Lead-free solid-state organic-inorganic halide perovskite

ARTICLES
PUBLISHED ONLINE: 4 MAY 2014 | DOI: 10.1038/NPHOTON.2014.82
Lead-free solid-state organic–inorganic halide
perovskite solar cells
Feng Hao1, Constantinos C. Stoumpos1, Duyen Hanh Cao1, Robert P. H. Chang2
and Mercouri G. Kanatzidis1 *
Lead-free solution-processed solid-state photovoltaic devices based on methylammonium tin iodide (CH3NH3SnI3)
perovskite semiconductor as the light harvester are reported. Featuring an optical bandgap of 1.3 eV, the CH3NH3SnI3
perovskite material can be incorporated into devices with the organic hole-transport layer spiro-OMeTAD and show an
absorption onset at 950 nm, which is signiﬁcantly redshifted compared with the benchmark CH3NH3PbI3 counterpart
(1.55 eV). Bandgap engineering was implemented by chemical substitution in the form of CH3NH3SnI3–xBrx solid solutions,
which can be controllably tuned to cover much of the visible spectrum, thus enabling the realization of lead-free solar cells
with an initial power conversion efﬁciency of 5.73% under simulated full sunlight. Further efﬁciency enhancements are
expected following optimization and a better fundamental understanding of the internal electron dynamics and
corresponding interfacial engineering. The reported CH3NH3SnI3–xBrx perovskite solar cells represent a step towards the
realization of low-cost, environmentally friendly solid-state solar cells.
T
he development of clean alternatives to current power generation methods is immensely important to preserving the
global environment and assuring sustained economic
growth1. The recent emergence of halide perovskites as light harvesters and hole-transport materials has revolutionized the emerging photovoltaic technologies2–14. Organic–inorganic hybrid
perovskite compounds based on metal halides adopt the ABX3 perovskite structure. This structure consists of a network of cornersharing BX6 octahedra, where the B atom is a metal cation (typically
Sn2þ or Pb2þ) and X is typically F2, Cl2, Br2 or I2. The A cation is
selected to balance the total charge and it can even be a Csþ or a
small molecular species15–17. Recent implementation of
CH3NH3PbX3 (X ¼ I, Cl, Br) perovskite absorbers with the
organic
hole
conductor
2,2′ ,7,7′ -tetrakis-(N,N-di-p-methoxyphenylamine)9,9′ -spirobiﬂuorene (spiro-OMeTAD) enabled
power conversion efﬁciencies (PCEs) greater than 15% (refs 8,18),
and has been recognized as the ‘next big thing in photovoltaics’19–22.
A planar heterojunction photovoltaic device incorporating
vapour-deposited perovskite (CH3NH3PbI32xClx) as the absorbing
layer has demonstrated overall PCEs of over 15% with a high opencircuit voltage of up to 1.07 V, further highlighting the industrial
application potential of this conﬁguration in the near future7.
Recent studies have indicated that mixed-halide organic–inorganic
hybrid perovskites can display electron–hole diffusion lengths of
over 1 mm, which is consistent with our reports of very high
carrier mobilities in these materials23 and supports our expectations
for highly efﬁcient and cheap solar cells using thick absorption
layers24,25. However, to realize commercial applications of this technology it is important to reach analogous optical and photovoltaic
performance using lead-free organic–inorganic compounds.
Here, we report a ﬁrst attempt using the lead-free perovskite of
methylammonium tin iodide (CH3NH3SnI3) as the light-absorbing
material to fabricate solution-processed solid-state photovoltaic
devices. Featuring an even lower optical bandgap of 1.3 eV than
the 1.55 eV achieved with CH3NH3PbI3 , devices with
CH3NH3SnI3 in conjunction with an organic spiro-OMeTAD
hole-transport layer showed an absorption onset of 950 nm.
Further chemical alloying of iodide with bromide provides efﬁcient
energetic tuning of the band structure of the perovskites, leading to a
PCE of 5.8% under simulated full sunlight of 100 mW cm22.
Building on this very promising initial result and with further
reduction of interfacial losses, we believe a substantial increase in
efﬁciency can be achieved.
As shown in Fig. 1a,b, CH3NH3SnI3 adopts the perovskite structure type, crystallizing in the pseudocubic space group P4mm at
ambient conditions. Unlike CH3NH3PbI3 , which has a lower symmetry at room temperature (b-phase), the Sn analogue adopts its
highest symmetry phase (a-phase), even at room temperature.
The corner-sharing [SnI6]42 polyhedra form an inﬁnite threedimensional lattice with Sn–I–Sn connecting angles of 177.43(1)8
and 1808 for the a- and c-axes, respectively. The deviation from
the ideal cubic (Pm–3m) structure arises from orientational polarization of the CH3NHþ
3 cation along the C–N bond direction, which is
imposed on the three-dimensional [SnI3]2 inorganic lattice
coinciding with the crystallographic c-axis23.
CH3NH3SnI3 is a direct-gap semiconductor with an energy gap
of 1.3 eV, as has been shown experimentally and theoretically23,26.
The optical bandgap Eg of the CH3NH3SnI3 compound (determined
from diffuse reﬂectance measurements) is shown in Fig. 1c. The
optical absorption coefﬁcient (a/S) is calculated using reﬂectance
data according to the Kubelka–Munk equation27, a/S ¼ (1 2 R)2/2R,
where R is the percentage of reﬂected light, and a and S are the
absorption and scattering coefﬁcients, respectively. At room temperature, the material displays a strong photoluminescence emission
at 950 nm, which corresponds to the onset of the absorption edge
(Fig. 1c). The photoluminescence intensity can act as a qualitative
measure of the number of photogenerated carriers in semiconductors, as it is proportional to the number of e2hþ pairs produced
by the incident light28. As depicted in Fig. 1d, its bulk electrical
conductivity s is 5 × 1022 S cm21 at room temperature,
1
Department of Chemistry, Northwestern University, 2145 Sheridan Road, Evanston, Illinois 60208, USA, 2 Department of Materials Science and Engineering,
and Argonne-Northwestern Solar Energy Research (ANSER) Center, Northwestern University, 2145 Sheridan Road, Evanston, Illinois 60208, USA. * e-mail:
[email protected]
NATURE PHOTONICS | VOL 8 | JUNE 2014 | www.nature.com/naturephotonics
© 2014 Macmillan Publishers Limited. All rights reserved.
489
ARTICLES
NATURE PHOTONICS
a
DOI: 10.1038/NPHOTON.2014.82
Sn
IBr
C
N
b
a
Intensity (a.u.)
b
c
10
20
30
40
50
60
70
80
2θ (deg)
1.2
1.4
1.6
1.8
2.0
Conductivity (S cm−1 × 10−2)
Absorbance (a.u.)
1.0
−40
d
7.0
CH3NH3SnI3
−60
6.5
6.0
−80
5.5
Seebeck coefficient (µV K−1)
CH3NH3SnI3
Photoluminescence emission (a.u.)
c
−100
300
320
340
360
380
Temperature (K)
Energy (eV)
Figure 1 | Crystal structure, XRD pattern, optical absorption and photoluminescence spectra, conductivity and Seebeck coefﬁcient of CH3NH3SnI3
perovskite. a, Perovskite crystal structure of the CH3NH3SnI32xBrx absorber materials. b, Experimental (red) and simulated (black) X-ray diffraction pattern
for CH3NH3SnI3. c,d, Optical absorption and photoluminescence spectra (c) and conductivity and Seebeck coefﬁcient (d) as a function of temperature for a
sample of CH3NH3SnI3 prepared using the solution method23.
corresponding to a Seebeck coefﬁcient of 2 60 mV K21 (n-type).
The compound has a low carrier concentration on the order of
1 × 1014 cm23 and high electron mobilities (me) on the order of
2,000 cm2 V21 s21, which is comparable or even superior to
most traditional semiconductors, including Si, CuInSe2 and CdTe,
which have comparable bandgap energies. The doping level of
CH3NH3SnI3 can be varied greatly depending on the preparation
method. Carrier concentrations up to 1 × 1019 cm23 have been
reported for CH3NH3SnI3 (ref. 29), showing a strong p-type character and a metallic behaviour suggestive of a heavily doped semiconducting behaviour. We attribute this large difference in the transport
properties to Sn4þ impurities that are inherently present in commercial SnI2 , which are readily detectable by a mass loss at
150 8C via thermal gravimetric analysis (Supplementary Fig. 1).
Therefore, when assembling the solar cells, care must be taken in
depositing ﬁlms of tin perovskite with low carrier concentration
to maximize the carrier mobility within the active perovskite
layer. This means that excessive Sn4þ in the sample must be avoided.
The valence band maximum (EVB) of the CH3NH3SnI3 compound was determined from ultraviolet photoelectron spectroscopy
(UPS) measurements. A representative UPS spectrum for the
CH3NH3SnI3 is shown in Supplementary Fig. 2, where the energy
is calibrated with respect to the He I photon energy (21.21 eV).
The valence band energy EVB is estimated to be 25.47 eV below
vacuum level, which is close to the reported value for
CH3NH3PbI3 (25.43 eV)5. From the observed optical bandgap,
the conduction band energy ECB of CH3NH3SnI3 was determined
to be at 24.17 eV, that is, slightly higher than the ECB for the
TiO2 anatase electrode (24.26 eV)5.
To fabricate the solid-state solar cells, mesoporous anatase TiO2
ﬁlms were prepared by spin coating a solution of colloidal anatase
particles (20 nm in size) onto a 30-nm-thick compact TiO2 underlayer30. The underlayer was deposited by atomic layer deposition
on a pre-patterned transparent-conducting-oxide-coated glass
490
substrate acting as the electric front contact of the solar cell.
Deposition of the perovskite light-absorbing layer was carried out
by spin coating in a nitrogen glove box to avoid hydrolysis and oxidation of the tin perovskite in contact with air. The triarylamine
derivative 2,2′ ,7,7′ -tetrakis-(N,N-di-p-methoxyphenylamine)-9,9′ spirobiﬂuorene (spiro-OMeTAD)31 was then applied as a holetransporting material (HTM) on top of the mesoporous TiO2 and
perovskite layer. Lithium bis(triﬂuoromethylsulphonyl)imide and
2,6-lutidine were added in the HTM solution as important
dopants to increase the hole mobility31. Figure 2 shows a representative cross-sectional scanning electron microscopy (SEM) image of
a typical solar cell device. The mesoporous TiO2 ﬁlm had an average
thickness of 350 nm and was inﬁltrated with the perovskite nanocrystals using the spin-coating procedure. The HTM penetrates into
the remaining pore volume of the perovskite/TiO2 layer and forms a
200-nm-thick capping layer on top of the composite structure. A
thin gold layer was thermally evaporated under high vacuum onto
the HTM layer, forming the back contact electrode of the device.
The solid-state device based on the CH3NH3SnI3 perovskite
shows a high mean short-circuit photocurrent density Jsc of
16.30 mA cm22, an open-circuit voltage Voc of 0.68 V and a moderate ﬁll factor (FF) of 0.48 under AM 1.5G solar illumination, corresponding to a PCE of 5.23% (Fig. 3a). This high current density was
achieved with submicrometre-thick TiO2 ﬁlms (that is, 350 nm)
because of the large optical absorption cross-section of the perovskite material and the well-developed interfacial pore ﬁlling by the
hole conductor (Fig. 2). More importantly, the incident photonto-electron conversion efﬁciency (IPCE) of the CH3NH3SnI3based device covers the entire visible spectrum and reaches a
broad absorption maximum of over 60% from 600 nm to 850 nm.
It is accompanied with a notable absorption onset up to 950 nm
(Fig. 3b), which is in good agreement with the optical bandgap of
1.30 eV. Integrating the overlap of the IPCE spectrum with the
AM 1.5G solar photon ﬂux yields a current density of
NATURE PHOTONICS | VOL 8 | JUNE 2014 | www.nature.com/naturephotonics
© 2014 Macmillan Publishers Limited. All rights reserved.
NATURE PHOTONICS
ARTICLES
DOI: 10.1038/NPHOTON.2014.82
x=0
Au electrode
Spiro-OMeTAD
x=1
Perovskite/TiO2
Blocking TiO2 layer
x=2
FTO layer
x=3
Figure 2 | Representative cross-sectional SEM view of a completed photovoltaic device with CH3NH3SnI3 perovskite. Individual layers are indicated on the
left. Scale bar, 200 nm. Real device images are shown on the right, indicating the colours of the photovoltaic devices made with CH3NH3SnI32xBrx as a
function of I/Br ratio.
16.60 mA cm22, which is in excellent agreement with the measured
photocurrent density. This conﬁrms that any mismatch between the
simulated sunlight and the AM 1.5G standard is negligibly small. It
is important to note that, although the obtained Jsc for the
CH3NH3SnI3 perovskite device is less efﬁcient than that for the
CH3NH3PbI3 device7,8,18,32, the maximum current density that can
be generated exceeds 30 mA cm22 when integrating the AM 1.5G
solar spectrum below the bandgap of CH3NH3SnI3 perovskite
(1.30 eV). To ﬁgure out the limiting factor for the fair photocurrent
density, devices with thinner active layer thickness (150 nm) were
constructed and tested. As shown in Supplementary Fig. 3, a Jsc of
12 mA cm22 was achieved, with a Voc of 0.74 V and a FF of
0.45, thus generating a PCE of 4.44%. Accordingly, the diffusion
Current density (mA cm−2)
a 20
CH3NH3SnI3
CH3NH3SnI2Br
15
CH3NH3SnIBr2
CH3NH3SnBr3
10
5
0
−5
0.0
0.2
0.4
0.6
0.8
1.0
1.2
Voltage (V)
b
CH3NH3SnI3
CH3NH3SnI2Br
80
CH3NH3SnIBr2
CH3NH3SnBr3
IPCE (%)
60
40
20
0
400
500
600
700
800
900
1,000
Wavelength (nm)
Figure 3 | Photovoltaic and IPCE characteristics for devices with
CH3NH3SnI32x Brx perovskites. a,b, Photocurrent density–voltage (J–V)
characteristics (a) and corresponding IPCE spectra (b) of devices based on
CH3NH3SnI32xBrx (x ¼ 0, 1, 2, 3) perovskites.
length of the tin perovskite might not be a main limiting factor
for device performance24,25. The ﬁlm morphology and quality of
the spin-coated CH3NH3SnI3 ﬁlm were then investigated by SEM
(Supplementary Fig. 4). Poor ﬁlm quality and coverage of the tin
perovskite on the mesoporous TiO2 electrodes were observed,
which has recently been recognized as an important factor determining perovskite solar cell performance8,13. Future device optimization will therefore focus on the improvement of perovskite ﬁlm
quality and interfacial recombination inhibition.
It has recently been observed that the charge accumulates in high
density in the perovskite absorber material rather than only in the
semiconducting TiO2 electrodes, making this type of photovoltaic
device fundamentally different from dye-sensitized solar cells33.
Thus, the Voc in a perovskite solar cell is not only related to the
energy difference between the HTM potential and the TiO2 conduction bandedge, but could also be correlated with the energy difference between the HTM potential and the conduction bandedge of
the perovskite itself. From the abovementioned band alignment it
can be inferred that the conduction bandedge ECB of
CH3NH3SnI3 is 0.24 eV lower than in CH3NH3PbI3 , thus
leading to a lower Voc for the CH3NH3SnI3 perovskite device.
Therefore, in an attempt to increase the Voc of these lead-free
devices, chemical substitution of the iodide atom with bromide
was applied in order to favourably tune the bandgap energetics32.
The CH3NH3SnI32xBrx compounds were prepared by mixing
stoichiometric amounts of CH3NH3X and SnX2 (X ¼ Br, I), ﬁnely
homogenized in a mortar in the nitrogen glove box. The resulting
solids were sealed in silica ampules under 1 × 1024 mbar vacuum
and heated to 200 8C to complete the reaction. As shown by the
X-ray diffraction (XRD) patterns in Fig. 4a, this series of compositions forms a continuous solid solution throughout the composition range, without displaying any structural transitions (at
room temperature), thus retaining the crystal structure of both
end members, crystallizing at the P4mm space group. The properties
of the solid solutions are clearly displayed by a continuous contraction of the lattice parameters from the CH3NH3SnI3 to
CH3NH3SnBr3 end members, which results in a widening of the
bandgap (Table 1).
To check the optical properties in the hybrid halide perovskite,
ultraviolet–visible diffuse reﬂectance spectra of CH3NH3SnI32xBrx
(x ¼ 0, 1, 2, 3) were measured and transformed to absorption
spectra, as mentioned above. Figure 4b shows that the absorption
onset of CH3NH3SnI32xBrx (x ¼ 0, 1, 2, 3) hybrid halide perovskites
can be tuned from 954 nm (1.30 eV for CH3NH3SnI3) to 577 nm
(2.15 eV for CH3NH3SnBr3), thus resulting in signiﬁcant colour
tunability for perovskite photovoltaic devices (as shown in the
NATURE PHOTONICS | VOL 8 | JUNE 2014 | www.nature.com/naturephotonics
© 2014 Macmillan Publishers Limited. All rights reserved.
491
ARTICLES
NATURE PHOTONICS
DOI: 10.1038/NPHOTON.2014.82
a
Intensity (a.u.)
CH3NH3SnBr3
CH3NH3SnIBr2
CH3NH3SnI2Br
CH3NH3SnI3
sim_CH3NH3SnI3
10
20
30
40
50
60
70
80
90
2θ (deg)
b
c
CH3NH3SnI3 (1.30 eV)
−2.05
CH3NH3SnI2Br (1.56 eV)
CH3NH3SnIBr2 (1.75 eV)
Energy (eV)
Absorbance (a.u.)
CH3NH3SnBr3 (2.15 eV)
−4.26
−4.17
1.30
−3.96
1.56
−3.78
1.75
−3.39
2.15
−5.22
−5.47 −5.52 −5.53 −5.54
1.0
1.5
2.0
2.5
3.0
Energy (eV)
TiO2
1
2
3
0
CH3NH3SnI3−xBrx
SpiroOMeTAD
Figure 4 | XRD patterns, absorption spectra and schematic energy-level diagram of CH3NH3SnI32x Brx compounds. a,b, XRD patterns (a) and absorption
spectra (b) of the CH3NH3SnI32xBrx (x ¼ 0, 1, 2, 3) perovskites. c, Schematic energy-level diagram of CH3NH3SnI32xBrx with TiO2 and spiro-OMeTAD HTM.
The valence band maxima ECB of the methylammonium tin halides were extracted from UPS measurements under high vacuum.
right panel of Fig. 2). The intermediate iodide/bromide hybrid perovskites of CH3NH3SnI2Br and CH3NH3SnIBr2 show absorption
onsets of 795 nm (1.56 eV) and 708 nm (1.75 eV), respectively.
The valence band energy EVB of the CH3NH3SnI32xBrx compounds
was also estimated from the UPS measurements. As illustrated in
Fig. 4c, the ECB increased from –4.17 eV below vacuum level for
CH3NH3SnI3 to 23.96 eV for CH3NH3SnI2Br and 23.78 eV for
CH3NH3SnIBr2 , and ﬁnally to 23.39 eV for CH3NH3SnBr3. It is
obvious from the band alignment diagram that the change in the
bandgap Eg of the CH3NH3SnI32xBrx compounds is mainly due
to the conduction band shift to higher energy, with the valence
band energy remaining practically unchanged. These changes in
energy levels allow for bandgap engineering and the tuning of energetics for more efﬁcient solar cell architectures.
Figure 3a,b presents representative photocurrent density–voltage
(J–V ) characteristics and IPCE spectra for devices constructed with
the CH3NH3SnI32xBrx perovskites as light harvesters. The photovoltaic parameters are summarized in Table 1. As demonstrated
in the right panel of Fig. 2, through the chemical compositional
control of CH3NH3SnI32xBrx , the corresponding device colour
can be tuned from black for CH3NH3SnI3 to dark brown for
CH3NH3SnI2Br and to yellow for CH3NH3SnBr3 with increasing
Br content. Notably, Jsc decreased from 16.30 mA cm22 for
CH3NH3SnI3 to 8.26 mA cm22 for CH3NH3SnBr3 with increasing
Br content, whereas Voc increased from 0.68 V to 0.88 V when
switching from the pure iodide to pure bromide perovskite. In
addition to the signiﬁcant improvement in Voc , an increase in FF
from 0.48 to 0.59 was also observed upon the incorporation of the
Br ions. Amongst the investigated CH3NH3SnI32xBrx perovskites,
the device with CH3NH3SnIBr2 delivered the highest PCE of
5.73%, with a Jsc of 12.30 mA cm22, a Voc of 0.82 V and a FF of
0.57. The reduction of Jsc with increasing Br content is directly
related to the blueshift of absorption onset, as indicated from the
IPCE spectra in Fig. 3b. Consistent with the bandgap tuning, the
onset of the IPCE spectra blueshifted from 950 nm for the iodide
perovskite to 600 nm for the pure bromide perovskite.
Integrating the overlap of these IPCE spectra with the AM 1.5G
solar photon ﬂux yields a current density Jcal that is similar to the
measured photocurrent density Jsc (Table 1). The improvement in
Voc can be attributed to the raised conduction bandedge ECB with
Table 1 | Optical bandgaps and reﬁned lattice parameters of the CH3NH3SnI32xBrx (x 5 0, 1, 2, 3) perovskites and
corresponding solar cell performance parameters.
Perovskites
CH3NH3SnI3
CH3NH3SnI2Br
CH3NH3SnIBr2
CH3NH3SnBr3
E g* (eV)
1.30
1.56
1.75
2.15
Lattice parameters (Å)
a ¼ 6.169(1) c ¼ 6.173(4)
a ¼ 6.041(1) c ¼ 6.053(4)
a ¼ 5.948(1) c ¼ 5.953(4)
a ¼ 5.837(1) c ¼ 5.853(4)
J †sc (mA cm22)
16.30+0.71
14.38+0.49
12.30+0.47
8.26+0.53
J ‡cal (mA cm22)
16.60
13.96
11.73
7.93
Voc (V)§
0.68+0.03
0.77+0.02
0.82+0.03
0.88+0.03
FF
0.48+0.03
0.50+0.02
0.57+0.02
0.59+0.02
PCE (%)}
5.23+0.18
5.48+0.15
5.73+0.23
4.27+0.18
R #s (V)
105.00
103.54
65.24
60.98
The photovoltaic parameters were the average of six devices in the same bath.
*Optical bandgap determined from the diffuse-reﬂectance measurements; †short-circuit photocurrent density; ‡calculated photocurrent density from the integration of representative IPCE curves shown in Fig. 3b
with the AM 1.5G solar spectrum; §open-circuit photovoltage; ﬁll factor; }power conversion efﬁciency; #series resistances derived from the J–V curves in Fig. 3a.
492
NATURE PHOTONICS | VOL 8 | JUNE 2014 | www.nature.com/naturephotonics
© 2014 Macmillan Publishers Limited. All rights reserved.
NATURE PHOTONICS
ARTICLES
DOI: 10.1038/NPHOTON.2014.82
increasing Br content in CH3NH3SnI32xBrx. However, the FF is signiﬁcantly lower than the reported values for high-efﬁciency perovskite solar cells7,8,18. It is well recognized that series resistance Rs is
one of the major factors inﬂuencing the FF of solar cells. This
arises mainly from three factors: (1) the active and interfacial
layer resistances, (2) electrode resistance and (3) contact resistance.
In the present work, variations in Rs among the cells with
CH3NH3SnI32xBrx perovskites will be mainly caused by differences
in the active layer resistance, because factors (2) and (3) will be
similar, being common in all of our devices. We estimated Rs
from the slope of the J–V curve at the open-circuit voltage point.
As shown in Table 1, the Rs value decreased from 105 V for a
device with CH3NH3SnI3 to 60.98 V for CH3NH3SnBr3 , which is
in accordance with the observed FF enhancement from 0.48 to
0.59. Another important issue that needs to be speciﬁed is that
insufﬁcient coverage of the perovskite ﬁlm (Supplementary Fig. 4)
might result in a high frequency of shunt paths allowing direct
contact between the HTM and the TiO2 compact layer, which will
act as a parallel diode in the device, deteriorating the FF and Voc
(ref. 34). Further FF improvements are expected from more efﬁcient
interfacial engineering to inhibit back electron recombination.
An important issue to note for the tin perovskite solar cell is the
well-known poor atmospheric stability of the Sn-based perovskite
compared to its lead analogy16,21,35. We carried out a preliminary
stability investigation of the CH3NH3SnI3 perovskite solar cell by
storing the devices in a nitrogen glove box after sealing with
Surlyn ﬁlms. Encouragingly, the devices retained almost 80% of
the initial performance in the ﬁrst 12 h (Supplementary Table 1).
Performance loss arises mainly from the decrease in photocurrent
density and FF, primarily due to the p-type doping via Sn2þ oxidation induced during the fabrication process. Far better stability
can be expected if more advanced sealing techniques are adopted.
In summary, methylammonium tin halide perovskites
(CH3NH3SnI32xBrx) have been used as lead-free light harvesters
for solar cell applications for the ﬁrst time. Featuring an ideal
optical bandgap of 1.3 eV, devices with CH3NH3SnI3 perovskite
together with an organic spiro-OMeTAD hole-transport layer
showed a notable absorption onset up to 950 nm, which is signiﬁcantly redshifted compared with its benchmark CH3NH3PbI3
counterpart
(1.55 eV).
The
bandgap
engineering
of
CH3NH3SnI32xBrx perovskites can be controllably tuned to cover
much of the visible spectrum, thus enabling the realization of
lead-free, colourful solar cells and leading to a promising initial
PCE of 5.73% under simulated full sunlight. Further efﬁciency
enhancements would be expected by the fundamental understanding of the internal electron dynamics and corresponding interfacial
engineering. The reported CH3NH3SnI32xBrx perovskites are
believed to represent a signiﬁcant step towards the realization of
low-cost, high-efﬁciency, environmentally benign, next-generation
solid-state solar cells.
Methods
Materials. Unless stated otherwise, all materials were purchased from Sigma-Aldrich
and used as received. Spiro-OMeTAD was purchased from Merck KGaA. CH3NH3I,
CH3NH3Br and SnI2 were synthesized and puriﬁed according to a reported
procedure23. CH3NH3SnI32xBrx compounds were prepared by mixing
stoichiometric amounts of CH3NH3X and SnX2 (X ¼ Br, I), ﬁnely homogenized in a
mortar in a nitrogen glove box. The resulting solids were sealed in quartz ampules
under 1 × 1024 mbar vacuum and heated to 200 8C to complete the reaction23.
Material characterization. Optical diffuse-reﬂectance measurements were
performed at room temperature using a Shimadzu UV-3101 PC double-beam,
double-monochromator spectrophotometer operating from 200 nm to 2,500 nm.
BaSO4 was used as a non-absorbing reﬂectance reference. Photoluminescence
spectra were measured with an OmniPV photoluminescence system, equipped with
a diode-pumped frequency-doubled Nd:YAG laser (500 mW power output, class 4)
emitting at 532 nm coupled with a bundle of eight 400-mm-core optical ﬁbres as an
excitation source. Resistivity measurements were made for arbitrary current
directions in the a–b plane using a standard point contact geometry. A homemade
resistivity apparatus was used that was equipped with a Keithley 2182A
nanovoltometer, Keithley 617 electrometer, Keithley 6220 Precision d.c. source and a
high-temperature vacuum chamber controlled by a K-20 MMR system. Seebeck
measurements were performed on the same homemade apparatus using Cr/Cr:Ni
thermocouples as electric leads, which were attached to the sample surface by means
of colloidal graphite isopropanol suspension. The temperature gradient along the
crystal was generated by a resistor on the ‘hot’ side of the crystal. The data were
corrected for the thermocouple contribution using a copper wire. SEM and energydispersive spectroscopy (EDS) measurements were performed with a Hitachi
SU8030 scanning electron microscope equipped with an Oxford X-max 80 SDD
EDS detector. Data were acquired with an accelerating voltage of 15 kV.
Device fabrication. A ﬂuorine-doped tin oxide-coated glass substrate (Tec15,
Hartford Glass) was patterned by etching with Zn metal powder and 2 M HCl
diluted in deionized water. The substrates were then cleaned by ultrasonication with
detergent, rinsed with deionized water, acetone and ethanol, and dried with clean
dry air. A 30-nm-thick TiO2 compact layer was deposited on the substrates by an
atomic layer deposition system (Savannah S300, Cambridge Nanotech) using
titanium isopropoxide (TTIP) and water as precursors. The mesoporous TiO2 layer
composed of 20-nm-sized particles was deposited by spin coating at 4,500 r.p.m. for
30 s using a hydrothermal-synthesized TiO2 paste diluted in ethanol (1:4, weight
ratio). After drying at 125 8C, the TiO2 ﬁlms were gradually heated to 500 8C, baked
at this temperature for 15 min, and then cooled to room temperature (25 8C). After
cooling to room temperature, the substrates were treated in a 0.02 M aqueous
solution of TiCl4 for 30 min at 70 8C, rinsed with deionized water, and dried at
500 8C for 20 min. Before use, the ﬁlms were again dried at 500 8C for 30 min.
CH3NH3SnI32xBrx was dissolved in N,N-dimethylformamide at a weight
concentration of 30% while stirring at 70 8C. The solution was kept at 70 8C during
the whole procedure. The mesoporous TiO2 ﬁlms were then inﬁltrated with
CH3NH3SnI32xBrx by spin coating at 4,000 r.p.m. for 45 s and dried at 125 8C for
30 min to remove the solvent. The HTM was then deposited by spin coating at
4,000 r.p.m. for 30 s. The spin-coating formulation was prepared by dissolving
72.3 mg spiro-OMeTAD, 30 ml 2,6-lutidine, 17.5 ml of a stock solution of
520 mg ml21 lithium bis(triﬂuoromethylsulphonyl)imide in acetonitrile in 1 ml
chlorobenzene. Finally, 100 nm of gold was thermally evaporated on top of the
device to form the back contact. The devices were sealed in nitrogen using a 30-mmthick hot-melting polymer and a microscope coverslip to prevent oxidation.
Device characterization. J–V characteristics were measured under AM 1.5G light
(100 mW cm22) using the xenon arc lamp of a Spectra-Nova Class A solar
simulator. Light intensity was calibrated using an NREL-certiﬁed monocrystalline Si
diode coupled to a KG3 ﬁlter to bring the spectral mismatch to unity. A Keithley
2400 source meter was used for electrical characterization. The active area of all
devices was 10 mm2, and an 8 mm2 aperture mask was placed on top of the cells
during all measurements. IPCEs were characterized using an Oriel model QE-PV-SI
instrument equipped with a NIST-certiﬁed Si diode. Monochromatic light was
generated from an Oriel 300 W lamp.
Received 14 January 2014; accepted 18 March 2014;
published online 4 May 2014
References
1. Bisquert, J. Photovoltaics: the two sides of solar energy. Nature Photon. 2,
648–649 (2008).
2. Chung, I., Lee, B., He, J., Chang, R. P. H. & Kanatzidis, M. G. All-solid-state dyesensitized solar cells with high efﬁciency. Nature 485, 486–489 (2012).
3. Kojima, A., Teshima, K., Shirai, Y. & Miyasaka, T. Organometal halide
perovskites as visible-light sensitizers for photovoltaic cells. J. Am. Chem. Soc.
131, 6050–6051 (2009).
4. Etgar, L. et al. Mesoscopic CH3NH3PbI3/TiO2 heterojunction solar cells. J. Am.
Chem. Soc. 134, 17396–17399 (2012).
5. Kim, H.-S. et al. Lead iodide perovskite sensitized all-solid-state submicron thin
ﬁlm mesoscopic solar cell with efﬁciency exceeding 9%. Sci. Rep. 2, 591 (2012).
6. Heo, J. H. et al. Efﬁcient inorganic–organic hybrid heterojunction solar cells
containing perovskite compound and polymeric hole conductors. Nature
Photon. 7, 487–492 (2013).
7. Liu, M., Johnston, M. B. & Snaith, H. J. Efﬁcient planar heterojunction
perovskite solar cells by vapour deposition. Nature 501, 395–398 (2013).
8. Burschka, J. et al. Sequential deposition as a route to high-performance
perovskite-sensitized solar cells. Nature 499, 316–319 (2013).
9. Kim, H.-S. et al. High efﬁciency solid-state sensitized solar cell-based on
submicrometer rutile TiO2 nanorod and CH3NH3PbI3 perovskite sensitizer.
Nano Lett. 13, 2412–2417 (2013).
10. Abrusci, A. et al. High-performance perovskite–polymer hybrid solar cells via
electronic coupling with fullerene monolayers. Nano Lett. 13, 3124–3128 (2013).
11. Edri, E., Kirmayer, S., Cahen, D. & Hodes, G. High open-circuit voltage solar
cells based on organic–inorganic lead bromide perovskite. J. Phys. Chem. Lett. 4,
897–902 (2013).
NATURE PHOTONICS | VOL 8 | JUNE 2014 | www.nature.com/naturephotonics
© 2014 Macmillan Publishers Limited. All rights reserved.
493
ARTICLES
NATURE PHOTONICS
12. Bi, D., Yang, L., Boschloo, G., Hagfeldt, A. & Johansson, E. M. J. Effect of
different hole transport materials on recombination in CH3NH3PbI3 perovskitesensitized mesoscopic solar cells. J. Phys. Chem. Lett. 4, 1532–1536 (2013).
13. Ball, J. M., Lee, M. M., Hey, A. & Snaith, H. J. Low-temperature processed mesosuperstructured to thin-ﬁlm perovskite solar cells. Energy Environ. Sci. 6,
1739–1743 (2013).
14. Laban, W. A. & Etgar, L. Depleted hole conductor-free lead halide iodide
heterojunction solar cells. Energy Environ. Sci. 6, 3249–3253 (2013).
15. Borriello, I., Cantele, G. & Ninno, D. Ab initio investigation of hybrid organic–
inorganic perovskites based on tin halides. Phys. Rev. B 77, 235214 (2008).
16. Mitzi, D. B. Synthesis, structure, and properties of organic–inorganic perovskites
and related materials. Prog. Inorg. Chem. 48, 1–121 (1999).
17. Kagan, C. R., Mitzi, D. B. & Dimitrakopoulos, C. D. Organic–inorganic hybrid
materials as semiconducting channels in thin-ﬁlm ﬁeld-effect transistors. Science
286, 945–947 (1999).
18. Lee, M. M., Teuscher, J., Miyasaka, T., Murakami, T. N. & Snaith, H. J. Efﬁcient
hybrid solar cells based on meso-superstructured organometal halide
perovskites. Science 338, 643–647 (2012).
19. Snaith, H. J. Perovskites: the emergence of a new era for low-cost, high-efﬁciency
solar cells. J. Phys. Chem. Lett. 4, 3623–3630 (2013).
20. Park, N.-G. Organometal perovskite light absorbers toward a 20% efﬁciency lowcost solid-state mesoscopic solar cell. J. Phys. Chem. Lett. 4, 2423–2429 (2013).
21. Hodes, G. Perovskite-based solar cells. Science 342, 317–318 (2013).
22. Bisquert, J. The swift surge of perovskite photovoltaics. J. Phys. Chem. Lett. 4,
2597–2598 (2013).
23. Stoumpos, C. C., Malliakas, C. D. & Kanatzidis, M. G. Semiconducting tin and
lead iodide perovskites with organic cations: phase transitions, high mobilities,
and near-infrared photoluminescent properties. Inorg. Chem. 52,
9019–9038 (2013).
24. Xing, G. et al. Long-range balanced electron- and hole-transport lengths in
organic–inorganic CH3NH3PbI3. Science 342, 344–347 (2013).
25. Stranks, S. D. et al. Electron–hole diffusion lengths exceeding 1 micrometer in an
organometal trihalide perovskite absorber. Science 342, 341–344 (2013).
26. Umari, P., Mosconi, E. & de Angelis, F. Relativistic solar cells. Preprint at
http://arxiv-web.arxiv.org/abs/1309.4895 (2013).
27. Gate, L. F. Comparison of the photon diffusion model and Kubelka–Munk
equation with the exact solution of the radiative transport equation. Appl. Opt.
13, 236–238 (1974).
28. Rein, S. Lifetime Spectroscopy: A Method of Defect Characterization in Silicon for
Photovoltaic Applications (Springer, 2004).
494
DOI: 10.1038/NPHOTON.2014.82
29. Takahashi, Y., Hasegawa, H., Takahashi, Y. & Inabe, T. Hall mobility in tin
iodide perovskite CH3NH3SnI3: evidence for a doped semiconductor. J. Solid
State Chem. 205, 39–43 (2013).
30. Ito, S. et al. Fabrication of thin ﬁlm dye sensitized solar cells with solar to electric
power conversion efﬁciency over 10%. Thin Solid Films 516, 4613–4619 (2008).
31. Bach, U. et al. Solid-state dye-sensitized mesoporous TiO2 solar cells with high
photon-to-electron conversion efﬁciencies. Nature 395, 583–585 (1998).
32. Noh, J. H., Im, S. H., Heo, J. H., Mandal, T. N. & Seok, S. I. Chemical
management for colorful, efﬁcient, and stable inorganic–organic hybrid
nanostructured solar cells. Nano Lett. 13, 1764–1769 (2013).
33. Kim, H.-S. et al. Mechanism of carrier accumulation in perovskite thin-absorber
solar cells. Nature Commun. 4, 2242 (2013).
34. Docampo, P. & Snaith, H. J. Obviating the requirement for oxygen in SnO2based solid-state dye-sensitized solar cells. Nanotechnology 22, 225403 (2011).
35. Mitzi, D. B., Dimitrakopoulos, C. D. & Kosbar, L. L. Structurally tailored
organic–inorganic perovskites: optical properties and solution-processed
channel materials for thin-ﬁlm transistors. Chem. Mater. 13, 3728–3740 (2001).
Acknowledgements
The authors thank T. Marks for use of the solar simulator and IPCE measurement system.
Electron microscopy and elemental analysis were carried out at the Electron Probe
Instrumentation Center (EPIC) at Northwestern University. This research was supported
as part of the ANSER Center, an Energy Frontier Research Center funded by the US
Department of Energy, Ofﬁce of Science, Ofﬁce of Basic Energy Sciences (award no.
DE-SC0001059) and ISEN at Northwestern University.
Author contributions
M.G.K. conceived the experiments and directed the study. F.H. and C.C.S. carried out the
material synthesis, device fabrication and performance measurements. D.H.C. prepared the
TiO2 blocking layer for the electrodes. R.P.H.C. contributed to the revision of the
manuscript. All authors discussed the results and commented on the manuscript.
Additional information
Supplementary information is available in the online version of the paper. Reprints and
permissions information is available online at www.nature.com/reprints. Correspondence and
requests for materials should be addressed to M.G.K.
Competing ﬁnancial interests
The authors declare no competing ﬁnancial interests.
NATURE PHOTONICS | VOL 8 | JUNE 2014 | www.nature.com/naturephotonics
© 2014 Macmillan Publishers Limited. All rights reserved.

Download Report