Tailoring the physical properties of hybrid lead metal halide APbX3 perovskites by means of compositional engineering is one of the key factors contributing to the development of highly efficient and stable perovskite solar cells. While the beneficial effects of partial ionic replacement at the A- and X-sites are largely demonstrated, partial replacement of Pb2+ is less explored. Here, we developed a solution-based procedure to prepare thin films of mixed-metal MAPb1-aMnaI3 perovskites. Although Mn2+ ions have a size that can potentially fit in the B-sites of MAPbI3, using a combination of structural and chemical analysis, we show that only less than 10% of Pb2+ can be replaced by Mn2+. A 3% replacement of Pb2+ by Mn2+ leads to an elongation of the charge carrier lifetimes as concluded from time-resolved PL measurements. However, by analysis of the time-resolved microwave conductance data, we show that the charge carrier mobilities are largely unbalanced, which is in accordance with density functional theory (DFT) calculations indicating that the effective mass of the hole is much higher than that of the electron. Increasing the concentration of Mn2+ in the precursor solution above 10% results in formation of amorphous Mn-rich domains in the film, while the perovskite lattice becomes depleted of Mn2+. These domains negatively affect the charge carrier mobilities and shorten the lifetime of photogenerated carriers. The resulting reduction in charge carrier diffusion lengths will severely limit the photovoltaic properties of solar cells prepared from these mixed metal halide perovskites.

The extraordinarily rapid rise in the power conversion efficiency of metal halide perovskite solar cells (PSCs)1,2 has in the last years been accompanied by improvements in their stability.3 While the first breakthroughs in the field of PSCs were due to changes in the device architecture4,5 and processing conditions,6 compositional engineering of the perovskite ABX3 crystal structure has recently played a crucial role in the achievement of highly stable and efficient PSCs.7 It has been shown that the properties of hybrid halide perovskites can be tailored and optimized by (partially) replacing the ions at the A-sites and X-sites of the perovskite lattice,7–9 and the most efficient PSCs are currently realized using mixed-cation, mixed-halide perovskites.10 By contrast, blending the divalent metal cations at the B-sites to obtain mixed-metal perovskites is a relatively less explored path. Most of the efforts in this direction involved the partial replacement of Pb2+ with Sn2+ and were initially motivated by the necessity of mitigating the toxicity of lead.11–14 MAPb1-aSnaX3 perovskites (where MA = CH3NH3 and X = I, Br, or Cl) have been synthesized for the full compositional range 0 ≤ a ≤ 111 and used in photovoltaic devices with efficiencies as high as 17.6%,15 although the ease with which Sn2+ oxidizes to Sn4+ severely limits the processing of this class of perovskites.16,17 Recently, Klug and co-workers demonstrated that partial replacement of lead is possible using transition or alkaline earth metals.18 In those cases, the photovoltaic performance significantly dropped if more than a few atomic percent of lead was replaced. Very recently, positive effects on the stability and performance were reported on replacing 3% of the Pb2+ by Zn2+, Mn2+, or Ni2+.19 

In this paper, we investigate how partial replacement of Pb2+ with Mn2+ affects the optoelectronic properties of MAPbI3. Manganese is one of the most abundant metals in the Earth’s crust and has a low toxicity.20 Besides the possibility of using the mixed metal perovskite MAPb1-aMnaI3 in photovoltaic devices,18,19,21 its magnetic properties can be used in low-power, light-controlling applications such as magneto-optical data storage devices.22 Recently, partial replacement of Pb2+ with Mn2+ has been demonstrated for lead halide perovskite nanocrystals and nanoplatelets.23–25 

We prepared thin films of MAPb1-aMnaI3 from solution, and we characterized the chemical composition, crystal structure, and optoelectronic properties. Although the elemental analysis confirms the presence of Mn2+ and the material retains the perovskite crystal structure even up to 50% replacement, we demonstrate that only less than 10% of Pb2+ can be replaced by Mn2+ at the B-sites. The excess manganese is located in amorphous, Mn-rich domains.

We characterized the optoelectronic properties of the mixed metal perovskite thin films by means of UV-vis absorption spectroscopy, time-resolved photoluminescence (TRPL), and time-resolved microwave conductance (TRMC). If less than 10% of Pb2+ is replaced by Mn2+, charge carrier lifetimes increase; however, the charge transport becomes largely unbalanced: The mobility of holes is more than one order of magnitude lower than the mobility of electrons, which is in accordance with density functional theory (DFT) calculations revealing the effective masses. Segregation of Mn2+ in domains outside the perovskite crystals restores balanced transport; however, the mobilities are significantly lower and the recombination rate is faster than in MAPbI3.

Details regarding the experiments are summarised below. Equimolar amounts of hydroiodic acid (HI, 57 vol% in water) and methylamine (CH3NH2, 33 wt. % in ethanol) were added in a round-bottom flask and let to react for 2 h in an ice bath while stirring. The precipitate of the solution was collected by rotary evaporation at 50 °C, then washed with diethyl ether, and dried at 60 °C in a vacuum oven for at least 24 h before use.

MAI, PbAc2, and MnCl2 were dissolved in dimethylformamide (DMF) inside of a nitrogen-filled glovebox. Various solutions with different PbAc2:MnCl2 molar ratios have been prepared, keeping the molar ratio of MAI to the total concentration of inorganic salt constant (3:1). The solutions were heated on a hot plate at 70 °C and magnetically stirred for 1 h. Subsequently, the solutions were allowed to cool down to room temperature and filtered (polypropylene filters, 0.45 μm) before use.

Samples for XRD, UV-vis absorption, TRPL, and TRMC measurements have been deposited on quartz substrates; samples for SEM and EDX have been deposited on FTO-coated glass substrates. The quartz substrates were cleaned by O2 plasma; for FTO-coated glass substrates, a thorough cleaning by using an ultrasonic bath in acetone and isopropanol has been performed prior to O2 plasma cleaning. The cleaned samples have been transferred in the glovebox, and perovskite films have been deposited on them by spin coating (2000 RPM, 2000 RPM/s, 45 s). The deposited films were annealed at 100 °C (15 min for MAPbI3 films, 2 h for films containing Mn2+) with a thickness of about 200 nm.

XRD measurements were performed using a Bruker D8 Advance diffractometer in the Bragg-Brentano geometry. The monochromatic Co Kα X-ray source was operated at 40 kV and 40 mA.

The fraction of absorbed light was measured with a Perkin-Elmer UV/VIS/NIR Lambda 1050S spectrometer containing a 150 mm Spectralon integrating sphere. The perovskite thin films were measured in the wavelength range 400–850 nm.

TRPL measurements were carried out using an Edinburgh LifeSpec spectrometer equipped with a single-photon counter. The perovskite films were excited at 405 nm with a picosecond pulsed diode laser. A high pass filter was used to prevent photons with wavelength below 450 nm to reach the detector. PL emission spectra of all samples were measured in the range 650–850 nm. PL lifetimes were recorded at the maximum emission wavelength.

Samples were sealed in a microwave cavity within a nitrogen-filled glovebox. The fraction of absorbed microwave (frequency ca. 8.5 GHz) power (ΔP/P) was measured after pulsed optical excitation (repetition rate 10 Hz, λ = 550 nm). The rise of ΔP/P is limited by the width of the laser pulses (3.5 ns) and by the response time of the microwave cavity (18 ns). A series of neutral density filters was used to perform light-intensity dependent TRMC measurements.

Top-view SEM secondary electron images were taken with a JEOL-JSM-610LA microscope using 10 kV voltage energy, 10 mm working distance, and ×3000 magnification. Backscattered electron images were taken using 20 kV voltage, 10 mm working distance, and ×200 magnification.

EDX elemental maps and point analysis were carried out using the EDX detector of the JEOL-JSM-610LA microscope, with settings of 20 kV voltage, 10 mm working distance, and ×200 magnification.

The structural optimization of MAPbI3, MAPb0.9375Mn0.0625I3, and MAPb0.875Mn0.125I3 using 12 (unit cell), 192, and 96 atoms (super cells) was performed using DFT within the Perdew, Burke, and Ernzerhof (PBE)26 functional implemented in the Vienna ab initio simulation package (VASP).27,28 The outermost s, p, and d (in the case of Pb and Mn) electrons are treated as valence electrons whose interactions with the remaining ions are modeled by pseudopotentials generated within the projector-augmented wave (PAW) method.29,30 An energy cutoff of 500 eV and 6 × 6 × 6, 2 × 3 × 3, and 3 × 3 × 3 k-point meshes were used to achieve energy and force convergence of 0.1 meV and 2 meV/Å for 12, 192, and 96 atoms cells, respectively. The spin-orbit coupling effect was taken into account for the band structure calculations.

Recently, the lead-free compound methylammonium manganese iodide (MAMnI3) has been synthesized and characterized; such a compound is amorphous and does not form a cubic perovskite structure.31 We started our study by investigating how much Pb2+ can be replaced by Mn2+ before disrupting the perovskite crystal structure. The most commonly used predictor for the perovskite phase is the Goldschmidt tolerance factor,32 

t=rA+rX2rB+rX,
(1)

where rA, rB, and rX are the ionic radii of monovalent cation (A), divalent cation (B), and halogen anion (X), respectively [Fig. 1(a)]. It has been empirically observed that the perovskite structure can be formed if the tolerance factor is in the range 0.875 < t < 1.06.33 Besides the tolerance factor, the octahedral factor is used to assess if the divalent cation B has the right size to fit in the X6 octahedron,34 

μ=rBrX.
(2)

If μ < 0.41, the BX6 octahedra cannot form, which excludes the formation of the perovskite structure.34 By combining the perovskite formation criteria for t and μ, one can define a perovskite formation region on the t-μ structure map [Fig. 1(b)]. We calculated t and μ for MAPb1-aMnaI3 with 0 < a < 1; for the mixed-metal compounds, we approximated the radius of the divalent cation as the weighted average of the cationic radius of Pb2+ (rPb) and of Mn2+ (rMn),

rB=arMn+1arPb.
(3)

For each composition considered, a point is placed on the t-μ structure map in Fig. 1(b). The replacement of Pb2+ with the smaller Mn2+ cations results in a reduction in μ, which drops below the empirical limit for a > 0.5.

FIG. 1.

(a) Crystal structure of ABX3 perovskite materials; (b) t-μ structure map for MAPb1-aMnaI3. The black circles represent t-μ combinations calculated for different values of a. The shaded area indicates the region in which the perovskite structure does not form. The cationic and anionic radii for the calculation of t and μ have been taken from Ref. 33.

FIG. 1.

(a) Crystal structure of ABX3 perovskite materials; (b) t-μ structure map for MAPb1-aMnaI3. The black circles represent t-μ combinations calculated for different values of a. The shaded area indicates the region in which the perovskite structure does not form. The cationic and anionic radii for the calculation of t and μ have been taken from Ref. 33.

Close modal

The above analysis suggests that it is thus possible to replace half of the Pb2+ with Mn2+ without disrupting the perovskite structure. To verify this prediction, we prepared mixed-metal compounds using a one-step solution process: methylammonium iodide (MAI), lead acetate (PbAc2), and manganese chloride (MnCl2) were dissolved in DMF and subsequently spin cast to form thin films. The molar ratio of MAI to the total amount of inorganic salt content was kept constant at 3:1. We define the amount of Mn2+ in the precursor solution (asol) as the percentage of MnCl2 with respect to the total amount of bivalent metal salts,

asol=nMnCl2nMnCl2+nPbAc2×100%,
(4)

where ni is the number of moles of the salt i in the precursor solution. For the remaining part of this paper, a film with asol = 10% denotes a film that is prepared from a precursor solution with a fraction Mn2+ of 0.1.

After spin casting the films, the elemental composition was measured by means of energy-dispersive X-ray spectroscopy (EDX). Quantitative analysis of the data provides the atomic fraction of each element in the sample (Table SI in the supplementary material); the percentage of Mn2+ with respect to the total amount of bivalent metals is given by

afilm=fMnfMn+fPb×100%,
(5)

where fi is the atomic fraction of the element i in the film.

Except for the film with asol = 50%, all percentages found for afilm are somewhat smaller than those for asol (see Table I). Moreover, from the backscattered electron images and elemental mapping of Mn (Fig. S1), it is evident that for films with asol > 30%, the films are not uniform and that the lateral distributions of Pb and Mn are not homogenous.

TABLE I.

Comparison of the Mn2+ fraction in the precursor solutions (asol) with that in the films (afilm).

asol (%)afilm (%)
10 8.93 
20 16.44 
30 19.56 
40 31.41 
50 51.86 
asol (%)afilm (%)
10 8.93 
20 16.44 
30 19.56 
40 31.41 
50 51.86 

Although elemental analysis confirms the presence of Mn2 in the films (with roughly the same fraction as in the precursor solutions), this does not imply whether Mn2+ ions occupy the B-sites of the perovskite lattice. The X-ray diffraction (XRD) patterns for films with different concentrations of Mn2+ are shown in Fig. 2(a). The diffraction pattern of the Mn-free film presents the characteristic peaks of MAPbI3.35 We note here that the (110) peaks of our samples are located at 2ϑ ≈ 16.35°, in contrast to the more frequently reported value of 2ϑ ≈ 14.20°. The discrepancy is due to the Co Kα X-ray source used in our experiment, which emits X-ray at longer wavelengths than the more commonly used Cu Kα source. Interestingly, independent of the values of asol, the perovskite structure is preserved. Replacement of Pb2+ ions with the smaller Mn2+ ions is expected to cause a reduction of the perovskite lattice, resulting in a shift of the XRD reflections toward larger angles. For samples with a low content of Mn, the (110) peak indeed shifts toward larger 2ϑ [Figs. 2(b) and 2(c)], which is in agreement with the recent literature.19 Remarkably, the shift is not monotonously increasing with higher Mn concentrations. Instead, for films with asol > 3%, the (110) peak shifts back toward the position of the (110) peak of MAPbI3. We explain this as follows: while limited amounts of MnCl2 in the precursor solution results in the replacement of Pb2+ with Mn2+ at the B-sites of the lattice, larger concentrations of MnCl2 yield the formation of Mn-rich domains, depleting the perovskite lattice with Mn2+. Hence, the XRD pattern reverts back to that of MAPbI3. We can also rule out the formation of Mn2+ interstitials as this should result in an expansion of the perovskite lattice and consequently a shift of the XRD peaks toward lower angles, as recently observed by Bai and co-workers for Mn-alloying of CsPbI2Br.36 However, we do not observe any expansion of the MAPbI3 lattice even at the highest Mn2+ concentrations (Table SII of the supplementary material).

FIG. 2.

(a) XRD patterns (Co Kα radiation, λ = 1.79 Å) of mixed-metal perovskite films prepared from solutions containing various concentrations of MnCl2 as indicated by the percentage; (b) local region of the XRD patterns around the strongest cubic diffraction peak (110). The dashed lines represent Gaussian fits of the peaks; (c) 2ϑ value of the (110) peak centers as a function of the percentage MnCl2 in the precursor solution.

FIG. 2.

(a) XRD patterns (Co Kα radiation, λ = 1.79 Å) of mixed-metal perovskite films prepared from solutions containing various concentrations of MnCl2 as indicated by the percentage; (b) local region of the XRD patterns around the strongest cubic diffraction peak (110). The dashed lines represent Gaussian fits of the peaks; (c) 2ϑ value of the (110) peak centers as a function of the percentage MnCl2 in the precursor solution.

Close modal

When asol exceeds 30%, new features appear in the XRD patterns (Fig. S2). These findings agree with the SEM images of these films (Fig. S1), showing the appearance of new, non-amorphous phases (particularly evident for asol = 40% and 50%). Therefore, we decided to not further investigate films with asol > 30%.

For films with asol = 30%, presence of Cl is detected by means of EDX (Table SI). Similarly as discussed above for Mn2+, we would expect a shift of the XRD peaks toward larger angles if Cl replaced the larger I ions in the perovskite structure; the absence of such a shift indicates that Cl is also segregated in Mn-rich domains.

Figure 3(a) shows the absorptance measured for thin films of mixed-metal perovskite. The absorption onset does not shift significantly on increasing the Mn2+ percentage, indicating that the bandgap of MAPbI3 is virtually unaffected by the presence of Mn2+. The fact that the optical absorption is not significantly modified by the presence of Mn-rich domains implies that these have a large bandgap. Careful analysis of the PL emission spectra presents a small red-shift of the PL peak position [Fig. 3(b)] with respect to the PL spectrum of MAPbI3. The peak shift reaches a maximum for films with asol = 3% and decreases with higher concentrations of Mn2+, at which the PL peak position is almost identical to that of MAPbI3 [Fig. 3(c)]. The dependency is similar to what we observed in Figs. 2(b) and 2(c) for the XRD data, which strengthens the idea that for films with asol > 3% the Mn2+ is located in Mn-rich domains and not in the perovskite lattice.

FIG. 3.

(a) Absorptance of thin mixed-metal perovskite films prepared from solutions containing various concentrations of MnCl2; (b) normalized PL spectra measured on optical excitation at 405 nm; (c) position of the PL peak vs MnCl2 percentage; (d) TRPL traces measured at the wavelength at which the PL is maximum.

FIG. 3.

(a) Absorptance of thin mixed-metal perovskite films prepared from solutions containing various concentrations of MnCl2; (b) normalized PL spectra measured on optical excitation at 405 nm; (c) position of the PL peak vs MnCl2 percentage; (d) TRPL traces measured at the wavelength at which the PL is maximum.

Close modal

Apart from the shift in the PL peak position, the TRPL lifetime [Fig. 3(d)] is influenced by the presence of Mn2+. We performed mono-exponential fits on the TRPL traces (Fig. S3) and collected the TRPL lifetimes in Table SIII. The TRPL lifetime for the MAPbI3 sample is 70 ns. For the film with asol = 3%, the TRPL lifetime is a factor 2 larger (140 ns). This enhancement is in agreement with that measured in Ref. 19 by means of transient absorption experiments. On the contrary, the formation of Mn-rich domains causes a reduction in the TRPL lifetime. This occurs to some extent in the film with asol = 10% and becomes particularly evident for even higher Mn2+ concentrations. Furthermore, the PL emission of the film with asol = 30% is significantly quenched in comparison with the other samples (Fig. S4). These results indicate that for large Mn2+ concentrations non-radiative recombination pathways dominate the charge dynamics, which will be discussed in more detail below.

We further investigated the charge dynamics of light induced excess carriers by TRMC.37 In short, photoexcitation of the film by means of a nano-second laser pulse at 550 nm results in the generation of free charges, which leads to an increase in the conductance of the sample. The photo-conductance (ΔG) is probed by measuring the reduction in reflected microwave power in time. After an initial rise corresponding to photogeneration of free charges, ΔG decays as a result of charge recombination and/or immobilization of charges in deep trap states. The TRMC traces have been corrected for the fraction of absorbed photons, Fa, which eliminates the effect of optical inhomogeneities. Intensity-normalized TRMC traces are displayed in Figs. 4(a)–4(d) for films with different concentrations of Mn. For the film with asol = 3%, the TRMC lifetimes are slightly longer [compare Fig. 4(a) with 4(b)], in line with the TRPL observations. However, for films with asol > 3%, the charge carrier lifetimes drop significantly.

FIG. 4.

[(a)–(d)] TRMC traces for mixed-metal perovskite thin films with different concentrations of Mn, recorded upon pulsed laser excitation at 550 nm at room temperature. The change in photoconductance, ∆G, is normalized for the number of incident photons I0 and for the fraction of absorbed photons Fa. Solid lines represent experimental data, and dashed lines are calculated using the model presented in Ref. 38; (e) schematic representation of the kinetic model used for the modeling. The meaning of the symbols is reported in Table II; (f) sum of electron and hole mobilities, ∑μ, as function of the percentage MnCl2 in the precursor solution.

FIG. 4.

[(a)–(d)] TRMC traces for mixed-metal perovskite thin films with different concentrations of Mn, recorded upon pulsed laser excitation at 550 nm at room temperature. The change in photoconductance, ∆G, is normalized for the number of incident photons I0 and for the fraction of absorbed photons Fa. Solid lines represent experimental data, and dashed lines are calculated using the model presented in Ref. 38; (e) schematic representation of the kinetic model used for the modeling. The meaning of the symbols is reported in Table II; (f) sum of electron and hole mobilities, ∑μ, as function of the percentage MnCl2 in the precursor solution.

Close modal

Hutter et al. modelled TRMC traces of MAPbI3 films using a kinetic model representing a p-type semiconductor including band-to-band recombination and recombination via deep trap states [Fig. 4(e)].38 The trapping process in this model is assumed to be a second-order process. More details about the method and the set of coupled differential equations which describe the time evolution of the populations of free and trapped charges are given in the supplementary material. We used the same model to reproduce the TRMC data presented in Figs. 4(a)–4(d). Note that this fully mathematical model also holds for the opposite situation, i.e., an n-type perovskite with trap states for holes and additional dark CB electrons.

Using the parameters listed in Table II, the model reproduces the experimental trends fairly well [dashed lines in Figs. 4(a)–4(d)]. We note that for MAPbI3 without Mn2+ we obtain similar kinetic values as reported earlier.38 The film with asol = 3% is characterized by a slightly smaller k2 as compared to MAPbI3. This is well in line with the longer TRPL lifetime found for this sample. In addition, the presence of Mn2+ in the perovskite lattice has a large effect on the mobilities of both charges, which become greatly unbalanced. We note here that the TRMC does not allow distinguishing between electrons and holes since mobile carriers interact with the microwave radiation independently on the sign of their charge. Nevertheless, the kinetic model reported in the supplementary material makes a distinction between the majority and the minority charges. In order to reproduce the TRMC data for the sample with asol = 3%, we have to set the mobility of majority charges one order of magnitude larger than the mobility of minority charges.

TABLE II.

Parameters used for the modeling of the TRMC traces in Figs. 4(a)–4(d).

Parameter (units)Description0% Mn3% Mn10% Mn30% Mn
μ> (cm2 V−1 s−1Majority carrier mobility 20.1 11.1 6.1 1.7 
μ< (cm2 V−1 s−1Minority carrier mobility 17.0 0.5 2.1 1.4 
μ</μ> (−) Mobility ratio 0.85 0.05 0.34 0.82 
k2 (× 10−10 cm3 s−1Band-to-band recombination rate 6.9 6.4 7.2 15.0 
kT (× 10−10 cm3 s−1Trap filling rate 1.5 3.6 6.0 7.0 
kR (× 10−11 cm3 s−1Trap recombination rate 1.4 5.8 7.0 7.0 
NT (× 1015 cm−3Trap density 5.4 5.5 6.0 5.5 
d0 (× 1015 cm−3Background density 4.2 3.5 9.0 7.0 
Parameter (units)Description0% Mn3% Mn10% Mn30% Mn
μ> (cm2 V−1 s−1Majority carrier mobility 20.1 11.1 6.1 1.7 
μ< (cm2 V−1 s−1Minority carrier mobility 17.0 0.5 2.1 1.4 
μ</μ> (−) Mobility ratio 0.85 0.05 0.34 0.82 
k2 (× 10−10 cm3 s−1Band-to-band recombination rate 6.9 6.4 7.2 15.0 
kT (× 10−10 cm3 s−1Trap filling rate 1.5 3.6 6.0 7.0 
kR (× 10−11 cm3 s−1Trap recombination rate 1.4 5.8 7.0 7.0 
NT (× 1015 cm−3Trap density 5.4 5.5 6.0 5.5 
d0 (× 1015 cm−3Background density 4.2 3.5 9.0 7.0 

The formation of Mn-rich domains for films with asol > 3% leads to faster decay kinetics, which translates into an increase in k2. In addition, the TRMC data clearly show that alloying Mn2+ in MAPbI3 negatively affects the mobilities. Figure 4(f) shows the sum of electron and hole mobilities, ∑μ, used to calculate the TRMC traces in Figs. 4(a)–4(d). For the film with asol = 3%, a reduction of ∑μ by a factor 3 is found. With higher concentrations of Mn2+, an additional reduction in mobility, in combination with an increase in the trap density, trap filling, and trap recombination rate constants are observed.

To verify how the presence of Mn2+ changes the charge dynamics for electron and holes, we carried out DFT calculations including spin-orbit coupling to reveal the band structure of MAPbI3, MAPb0.9375Mn0.0625I3, and MAPb0.875Mn0.125I3 (Fig. 5). The choice of these compositions is motivated by the computational constrains of the DFT simulations. We have calculated the effective masses of electrons and holes, listed in Table III by using the parabolic fitting method. The effective masses of MAPbI3 are in good agreement with reported theoretical studies.39,40 Interestingly, the curvature of the valence band (VB) of MAPb0.9375Mn0.0625I3 is much wider than the curvature of the conduction band (CB) due to interaction of Mn with the rest of the Pb1-aI3 lattice (hybridization of s and d states of Mn with s state of Pb and p state of I). This means that the effective mass of the holes is much larger than the effective mass of the electrons (Table III), and hence the electrons are more mobile than holes. Another interesting point to note is that the effective mass of holes of the MAPb0.9375Mn0.0625I3 is significantly larger than that of MAPbI3, while the effective mass of the electrons is similar in the two samples. Given the effective masses calculated by DFT, we argue that the electrons in the mixed metal perovskite have a substantial larger mobility than the holes. Combining TRMC and DFT allows us to conclude that, in the sample with asol = 3%, the electrons are the majority carriers and that therefore the material is most likely n-type. For samples with a larger concentration of Mn2+, we keep the assumption of the n-type material.

FIG. 5.

Band structure (VBM and CBM in bold) of cubic (a) MAPbI3, (b) MAPb0.9375Mn0.0625I3, and (c) MAPb0.875Mn0.125I3 as calculated by DFT with spin orbit coupling.

FIG. 5.

Band structure (VBM and CBM in bold) of cubic (a) MAPbI3, (b) MAPb0.9375Mn0.0625I3, and (c) MAPb0.875Mn0.125I3 as calculated by DFT with spin orbit coupling.

Close modal
TABLE III.

Calculated effective masses of electrons and holes expressed in units of the free-electron mass, (m0) for MAPbI3, MAPb0.9375Mn0.0625I3, and MAPb0.875Mn0.125I3.

mn*mp*mp*/mn*mn*mp*mp*/mn*
MAPbI3 0.23 (R-Γ) 0.25 (R-Γ) 1.09 0.33 (R-M) 0.51 (R-M) 1.55 
MAPb0.9375Mn0.0625I3 0.25 (Γ-M) 2.79 (Γ-M) 11.16 0.35 (Γ-R) 1.37 (Γ-R) 3.91 
MAPb0.875Mn0.125I3 0.51 (Γ-R) 1.10 (Γ-R) 2.16 0.68 (Γ-M) 2.16 (Γ-M) 3.18 
mn*mp*mp*/mn*mn*mp*mp*/mn*
MAPbI3 0.23 (R-Γ) 0.25 (R-Γ) 1.09 0.33 (R-M) 0.51 (R-M) 1.55 
MAPb0.9375Mn0.0625I3 0.25 (Γ-M) 2.79 (Γ-M) 11.16 0.35 (Γ-R) 1.37 (Γ-R) 3.91 
MAPb0.875Mn0.125I3 0.51 (Γ-R) 1.10 (Γ-R) 2.16 0.68 (Γ-M) 2.16 (Γ-M) 3.18 

The effective masses of electrons and holes in MAPb0.875Mn0.125I3 are larger than those of MAPbI3. This is in agreement with the incorporation of a transition metal, such as Ag, Cu, or Au in MA2BB′I6 (B and B′ are metal cations). The corresponding effective masses are generally larger than those of lead halide perovskites.41 However, in view of the XRD data previously shown and the conclusions about the segregation of Mn2+ out of the perovskite structure for films with asol > 3%, we do not consider the calculated DFT effective masses of MAPb0.875Mn0.125I3 to be related to the lower TRMC mobilities observed in Figs. 4(c) and 4(d). Instead, at higher Mn2+ concentrations, it seems likely that the reduction in ∑μ is, at least partially, caused by a reduction in the size of crystalline domains in the film and the formation of more grain boundaries that would slow down the motion of free charges. Secondary electron SEM images (Fig. S5) show that the perovskite crystalline domains are smaller for the film with asol = 30% than those for the film with asol = 10%. Thus, the change in domain size is likely to lead to a reduction in the charge mobilities when the Mn2+ concentration is increased from asol = 10% to asol = 30%. In addition, the amorphous Mn2+ rich phases formed in films with asol > 10% are expected to enhance to charge decay as observed in the TRPL and TRMC measurements.

In this work, we partially replaced Pb2+ with Mn2+ in MAPbI3 and studied the change in structural and electronic properties of the resulting films. From structural and elemental analysis, we conclude that only a limited amount of the Pb2+ atoms at the B position of the perovskite structure can be replaced by Mn2+. Using precursor solutions with higher Mn2+ concentrations leads to the formation of additional amorphous Mn-rich domains. At 3% replacement, the band-to-band recombination rate slightly reduces yielding somewhat longer charge carrier lifetimes as determined from TRPL and TRMC measurements. However, this is accompanied by a substantial increase in the effective mass of the holes, resulting in a strongly unbalanced charge transport. According to the recent results of Zheng et al., the increase in charge carrier lifetimes still positively affects the photovoltaic performance of devices made from this composition.19 For higher concentrations of Mn2+, formation of Mn-rich domains restores the balanced transport, although both electron mobility and hole mobility become ca. 1 order of magnitude smaller than in MAPbI3 due to enhanced charge scattering. Besides the lowering of the mobilities, a reduction in the charge lifetime is found. At higher concentrations of Mn2+, the reduced charge carrier diffusion lengths will be severely limiting the photovoltaic properties of solar cells prepared from these mixed metal halide perovskites. These findings are in line with the lower efficiencies observed in MAPbI3 solar cells in which more than 5% of the Pb2+ was replaced by Mn2+.18,19 Combination of experimental and theoretical data indicates that alloying MAPbI3 with small amounts of Mn2+ is a unique handle to manipulate the dynamics of the carriers.

See supplementary material for the differential equations used for modeling the TRMC data and for additional data.

This research was carried out under Project No. F71.4.15562a in the framework of the Partnership Program of the Materials innovation institute M2i (www.m2i.nl) and the Foundation of Fundamental Research on Matter (FOM) (www.fom.nl), which is part of the Netherlands Organization for Scientific Research (www.nwo.nl). S.T. and J.J. acknowledge funding by the Computational Sciences for Energy Research (CSER) tenure track program of Shell, NWO, and FOM (Project No. 15CST04-2). D.B. and T.J.S. are thankful to Valentina Caselli and Duco Bosma for the collection of SEM data and Dengyang Guo for discussion on the TRMC data interpretation.

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Supplementary Material