law.
The inset is the sketch of electron-only device configuration. (c) Schematic diagram of the difference in Jsc and PCE of solar cells caused by the
anisotropy of the transport performance of the oriented CsPbBr3 film.
optical absorption are not the main reason for the Jsc
improvement of HBr-30. According to previous literature,46
the facet-dependent optoelectronic anisotropy was confirmed
in CsPbBr3. Thus, the significantly increased Jsc for HBr-30
with (121) orientation may be derived from its higher carrier
transport than HBr-0 with (101) orientation. This verification
will be studied in detail in the next section. The HBr-50-based
device delivers a PCE of 4.83%, along with a Jsc of 5.59 mA
cm−2, a Voc of 1.29 V, and a FF of 0.67. The higher Voc and FF
are also attributed to obtain a high coverage of the CsPbBr3
film. The reproducibility of as-obtained CsPbBr3 films based
PSCs is evaluated by testing 30 individual devices from one
batch, and the resulting PV parameters of Jsc, Voc, FF, and PCE
are summarized in Figure 5c. As can be seen from these box
charts, the PV parameters of all samples show good
reproducibility.
Furthermore, the J−V curves under reverse and forward
scans of the champion device based on HBr-30 are
investigated, and the results are shown in Figure 5d; the
obvious hysteresis effect appear, which probably originate from
the large energy-level offset between carbon and CsPbBr3, as
well as the interface defects between c-TiO2 and CsPbBr3.62−64
Moreover, the steady-state photocurrent outputs and the
photon-to-current conversion efficiency (IPCE) and integrated
Jsc curves for characterizing the reliability of J−V measurements of champion device based on HBr-30 are provided in
Figure 5e,f. It can be seen that a steady-state current density of
6.08 mA cm−2 and PCE of 6.71% are obtained (measured at
bias voltage of 1.08 V), close to champion PCE tested by the
J−V curve under reverse scans. The IPCE spectra of HBr-30based device exhibits narrow light harvesting because of the
large band gap (2.3 eV). In a wavelength region of 300−550
nm, the maximal IPCE value is ∼85%. The integrated Jsc values
calculated from the IPCE curves are 6.93 mA cm−2, which is
close to the Jsc extracted from the J−V curve. The stability is
another key problem to be resolved before the commercial
applications of the cells. Thus, the stability measurement of the
unencapsulated devices can be carried out and the normalized
PCEs in the devices based on HBr-0 and HBr-30 are plotted in
Figure S6. When stored in atmosphere with a relative humidity
of 50−60% and a temperature of 25 °C for 25 days, the HBr-0based and HBr-30-based PSCs can retain 93.3 and 93.6% of
the initial efficiency, respectively. Our result indicated that the
CsPbBr3 PSCs have a relatively good stability.
3.4. Anisotropic Carrier Mobility of Oriented CsPbBr3
Films. To gain more insights into the carrier-transport
properties of (121) and (101) orientation of CsPbBr3, the
carrier effective mass and carrier mobility (μ) were obtained
theoretically and experimentally. Based on the density
functional theory (DFT), the carrier effective masses
(referenced to the electron rest mass m0) are calculated by
eq 2
ℏ2
m* =
2|α|
(2)
where ℏ is the reduced Planck’s constant and α is the second
derivative of the curve in the conduction band (for electrons)
and valence band (for holes).65 To obtain parameter α, the
band structures will be calculated. Figure 6a shows band
structures calculated by the conventional (1 × 1 × 1) cells; the
conduction band minimum (CBM) and valence band
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maximum (VBM) are located at the Γ point. The carrier
effective masses along some directions of the orthorhombic
CsPbBr3 are calculated and listed in Table 2. It can be seen
films are successfully fabricated. The morphology of PbBr2
films is demonstrated to be the key factor that affects the
subsequent crystal growth manner of CsPbBr3, thereby leading
to the consequent thin films with different orientations. The
dense PbBr2 film with less voids results in the (101)-oriented
CsPbBr3 film, while the PbBr2 film with abundant nanochannels adjusted by the dosage of HBr additive leads to the
(121)-oriented film. A higher carrier mobility is observed in
the (121)-oriented CsPbBr3 film, which is consistent with the
Jsc results that a much higher value is observed for the (121)oriented CsPbBr3 film (7.47 mA cm−2) than the (101)oriented one (4.44 mA cm−2). Finally, (121)-oriented CsPbBr3
PSCs exhibit a considerable enhancement in photovoltaic
performance in comparison with the (101)-oriented one, and
the champion PCE increases from original 2.56 to 6.91% due
possibly to the higher surface coverage and carrier mobility.
This study provides a facile yet effective method for preparing
compact preferentially oriented CsPbBr3 films with beneficial
charge carrier transport and also points out the importance of
crystal orientation engineering in design and fabrication of
high-performance optoelectronic devices.
Table 2. Effective Masses of Electrons and Holes (Relative
to the Electron Rest Mass) along Different Directions in the
Brillouin Zone of Orthorhombic CsPbBr3
direction
mh*/m0
me*/m0
[100]
[010]
[001]
[101]
[121]
1.07
0.90
1.02
0.28
0.07
4.37
0.43
4.56
1.13
0.18
that the effective masses of electrons and holes along the [121]
direction are lower than those along the other crystal
directions. Taking advantage of the relationship between
carrier effective masses and their mobility
i q y
μ = jjj zzzτ ̅
(3)
k m* {
where q is the elementary charge and τ̅ is the average scattering
time.66 It can be concluded that the carrier mobility along the
[121] direction is larger than those along the other crystal
directions. Therefore, the HBr-30 film with (121) orientation
has better carrier transport properties than one with (101)
orientation.
The carrier transport properties of HBr-0 and HBr-30 films
were further verified by experimental measurement. The
electron-transport properties in HBr-0 and HBr-30 films
were determined by the SCLC measurement with a simple
structure of FTO/c-TiO2/CsPbBr3/PC61BM/C, as shown in
Figure 6b. In the dark J ≈ Vn curves, the Ohmic (n = 1), trapfilling (n > 3), and child (n = 2) regions were all observed with
the increase of applied electric bias. The trapping state
densities were calculated using the equation nt = 2εε0VTFL/
(eL2), where ε (=16.46) is the relative dielectric constant of
CsPbBr3, ε0 is the vacuum permittivity, e is the elemental
charge, and L is the thickness of the CsPbBr3 film.12,67,68 The
electron-trap density in HBr-30 is estimated to be around 1.08
× 1016, close to 1.19 × 1016 cm−3 in HBr-0. The electron
mobility is estimated from the child region according to Mott−
Gurney’s equation μ = 8JDL3/9εε0V2, where JD is the current
density and V is the applied voltage. According to the equation,
a higher electron mobility of 6.01 × 10−3 cm2 V−1 s−1 is
derived in HBr-30, whereas it is only 1.05 × 10−3 cm2 V−1 s−1
for HBr-0, showing an obvious anisotropy of electron-transport
properties, which is in good agreement with the above
theoretical results. Therefore, based on the investigation
results of the above CsPbBr3 transport properties, it can be
concluded that the increase in the Jsc of HBr-30-based PSCs is
mainly because of the (121)-oriented films having higher hole
and electron mobility than that of (101) orientation, thereby
effectively improving the PCE of the device, as shown in Figure
6c.
■
Article
ASSOCIATED CONTENT
sı Supporting Information
*
The Supporting Information is available free of charge at
https://pubs.acs.org/doi/10.1021/acsaelm.0c00909.
Reaction scheme, simulated XRD pattern, void-labeled
SEM images, estimated porosity, cross-sectional SEM
images, UV−vis absorption spectra, and long-term
stability measurements (PDF)
■
AUTHOR INFORMATION
Corresponding Author
Tao Yu − National Laboratory of Solid State Microstructures,
Nanjing University, Nanjing 210093, P. R. China;
Ecomaterials and Renewable Energy Research Center
(ERERC) at School of Physics, Collaborative Innovation
Center of Advanced Microstructures, and Jiangsu Provincial
Key Laboratory for Nanotechnology, Nanjing University,
Nanjing 210093, P. R. China; orcid.org/0000-00031981-3469; Email: yutao@nju.edu.cn
Authors
Xiaopeng Han − National Laboratory of Solid State
Microstructures, Nanjing University, Nanjing 210093, P. R.
China; Ecomaterials and Renewable Energy Research Center
(ERERC) at School of Physics, Nanjing University, Nanjing
210093, P. R. China
Xin Wang − National Laboratory of Solid State
Microstructures, Nanjing University, Nanjing 210093, P. R.
China; Ecomaterials and Renewable Energy Research Center
(ERERC) at School of Physics, Nanjing University, Nanjing
210093, P. R. China; orcid.org/0000-0002-4301-9786
Jianyong Feng − National Laboratory of Solid State
Microstructures, Nanjing University, Nanjing 210093, P. R.
China; College of Engineering and Applied Sciences, Nanjing
University, Nanjing 210093, P. R. China; orcid.org/
0000-0003-4275-8306
Huiting Huang − National Laboratory of Solid State
Microstructures, Nanjing University, Nanjing 210093, P. R.
China; College of Engineering and Applied Sciences, Nanjing
4. CONCLUSIONS
The manipulation of crystal orientation in CsPbBr3 films
enabled us to investigate the crystal anisotropic effect on
optoelectronic properties and optimize PV performance by
crystal orientation engineering. By using the HBr additive in a
two-step solution method, highly (121)-oriented CsPbBr3
381
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University, Nanjing 210093, P. R. China; orcid.org/
0000-0001-7906-6122
Zhi Zhu − National Laboratory of Solid State Microstructures,
Nanjing University, Nanjing 210093, P. R. China;
Ecomaterials and Renewable Energy Research Center
(ERERC) at School of Physics, Nanjing University, Nanjing
210093, P. R. China
Zhaosheng Li − National Laboratory of Solid State
Microstructures, Nanjing University, Nanjing 210093, P. R.
China; College of Engineering and Applied Sciences, Nanjing
University, Nanjing 210093, P. R. China; orcid.org/
0000-0001-8114-0432
Zhigang Zou − National Laboratory of Solid State
Microstructures, Nanjing University, Nanjing 210093, P. R.
China; Ecomaterials and Renewable Energy Research Center
(ERERC) at School of Physics, Collaborative Innovation
Center of Advanced Microstructures, and Jiangsu Provincial
Key Laboratory for Nanotechnology, Nanjing University,
Nanjing 210093, P. R. China
Complete contact information is available at:
https://pubs.acs.org/10.1021/acsaelm.0c00909
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Author Contributions
The manuscript was written through contributions of all
authors. All authors have given approval to the final version of
the manuscript.
Notes
The authors declare no competing financial interest.
■
ACKNOWLEDGMENTS
This work was supported primarily by the National Key
Research and Development Program of China
(2018YFA0209303) and the National Natural Science
Foundation of China (61377051, U1663228, 51902153, and
51972165). We thank Qingxiao Meng, Jincheng Li, and
Wenjing Su for informative discussion and experimental and
technical assistance.
■
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