Low-Temperature Solution-Grown CsPbBr3 Single Crystals and Their Characterization

Article abstract of DOI:10.1021/acs.cgd.6b00764

Cesium lead bromide (CsPbBr₃) was recently introduced as a potentially high performance thin-film halide perovskite (HaP) material for optoelectronics, including photovoltaics, significantly more stable than MAPbBr₃ (MA = CH₃NH₃⁺). Because of the importance of single crystals to study relevant material properties per se, crystals grown under conditions comparable to those used for preparing thin films, i.e., low-temperature solution-based growth, are needed. We show here two simple ways, antisolvent-vapor saturation or heating a solution containing retrograde soluble CsPbBr₃, to grow single crystals of CsPbBr₃ from a precursor solution, treated with acetonitrile (MeCN) or methanol (MeOH). The precursor solutions are stable for at least several months. Millimeter-sized crystals are grown without crystal-seeding and can provide a 100% yield of CsPbBr₃ perovskite crystals, avoiding a CsBr-rich (or PbBr₂-rich) composition, which is often present alongside the perovskite phase. Further growth is demonstrated to be possible with crystal seeding. The crystals are characterized in several ways, including first results of charge carrier lifetime (30 ns) and an upper-limit of the Urbach energy (19 meV). As the crystals are grown from a polar aprotic solvent (DMSO), which is similar to those used to grow hybrid organic-inorganic HaP crystals, this may allow growing mixed (organic and inorganic) monovalent cation HaP crystals.

Full text of DOI:10.1021/acs.cgd.6b00764

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Article
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Low-Temperature Solution-Grown CsPbBr
Single Crystals and Their Characterization
Yevgeny Rakita, Nir Kedem, Satyajit Gupta, Aditya Sadhanala, Vyacheslav Kalchenko,
Marcus L. Böhm, Michael Kulbak, Richard H. Friend, David Cahen, and Gary Hodes
Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.6b00764 • Publication Date (Web): 22 Aug 2016
Downloaded from http://pubs.acs.org on August 23, 2016
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LowꢀTemperature SolutionꢀGrown CsPbBr3
Single Crystals and Their Characterization
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Yevgeny Rakitaa, Nir Kedema, Satyajit Guptaa, Aditya Sadhanalab, Vyacheslav
Kalchenkoc, Marcus L. Böhmb, Michael Kulbaka, Richard H. Friendb, David Cahena*
and Gary Hodesa*
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a. Department of Materials and Interfaces, Weizmann Institute of Science, Rehovot, 7610001, Israel
b. Cavendish Laboratory, JJ Thomson Avenue, Cambridge CB3 0HE, United Kingdom
c. Departments of Veterinary Resources; Weizmann Institute of Science, Rehovot, 7610001, Israel
Keywords: Halide Perovskite, Single Crystal, Photovoltaics, Urbach energy
Abstract:
Cesium lead bromide (CsPbBr3) was recently introduced as a potentially high performance
thinꢀfilm halide perovskite (HaP) material for optoelectronics, including photovoltaics,
+
significantly more stable than MAPbBr3 (MA=CH3NH3 ). Because of the importance of single
crystals to study relevant material properties per se, crystals grown under conditions
comparable to those used for preparing thin films, i.e. low-temperature solution-based growth,
are needed. We show here two simple ways: antiꢀsolventꢀvapor saturation or heating a solution
containing retrograde soluble CsPbBr3, to grow single crystals of CsPbBr3 from a precursorꢀ
solution, treated with acetonitrile (MeCN) or methanol (MeOH). The precursor solutions are
stable for at least several months. Millimeterꢀsized crystals are grown without crystalꢀseeding
and can provide a 100% yield of CsPbBr3 perovskite crystals, avoiding a CsBrꢀrich (or PbBr2ꢀ
rich) composition, which is often present alongside the perovskite phase. Further growth is
demonstrated to be possible with crystalꢀseeding. The crystals are characterized in several
ways, including first results of charge carrier lifetime (30 ns) and an upperꢀlimit of the Urbach
energy (19 meV). As the crystals are grown from a polar solvent (DMSO), which is similar to
those used to grow hybrid organicꢀinorganic HaP crystals, this may allow growing mixed
(organic and inorganic) monovalent cation HaP crystals.
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* Gary Hodes: gary.hodes@weizmann.ac.il
David Cahen: david.cahen@weizmann.ac.il
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Introduction:
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In the last few years, halide perovskites (HaPs) with AMX3 stoichiometry (where A is
methylammonium (MA), formamidinium (FA) or cesium (Cs) ; M=Pb or Sn ; X=Cl, Br, I)
have attracted attention because of their remarkable optoelectronic (especially photovoltaic)
properties, compared to other competitive, lowꢀcost, optoelectronic materials (e.g., various
organic molecules and polymers, CdX, PbX (X=S, Se, Te), Cu(In,Ga)Se2).1,2 Some
compositions of AMX3 yield small area photovoltaic lightꢀtoꢀelectrical conversion efficiencies
above 21%,2 and also relatively efficient and tunable lightꢀemitting diodes,3,4
photodetectors/resistors5,6 and highꢀenergy radiation detectors7. To help understand the
remarkable optoelectronic properties of these materials, one would like to be able to study bulk
material of controlled quality, prepared under conditions that are not too different from those
used for making films employed in most HaPꢀbased devices. Clearly single crystals are a
desirable form of materials for this purpose.
CsPbBr3, has recently gained attention as a stable HaP,8,9 but hitherto no method was
available to prepare macroscopic crystals by the low temperature solution methods that are
used to make films for photovoltaic cells. For the hybrid organic–inorganic halide perovskites
(where A is usually MA or FA), it was already demonstrated that single crystals with mm to cm
dimensions can be grown using different solutionꢀbased methods at lowꢀtemperature (around or
below 100°C): antiꢀsolvent vapor saturation,10,11 slow heating of a solution with a retrogradeꢀ
soluble compound,12 or topꢀseeded slowꢀcooling in a precursorꢀsolution.13,14 However, organic
elements in these compositions can act as ‘weak links’ and decrease the stability (and thus
reliability) of the entire structure, as was demonstrated in a comparison between the MAPbBr3
and CsPbBr3 HaPs.8 It was also shown that exchanging to the fully inorganic compound (i.e.
CsPbBr3) does not necessarily exact a price in photovoltaic performance.9 To be able to use
single crystals to study CsPbBr3 and compare it with its organicꢀinorganic hybrid analogs, the
crystals should preferably be prepared under conditions that are comparable to those of the thin
films used in device work. While the hybrid HaPs can indeed be grown as macroscopic crystals
from solution, no such process was available for CsPbBr3 crystals. Here we report on this
missing link.
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At standard temperature and pressure, the perovskite structure is the thermodynamically
mostꢀstable structure of CsPbCl3 and CsPbBr3 (but not of CsPbI3).15 The only reported
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procedure of making mmꢀsize single crystals of these materials is by Bridgman growth, where
precursors are melted (above 600°C) in a quartz tube and passed through a multiꢀtemperature
zone tube furnace).7,16
For hybrid organicꢀinorganic perovskites, the known compositions that can be created have
a single stoichiometric composition (i.e. AMX3) regardless of precursor ratio and present a
retrogradeꢀsolubility in a variety of organic solvents,17 which allow an efficient and easy crystal
growth. For the fully inorganic analog (for our particular interest: CsPbBr3), on the other hand,
the situation is more challenging. First, similar to compositions with other alkaline A atoms,
such as RbBr:PbBr2, in the CsBr:PbBr2 system 1:2 and 4:1 compositions, as well as the 1:1
perovskite one, are known.18 Based on the phaseꢀdiagram of PbBr2:CsBr compositions,19 it is
difficult to prepare a pure CsPbBr3 compound, since either the CsBrꢀrich Cs4PbBr6 or PbBr2ꢀ
rich CsPb2Br5 are expected to be formed at slight deviation from a perfect 1:1 PbBr2:CsBr ratio.
In addition, there is a large difference in the solubility of PbBr2 and CsBr in aprotic solvents,
making a big difference in the reactivity of the dissolved species, which results in precipitation
of CsBrꢀ or PbBr2ꢀrich compounds alongside the desired CsPbBr3 perovskite. Since, usually,
CsBr is less soluble than PbBr2 in aprotic solvents like dimethylformamide, gammaꢀ
butyrolactone, dimethylsulfoxide (generally used for growing hybridꢀhalide perovskites
crystals), CsBrꢀrich products (e.g., CsBr:PbBr2 = 4:1; Cs4PbBr6) tend to form. Thus, growing
the 1:1 (CsPbBr3) composition presents a serious challenge.
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Although Cs4PbBr6, was usually evident as the main or side product, we found conditions
that allow growth of pure CsPbBr3 crystals, without the presence of additional phases. Two
lowꢀtemperature (< 100ºC, down to roomꢀtemperature) methods were used:
ꢀ1ꢀ slow antiꢀsolvent vapor saturation of a precursor solution and
ꢀ2ꢀ slow heating of a solution containing CsPbBr3 as a retrogradeꢀsoluble compound.
Here we describe, for the first time, procedures to grow pure CsPbBr3 crystals with dimensions
of several mm from a precursor solution. The precursor solution was found to be stable for
several months, and larger crystals (using CsPbBr3 crystal seeds) could be further grown with
it. We use several characterization techniques to assess the potential of these solutionꢀgrown
crystals for optoelectronics and photovoltaics, including first results of charge carrier lifetime
and of the Urbach energy.
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Protocol for crystal growth
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Precursors
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Powders of cesium bromide (CsBr) (Aldrich, 99.999%) without further purification and
lead bromide (PbBr2) (Aldrich, ≥98%) that was dried in a vacuum oven overnight, were
dissolved in dimethylsulfoxide (DMSO) (Aldrich, ≥99.9%). Acetonitrile (MeCN) (BioLab
LTD, HPLC grade, 99.97%) and methanol (MeOH) (BioLab LTD, HPLC grade, 99.95%) were
used as received.
Precursor Solution preparation
A 0.45 M solution (slightly below the ~ 0.5 M solubility limit) of the perovskite precursors
(equimolar amounts of CsBr and PbBr2) in the same solution of DMSO was prepared in
ambient air (ca. 45% RH) under continuous stirring at ~50oC, until no powder is observed. It is
important to mix the precursors in the same volume, because the room temperature solubility
limit of CsBr in DMSO is ca. 0.25 M when it is dissolved separately. After cooling to room
temperature, the DMSO solution was titrated (dropwise under continuous stirring) with MeCN
or MeOH. During the titration, a yellowꢀorange precipitant appeared with addition of each drop
and quickly reꢀdissolved. As the system gets closer to the saturation point (more pronounced
when MeCN is added), a permanent white solid precipitates (Figure S1(a)). At MeCN : DMSO
and MeOH : DMSO ratios of 1.1 : 1 and 0.55 : 1, respectively, the yellowꢀorange precipitate no
longer reꢀdissolved. These saturated solutions were thoroughly sealed (to prevent loss of the
volatile MeCN or MeOH) and stirred for 24 hours at 50 °C. After heating for 24 h, a pale
greenꢀyellow stronglyꢀfluorescent precipitant was clearly observed alongꢀside the other
precipitated species (Figure S1(b)). The saturated solutions can be further stored (best in the
dark) for at least several months, until their use for crystal growth. Before crystal growth, the
saturated precursor solutions are filtered with PTFE 0.2 ꢁm poreꢀsize syringe filters. No
noticeable differences in crystal growth or kinetics were observed between properly stored or
freshly prepared solutions.
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Crystal growth
-1- slow Vapor-Saturation of an Anti-solvent (VSA)
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The concept of slow saturation with the vapor of an antiꢀsolvent was shown earlier to be a
possible method to grow different hybrid organicꢀinorganic HaPs.10,11 A scheme of the
crystallization setup is presented in Figure S2(a). The filtered precursor solutions were placed
in a clean crystallization flask and covered with a filter paper and a glass petri dish on top to
limit antiꢀsolvent vapor diffusion. The covered crystallization flask was then placed inside a
deeper, flatꢀbottomed, glass dish, which contained MeCN or MeOH (antiꢀsolvent). The antiꢀ
solvent was chosen to be similar to the one previously used for saturating (titrating) the DMSO
precursorꢀsolution. Sealing the outer glass flask was done using a filter paper and glass top. In
this case, the filter paper was used to create a balanced antiꢀsolvent atmosphere and to prevent
the antiꢀsolvent from condensing and dripping on the crystallization innerꢀsetup. While crystal
formation can occur at room temperature, heating the antiꢀsolvent bath (to allow a higher antiꢀ
solvent residual vapor pressure) will accelerate the crystal growth. This can be done by placing
the crystallization setup on a hot plate with the addition of an elevating element for the inner
crystallization flask to minimize a nonꢀuniform heat dissipation coming from the interface with
the hotplate.
We note that H2O can also be used as an antiꢀsolvent to grow CsPbBr3 crystals. However,
using H2O causes the orange crystals to bleach if the amount of the antiꢀsolvent is not carefully
controlled (see Figure S3). For that reason, crystals that were grown by using H2O as an antiꢀ
solvent were not further investigated.
Figure 1(a) shows the results of crystal growth from MeCNꢀ and MeOHꢀ saturated
solutions. A 48 hour growth from MeCNꢀsaturated solution showed wellꢀfaceted rectangular
prisms (Figure 1(a,i)), while those grown from MeOH showed a terraceꢀlooking morphology
(Figure 1(a,ii)). If the same procedure is performed without the preꢀsaturation step, paleꢀ
yellow and strongly fluorescent precipitant, relates to Cs4PbBr6,20 appear, followed by growth
of CsPbBr3 orange crystals on top of the Cs4PbBr6 crystallites (Figure 1(a,iii)).
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-2- Slow Heating of a solution with a Retrograde-Soluble compound (HRS)
The slow heating of a solution with a retrogradeꢀsoluble (HRS) concept was demonstrated
before as an additional method to grow different hybrid organicꢀinorganic HaPs.5,17,21 To grow
CsPbBr3 crystals the procedure is as follows: clean 4 mL or 20 mL glass vials were filled with
filtered precursorꢀsolutions and heated to the desired temperature. An unsaturated DMSO
precursorꢀsolution showed no retrogradeꢀsoluble compounds. When saturated with MeCN, and
especially, with MeOH, retrogradeꢀsoluble compounds (mainly, but not only, CsPbBr3) appear
upon heating. To eliminate the appearance of undesirable precipitants (mostly the strongly
yellowꢀgreen fluorescing Cs4PbBr6) a twoꢀstep heating cycle was required. First, the solution
was heated to the desired temperature for 4 h and left to cool to roomꢀtemperature under
continuous stirring. After all the orange precipitant dissolved, a mixture of Cs4PbBr6 and some
unknown white precipitate was still present at the bottom of the flask. The solution was then
filtered and placed in a new vial. The second heating cycle showed no additional retrograde
soluble compounds apart from the orange CsPbBr3 crystals (Figure S1(c)). From the MeCNꢀ
saturated solution, CsPbBr3 crystals start to appear clearly only above ca. 120 ºC. From the
MeOHꢀsaturated solution, however, this occurs already at ca. 40º C. To avoid crystalꢀgrowth on
the sides of the glass vial, one can heat only the bottom of the vial. Due to the very low yield of
crystal growth obtained from the MeCNꢀsaturated solution, we further explored the HRS
method with the MeOHꢀsaturated solution. Figure S2(b) shows a scheme of the HRS
crystallization setup.
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Figure 1(b, i & ii) show CsPbBr3 crystals grown from a MeOHꢀsaturated solution with
and without crystal seeding with a growth rate of 10°C /hour up to 80°C. If the same procedure
is performed with only one heating cycle (Figure 1(b,iii)), a strongly fluorescent CsBrꢀrich
compound (Cs4PbBr6) is clearly evident together with the CsPbBr3 crystals.
MeCN and MeOH were chosen from among 9 volatile commonlyꢀused antiꢀsolvents viz.
dibromoethane, dibromomethane, ethyl acetate, 2ꢀpropanol, tetrahydrofuran, acetone, H2O,
MeCN and MeOH. All of them, except H2O, MeCN and MeOH, showed rapid precipitation of
the CsBrꢀrich compound, without obvious signs of CsPbBr3 orange crystal growth. Therefore,
they were not further investigated. Details about the different antiꢀsolvents that did not yield
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CsPbBr3 crystals from DMSO solution (and for H2O) can be found in the last section of the
Supplementary Information.
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The following equations provide a general description of formation of the various compounds
we describe in this paper
:
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(1)
(2)
(3)
ꢀꢁꢂꢃꢄꢅꢆꢇ ꢈ ꢉꢊꢋꢌ → ꢀꢁꢂꢃꢄ
ꢑꢒꢂꢃ ꢈ ꢉꢊꢋꢌ → ꢑꢒꢂꢃꢅ
ꢅ
ꢇ
ꢍꢎꢏꢐ
ꢇ
ꢍꢎꢏꢐ
ꢅꢆꢇ
ꢕꢅ ꢇ ꢖ
ꢓ ꢀꢁꢂꢃꢄ ∙ ꢔ ꢑꢒꢂꢃ
ꢍꢎꢏꢐ
ꢅ ꢇ
ꢓ ∙ ꢀꢁꢂꢃꢄ
ꢈ ꢔ ∙ ꢑꢒꢂꢃꢅ
→
ꢇ
ꢍꢎꢏꢐ
ꢅ
ꢇ
ꢍꢎꢏꢐ
ꢘꢙꢚꢛꢜꢆꢝꢞꢟꢠꢙꢚ/ꢡ
ꢕꢅ ꢇ
ꢅ ꢇ
ꢅ ꢇ
ꢖ
ꢅ
ꢇ
(4)
(5)
ꢓ ꢀꢁꢂꢃꢄ ∙ ꢔ ꢑꢒꢂꢃ ꢍꢎꢏꢐ ꢈ 4ꢓ ꢗ ꢔ ∙ ꢑꢒꢂꢃꢅ
ꢢꢣꢣꢣꢣꢣꢣꢣꢣꢣꢣꢤ ꢓ ∙ ꢑꢒꢥꢀꢁꢂꢃꢦ ꢆ
ꢇ
ꢅ ꢇ
ꢍꢎꢏꢐ
ꢘꢙꢚꢛꢜꢆꢝꢞꢟꢠꢙꢚ/ꢡ
ꢢꢣꢣꢣꢣꢣꢣꢣꢣꢣꢣꢤ ꢓ ∙ ꢑꢒꢀꢁꢂꢃꢧ ꢆ
ꢕꢅ ꢇ
ꢖ
ꢅ
ꢇ
ꢓ ꢀꢁꢂꢃꢄ ∙ ꢔ ꢑꢒꢂꢃ ꢍꢎꢏꢐ ꢈ ꢓ ꢗ ꢔ ∙ ꢑꢒꢂꢃꢅ
ꢇ
ꢅ ꢇ
ꢍꢎꢏꢐ
Eq. 1 and 2 represent the precursor dissolution. Eq. 3 is a general representation of a
complex formation, where further investigation is required to reveal the exact formula of
the complex. Eq. 4 represents the immediate reaction when an antiꢀsolvent is added. Eq. 5
represents the CsPbBr3 perovskite precipitation when an antiꢀsolvent is added to an
equilibrated CsBrꢀdepleted solution.
Figure 1. (a) CsPbBr3 crystals grown by a VSA method: (i) growth from MeCNꢀsaturated solution on a
50ºC hotplate; (ii) growth from MeOHꢀsaturated solutions at RT. Both were grown for ca. 48 hours; (iii)
growth from an unsaturated DMSO precursor solution, using MeCN as an antiꢀsolvent – showing
(LEFT) a yellow precipitate on top of which grew orange CsPbBr3. When illuminated with a UV light
(RIGHT) the yellow precipitate shows a strong green fluorescence, indicating that it contains the CsBrꢀ
rich (Cs4PbBr6) compound. (b) CsPbBr3 crystals grown by an HRS method: growth at a heating rate of
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10°C/hour up to 80°C (i) without and (ii) with seeding of carefullyꢀselected freshlyꢀgrown crystals; (iii)
before and after heating to 80 °C for the first time (i.e., without a preliminary heating/cooling cycle),
showing the undesired Cs4PbBr6 fluorescent compound as a side product.
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The following section deals with the characterization of the CsPbBr3 crystals. Except for
structural verification via XRD (Figure 2) that was done for crystals grown by all the methods
mentioned above, we further investigated crystals that were selected based on their (large) size
and clearly developed facets. These were obtained from the VSAꢀgrowth method using a
MeCNꢀsaturated solution (as shown in Figure 1(a,i)). Since seeding could induce undesirable
defects at the seed boundary, we did not use crystals that resulted from seeding.
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Characterization results:
For characterizing the crystals that were grown, we used:
1. powder xꢀray diffraction (XRD) for structural information
2. energyꢀdispersive xꢀray fluorescence spectroscopy (EDS) for elemental
composition
3. thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC) for
thermal behavior (phase transition and melting point);
4. Secondꢀharmonic generation (SHG) spectroscopy to characterize the (aꢀ)symmetry
of the crystals.
5. photothermal deflection spectroscopy (PDS), and photoluminescence spectroscopy
(PL) for optoelectronic transitions (excitation and emission, respectively) and
Urbach energy determination (PDS);
6. timeꢀresolved photoluminescence (TRPL) for characterizing the lifetime of
opticallyꢀexcited carriers;
7. impedance spectroscopy (IS) for determining the bulk resistivity and relative
permittivity, as well as measuring photoꢀresponsivity;
Nanoꢀindentation experiments to determine the elastoꢀmechanical properties of the crystals
were previously reported.22
A brief description of the different methods can be found in a dedicated section in the
supporting information. Whenever it was possible (based on available literature), we compared
results with those of thermally grown CsPbBr3 crystals.
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The powder XRD pattern of the pulverized orange crystals agrees well with patterns
obtained from meltꢀgrown CsPbBr3 and the simulated pattern of CsPbBr3 (Figure 2).7 Placing
the nativelyꢀdeveloped faces of a single crystal parallel to the scanned surface in a θ/2θ scan,
showed it to be the (101) and (002) orthogonal crystallographic planes (Figure 2(i)), where
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(101) was found to be the more developed face.
To assess the single crystallinity of the grown crystals, we performed a θꢀscan to get a rockingꢀcurve
(Figure S4(a)) on a ~ 2x1 mm2 crystal surface (Figure S4(c)) lying closely parallel to the (101)
crystallographic plane (fixed on 2θ=64.0, which is the (404) diffraction angle). By scanning the crystal
in 4 different angles (0˚, 90˚, 180˚, 270˚), we verified that the reflection corresponds with the (404)
peak. Although we scanned using a wide receiving slit, which allowed collection from ꢂ2θ of ~1.4˚
(Figure S4(b)), we could not detect the (080) and (414) peaks, which appear in a powder pattern and
could theoretically be detected within the range of ±1.4˚ from 2θ=64.0˚. Their absence confirms the
single crystal nature of the sample and the ability to grow single crystals using this method. We note
that among the samples that were checked in this way there were some that showed additional
diffraction peaks at θ=31.2˚ (~(080)) as well as at θ=32.0 (~(404)), indicating that further improvement
is needed to increase the yield of a singleꢀdomain crystals.
Since powder XRD might be not sensitive enough to reveal slight structural distortions, we
used SHG spectroscopy to check if the solutionꢀgrown crystals retain their centrosymmetry, as
expected from the mmm point group (orthorhombic, Pnma space group), to which CsPbBr3
crystals were reported to belong.7 The SHG spectroscopy was performed on a coarseꢀgrained
CsPbBr3 crystal to avoid loss of signal due to phaseꢀ(mis)matching between the first and the
second generated harmonics and to expose different crystallographic orientations.23 We
observed only two photon (Ex: 860 nm) fluorescent emission (at 537±2 nm) but no SHG at half
the excitation photon wavelength; Em: 430 nm) (Figure 3). We conclude that, indeed, the
crystal is centrosymmetric, agreeing with the reported crystal symmetry.
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Figure 2. Roomꢀtemperature xꢀray diffraction pattern of (i) solutionꢀgrown CsPbBr3 single crystal that
its nativelyꢀgrown (101) face was parallel to the scanning plane in a θ/2θ scan; (ii) powder of a
pulverized solutionꢀgrown crystal;(iii) a simulated pattern based on thermallyꢀgrown crystals’ CIF.7 The
crystal orientations shown in bold correspond to the nativelyꢀexposed faces of the solutionꢀgrown
crystal.
EDS yields a Cs:Pb:Br atomic ratio of (1.00±0.04):(1):(3.4±0.1) (taking the Pb quantity as
unity). The excess in the Br signal may come from reabsorption of Pb and Cs xꢀray emissions
by Br (the EDS lines that were used for quantification are 1.48, 2.34 and 4.28 keV, for the Br,
Pb and Cs respectively). The effect is only seen for the Br, probably as it is the most abundant
element as well as its ability to absorb both the Cs and Pb xꢀrays. Pb, in contrast, may only
absorb xꢀrays emitted by the Cs. Based on the EDS lineꢀspectrum (Figure S5), no other
elements were detected.
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Figure 3. (a) CsPbBr3 solutionꢀgrown single crystal absorption (from PDS) and twoꢀphoton
photoluminescence (using 860 nm laser) spectra. The Urbach energy is calculated from the absorbance
data (Figure S7). A higher resolution spectrum (inset) shows no signs of SHG at 430 nm. (b) Two 860
nm photonꢀexcited photoluminescence images of a crystal (top) and of a dispersed powder, obtained
from grinding a crystal (bottom).
Thermal analysis results (Figure 4) are similar to what has been reported for meltꢀgrown
CsPbBr3 crystals7,24 and to the TGA analysis of spinꢀcoated CsPbBr3 (i.e. orthorhombic
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(Pnma) ↔ cubic (Pm3m) phase transformation at 131 °C24 and melting at 567 °C7). The
reported transition between the orthorhombic (Pnma) to the tetragonal (P4/mbm) symmetries
(at 88 ºC) was not observed in our analysis, possibly due to insufficient sensitivity in the
measurement.7,24 Thermal analysis performed on CsPbBr3 crystals (Figure S6) sometime
o
showed a sharp change (both in TGA and the first heating cycle in DSC) at ~230 C, which
may be related to trapped solvent (DMSO; bp:189 °C).
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Figure 4. (a) DSC and (b) TGA and DTA of powdered samples, obtained from pulverizing solutionꢀ
grown CsPbBr3 crystals. At 131 °C (inset in (a)) a phase transition between tetragonal and cubic
structure occurs.24,25 At 567 °C (differentialꢀthermogravimetricꢀanalysis (DTA) in (b)) the crystal melts
– in agreement with literature values for meltꢀgrown CsPbBr3 crystals and with results from the TGA
analysis obtained from spinꢀcoated CsPbBr3.7,8 The green line in (a) indicates that there is no sharp
change during the first heating cycle, as sometimes occurred when a crystal was analyzed (see Figure
S6 and further discussion in the text).
From optical absorption data, obtained by PDS (Figure 3 and Figure S7), an Urbach
energy of Eu=19 meV was derived from the slope of the logarithmic plot of the absorbance vs.
the photon energy (ꢨꢩꢪꢫꢬꢩꢨꢭꢮꢯ ∝ ꢰꢱꢲꢕꢳꢳ ꢖ). While this is higher than the 15 meV value
ꢴ
found for MAPbI3, it is small compared to that of e.g., Cu(In,Ga)Se2 (Table 2 in Ref. 26). We
also find that a thin film that was prepared by spin coating from a similar DMSO solution
shows a higher Urbach energy ( ~36 meV; Figure S7(b)), which implies that the crystal is of
high quality. The CsPbBr3 thin film fabrication procedure is reported elsewhere.9
Impedance spectroscopy (IS) between two parallel faces of a single crystal, using carbonꢀ
pasted contacts, was used to measure the bulk resistivity and relative permittivity (Figure 5 and
Figure S8). The average resistivity, ρ, of the solutionꢀgrown crystals was found to be ~ 0.10 ±
0.04 Gꢃꢄcm and relative permittivity εr ≈ 41 ± 4. The resistivity is an order of magnitude lower
than that of crystals obtained by meltꢀbased synthesis (~1 Gꢃꢄcm).7 εr, being a static (or low
frequency) dielectric constant is clearly higher than the estimated optical one (εr (optical) ~4).27
Similarly, the optical dielectric constants of MAPbBr3 and MAPbI3 are found to be ~4.5 and
6.5,28 while their GHz (microwave) dielectric constants are ~25 and 30, respectively21. When
illuminating (ca. 5 mW/cm2 of a white LED) during the IS measurement, we find that the
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resistivity decreases by ~1 order of magnitude (Figure 5). Since only a small fraction of the
crystal thickness actually absorbs the light (except for a very small amount of light in the lower
energy Urbach tail region– see Figure S7) and there is shading by the carbon contact, the
overall observed photoconductivity seems to be reasonable.
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Two different excited carrier lifetimes were found by TRPL (Figure 6): 4.4 ± 0.1 ns and
30 ± 3 ns, which are speculated to be the trapꢀassisted and free charge carrier recombinationꢀ
based decays, respectively.29 Both lifetime regimes are significantly shorter than those reported
for MAPbI3 and MAPbBr3 single crystals.10 Other fullyꢀinorganic photovoltaicꢀgrade materials,
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CIGS and CdTe show lifetimes of ~50 and ~150 ns , respectively. In addition, a ~50 meV
redꢀshift, which is observed in the PL decay over time, may indicate photon recycling as was
recently suggested32,33 and subsequently reported for MAPbI3 34. An alternative reason might
be diffusion and relaxation of the excited carriers to lower energy levels as the PL decays over
time.
Figure 5. IS (in air at 48% RH) between two parallel faces of a solutionꢀgrown CsPbBr3 single crystal in
the dark and under white LED illumination (5ꢀ7 mW/cm2 – calibrated with a Si diode).
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Figure 6. TRPL spectroscopy of solutionꢀgrown CsPbBr3 crystals. (a) A full PL spectrum
representation with respect to time and intensity – shows that the maximum intensity shifts with time
towards longer wavelength (white dashed line). (b) PL intensity vs time at 532 nm from which the short
and long lifetimes are extracted. τ1 and τ2 are speculated to be the trapꢀassisted and freeꢀcharge carrierꢀ
recombination based decays, respectively.29
We observed a strongly photoluminescent phase as a sideꢀproduct during the crystal
formation, which we attribute to Cs4PbBr6.20 Since the green luminescence could also be
attributed to confined CsPbBr3 nanocrystals in a CsBr matrix,35 its identity was further
investigated. We extracted the yellow, strongly fluorescent, precipitate and compared it to
pulverized solutionꢀgrown CsPbBr3 crystals via PL and powder XRD (Figure 7). A clear
qualitative difference in the PL intensity is shown in the photograph (Figure 7(a)). A more
quantitative analysis shows (under identical excitation and collection conditions) 104 times
stronger PL intensity (Figure 7(b)) in favor to Cs4PbBr6. By increasing the signal to noise PL
from the pulverized CsPbBr3 by opening the excitation and emission (collection) slits, and
normalized the two spectra (Figure 7(c)). We see that although the PL emission of two samples
is in the green region, they are clearly distinct from each other by ~ 20 nm (0.1 eV). A further
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identification of the luminescent phase was carried out via powder XRD (Figure 7(d)), which
shows very good agreement with the Cs4PbBr6 pattern.36
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Figure 7. (a) Pictures of the extracted luminescent precipitate and pulverized CsPbBr3 under room light
and UV (365 nm) lamp. (b) PL spectra of the two samples shown in (a) using identical excitation and
emission slit widths. Excitation wavelength ꢀ 365nm. (c) Normalized PL spectrum of the two samples
shown in (a) under different slitꢀwidth conditions. Excitation wavelength – 365nm. (d) Powder XRD
pattern of the powders shown in (a). The labeled red crystallographic (hkl) planes of Cs4PbBr6.36 Some
observed peaks in the Cs4PbBr6 diffraction pattern correspond to CsPbBr3, indicating that the
luminescent phase contain small amounts of CsPbB3.
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Crystallization process
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While unraveling the crystallization mechanism of solutionꢀgrown CsPbBr3 crystals awaits
further study, we draw attention to several qualitative observations.
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ꢀ1ꢀ The maximum solubility of the components at roomꢀtemperature was found to be ca.
0.5 M, which was achieved only after mixing the two components. Solubility tests of PbBr2 and
CsBr precursors individually showed that, although 0.183 g/mL (0.50 M) PbBr2 completely
dissolved in DMSO (still below the solubility limit), only 0.055 g/mL (0.26 M) of CsBr could
be dissolved in DMSO. The remaining 0.24 M seem to form a soluble [(x)CsBr∙(y)PbBr2]
complex with DMSO, where (x) and (y) are stoichiometric factors. To test the effect of adding
an antiꢀsolvent to a CsBrꢀdepleted solution, we prepared a precursor solution of 0.25 M CsBr
and 0.5 M PbBr2. Upon slow addition of MeCN, we found that before reaching the saturation
point (where the orange precipitant stops reꢀdissolving) the CsBrꢀdepleted solution stays clear,
unlike the case with the equimolar solution (see Figure S1(a)). This supports the existence of
CsBr excess as an available species for crystallization. These examples further suggest that it is
possible to find an optimal concentration balance between the precursors (which might be, as in
this case, not 1:1) in which the addition of an antiꢀsolvent will result in precipitation of a
desired product. Titrating up to the saturation point of the DMSO precursor solution and further
filtering it, forced the system to get close to this optimal ratio. This method might have
potential as a generic way of extracting a desired product out of similarlyꢀchallenging systems.
ꢀ2ꢀ Although finding a proper antiꢀsolvent might seem to be a matter of luck, we note that
among the 9 tested antiꢀsolvents, those with the highest dielectric constants (MeCN (ε ≈ 37),
MeOH (ε ≈ 32) and H2O (ε ≈ 78)) showed positive results of CsPbBr3 formation, where those
of MeCN and MeOH are the closest to that of DMSO (~42).37 The solvation of ions in polar,
aprotic solvents (such as DMSO) is usually via the first solvation shell of the ions in ionꢀ
solvent interactions38 (i.e. complexation). The high dielectric constant, therefore, plays an
important role in screening the interactions between the solvated salts and the solvent (DMSO).
We noticed that addition of antiꢀsolvents with low dielectric constant usually results in
preferred precipitation of the Cs4PbBr6 compound. However, when the antiꢀsolvent has a higher
dielectric constant than the solvent (compare between H2O and DMSO), much less of the (Csꢀ
rich) Cs4PbBr6 precipitates. After adding a certain amount of H2O, a white precipitate appears
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(see Figure S3), which is attributed (based on Ref. ) to CsPb2Br5 (PbBr2ꢀrich) phase. This
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suggests that the unique precursor ratio (described in the previous point) should be coupled to a
specific antiꢀsolvent, the dielectric constant of which plays a significant role.
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Crystal characterization:
The solution–based path described here is found to yield good quality CsPbBr3 single
crystals. Based on results from XRD, EDS, TGA and DSC, these are similar to those obtained
for crystals grown at melting temperatures (> 600°C; Bridgman method). The crystals are
grown in a matter of a few hours (using the HSR method) up to a few days (using the VSA
method). For future studies and possible applications, the VSA method with MeCN as an antiꢀ
solvent was found to be the most suitable one, as it provides the most controllable and
reproducible way of growing crystals. Nevertheless, one may find the other antiꢀsolvents or,
especially, the HSR method more appealing – mostly due to its comparative simplicity and
speed of preparation.
Despite the similarity, the solutionꢀgrown crystals differ from the thermally grown ones in
some respects, such as their conductivity. It is reasonable to assume that, due to the presence of
a solvent and the lower temperature, the crystallization kinetics will be different and, thus, the
defect density and resulting properties may be different. Therefore, apart from the simplicity,
which will make the crystals more readily accessible for research purposes or for optoelectronic
uses (e.g. xꢀray or γꢀray detection as was previously demonstrated7), crystals from lowꢀ
temperature solutionꢀbased growth will likely be closer in their properties to those of solutionꢀ
deposited HaP films, which are used in solar cells and LEDs.
Another issue that arises due to the use of a solution medium for crystal growth is
occluded solvent molecules, which can lead to defects in the crystal. Although thermal analysis
of the powder did not show clear signs of trapped solvent, similar analysis for single crystals
showed features that may be attributed to trapped solvent (Figure S6). More work is required to
resolve this point.
For a more detailed investigation of possible distortions or defects in the solutionꢀgrown
crystal, SHG and Urbach energy measurements were carried out. SHG spectroscopy showed
(Figure 3) that growth from solution does not lead to a breaking of the symmetry to a nonꢀ
centrosymmetric space group (as might be the case because of defects in the crystal). Some
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very weak, but nonꢀzero, SHG signal that came from a thermallyꢀgrown CsPbBr3 single
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crystals was attributed to surfaceꢀ or local defectꢀrelated symmetry breaking, a signal that
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was not detected in our measurements. It was also found that the Urbach energy, a parameter
that can be correlated with structural and/or electronic disorder,40 is only 19 meV, lower than
what has been reported for photovoltaicꢀquality CIGS thin films.26 Due to the crystal thickness,
this value of the Urbach energy should also be taken as an upperꢀlimit and might actually be
lower (for further explanation see Figure S7 and ‘Characterization methods’ section in the
supporting information). Still, the lower resistivity than that of meltꢀgrown CsPbBr3, the
shorter lifetime than that found for other solutionꢀgrown HaP crystals, and the ~50 meV red
shift in the PL with relaxation time might indicate that further purification of the crystals may
be possible.
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The hybrid organicꢀinorganic MAPbBr3 HaP shows lower longꢀterm stability than the fully
inorganic analog, but has similar mechanical properties and gives comparable photovoltaic
performance.8,9,22 The Young’s modulus (15.8 ± 0.6 GPa) and hardness (~0.34 ± 0.02 GPa) of
these solutionꢀgrown crystals are quite similar to those of MAPbBr3 (~19.6 ± 0.03 GPa and
~0.36 ± 0.02 GPa, respectively)22 ꢀ an indication for the lack of importance of the A group (Cs+
vs. MA+) in the dominating interatomic bonds. Since solutionꢀgrown CsPbBr3 crystals show
photoconductivity (Figure 5) and nonꢀlinear absorption (Figure 3(b)), Cs can replace the
organic cation to give a more reliable material for (lowꢀcost) single crystalꢀrelated applications
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(e.g., photodetectors, as was demonstrated for MAPbI3 , MAPbCl3 or nonꢀlinear absorbers for
ultraꢀfast photonics, as demonstrated for MAPbBr3 41).
Regarding the strongly fluorescent precipitate, we verified it is Cs4PbBr6 or, at least,
mostly Cs4PbBr6. It fluoresces at ~0.1 eV higher energy than CsPbBr3 crystals. It was
previously reported that Cs4PbBr6 can be luminescent at the green region (~545nm),42 but also
that confined CsPbBr3 quantum dots in a CsBr matrix can show a similar effect.35 It was found
that the emission of confined CsPbBr3 crystals in CsBr matrix can be blueꢀshifted by ~0.1 eV
from the bulk crystal, which is very similar to our observation.35 However, in order to get a 0.1
eV blueꢀshift due to quantum confinement, the CsPbBr3 particles need to be at most 10 nm in
diameter.43 Based on the Scherrer equation, which allows to estimate the particle size based on
the peak broadening, and a closer look on the (121) peak FWHM of the CsPbBr3 (ꢂ(2θ)~0.2˚)
in the Cs4PbBr6 pattern (Figure 7(d) at 2θ=21.4˚), the CsPbBr3 particles are much larger than
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10 nm. Therefore it is very unlikely that the strong luminescence originates from confined
CsPbBr3 crystals in a Cs4PbBr6 matrix, To exclude the possibility that the luminescence is
correlated to a solventꢀeffect with the Cs4PbBr6 or CsPbBr3, we stored the samples in ambient
atmosphere at room temperature or moderately heated to remove residues of the solvent. No
clear PLꢀintensity loss was observed.
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Summary:
Simple and reproducible procedures to grow high quality CsPbBr3 single crystals from
solution without the usuallyꢀobserved related byproducts: Cs4PbBr6 or CsPb2Br5, were
described. Two separate methods were found to be suitable: ꢀ1ꢀ slow VaporꢀSaturation of an
Antiꢀsolvent (VSA) and ꢀ2ꢀ Slow Heating of a solution with a RetrogradeꢀSoluble compound
(HRS). For both methods a precursor (PbBr2 and CsBr) DMSO solution was first titrated with
one of the following antiꢀsolvents: MeCN or MeOH. The preꢀsaturation step was important to
prevent precipitation of the undesired Cs4PbBr6 alongside to the desired CsPbBr3. For the VSA
method, MeCN was found to be more suitable than MeOH, while for the HRS method, MeOH
is preferable.
The low temperature solutionꢀgrown method presented here provides a readily available
alternative to obtain high quality, mmꢀsize, CsPbBr3 single crystals to the thermallyꢀgrown
method (> 600°C). The absence of symmetryꢀbreaking (based on SHG spectroscopy) and the
19 meV upperꢀlimit of the Urbach energy suggest high quality of the solutionꢀgrown crystals.
The dielectric constant (at subꢀMHz dielectric response) was found to be ~ 40.
The order of magnitude lower resistivity, compared to thermally grown CsPbBr3, and the
low excited carrier lifetime, compared to other hybrid HaPs, suggest the presence of
appreciable dopant (charged defect) density in the CsPbBr3 crystals formed by this growth
method. If true, this implies the possibility to control the crystal doping: intrinsically (by
modifying the growth conditions) or extrinsically (by adding known impurities to the solution
medium). The fact that CsPbBr3 is clearly more stable than MAPbBr3, but with similar
mechanical and optoelectronic properties, suggests that the methods described in this paper may
allow production of more reliable optoelectronic components, with respect to analogous hybrid
HaPs.
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Because high dipole moment aprotic solvents (e.g., DMSO, DMF, GBL) are commonly
used in synthesis of HaPs (for both thin films and single crystals), the presented growth method
is a potential path to grow perovskite structures with both organic and inorganic A group,
similar to recently reported Cs0.2FA0.8PbI2.84Br0.16 thin films.44 A combination of organic (e.g.
MA and FA) and inorganic (e.g. Cs, Rb) monovalent cations may lead us to better
understanding of the organic cation’s role in halide perovskites. The present work is an
important step towards this goal.
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Supporting information:
The Supporting Information is available free of charge on the ACS Publications website at
DOI: XXXXXXX.
Images of reaction systems during or after different reaction steps (Figures S1 and S3);
schematic view of the VSA setup and the HRS sequence (Figure S2); Rockingꢀcurve of a
single crystal (Figure S4); EDS spectrum (Figure S5); thermal analysis (TGA, DTA and
DSC) of nonꢀpulverized crystals (Figure S6); PDS results from a single crystal and a thin
film of CsPbBr3 (Figure S7); impedance spectroscopy (Figure S8); detailed description of
characterization methods and equipment; list of other tested antiꢀsolvents.
Acknowledgments:
The authors thank Dr. Isai Feldman, Dr. David Ehre, Dr. Thomas M. Brenner and Arava Zohar from the
Weizmann Institute of Science for fruitful discussions and technical support. D.C. holds the Sylvia and
Rowland Schaefer Chair in Energy Research.
Funding:
G.H and D.C acknowledge the Israel Ministry of Science and the Israel National NanoꢀInitiative for
partial support and G.H. acknowledges the Leona M. and Harry B. Helmsley Charitable Trust.. A.S.,
M.L.B. and R.H.F would like acknowledge EPSRC for their support. A.S. would also like to
acknowledge support from an IndiaꢀUK APEX project. D.C. holds the Sylvia and Rowland
Schaefer Chair in Energy Research.
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For Table of Contents Use Only
Manuscript title: “LowꢀTemperature SolutionꢀGrown CsPbBr3 Single Crystals and Their
Characterization”
Author list: Rakita, Yevgeny; Kedem, Nir; Gupta, Satyajit; Sadhanala, Aditya;
Kalchenko, Vyacheslav; Böhm, Marcus; Kulbak, Michael; Friend, Richard; Cahen,
David; Hodes,
TOC graphic:
TOC synopsis: As an alternative to a highꢀtemperature (>600˚C) meltꢀbased synthesis, a
simple solutionꢀbased growth provides millimeterꢀsized, pure CsPbBr3 perovskite singleꢀ
crystals without formation of CsBrꢀrich or PbBr2ꢀrich phases, which are often present
alongside the perovskite phase. Characterization of these crystals, which are of interest for
optoelectronics, includes charge carrier lifetime and an upperꢀlimit of the Urbach energy.
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