ARTICLESPUBLISHED ONLINE: 18 MAY 2015 | DOI: 10.1038/NPHOTON.2015.78

Thin-ﬁlm Sb2Se3 photovoltaics with oriented
one-dimensional ribbons and benign
grain boundaries
Ying Zhou1,2, Liang Wang1,2, Shiyou Chen3, Sikai Qin1,2, Xinsheng Liu1,2, Jie Chen1,2, Ding-Jiang Xue1,2,Miao Luo1,2, Yuanzhi Cao1, Yibing Cheng1, Edward H. Sargent4 and Jiang Tang1,2*
Solar cells based on inorganic absorbers, such as Si, GaAs, CdTe and Cu(In,Ga)Se2 , permit a high device efﬁciency andstability. The crystals’ three-dimensional structure means that dangling bonds inevitably exist at the grain boundaries(GBs), which signiﬁcantly degrades the device performance via recombination losses. Thus, the growth of singlecrystalline materials or the passivation of defects at the GBs is required to address this problem, which introduces anadded processing complexity and cost. Here we report that antimony selenide (Sb2Se3)—a simple, non-toxic and low-costmaterial with an optimal solar bandgap of ∼1.1 eV—exhibits intrinsically benign GBs because of its one-dimensional
crystal structure. Using a simple and fast (∼1 μm min–1) rapid thermal evaporation process, we oriented crystal
growth perpendicular to the substrate, and produced Sb2Se3 thin-ﬁlm solar cells with a certiﬁed device efﬁciency of 5.6%.
Our results suggest that the family of one-dimensional crystals, including Sb2Se3 , SbSeI and Bi2S3 , show promise inphotovoltaic applications.

T

he urgent need for high-efﬁciency, low-cost solar cells drives thesustained research on new absorber materials for thin-ﬁlm
photovoltaics. Copper zinc tin sulﬁde (CZTS) and organic–
inorganic metal halide perovskites (for example, CH3NH3PbI3)
are emerging and promising absorber materials that have achieved
impressive certiﬁed device efﬁciencies1,2. However, the complexityof defect physics associated with CZTS3,4 appears to limit further
signiﬁcant efﬁciency improvements in this system, and concerns
over stability and the reliance on lead in CH3NH3PbI3 (ref. 5)creates the need for new, stable, Pb-free materials.
These and other widely explored photovoltaic absorber materials(such as Si, GaAs, CdTe, InP and Cu(In,Ga)Se2 (CIGS)) exhibit athree-dimensional (3D) crystal structure6–8, that is, they are boundby covalent and/or ionic bonds in all three spatial dimensions. The
3D crystal structure guarantees an isotropic carrier transport andpermits a relaxed orientation control. However, at discontinuities
such as GBs, dangling bonds typically act as recombination centres,
and require further steps to be remedied 9–11. For example, in CdTe
(Fig. 1a) recombination loss through dangling bonds at the GBs
lowers the open-circuit voltage (VOC) even when the best-available
passivation methods are employed12.
We hypothesized that materials possessing a one-dimensional(1D) crystal structure (Fig. 1b), such as the ribboned compounds
Sb2Se3 and Bi2S3 , could, if suitably oriented, offer a compelling performance. In Sb2Se3 , for example, (Sb4Se6)n ribbons stack along the[001] direction through strong covalent Sb–Se bonds, whereas in the[100] and [010] directions the (Sb4Se6)n ribbons are held together byvan der Waals forces. Rarely have these materials been explored forphotovoltaics because of concerns about the poor carrier transport
between the ribbons, and worries about the mechanical stability of
the ﬁlms.

We posited that a new materials-processing strategy designed toorient ribbons vertically on the substrate could permit photogenerated carriers to travel efﬁciently along the covalently bonded 1D
axis. In the orthogonal direction, the parallel-stacked ribbons
would substantially provide no dangling bonds, even at GBs, andthereby minimize recombination losses, which is an important
advantage in photovoltaic applications because recombination
losses at GBs are one of the major limiting factors for high-efﬁciency
thin-ﬁlm solar cells.
We began with ﬁrst-principles simulations (VASP code) tostudy the structural relaxation and electronic structure of Sb2Se3
surfaces. The available terminations of (100) and (010) surfaces
in Sb2Se3 require no breaking of the Sb–Se bonds (Fig. 1c) andthus produce no dangling bonds. Other surfaces parallel to the[001] direction also have no dangling bonds, for example, the(110) and (120) surfaces. Their surface energy is therefore lower
than those of surfaces with dangling bonds, such as the (001),(211) and (221) surfaces. Calculation reveals that, as expected,
the surface energies for the (100), (010), (110), (120), (001),(211) and (221) surfaces, respectively, are 0.44, 0.25, 0.33, 0.32,
0.46, 0.56 and 0.53 J m–2. This indicates that the most-abundantsurfaces in Sb2Se3 samples will be the (010), (110) and (120)surfaces, that is, those with the lowest formation energies andno breakage of the covalent bonds. Density functional theory
conﬁrms quantitatively that no extra states are introduced inside
the bandgap by these terminations. A clean bandgap is seen in
the calculated density of states (DOS) of the four Sb2Se3 surfaces(Fig. 1d) and the gap is comparable to that of bulk Sb2Se3.
Furthermore, there is no appreciable change in the calculated
DOS over a wide energy range (−10 to 3 eV), which indicatesthat there is no signiﬁcant reconstruction on these surfaces. In

1

Wuhan National Laboratory for Optoelectronics (WNLO), Huazhong University of Science and Technology (HUST), Wuhan 430074, China. 2 School of
Optical and Electronic Information, Huazhong University of Science and Technology (HUST), Wuhan 430074, China. 3 Key Laboratory for Polar Materials
and Devices (MOE), East China Normal University, Shanghai 200241, China. 4 Department of Electrical and Computer Engineering, University of Toronto,
Toronto, Ontario M5S 3G4, Canada. * e-mail: jtang@mail.hust.edu.cn

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ARTICLES

NATURE PHOTONICS DOI: 10.1038/NPHOTON.2015.78

a
+
+
−
−
Te

n-type
TCO

Cd
GBs
[100]

[001]

b

+

+

−

−

n-type
TCO
Sb

Se

GBs

RibboncdBulk

DOS (a.u.)

(010)

(100)

(110)

(120)
(010)

(100)

−10

−8

−6

−4

−2

0

2

Energy (eV)Figure 1 | Crystal structure, DOS and recombination loss at the GBs in CdTe and Sb2Se3 solar cells. a, CdTe possesses a 3D crystal structure and has
dangling bonds (shown as red rods) at the GBs, which act as defects that cause a recombination loss of the photogenerated carriers. b, Sb2Se3(orthorhombic) is composed of (Sb4Se6)n ribbons stacked in parallel in the [001] direction. All the atoms at the edge of these ribbons are saturated(highlighted as red spheres) and introduce no recombination loss at the GBs once they are oriented vertically onto the substrates. For clarity, we omitted
from the image those (Sb4Se6)n ribbons that are present underneath this layer, but in a staggered manner. c, Atomic conﬁguration of the (010) and (100)surfaces (marked by the dashed black grids) in the Sb2Se3 crystal. d, DOS of the bulk Sb2Se3 and of four surfaces that are parallel to the [001] direction.
The forbidden gap is the empty zone near 0 eV. TCO, transparent conducting oxide.

sum, the ribbons represent the fundamental building block of the
Sb2Se3 structure, and surface reconstruction is negligible as longas this basic repeat unit is not broken.
Overall, this computational study conﬁrms that, as long as the
Sb2Se3 ribbons are suitably oriented, the GBs will be terminatedby the intrinsically benign surfaces (for example, (100), (010),(110) and (120) planes) and so the recombination loss minimized.
This is in striking contrast with most known photovoltaic absorbers,
in which the breakage of covalent bonds introduces defect states andrecombination centres at the GBs13–15.
A number of additional properties make Sb2Se3 particularly
worthy of experimental investigation as high-efﬁciency, low-costthin-ﬁlm solar cells16. Like CdTe, but in contrast with CIGS andCZTS, Sb2Se3 is a simple binary compound with a ﬁxed phaseand stoichiometry. It has a very strong absorption coefﬁcient(>105 cm−1 at short wavelengths) and its bandgap is ∼1.1 eV (refs
17,18), optimal for single-junction solar cells. The constituents of2

Sb2Se3 are non-toxic and low in cost (Sb has a similar cost to Cu),
and, as we proceed to show herein, Sb2Se3 ﬁlms are produced
using minimal energy, which, in principle, enables a short energypayback time for a solar cell19. All of these considerations motivated
the further study of oriented Sb2Se3 ﬁlms and the devices herein.
Sb2Se3 has a low melting point of 608 °C and a high saturated
vapour pressure (∼1,200 Pa at 550 °C)20, which enables us to carryout ﬁlm deposition using rapid thermal evaporation (RTE) in a tube
furnace with a high ramp rate (Supplementary Fig. 1). Sb2Se3
powder was directly applied via evaporation under a low vacuum
pressure (∼8 mtorr) maintained by a mechanical pump. Onceheated, Sb2Se3 powder evaporated and, because of the temperature
gradient, condensed on the substrate to form the Sb2Se3 ﬁlm. The
distance between the evaporating source and the substrate was kept
at the low value of 0.8 cm to enable a high material usage and fast
deposition rate. The deposition rate was as high as 1 μm min–1,
much greater than that of regular thermal evaporation (typically

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ARTICLES

NATURE PHOTONICS DOI: 10.1038/NPHOTON.2015.78
a

Au 100 nm
Thermal evaporation

I

p-Sb2Se3 390 nm
RTE-deposited
Vacuum, ~1 Pa
Deposition rate, 1 μm min−1

II

III

n-CdS 60 nm
Chemical bath depositon

FTO

100 nmbcd~0.320 nm~0.320 nm
(211)

(211)

~0.318 nm

(211)

(211)

~0.318 nm~0.318 nm(211) .

~0.320 nm(211).

Figure 2 | Structure and TEM analysis of Sb2Se3 ﬁlms and devices. a, Cross-sectional TEM image of the certiﬁed device. The nature of each layer and thecorresponding preparation procedures are given in the right column. The inset also shows the picture of a ﬁnished device; the nine yellow squares are thegold electrodes. HRTEM scans were performed at points I, II and III, and the corresponding lattice fringes are shown in b–d.

0.01–0.1 μm min–1) or sputtering (typically 0.01–0.05 μm min–1) andcomparable to conﬁned-space sublimation (CSS)21, a key technologythat has enabled the high manufacturing throughput and commercial
success of CdTe solar cells. The RTE process is distinct from CSS
because, in the RTE performed herein, Sb2Se3 melts and evaporates
from the liquid phase, in contrast to a direct sublimation from thesolid. Using this simple and fast technology, phase-pure Sb2Se3 ﬁlm(Supplementary Fig. 2) and solar cells were readily fabricated withhigh reproducibility.
We ﬁrst analysed, using transmission electron microscopy(TEM), speciﬁc regions of the Sb2Se3 ﬁlms made in actual solar
devices. The full device stack consisted of a ﬂuorine-doped tinoxide (FTO) substrate, a thin (∼60 nm) CdS layer produced
through chemical bath deposition, an approximately 390 nm thick
Sb2Se3 absorber layer deposited by the RTE process and top Au electrodes from thermal evaporation. For TEM characterization, thesample was prepared by cross-sectioning from the certiﬁed device
using a focused ion beam. The cross-sectional TEM image showedthat the Sb2Se3 ﬁlm was compact and composed of large Sb2Se3
grains of size equal to the ﬁlm thickness (Fig. 2a). To check
whether the grains were single crystalline, we applied highresolution transmission electron microscopy (HRTEM) to analyse
three arbitrarily selected points (I, II and III in Fig. 2a) with thecorresponding lattice fringes shown in Fig. 2b–d. The distances
between lattice lines were measured as 0.320 nm and 0.318 nm,

which corresponds to the separation of the (211) and (211) planes
in orthorhombic Sb2Se3. The crystal planes extend continuously
from the top to the bottom of the active region of the device,
which conﬁrms that this grain was single crystalline. This trend
was conﬁrmed with further studies of multiple grains(Supplementary Fig. 3), reinforcing the picture that the Sb2Se3
ﬁlms are made up of single-crystalline grains.
Orientational control of the Sb2Se3 ﬁlm is predicted to be crucial
to realize an efﬁcient carrier transport and benign GBs. We investigated, therefore, the possibility of a correlation between Sb2Se3 ﬁlm
orientation and photovoltaic device performance. Different substratetemperatures (Tsub) during the RTE process were explored with an
identical Sb2Se3 ﬁlm thickness maintained. Devices of Class A were
deposited onto the substrate kept at 300 °C, and devices of Class B
were deposited onto a 350 °C substrate. Both devices had an area of
0.095 cm2 deﬁned by the gold electrodes. Typical device performance(Fig. 3a) measured under 100 mW cm–2 illumination (Class 3A solar
simulator) revealed a short-circuit current density (JSC) for device
Class A of 27.2 mA cm–2, ﬁll factor (FF) of 53% and series resistance(RS) of 40 Ω, whereas the corresponding values in device Class B were18.4 mA cm–2, 47% and 89 Ω. We ﬁrst applied energy-dispersive
X-ray spectroscopy (Supplementary Fig. 4 and Supplementary
Tables 1–3) and capacitance voltage measurement (Supplementary
Fig. 5) to analyse our devices and found out that both the Sb2Se3
ﬁlms deposited at substrate temperatures of 300 °C and 350 °C were

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NATURE PHOTONICS DOI: 10.1038/NPHOTON.2015.78

a 30

(221)

Tsub = 350 °C(211)

Normalized intensity (a.u.)

Current density (mA cm−2)b

Tsub = 300 °C

25
20
15
10
5
0

300 °C

(120)

*
*

*

350 °CJCPDS 15-0861

−5
−10
0.1

0.2

0.3

0.4

10

20

30

Voltage (V)8Rs (Ω)

Efficiency (%)

6
5

Top

2.0
1.5
1.0

Se
Sb

0.5

+

+

0.0
0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2
TC(120)

4

60

d

2.5

160
140
120
100
80
60
40
20

RSh (kΩ)

7

50

1]

c

40
2θ (°)

[00

0.0

+

[001]

3

[120] orientation
0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

1.8

2.0

[221] orientation

Substrate

2.2

TC(120)Figure 3 | Device performance and its correlation with crystalline orientation. a,b, Solar cells with different photovoltaic performance (a) and thecorresponding XRD patterns (b). The standard diffraction pattern for Sb2Se3 (JCPDS 15-0861) is included for reference. The asterisks mark the peaks indexed
to the FTO substrate. The peak at 13° is associated by indexing to the excess of the trace element Se. c, Statistics of the TC(120) peaks with the device
efﬁciency, RS (left axis in the inset) and RSh (right axis in the inset) of the corresponding devices. To obtain the statistics presented, 140 devices were
included. d, Atomic structures of [120]- and [221]-oriented grains in Sb2Se3. The substrate is the CdS buffer layer on which the Sb2Se3 ribbons grow. On
photoexcitation, charge carriers must hop between ribbons in the [120]-oriented grain (dashed red arrows), but are able to move along the [221]-oriented
grain smoothly (solid red arrows).

slightly Se rich (the molar ratio Se:Sb was 1.51) and showed a very
similar doping density. This indicates that the ﬁlm compositionand consequent doping density probably did not account for theobserved device-efﬁciency difference. The large improvement fromdevice Class B (3.2% efﬁciency) to Class A (5.6% efﬁciency) cannot
be accounted for by stoichiometry and/or compositional differences,and we instead ascribed it to different ﬁlm orientations. Indeed,for device Class A, the diffraction intensity associated with the(120) peak was much weaker than that for the (211) peak, whereas
in device Class B the intensities of these two peaks were comparablewith each other (Fig. 3b).
To quantitate the orientation effects, we calculated the texture
coefﬁcient of the (120) (TC(120)) peaks from 140 devices andplotted the corresponding device efﬁciency versus texture coefﬁcient(TC) in Fig. 3c. The TC measures ﬁlm orientation (Supplementary
Fig. 6), with large TCs for a diffraction peak indicating a preferredorientation along this direction22. There existed a strong correlation
between device performance and ﬁlm orientation: device efﬁciency
monotonically decreased when the TC(120) orientation increased.
Further analysis revealed that for all of these devices, the shunt
resistance (RSh) remained constant at approximately 1,600 Ω,
which conﬁrms a similar junction quality, whereas RS increased
monotonically with an increased value of the TC(120). We can
explain this observation (that the preferred orientation along the[120] direction resulted in a signiﬁcantly increased RS and decreased
device efﬁciency) by the fact that the [120]-oriented grain consists of(Sb4Se6)n ribbons horizontally stacked in parallel with the substrate
4(Fig. 3d). In contrast, a [221]-oriented grain consists of tilted(Sb4Se6)n ribbons stacked vertically on the substrate. A [211]oriented grain is also composed of tilted (Sb4Se6)n ribbons, but toa different angle, and thus we limit our transport discussion to
the [221]-oriented grains. Naturally, carrier transport in the [211]oriented grains should be much easier than in the [120]-oriented
grains because in the former carriers travel within the covalently
bonded (Sb4Se6)n ribbons and in the latter carriers are required tohop between ribbons held together by van der Waals forces.
Furthermore, GBs of the [211]-oriented grains are composed ofplanes of low surface energy (hk0), such as the (100), (010), (110),(120) planes, which are free of dangling bonds and should cause
little recombination loss. In sum, we explain the lower RS , betterFF and higher JSC as the result of the improved transport andlowered recombination loss observed in devices with a preferred[211] orientation of the Sb2Se3 active layers.
For polycrystalline thin-ﬁlm solar cells, the passivation of GBs tosuppress strong carrier recombination is mandatory to achieve highefﬁciency devices. Typical examples are the carefully engineered Cudeﬁcient GBs in CIGS solar cells23,24 and a high-temperature CdCl2
treatment for CdTe solar cells24,25. We studied the properties of GBs
in our Sb2Se3 solar cells using Kelvin probe force microscope(KPFM) and electron-beam-induced current (EBIC) measurements. Two-dimensional topography spatial maps and the corresponding surface potentials of Sb2Se3 thin ﬁlms (Fig. 4a,b) revealthat there is no correlation between GBs (identiﬁable by notablechanges in surface topography) and the substantial potential

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ARTICLESa

250 nm/ c
200 mVb

0.5 μm

2
40

0
−2

30

−4
20

−6
−8

10

−10
−12

0

−250 nm/−200 mV

100

200

300

400

500

600

x (nm)

f 33,780

e

d

50

4

Height (nm)0.5 μmContact potential difference (mV)

NATURE PHOTONICS DOI: 10.1038/NPHOTON.2015.78

A

A

B
C
D

B
C
D

Intensity (a.u.)

33,76033,74033,72033,7000.5 μm33,6800.5 μm

0.0

0.3

0.6

0.9

1.2

1.5

Distance (μm)

Figure 4 | Surface potential at Sb2Se3 GBs and EBIC images from a crystallographically well-oriented device (deposited onto 300 °C substrates).
a,b, Atomic force microscopy (a) and scanning KPFM image (b) of the device. The scanning area is 2.5 μm × 2.5 μm. c, Contact potential difference (red)and height (blue) along the solid black lines drawn in a and b. The valleys of the height curve correspond to the GBs. d,e, Cross-sectional SEM (d) and EBIC
images (e) for the cleaved Sb2Se3 devices. For the coloured legend, A is Au, B is Sb2Se3, C is CdS and D is FTO. White parts indicate the zones in which
carriers induced by the electron beam were efﬁciently collected and grey areas indicate the Sb2Se3 zones that contributed less photocurrent. f, The intensity
proﬁling along the dashed blue line marked in e. The scattering in signal intensity is negligible in the scanned 1.5 μm length, which suggests a uniform
photocurrent-collecting efﬁciency even across GBs.

variation in the KFPM image. Overall, the average roughness ofan Sb2Se3 ﬁlm is 23 nm and the average surface potential
difference is a very low 9.1 mV (much below the thermal energy,kBT). The surface-potential ﬂuctuations are extremely small
compared to those of CIGS and CZTS ﬁlms (generally >100 mV
in a 2.5 μm × 2.5 μm zone9,25). In an illustrative line scan crossing
the GBs (Fig. 4c), the surface potential difference between two
grains is as low as 10 mV, which indicates a lack of signiﬁcantband bending and surface defects in the Sb2Se3 ﬁlms.
Further scanning electron microscopy (SEM) and EBIC analysesof the Sb2Se3 solar cells are reported in Fig. 4d,e, respectively. The
thicknesses of the Sb2Se3 and CdS layers were 390 nm and 60 nm,
respectively, which is consistent with the results obtained from theTEM characterization (Fig. 2a). Bright areas in an EBIC image indicate regions of high collection efﬁciency for minority carriers.
Whereas in CdTe and CIGS a high contrast indicates a nonuniform collection efﬁciency26, the Sb2Se3 solar cell exhibits noappreciable variation in the EBIC signal across the absorber layer
close to the CdS layer (Fig. 4e) and, more importantly, also acrossa 2 μm long line scan (Fig. 4f ). As we have already shown, thesame materials processing led to ∼300 nm grain sizes, results thatindicate the EBIC signal, and hence collection efﬁciency, is
uniform over multiple grains, excluding severe carrier recombination or an efﬁcient current-collecting path presented at GBs. The
EBIC ﬁndings echo the KFPM results above and clearly demonstratethat the GBs and grain interior are indistinguishable in our Sb2Se3
ﬁlm, an expected consequence of GBs free of dangling bonds.
We obtained certiﬁcation of the solar cells reported herein(Fig. 5a,b, Newport Certiﬁcate Number (NCN) 1017). In all thedevices reported herein, illumination is from the FTO side andphotogenerated carriers separate at the p-Sb2Se3/n-CdS interface.
Electrons are injected into the CdS and collected by the FTO, andholes travel through the Sb2Se3 absorber and are collected by theback Au ohmic contact. The certiﬁed device (0.095 cm2) exhibiteda JSC of 25.1 mA cm–2, a VOC of 0.40 V and an FF of 55.7%, which corresponds to a power-conversion efﬁciency (PCE) of 5.6% (Fig. 5a).
This agrees with the in-house measurements, and in each case no
hysteresis between the forward and reverse scans was observed(Supplementary Fig. 7). We also built large-area (1.08 cm2) devices
and obtained a device efﬁciency of 5.4% measured with the aid ofaperture and mechanical scribing (Supplementary Figs 8 and 9).
This device efﬁciency surpasses that of many other emerging
inorganic solar cells based on Cd-free active absorbers, such as SnS(ref. 27) and FeS2 (ref. 28). Measurement of the external quantum
efﬁciency (EQE) showed that the photocurrent contribution
extended to 1,070 nm, consistent with the measured optical bandgap. By dividing the EQE spectrum by the optical absorption, we
obtained an internal quantum efﬁciency (IQE) close to 100% in
the 550–600 nm solar spectrum. The high IQE is again consistent
with the low recombination loss assigned to the benign GBs
studied in detail herein. The photovoltaic performance obtained
across a study of 100 devices (Fig. 5c) demonstrates a tight distribution with an average efﬁciency of 5.24% and a standard deviationof 0.34%. The device efﬁciency (Fig. 5d) increased by about
0.3 power points (from 5.32% to 5.66%) when stored in ambient
air for up to 45 days without encapsulation, in good agreementwith our previous reports of thermally evaporated Sb2Se3 solar
cells19,29. Damp-heat testing (85 °C and 85% humidity, no encapsulation) of the same device showed a slightly reduced device performance, from an initial 5.67% efﬁciency to 5.39% after 10 hours of
testing and to 5.16% after 100 hours of testing. This amountof degradation is considerably lower than that of many next-generation solar cells, such as organic–inorganic halide perovskite andpolymer solar cells5.

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ARTICLES
I–V NCN 1,017
2.5

b 100

100

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60

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40

20

20

1.5
1.0
0.5

EQE (%)VOC = 0.40 V
JSC = 25.14 mA cm−2
FF = 55.7%
η = 5.6%
Area = 0.095 cm2

0.0
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Voltage (V)

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25 days1h
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100 h

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PCE (%)

Number of devices

IQE (%)

Current (mA)a

NATURE PHOTONICS DOI: 10.1038/NPHOTON.2015.78

20
15

3
2

10

25 °C
in ambient air

85 °C
85% humidity

1

5
0

0
4.6

4.8

5.0

5.2

5.4

5.6

Time

5.8

PCE (%)

Figure 5 | Device performance (certiﬁed) and stability. a,b, Current–voltage (I–V) (a) and EQE and IQE spectra (b) of the FTO/CdS/Sb2Se3/Au solar cell
independently certiﬁed by the Newport Cooperation (Calibration Certiﬁcate No. 1017). c, Histogram of device efﬁciencies obtained from over 100 individuallyfabricated devices: VOC = 0.385 ± 0.012 V, JSC = 25.47 ± 1.75 mA cm–2 and FF = 53.37 ± 3.72%, with efﬁciencies of 5.24 ± 0.34%. d, The stability of a typicaldevice without package stored under regular laboratory conditions (ambient air, no shading) and then subjected to a damp–heat measurement (85 °C, 85%
humidity, in the dark).

The path to improved solar PCEs requires major progress on
the current (from 25 mA cm–2 towards the maximum available
42 mA cm–2 for the 1.1 eV bandgap), the voltage (from 0.4 to
0.75 V for this bandgap) and the FF (from a 50% range into an
80% range). The low EQE at short wavelengths (Fig. 5b) suggeststhat the CdS front layer should be replaced by a more desirable
junction-forming material with a larger bandgap and an environmentally friendly composition, such as TiO2 or ZnO. Although
the ambipolar transport is achieved efﬁciently over ∼150 nm(>80% EQE at a wavelength of 600 nm at which the absorption
length is ∼150 nm), these materials and devices require an indepth study to determine whether the minority carrier diffusion
length requires extending, or whether an electron-blocking backcontact (back surface ﬁeld) is the major requirement for an
improved current. The modest VOC and FF combined with the evidence of benign GBs suggests that electrical interfaces to Sb2Se3(for an efﬁcient charge extraction, maximal quasi-Fermi-level splitting and the blocking of one carrier at each interface) deserves
signiﬁcant attention.
In sum, this work presents a non-toxic, stable and manufacturable method of solar-cell fabrication that, in the initial trials, led
to a certiﬁed device efﬁciency of 5.6%. The work reveals thatmaterials with a 1D crystal structure, little explored in photovoltaics,
offer promise as absorbers once the ﬁlms are oriented suitably fortransport. We have shown that these materials can simultaneously
sustain excellent transport along one axis and minimal recombination (owing to the benign GBs) along the orthogonal axes. The
family of 1D materials includes other V2–VI3 compounds, such as
Bi2S3 , and V–VI–VII compounds, such as SbSeI, which all
possess attractive material and optoelectronic properties.
Analogously, 2D binary layered materials, such as MoSe2 and6

WSe2 , when similarly aligned to a substrate are expected to
exhibit sufﬁcient transport and manifest the beneﬁts of covalently
bonded planes and partially benign GBs.

Methods
Methods and any associated references are available in the online
version of the paper.
Received 11 November 2014; accepted 14 April 2015;
published online 18 May 2015

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Acknowledgements
This work is supported by the Director Fund of WNLO, the National 1000 Young Talents
project, the National Natural Science Foundation of China (NSFC 61274055, 91233121,
91433105, 21403078) and the 973 Program of China (2011CBA00703). The authors thankthe Analytical and Testing Center of HUST, the Center for Nanoscale Characterization andDevices of WNLO and the Suzhou Institute of Nano-Tech and Nano-Bionics for thecharacterization support. Y. Yan at the University of Toledo and H. Zhong at the Beijing
Institute of Technology are acknowledged for helpful discussions.

Author contributions
Y.Z. and J.T. conceived the idea, designed the experiments and analysed the data. Y.Z. and
L.W. carried out most of the characterizations and device optimizations. S.C. performed thetheoretical simulations and analysed the results. S.Q. and X.L. initialized the RTE process.
J.C., D-J.X., M.L. and Y.C. participated in the device optimization and data analysis.
Y.Ch. helped with the manuscript preparation. E.H.S. and J.T. wrote the paper; all theauthors commented on the manuscript.

Additional information
Supplementary information is available in the online version of the paper. Reprints andpermissions information is available online at www.nature.com/reprints. Correspondence andrequests for materials should be addressed to J.T.

Competing ﬁnancial interests

The authors declare no competing ﬁnancial interests.

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7

ARTICLES

NATURE PHOTONICS DOI: 10.1038/NPHOTON.2015.78

Methods
Solar cell fabrication. All devices were deposited on glass-coated SnO2:F suppliedby Kaivo (Zhuhai, China) that had been cleaned using a detergent, acetone,
isopropanol and water rinse in sequence. Chemical bath deposition (CdSO4 , 65 °C,
16 minutes) was used to deposit a CdS layer of about 60 nm thickness. A CdCl2 posttreatment was applied to the CdS buffer layer: (1) a CdCl2 (99%, Aladdin) anhydrous
methanol saturated solution (about 20 mg ml–1) was spin coated onto the CdS layer
and then rinsed with methanol; (2) the treated CdS layer was baked on a hot plate at
400 °C for ﬁve minutes in air. An Sb2Se3 layer was deposited by RTE in a tube
furnace (MTI, Hefei, China). The set-up is shown schematically in Supplementary
Fig. 1. Sb2Se3 powder was placed on the AlN plate inside the quartz boat, and the
FTO/CdS substrate was placed on top of the quartz boat (0.8 cm distance from the
Sb2Se3 powder) with the CdS side facing down. When the pressure was reduced to
below 10 mtorr, the ﬁlm deposition was initiated. The deposition process was topreheat the source and substrate at 300 °C for 15 minutes, then ramp thetemperature up to approximately 550 °C within 30 seconds, maintain thistemperature for 30 seconds and then turn off the heating and allow the ﬁlm to cool
naturally to ∼150 °C. Films were then taken out for Au deposition. Gold backcontacts (0.095 cm2) were thermally evaporated using the electron beam and a
resistance evaporation thin-ﬁlm deposition system (Beijing Technol Science Co. Ltd)
under a vacuum pressure of 5 × 10−3 Pa.
Solar-cell performance measurement. The devices were measured under asimulated 100 mW cm–2 AM1.5G irradiation, which was generated through a 3A
solar simulator with a Xe light source (450 W, Oriel, Model 9119). A Keithley 2400
was applied to measure the current–voltage characteristics in air. The device area
(0.095 cm2) was deﬁned by the size of the gold electrodes. For the in-housemeasurements, no temperature control was applied. The scan rate was 0.1 V s–1, andno hysteresis was observed between the forwards (JSC to VOC) and reverse (VOC to JSC)
scan (Supplementary Fig. 7). We also applied aperture (using a metal mask to deﬁne
the incident light) and mechanical scribing to measure some of the devices andobtained comparable device efﬁciencies (Supplementary Figs 8 and 9). The device
was further independently certiﬁed by the Newport Cooperation (American
Certiﬁcate Number 1017 issued on 16 May 2014, without aperture). The ageing of
the devices below 85 °C with 85% humidity was applied in a desktop constanttemperature and humidity testing machine, QA-HPZ-30 (Dongguan Cheng Yang
Instruments Industry Co. Ltd).
XRD. XRD measurements were performed using a Philips diffractometer(Xpert PRO MRD) with a step of 0.017 ° and a step time of 10.16 seconds. Thelines used were Cu Kα1 and Cu Kα2 with wavelengths of 1.54060 nm and1.54443 nm, respectively.

TEM and selected area electron (SAED) characterization. Samples for TEM and
SAED analyses were prepared by ablating the certiﬁed device using a FEI Quanta 3D
FEG-FIB. A thin Pt layer was ﬁrst deposited on top of the device for protection. TEM
images and SAED measurements were taken using a FEI Titan G2 60–300 Probe Cs
Corrector STEM operating at 300 kV.
KPFM. KFPM was performed on a grounded Sb2Se3 sample using a NT-MDT
Dimension 3100 system. The topography and surface potential were simultaneously
measured using a Pt-coated silicon probe (NT-MDT, HQ: NSG11/Pt, f0 ≈ 150 kHz,k ≈ 5.5 N m–1 and sampling rate 0.25 Hz). The surface potential was calculated fromthe electrostatic force between the probe tip and the surface of the sample using a
lock-in ampliﬁer. The scanning area was 2.5 μm × 2.5 μm of 512 × 512 pixels. The
minimal distance between two scanned points was about 5 nm, and the potential
measurement limit and resolution were 10 mV and 1 mV, respectively.
EBIC measurement. The device was cleaved by hand after scribing on the back andloaded into an SEM vacuum chamber for EBIC measurements. SEM images weretaken on a Quanta 400 FEG FEI microscope equipped with a SmartEBIC (Gatan,Inc.). The gold back contact was contacted to the preampliﬁer by a small
micromanipulator to enable current measurements. The electron-beam voltage was
20 kV and the working distance was about 10.5 mm. Images were processed using
the Digital micrograph software.
Simulation methods. The structural relaxation and electronic structure of Sb2Se3
were calculated within the density functional formalism as implemented in the
VASP code30. For the exchange-correlation potential, we used the non-local vdWDF (van der Waals density functional) method proposed by Dion et al.31 to describe
the van der Waals interaction. The interaction between the core electrons and thevalence electrons is included in the frozen-core projector augmented-wave
pseudopotentials. An energy cut-off of 400 eV was applied in all cases. For Brillouinzone integration, we used the 4 × 4 × 12 Monkhorst-Pack k-point mesh for the
20-atom primitive cell, and 2 × 6 × 1 (2 × 2 × 1) meshes for the (100) [(010)] surface
slab models. The lattice vectors of the primitive cell and the atomic positions were
fully relaxed by minimizing the quantum mechanical stresses and forces.

References
30. Kresse, G. & Furthmüller, J. Efﬁcient iterative schemes for ab initio total-energy
calculations using a plane-wave basis set. Phys. Rev. B 54, 11169–11186 (1996).
31. Dion, M., Rydberg, H., Schröder, E., Langreth, D. C. & Lundqvist, B. I. Phys.
Rev. Lett. 92, 246401 (2004).

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