Copper-incorporation for polytypism and bandgap engineering of MAPbBr₃ perovskite thin films with enhanced near-Infrared photocurrent-response

Abstract

The optoelectronic properties of lead-based halide perovskites can be enhanced through B-site engineering. Here, we studied the B-site alloying of MAPbBr₃ thin films with copper (Cu²⁺). The alloyed perovskite thin films were characterized by a dark color, enlarged average grain boundary, and lowering of the optical bandgap from 2.32 eV for pristine MAPbBr₃ to 1.85 eV for 50% Cu-substituted MAPbBr₃. Various characterization methods revealed that the Cu-incorporation leads to the appearance of a Cu-rich secondary phase. The conductivity increased over three orders of magnitude upon alloying. Temperature-dependent conductivity measurements at temperatures ranging from 110 K to 300 K revealed the occurrence of two phase-transitions in Cu-substituted perovskite, and only one transition in pristine MAPbBr₃. Photocurrent measurements of the alloyed perovskites showed that band-carrier generation occurred upon excitation in the near-infrared region. First-principles point defect calculation shows the likelihood of compensating Br vacancy formation with high Cu-substituting concentrations. Calculation of atomic orbital projected density of states (Cu_Pb + v_Br defect complex) revealed the presence of localized defect states within the pristine bandgap, explaining the observed sub-bandgap absorption. The results provide an insight into the alloying importance in phase-modulation and tailoring the optoelectronic properties of perovskites for a wide range of efficient optoelectronic devices.

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Introduction

Lead-based perovskites have been widely utilized in optoelectronic devices, owing to their attractive features of strong absorption coefficient,^1–3 high charge carrier mobility,^4–8 tunable bandgap,^9–12 long diffusion length,^13–15 and low exciton binding energies.^16–19 One of the fundamental approaches to developing efficient perovskite-based optoelectronic devices is doping/alloying. Via doping, defect/impurity levels, introduced by dopants, overlap with the valence band (or the conduction band) of the host perovskite. As a result, different perovskites with versatile optical and electrical properties can be manipulated. Three types of MAPbX₃ doping have (in general) been investigated in numerous studies: (1) A-site doping where a methylammonium MA⁺ cation is doped with A⁺ (A⁺ = FA⁺, Rb⁺, Cs⁺, or K⁺).^20–30 (2) B-site doping where lead Pb²⁺ is doped with different metal cations Mⁿ⁺ (Mⁿ⁺ = Sn²⁺, Ca²⁺, Cd²⁺, Sr²⁺, Bi³⁺, Au³⁺, Ba²⁺, Mn²⁺, Ni²⁺, In³⁺, Cu²⁺, or Er³⁺).^31–44 (3) Halide anion X⁻ doping with different anions (Cl⁻, Br⁻, I⁻, or Se²⁻).^45–49

MAPbBr₃ perovskite has considerable potential for laser applications^50–53 and photodetection.^54,55 Various investigations have focused on enhancing the optoelectronic properties of this material via doping. Deng et al. have studied NH₄⁺-doped MAPbBr₃ perovskite quantum dots.⁵⁶ They found that NH₄⁺ inhibits non-radiative recombinations and reduces the defect density of MAPbBr₃. Li et al. have proposed local lattice strain compensation of Zn²⁺-doped MAPbBr₃ single crystal for photo-restriction under light illumination.⁵⁷ Furthermore, through Er³⁺-doping, the responsivity and external quantum efficiency of MAPbBr₃ single crystal photodetector increased five-fold relative to that of the pristine crystal.⁴⁴ Moreover, Bi³⁺-doped MAPbBr₃ single crystals have been extensively investigated. Early investigations by Abdelhady et al. suggested that Bi-doping-induced changes in the absorption spectrum resulted from bandgap modulation.³⁹ However, subsequent studies showed that Bi-doping increases the trap states within the bandgap of the MAPbBr₃ host without any effect on the bandgap and the exciton binding energy.^{38,40,42,43,58}

Polytypism is one example of polymorphism where the crystal structure is identical along two dimensions but different along the third dimension. Polytypism of hybrid lead-based perovskites has been reported in a few papers.^59–61 The polytypes of hybrid perovskite structure ABX₃ depend on the stacking sequence of the close-packed AX₃ layers. The most common polytype of hybrid perovskite is the 3C polytype with BX₆ corner-sharing octahedra, which is favorable for perovskites with tolerance factors of t ≤ 1.⁶⁰ The label of 3C resembles the three close-packed layers of the aristotype cubic cell (C).⁶² The other most common polytypes of perovskite oxides are hexagonal polytypes, which are favorable for perovskites with tolerance factors of t > 1, which can be obtained through using large A-cations or small B-cations.⁶¹ A 2H hexagonal polytype with BX₆ face-sharing octahedra leads to a one-dimensional (1D) array.⁵⁹ However, the other 4H and 6H hexagonal polytypes, which have three-dimensional (3D) frameworks, are composed of BX₆ face-sharing octahedra (h) and BX₆ corner-sharing octahedra (C).⁶⁰

In this study, we have investigated the Cu²⁺-chemical alloying effect on MAPbBr₃ as a host material. Upon substitution, MA(Pb:Cu)Br₃, thin films exhibit color darkness with increasing size of the average grain boundaries. Optical measurements indicate a sharp bandgap reduction with 50% replacement. X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS) and energy dispersive spectroscopy (EDS) mapping measurements indicate that Cu²⁺ ions are incorporated into the MAPbBr₃ crystal structure with the appearance of a new phase related neither to the crystal structure peaks of MAPbBr₃ nor the initial precursor. Theoretical simulations were utilized to further explore the Cu-incorporation mechanism. Photocurrent measurements of the doped perovskites revealed band carrier generation at an excitation photon energy below the bandgap, suggesting considerable potential for a wide range of optoelectronic applications.

Results and discussion

Anti-solvent assisted crystallization of methylammonium lead bromide perovskites plays the most significant role in DMF, at the crystallization temperature (70 °C).³⁹ In this work, pristine and Cu-alloyed MACu_xPb_(1−x)Br₃ (where x = 0%, 10%, 20%, 30%, 40%, and 50%) polycrystalline thin films were prepared via spin-coating of the precursor solution over a glass substrate with the assistance of toluene as an anti-solvent, as shown in Fig. 1 The perovskite films with different Cu²⁺ concentrations differ in color (see Fig. 1b) from yellowish orange for the non-alloyed MAPbBr₃ to almost dark brown for the 50% Cu-substituted MAPbBr₃.

Fig. 1 (a) Graphical illustration of perovskite thin film fabrication via spin-coating. (b) Picture of pristine (0%) and (10, 20, 30, 40, and 50%) Cu-substituted MAPbBr₃ thin films obtained from use of (0, 10, 20, 30, 40, and 50%) [CuBr₂]/([CuBr₂] + [PbBr₂]) mole percentages.

The top-view scanning electron microscopy (SEM) images presented in Fig. 2a–f show the morphology of the perovskite thin films. Upon Cu²⁺ alloying, the average grain boundary area increases from 0.042 μm² for pristine MAPbBr₃ to 0.71 μm² for 40% Cu-substituted MAPbBr₃, as shown in Fig. S1 (ESI†). However, an SEM image of the 50% Cu-substituted MAPbBr₃ thin film revealed that large micro-rods are formed as a side product. Energy dispersive scanning (EDS) elemental mapping results reveal the presence of homogeneously distributed Pb (blue-colored images in Fig. S2, ESI†). However, the Cu-elemental mapping results (red-colored images in Fig. S2, ESI†) show the gradual increase in Cu alloying. Cu is homogeneously distributed until an alloying level of 20% is reached. After that, a Cu-rich new phase appears in the film which is clearly observed at an alloying level of 50%.

Fig. 2 SEM images of undoped (a), 10% (b), 20% (c), 30% (d), 40% (e), and 50% (f) Cu-alloyed MAPbBr₃ thin films.

Fig. 3a shows the XRD pattern of the pristine and Cu-substituted MAPbBr₃ thin films. The alloyed samples maintain the cubic perovskite structure of MAPbBr₃ with three main diffraction peaks at (100), (200), and (300).⁶³ However, with incorporation of Cu into the material, new XRD peaks occur at 2θ = 11.2° (‡) and 2θ = 11.5° (§). Regarding to XRD data within 2theta range 14.6°–15.3° (As depicted in Fig. 3b), the main diffraction peak (100) of MAPbBr₃ cubic structure is shifted towards lower theta with increasing of copper alloying from pristine MAPbBr₃ into 20% alloyed one. This reveals MAPbBr₃ crystal lattice expansion owing to the copper occupation of the interstitial positions through MAPbBr₃ cubic structure rather than Cu/Pb metal substitution. This is consistent with the large difference in effective ionic radius of Pb²⁺ (119 pm) and Cu²⁺ (73 pm). In addition, Cu perovskites show Jahn–Teller distortion which stereochemically hinders the upkeeping of 3D perovskite structure. With further Cu-alloying, the XRD peak is shifted-back to its original position. This might be accounted for the preference of Cu-atoms to form the Cu-rich secondary phases rather than the occupation of the MAPbBr₃ interstitial positions. A plot of the logarithm (intensity) vs. 2θ, as shown in Fig. 3c, reveals four regions. These include the purple regions of 3D-MAPbBr₃ and peaks corresponding to crystalline residual metal halide salt in the perovskites with Cu-concentration ranging from 30% to 50% (blue and green regions represent CuBr₂ and PbBr₂, respectively). Interestingly, the orange region reveals a new diffraction peak, which occurs for 20% Cu-substituted MAPbBr₃ (‡). With further substitution, the intensity of this peak decreases while the intensity of another peak (occurring at a higher angle) increases (§). These extra XRD peaks may have resulted from two factors. One factor is that the formation of Pb-substituted 2D Cu-based perovskite secondary phase MA₂CuBr₄ may have contributed to their emergence. However, this factor can be excluded as the main peak of MA₂CuBr₄ perovskite should be appeared near 2θ = 9.2° not in the range (11°–12°) whatever the concentration of lead alloyed, as depicted in Fig. S3 (ESI†).⁶⁴ The other possible factor is the formation of a Pb-based perovskite secondary phase, owing to Cu-alloying. According to literature,⁶⁰ Pb-based perovskite, namely (FAPbI₃)_x(MAPbBr₃)_1−x, exhibits a sequential hexagonal polytypism (i.e., 2H (delta phase)-4H-6H-3R(3C)) during the crystallization process of its perovskite thin film (H, R, and C represent hexagonal, rhombohedral, and cubic phases, respectively). Moreover, the hexagonal phases have one dimensionality (1D) structure with XRD fingerprint peaks for 2θ ranging from 11.6° to 14.1°, and therefore we hypothesize that both peaks, i.e., ‡ and §, correspond to new phases, which are neither related to the crystal structure of 3D-MAPbBr₃ nor the initial precursors. These peaks may, however, be related to 1D hexagonal polytypes of Cu-MAPbBr₃ perovskite, which would be consistent with the formation of a rod-shaped phase in the 50% Cu-substituted thin film. Regarding to these observations, we speculate that hexagonal polytypes of Cu-substituted MAPbBr₃ perovskite. This might be attributed to the presence of MAPbBr₃ cubic perovskite which induces the excess copper to form polytypes of the most stable 1D microrods rather than 2D sheets.

Fig. 3 (a) XRD patterns of Cu-alloyed MAPbBr₃ perovskite thin films. (b) XRD peak 〈100〉 shift upon copper addition. (c) XRD patterns (log scale) for 2θ = 10–17°. (d) Energy above convex hull and bandgap for different phases of pure and Cu-alloyed MAPbBr₃ where C, 2H, 4H, and 6H represent cubic, hexagonal-2H, hexagonal-4H, and hexagonal-6H phases of MAPbBr₃, respectively.

To verify the validity of the assumption that hexagonal polytypes are formed in Cu-substituted MAPbBr₃, we calculated the energy above the convex hull (E_hull) of different phases (namely Cubic, 2H, 4H, and 6H (only for 50% alloying)) with varying substituted concentrations (MACu_xPb_1−xBr₃; x = 0, 0.25, 0.5, and 0.75). The possible crystal structures are presented in Fig. S4–S7 (ESI†). Fig. 3d shows E_hull calculation results where all possible secondary phases are taken from the Open Quantum Materials Database (OQMD).⁶⁵E_hull can be seen as a measure of comparative stability between different phases. It should be noted that E_hull is calculated at T = 0 K and a finite temperature might have some effect on comparative stability. For the pristine MAPbBr₃ perovskite, almost the same E_hull is obtained for the aforementioned four phases. However, for the Cu-alloyed compounds, the E_hull of the hexagonal phases differs modestly (by ∼10 meV per atom) from that of the cubic phase. This might explain some secondary hexagonal phase formation with Cu alloying as observed experimentally. Furthermore, the bandgap differs among the phases of the Cu-alloyed perovskites with the cubic phase showing the lowest bandgap in all the alloying concentrations. For the hexagonal phase bandgap decreases when we go from 2H > 4H > 6H.

In order to study the effect of Cu²⁺ incorporation on the elemental composition of Cu-MAPbBr₃ films, X-ray photoelectron spectroscopy (XPS) was performed, as depicted in Fig. 4 and Fig. S8 (ESI†). The Pb 4f spectrum for the MAPbBr₃ film reveals 4f_5/2 and 4f_7/2 peaks at 143.4 and 138.6 eV, respectively, corresponding to the Pb²⁺ cations⁶⁶ (Fig. 4a). The Br 3d spectrum (Fig. 4b) shows 3d_3/2 and 3d_5/2 peaks at 69.7 and 68.5 eV, respectively, corresponding to the Br⁻ anions. The C 1s spectrum for the MAPbBr₃ film presents C–N and C–C peaks at 286.2 and 284.7 eV, respectively, corresponding to the methylammonium CH₃NH₃⁺ cations (Fig. 4c). The non-shifting of Pb 4f, Br 3d, and C 1s peaks, with Cu-alloying, is consistent with the formation of MAPbBr₃ film as a major phase. In addition, Cu²⁺ cations embedded in the film are evident with the presence of Cu 2p_3/2 at 932.9 eV⁶⁷ for 10% and 20% samples which are slightly shifted to lower binding energy at 932.7 eV for higher concentration Cu-alloyed perovskites (Fig. S9a, ESI†), revealing the difference in chemical structure nature of Cu²⁺ alloyed in perovskite which is consistent with XRD results. The formation of Cu-satellite structure at higher Cu-alloyed samples (Fig. 4d) indicates the presence of Cu²⁺ species, as expected.⁶⁸ Gradual increasing of Cu 2p_3/2 peak area is a clear indication of increasing of copper concentration embedded in perovskite thin films upon alloying (Fig. S9b, ESI†). Notably, after incorporation of 50% Cu, the Br 3d, and C 1s peaks area is clearly changed which might be attributed to the Cu-rich secondary phase formation. Furthermore, new peaks at 141.4 eV and 136.6 eV are formed through the Pb 4f spectrum which might be assigned to the formation of Pb metal or the Pb occupation of the Cu-rich secondary phase interstitial positions.⁶⁹ The changes in peak area behavior for Pb and Br are depicted in Fig. S9c and d (ESI†).

Fig. 4 XPS spectra of Pb4f (a), Br3d (b), C1s (c), and Cu2p (d), and Br3d (e) of Cu-alloyed MAPbBr₃ perovskite thin films.

The optical properties of the samples are evaluated by obtaining UV-Vis absorption spectra of the pristine and Cu-substituted MAPbBr₃ thin films (as illustrated in Fig. 5a). The absorption edges undergo a slight red-shifting with alloying of up to 40% Cu, and subsequently, a sudden and considerable red-shift (near-IR range) occurs for the 50% Cu-substituted sample. Tauc plots, Fig. 5b, and Fig. S10 (ESI†), show that the optical direct bandgaps of the perovskites are 2.32 eV for pristine MAPbBr₃, 2.31 eV for 10%, 2.31 eV for 20%, 2.29 eV for 30%, 2.28 eV for 40%, and 1.85 eV for 50% Cu-substituted MAPbBr₃ thin films. Tuning of the optical bandgap energies is consistent with the perovskite thin film color tuning (from yellowish orange to dark brown) with increasing Cu content. Moreover, the trend observed for the absorption spectrum of the 50% Cu-substituted film is similar to that observed for Cu-based perovskite materials.⁷⁰ This may have resulted from the higher concentration of Cu-rich secondary phase in the 50% Cu-substituted film compared with the concentrations occurring in the other films. The inhomogeneous broadening of PL emission peaks with Cu-doping, as shown in Fig. S11a (ESI†), further confirms the heterogeneity of the doped films. The results revealed that 10% doping enhances room temperature PL emission of MAPbBr₃ (76% increase in PL peak area), as shown in Fig. S11b (ESI†). This may have resulted from the homogeneously distributed low-Cu-concentration doping effect of the 10% Cu perovskite thin film without secondary phase formation. However, further Cu-doping (of up to 50% doping) leads to quenching of the PL emission and 85% recovery of the pristine MAPbBr₃ PL emission (Fig. S11b, ESI†). Normalized PL spectra (Fig. S11c, ESI†) revealed that, with 50% Cu-doping, the PL shifts slightly (0.2 nm) to longer wavelength than those observed at other doping levels (Fig. S11d, ESI†).

Fig. 5 (a and b) Absorption spectra, and direct bandgap estimation: Tauc plot, respectively, of Cu-alloyed MAPbBr₃ thin films.

Using the four-probe method, we measured the electrical conductivity of perovskite thin films as a function of Cu alloying, as shown in Fig. 6a. The conductivity increases significantly with alloying and the highest value is obtained for the 40% Cu-substituted thin film, owing possibly to the presence of homogeneously distributed secondary phases. Subsequently, the conductivity decreases slightly for the 50% Cu substituted material, consistent with the heterogeneously distributed Cu-rich new phase formed in this material (as depicted in Fig. 2f and Fig. S2, ESI†). The 40% Cu-substituted MAPbBr₃ thin film exhibits the highest dark electrical conductivity (10.77 × 10⁻⁴ S cm⁻¹), which is nearly three orders of magnitude higher than that of the pristine MAPbBr₃ thin film (4.10 × 10⁻⁷ S cm⁻¹). To evaluate the secondary phase stability, we measure the conductivity under the effect of temperature (temperature range: 300 K–110 K) for one cooling/heating (C/H) cycle. We first measure the conductivity upon cooling from room temperature (300 K) to 110 K, and then upon heating back from 110 K to 330 K. Fig. 6b shows the results of temperature-dependent conductivity measurements performed on pristine MAPbBr₃ where the conductivity decreases from 7.05 × 10⁻⁸ S cm⁻¹ at 300 K to 2.72 × 10⁻⁸ S cm⁻¹ at 110 K. The reverse behavior is observed upon heating, i.e., increasing conductivity with slight hysteresis at temperatures ranging from 200 K to 220 K. This behavior may have resulted from a phase transition where the tetragonal structure at T < 200 K transforms into a cubic structure at T > 220 K.⁷¹Fig. 6c shows the results of the conductivity measurements performed on 40% Cu-MAPbBr₃. The conductivity decreases sharply from 1.16 × 10⁻³ S cm⁻¹ at 300 K to 2.06 × 10⁻⁷ S cm⁻¹ at 270 K and then decreases to 4.47 × 10⁻⁸ S cm⁻¹ at 110 K. Interestingly, upon heating, the 40% Cu sample behave like pristine MAPbBr₃ where the conductivity increases slightly to 2.66 × 10⁻⁷ S cm⁻¹ at 330 K. The cooling/heating cycle behavior of the 40% Cu samples can be attributed to a phase transition from the secondary hexagonal phase at room temperature (300 K) to the cubic phase at 270 K. The inset image shown in Fig. 6c reveals that, at temperatures of 200–220 K, the same hysteresis occurs as the pristine perovskite (Fig. 6b), thereby confirming the occurrence of another phase transition to the tetragonal structure at T < 200 K. The Cu-alloyed perovskite thin film recovers its original conductivity value (1.08 × 10⁻³ S cm⁻¹) after three days (Fig. 6d). This indicates that, compared with the pristine material, the 40% Cu perovskite requires more time for re-formation of its more stable hexagonal secondary phase with its original higher conductivity value. The activation energy (E_a) can reflect the energy difference between the conduction band and the Fermi energy (or defect levels). E_a required for charge transport can be obtained from conductivity measurements based on the Arrhenius equation S = S₀ exp(−E_a/k_BT) where S, S₀, k_B, and T are the conductivity, a constant, the Boltzmann constant, and the absolute temperature, respectively. We obtained an E_a of 29 meV for pristine MAPbBr₃ and a considerably lower estimated value of 2 meV for the 40% Cu-substituted perovskite, as depicted in Fig. S12 (ESI†). The reduced E_a required for the alloyed perovskite may correspond to the change in Fermi energy or arising defect levels. This suggests that the carrier density increases with Cu-alloying where the presence of hexagonal secondary phases may play a crucial role in E_a reduction. For additional insight into the temperature-dependent conductivity behavior, the phase transitions in the 40% Cu-substituted perovskite are further tracked during three consecutive days of XRD measurements. The main cubic MAPbBr₃ cubic structure of the material is retained after one day of the (C/H) cycle (Fig. 6e). However, the XRD peak of the hexagonal secondary phase is shifted to lower 2θ than the initial values (Fig. 6f). This shifting is accompanied by the appearance of a peak at 2θ = 9.2°, which corresponds to two-dimensional (2D) MA₂CuBr₄ perovskite.⁶⁴ Moreover, the XRD peak at 2θ = 14.3°, which is associated with the unreacted CuBr₂, disappears after the (C/H) cycle, thereby further confirming the formation of 2D Cu-based perovskite. This formation may have led to the aforementioned sudden decrease in the conductivity of the alloyed sample after the one-day (C/H) cycle. With the progression of the measurement, the XRD peak intensity of the MA₂CuBr₄ perovskite decreases, and the peak intensity of the hexagonal secondary phase increases. The direct dependence of the XRD peak intensity ratio (Iphase transitions/IMA₂CuBr₄) on time reveals that the hexagonal secondary phases are restored at the expense of 2D MA₂CuBr₄ perovskite (as depicted in Fig. S13, ESI†). Consequently, the original conductivity of the 40% Cu-substituted perovskite is restored after the three-day measurement period.

Fig. 6 (a) Conductivity of Cu-MAPbBr₃ perovskite thin films at room temperature. (b and c) Temperature-dependent conductivity of pristine MAPbBr₃ and 40% Cu-MAPbBr₃ perovskite thin films, respectively for one cooling/heating cycle. The inset in (c): zoom for low conductivity values. (d and e) Time-dependent conductivity and time-dependent XRD patterns, respectively, of 40% Cu-MAPbBr₃ perovskite thin film after cooling/heating cycle effect. (f) Plots on log scale for 2θ = 8–17°.

To determine the potential of pristine/alloyed perovskite thin films for optoelectronic devices, photocurrent measurements were performed as a function of the excitation wavelength (i.e., the excitation photon energy). Two types of transitions are (in general) observed (Fig. 7a): type I is a band-to-band transition (i.e., from valence band to conduction band), which requires higher excitation energy than the bandgap energy for photocarrier generation. Type II, which is a transition between a defect level and either the valence or conduction band, occurs with excitation energy either higher or lower than the bandgap energy for photocarrier generation.⁷²Fig. 7b shows the normalized photocurrent spectra of both pristine and 40% Cu-substituted MAPbBr₃ thin films as a function of the excitation photon energy. When the energy of excitation with photons is higher than the bandgap, E_Ph > E_g, the photocurrent of each sample rises sharply (i.e., intrinsic photocurrent generation occurs) despite the difference in alloying levels. Thereafter, the photocurrent decreases with increasing excitation photon energy. This may have resulted from the reduction in the penetration-depth (where carriers are generated at the perovskite surface), owing to rapid deterioration induced by the surface effect. It should be noted that, at higher energy, pristine perovskite reveals a higher surface recombination rate than that of 40% alloyed one, as shown in Fig. 7b, which might be due to the increasing of penetration depth upon Cu-alloying. No photocurrent generation occurs in pristine MAPbBr₃ subjected to excitation at the near-infrared region. However, band-carrier generation increases with increasing Cu-substituting level. This strongly suggests that sub-bandgap absorption levels (defect levels) are present in the Cu-alloyed samples. Furthermore, defect level-related photocurrent is generated (i.e. Type II transition) above the threshold energy (E_T), i.e., 1.25 eV (a defect level is located at ∼1.25 eV below the conduction band (or above the valence band)). This defect level behavior has been observed previously.^72–75 The bandgaps of the perovskites are 2.24 eV for pristine MAPbBr₃, and 2.22 eV, 2.19 eV, 2.14 eV, 1.90 eV, and 1.78 eV for 10%, 20%, 30%, 40%, and 50% Cu-substituted MAPbBr₃ thin films, respectively (as depicted in Fig. S14, ESI†). These values correspond closely to the optical bandgaps obtained from the Tauc plots. Bandgap narrowing, with near IR carrier generation, makes Cu-alloyed MAPbBr₃ perovskite thin film a promising candidate for NIR Photodetectors⁷⁶ and other optoelectronic devices.⁷⁷ It should be hinted that the only MAPbBr₃ based photodetector, with narrow NIR-visible light dual-mode, was recently reported by Zhao et al.⁷⁶

Fig. 7 (a) Graphical illustration of photo-excitation processes in Cu-alloyed MAPbBr₃ perovskite and (b) Photocurrent spectra of pristine MAPbBr₃ and 40% Cu-alloyed MAPbBr₃ thin films at room temperature.

To gain further insight into Cu alloying of MAPbBr₃, first principles point defect simulations were performed (details about the formalism are presented in ESI†). Here, we designate cubic MAPbBr₃ as the parent phase and calculate the probability of Cu_Pb extrinsic-defect formation. We employ a high Cu alloying concentration of 12.5% (5.59 × 10²⁰ atoms per cm³) in the model. The findings reported in the literature regarding the oxidation state of Cu substituted into halide perovskites are inconsistent^72,78–81 (both +1 and +2 charge states have been reported). In our calculation, we consider both +1 (Cu_Pb⁻¹) and +2 (Cu_Pb⁰) oxidation states of Cu, for modeling the Pb-replacing perovskite octahedra. The simulation results (Fig. 8a and b) revealed that, in the highly p-alloyed system (when the Fermi level is close to VBM), the neutral state is more stable than the other states. However, in the n-alloyed samples, Cu behaves as an acceptor type defect. Moreover, Cu substitution creates a charge imbalance in the system and can therefore trigger the formation of intrinsic secondary donor defects that compensate for this imbalance. The results reported in the literature indicate that, in the case of MAPbBr₃, the most probable intrinsic donors are a Br vacancy (v_Br) and an MA interstitial (MA_i).⁸² Consequently, in this study, we consider Cu_Pb, v_Br, and MA_i individual point defects as well as Cu_Pb + v_Br and Cu_Pb + MA_i defect complexes with different charge states. With Cu alloying, the octahedral geometry changes substantially (mainly due to size mismatch between Cu and Pb), and as such we consider Br vacancies at two different (namely axial and equatorial) sites of the octahedron. We incorporate the effect of chemical growth environment by considering the chemical potentials of the constituent elements.

Fig. 8 Formation energies of simulated point defects with respect to Fermi level associated with two different growth environments (Pb rich/Br poor (a) and Pb poor/Br rich (b)). Defect charge states (q) are the slope of the curves (we performed simulations for q = −1, 0, and +1) and charge transition points are denoted by circles. Red- and purple-shaded regions represent the valence band (and below) and conduction band (and above), respectively. (c) Compositional phase diagram showing competing secondary phases of MAPbBr₃. (d) Zoomed view of the region enclosed in the box in (c).

To elaborate, we draw a compositional phase diagram of MAPbBr₃ (Fig. 8c and d) considering all the secondary phases from OQMD. We choose two corner points (chemical potential values to be used in defect formation energy calculation) of the diagram, which represent Pb rich/Br poor and Pb poor/Br rich conditions. The chemical potential of Cu is kept as rich as possible (for consistency with our experimental conditions) by considering elemental Cu and CuBr₂ phases (experimental precursor) (additional details on the chosen chemical potential values are presented in the experimental section). To validate our methodology, we compare our calculated formation energies and charge transition levels of two intrinsic defects (namely v_Br and MA_i) with values reported in the literature, and a reasonable agreement is realized.⁸³ The point defect diagrams (Fig. 8a and b) indicate that Cu alloying will be more likely under Pb poor/Br rich (p-type) growth conditions than under Pb rich/Br poor growth conditions. It is likely that the Fermi level pinning will be near CBM (VBM) under Pb rich/Br poor (Pb poor/Br rich) conditions. In this case, the formation energy of compensating intrinsic defects and defect complexes remains close to that of Cu_Pb in both chemical environments. This means that defect compensation is likely for Cu alloying under these conditions. From our analysis, v_Br turns out to be the most probable compensating defect. We also calculate the defect complex binding energy (E_b = E_form(Cu_Pb + v_Br) − E_form(Cu_Pb) − E_form(v_Br)) at different Fermi levels. Our calculated value lies between −0.41 eV (E_F at VBM) and −0.53 eV (E_F at CBM), indicating that formation of the Cu_Pb + v_Br defect complex is quite likely. However, for Cu_Pb + MA_i, E_b lies between 0.04 eV (E_F at VBM) and 0.26 eV (E_F at CBM), suggesting a moderate to low formation tendency of this complex. Moderate to deep-level (0.2–0.5 eV above VBM) charge transitions are observed for Cu_Pb and the associated defect complexes.

We evaluate the effect of Cu alloying (and possible defect complexes) on the electronic properties of MAPbBr₃ by plotting the orbital projected density of states for 50% Cu substituted MAPbBr3 with (and without) a Br vacancy, as shown in Fig. 9. With only Cu alloying at Pb sites, Fig. 9b, we observe that the effective gap decreases slightly (from 2.12 eV to 1.90 eV) compared with that of the pristine material, as depicted in Fig. 9a. In Fig. 9c, with the formation of Br vacancy (Cu_Pb + v_Br, defect complex), the presence of localized defect states can be observed within the pristine bandgap, formed by Cu-d and Br-p orbitals. This may explain the experimentally observed sub-bandgap absorption with 40% Cu alloying (Fig. 7b).

Fig. 9 Atomic orbital projected density of states for (a) pure MAPbBr₃, (b) 50% Cu alloyed MAPbBr₃, and (c) 50% Cu alloyed MAPbBr₃ with a Br vacancy.

Experimental

Characterization

Scanning electron microscopy (SEM) with energy dispersive spectroscopy (EDS; JEOL, JEM-2100F) was used to determine the Pb/Cu atomic weight percentages and to obtain the elemental mapping of the perovskite samples. The results confirmed that the number of grain boundaries increased, and a Cu-rich secondary phase formed during Cu doping. The photoluminescence of the perovskite was characterized via Raman spectroscopy (JASCO, NRS4500 NMDS) performed with 532-nm excitation (green) lasers. Absorption measurements performed via ultraviolet-visible near-infrared (UV-Vis-NIR) spectroscopy (JASCO, V-670) were used to estimate the bandgap from a Tauc plot. X-ray (Rigaku XRD) reflection data was measured at a wavelength of 1.5406 Å. X-ray photoelectron spectroscopy (XPS) was characterized on JEOL JPS-9030 with X-ray source of Mg K Alfa (Monochromatic) with 25 mA beam current and 12 kV. In addition, electrical measurements were conducted using a source meter (ADCMT, 6245) with a temperature controller installed in a four-probe station (Riko International Ltd.). The sample temperature was controlled by cooling the sample stage of the four-probe station using liquid nitrogen with a temperature controller under vacuum condition (∼5.0 × 10⁻² Pa). Photocurrent measurements at room temperature were performed in the station using a source meter (KEITHLEY, 2611A) with monochromatic light illumination. The wavelength of this illumination was controlled by a monochromator (BUNKOKEIKI, SM2-25).

Computational details

First-principles calculations were performed using Density Functional Theory (DFT)⁸⁴ within the Projector-Augmented Wave (PAW)⁸⁵ formalism as implemented in a GPAW⁸⁶ code combined with an Atomic Simulation Environment (ASE).⁸⁷ A plane wave cutoff energy of 800 eV and k-point density of 6 Å⁻¹ (12 Å⁻¹) were employed for geometry optimization (ground state calculations). Electronic structure calculations were performed using a Perdew–Burke–Ernzerhof (PBE)⁸⁸ exchange–correlation functional. In addition, a 2 × 2 × 2 (96 atoms) supercell of the MAPbBr₃ cubic unit cell was used for point defect calculations. A k-point density of 3.5 Å⁻¹ (7 Å⁻¹) was used for sampling the Brillouin zone for the supercells. Forces were converged up to 0.02 eV Å⁻¹ for the defect calculations, where the supercell size was kept fixed. Compositional phase diagrams were obtained using a CPLAP code.⁸⁹ The computational workflow was constructed using the Atomic Simulation Recipes (ASR)⁹⁰ and executed using the My Queue⁹¹ task scheduler frontend.

Conclusions

We have demonstrated that Cu alloying of MAPbBr₃ thin films leads to an increase in the grain boundary area with the formation of Cu-rich secondary phases. The formation of hexagonal perovskite secondary phases is proposed to be more likely in the Cu-alloyed materials than in the non-alloyed materials as supported by calculated convex hull energies. Bandgap decreases slightly with Cu alloying until 40% concentration. A substantially reduced bandgap is observed in 50% Cu-substituted thin film which can be attributed to the presence of Cu-rich secondary phase. Intriguingly, the activation energy of the 40%Cu-substituted perovskite is considerably lower than that of the pristine material, and the conductivity is significantly higher (10³×). In addition, photocurrent carriers are generated in the NIR region of the alloyed material. Based on point defect simulation results, we propose that a compensating Br vacancy will form with alloying and the associated defect complex may induce defect states inside the pristine bandgap, which may explain the experimentally observed optical signature.

Author contributions

Amr Elattar conceived the manuscript's concept and designed the experiments. Amr Elattar, Kodai Nakao, and Kosei Tsutsumi synthesized the materials and conducted structural, electrical, and optical characterizations. Theoretical part was performed and wrote by Jiban Kangsabanik. The first manuscript was written by Amr Elattar and. At various phases, the writers discussed the findings and provided feedback on the manuscript. Kristian S. Thygesen led the assessment of theoretical part. Yasuhiko Hayashi coordinated the entire project.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

A. Elattar acknowledges the financial support from the Ministry of Higher Education, Egypt for his PhD scholarship (EJEP). We would like to thank Enago (https://www.enago.com/) for their English language editing service.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d2qm00491g

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