Minimized Optical/Electrical Energy Loss for 25.1% Monolithic Perovskite/Organic Tandem Solar Cells

doi:10.21203/rs.3.rs-5175952/v1

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Minimized Optical/Electrical Energy Loss for 25.1% Monolithic Perovskite/Organic Tandem Solar Cells

https://doi.org/10.21203/rs.3.rs-5175952/v1

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Perovskite/organic tandem solar cells (PO-TSCs) exploit the advantages of cost-effective fabrication, orthogonal solvent processing for perovskite and organic absorber layers, and compatibility with low-temperature, high-throughput deposition techniques. However, their performance remains hampered by energy losses of subcells and interconnecting layers (ICLs). Here, an energy loss management strategy for PO-TSCs is proposed, focusing on the simultaneous regulation of defect states in perovskite front subcells and the reduction of optical and electrical losses in the ICL. The synergistic effect of hydrogen bonding and coordination interactions between the pyridinium bromide perbromide and perovskite layer effectively mitigates ion migration, thereby minimizing energy losses. The optimized V₂O₅-based ICL structure not only demonstrates excellent transmissivity for near-infrared photons but also allows for barrier-free extraction of charge carriers. This can provide a low-loss interface, facilitating light management within the bulk heterojunction, which effectively balances the current between the front and rear subcells. Taken together, the resulting PO-TSCs deliver a power conversion efficiency of 25.1% with a high open-circuit voltage of 2.10 V.

Physical sciences/Energy science and technology/Renewable energy/Solar energy/Photovoltaics/Solar cells

Physical sciences/Materials science/Materials for energy and catalysis/Solar cells

The pursuit of superior power conversion efficiency (PCE) has spurred the development of tandem solar cells (TSCs), which hold the potential to surpass the Shockley–Queisser (S–Q) limit inherent in single-junction solar cells^{1, 2}. Perovskite/organic tandem solar cells (PO-TSCs) typically consist of three components, including a wide bandgap (WBG) front perovskite subcell that absorbs high-energy photons, a narrow bandgap rear organic subcell that absorbs transmitted low-energy photons, and an optically transparent interconnecting layer (ICL) that establishes electrical contact between the two subcells. Monolithic PO-TSCs have garnered significant interest due to their unique advantages, including low-cost fabrication, the use of orthogonal solvents for the perovskite and organic absorbers, and compatibility with low-temperature, high-throughput manufacturing processes like roll-to-roll and blade coating^{3, 4}. However, despite these advantages, the practical implementation of PO-TSCs has been hindered by significant challenges, particularly related to energy losses in the WBG perovskite layers and optoelectronic losses in the ICL^{5, 6}.

Energy losses in WBG perovskites are a critical barrier to achieving high-efficiency PO-TSCs. WBG perovskite layers, essential for capturing high-energy photons, are susceptible to phase segregation and defect formation, leading to non-radiative recombination^7–9. This recombination results in significant energy losses, thereby reducing the overall efficiency of the PO-TSCs. To mitigate energy losses, various strategies have been conducted, primarily targeting the reduction of non-radiative recombination losses, crystallographic inhomogeneity, or strain within the bulk perovskite phase^{10, 11}. Enhancing the bulk properties of a perovskite subcell often involves controlling the crystallization process or adjusting the perovskite composition. Methodologies such as introducing crystallization agents, designing charge transporting layers with respectable stability and energy level alignment, using non-stoichiometric solutions, tuning the anions fraction on the X-site and cations fraction on the A-site, or incorporating hybrid 2D–3D perovskite heterojunctions at the top surface of 3D perovskites structures have all proven effective in reducing energy losses^12–15.

To electrically connect the front and rear subcells in series within PO-TSCs, an ICL is necessary, where a recombination layer joins the electron transporting layer (ETL) to the hole transporting layer (HTL). An effective ICL must enable efficient charge transfer between subcells, ensure optical transparency, and provide mechanical and chemical protection for the bottom subcell. Nevertheless, the designs of ICLs often suffer from optical and electrical losses, leading to unbalanced short-circuit current density (J_SC) between the front and bottom subcells. Early work typically used ultrathin metal layers (thermally evaporated Ag or Au) to render the ohmic contact^{16, 17}. However, even with an ultrathin metal layer with an equivalent thickness in the range of 0.5–1 nm, significant optical losses can still occur. The transparent metal oxides, such as InO_x and SnO_x, are also a great option for recombination layers¹⁸. Nonetheless, InO_x and SnO_x are prepared using atomic layer deposition (ALD) techniques that demand high energy. It still remains a significant challenge to construct interlayer structures with low photoelectric loss through low-energy consumption methods to meet the demands of burgeoning PO-TSCs.

Here, we fabricated high-efficiency PO-TSCs by minimizing the optical and electric energy losses in subcells and ICL layer. Pyridinium bromide perbromide (PBP) was introduced to the WBG single-junction perovskite solar cells (PerSCs). By leveraging the synergistic enhancement effects of hydrogen bonding and coordination between PBP and perovskites, halide ion migration was effectively restricted, and defect density was reduced. The resulting 1.79 eV WBG single-junction PerSCs exhibited a reduced energy loss of 0.46 eV, achieving a maximum PCE of 19.08% with an open-circuit voltage (V_OC) of 1.33 V. Furthermore, a high-performance ICL structure was developed. The V₂O₅-based ICL demonstrated excellent NIR transmittance, good conductivity, and appropriate energy level alignment, minimizing charge accumulation in the ICL layer and balancing the current between the front and bottom subcells. These combined effects enabled the PO-TSCs to achieve a V_OC of 2.10 V, a fill factor (FF) of 81.1%, a J_SC of 14.68 mA cm^− 2, and a PCE of 25.1%.

Reducing energy loss in WBG perovskite subcells

PBP consists of an anionic part where the central bromine atom is bonded to two adjacent bromine atoms and a cationic part from pyridinium salts (Fig. 1a). It can occupy a halide site or equivalently fill a halide vacancy, coordinating with the perovskite. As illustrated in Fig. 1a, the pure solution of PBP, the mixed solution of PBP and PbI₂, and the mixed solution of PBP and CsI exhibited the color of orange, red-brown, and brown, respectively. These noticeable color differences indicate strong interactions between PBP and the perovskite precursors. Proton nuclear magnetic resonance (¹H NMR) spectroscopy was employed to investigate the existing interaction. For clarity in analysis, we labeled the hydrogen atoms at different positions of the PBP as H_a, H_b, and H_c, respectively. As depicted in Fig. 1b, compared to the pure PBP solution, all proton signals shift upfield in the PBP and PbI₂ mixed solution. However, it remains unclear whether this shift is due to the formation of N − H···I hydrogen bonds, coordination between N − H and Pb²⁺ or a combination of both effects. To elucidate this, we compared the ¹H NMR differences between the PBP and CsI mixed solution and the PBP solution. Only the peaks of H_a and H_b shift upfield, while the peak of H_c remains almost unchanged. Since PBP and CsI interact mainly through hydrogen bonding, the upfield shift of H_a and H_b is attributed to the formation of N − H···I hydrogen bonds^{19, 20}. Therefore, the upfield shift of H_c in the mixed solution of PBP and PbI₂ is solely due to the deshielding effect from the coordination of N − H with Pb^2 + 21. Moreover, combining the effects of hydrogen bonding and coordination, the shift of H_a and H_b peaks in the mixed solution of PBP and PbI₂ is significantly greater than that in the mixed solution of PBP and CsI. These results reveal the existence of N − H···I hydrogen bonds and coordination between PBP and perovskite precursors.

Further confirmation of the interactions between PBP and perovskites was obtained through X-ray photoelectron spectroscopy (XPS). The Pb 4f XPS spectra (Fig. 1c) show that the Pb 4f_5/2 and 4f_7/2 peaks of the perovskite surface shift from 143.04 and 138.22 eV (untreated) to 142.85 and 138.00 eV (PBP-treated), suggesting alterations in the electron cloud density surrounding Pb. This change could be attributed to electron sharing between the − NH − group in PBP and Pb²⁺ in the perovskite²². Similarly, as shown in the I 3d XPS spectrum in Fig. 1d, we find that the I 3d_3/2 and 3d_5/2 peaks of the untreated perovskite surface are located at 630.44 and 618.97 eV, respectively. After PBP optimization, these peaks shift to 630.09 and 618.63 eV, indicating the formation of N − H···I hydrogen bonds between PBP and I⁻ in the perovskite²³. In contrast, the Cs 3d XPS spectra (Supplementary Fig. 1) exhibit no notable shift, indicating the unchanged chemical environment around Cs, which is consistent with results obtained from the ¹H NMR analysis.

As for the uniformity and defect distribution details of the untreated and PBP-treated perovskite films, photoluminescence (PL) mapping displays the corresponding information (Fig. 1e). From the statistical histograms depicting the spatial distribution of luminescence intensity (Supplementary Fig. 2), it can be seen that the PL intensity distribution in the untreated perovskite film is more dispersed with poor distribution uniformity, indicating a low and uneven luminescence intensity that suggests the presence of numerous defects or non-uniform regions²⁴. In contrast, the PL intensity distribution in the PBP-treated perovskite film is more concentrated, reflecting improved uniformity and higher film quality with fewer defects.

The influence of PBP on perovskite morphology was verified through scanning electron microscopy (SEM) and atomic force microscopy (AFM) measurements. As shown in the SEM images (Fig. 1f), it is evident that while both untreated and PBP-treated perovskite films exhibit distinct perovskite grains and grain boundaries, the grain size of untreated perovskite films is uneven, with sizes primarily ranging from 350 to 500 nm, showing a broad distribution. In contrast, the grain size of the PBP-treated perovskite is concentrated within a narrower range of 450 to 500 nm, indicating a more uniform internal structure of the film. This helps reduce grain boundary defects caused by size differences, facilitates smoother carrier transport between grains and reduces recombination losses at the grain boundaries. Figure 1g illustrates the AFM images of untreated and PBP-treated perovskite films, where the perovskite film after PBP treatment possessed slightly larger grain domain sizes, leading to the surface roughness increase from 13 to 18 nm, which is conducive to forming better physical contact and reduce defects.

To exclude the defect evolution of WBG perovskite, PL spectra under continuous illumination were performed (Fig. 2a). During the initial 10 min of illumination, the PL intensity of the pristine perovskite film exhibits a rapid decline, indicating an increase in recombination centers on the perovskite surface. Following this period, the PL intensity continues to decrease, albeit at a slower rate, implying a gradual accumulation of defect states within the perovskite film. In contrast, the perovskite film with PBP treatment shows no significant reduction in PL intensity throughout the entire duration of illumination.

The trap state density (n_t) and carrier mobility (µ) can be obtained by fitting the space-charge-limited current (SCLC) model based on electron-only and hole-only devices. The trap-filled limit voltage (V_TFL) provides information about the n_t, calculated using the formula²⁵:

V _TFL = \(\:\frac{\text{e}{L}^{2}{n}_{t}}{2{\epsilon\:}{{\epsilon\:}}_{0}}\) (1)

where ε is the relative permittivity of the perovskite material, ε₀ is the vacuum permittivity, e is the elementary charge, and L is the thickness of the perovskite active layer. For the electron-only devices (Supplementary Fig. 3a), the V_TFL values for untreated and PBP-treated perovskite devices are 0.25 and 0.16 V, respectively. The electron trap state densities are calculated to be 5.2 × 10¹⁵ and 3.3×10¹⁵ cm^− 3, showing a reduction of 1.9 × 10¹⁵ cm^− 3 after PBP treatment. For the hole-only devices (Supplementary Fig. 3b), the V_TFL values for untreated and PBP-treated perovskite devices are 0.64 and 0.26 V, respectively. The corresponding hole trap state densities are 1.3 × 10¹⁶ and 5.4 × 10¹⁵ cm^− 3, representing a decrease of 7.6 × 10¹⁵ cm^− 3 after PBP modification. The significant reduced n_t of the PBP-treated perovskite devices indicates that the introduction of PBP can effectively improve the crystalline quality and mitigate defects in the perovskite films (Fig. 2b). Next, the µ was calculated by using the Mott-Gurney formula:

µ = \(\:\frac{8}{9}\)J_Dεε₀\(\:\frac{{L}^{3}}{{V}^{2}}\) (2)

where J_D is the current density, and V is the applied voltage. For the single electron devices (Supplementary Fig. 4a), the electron mobility increased from 1.17 × 10^− 3 cm² V^− 1 s^− 1 for untreated perovskite devices to 1.70 × 10^− 3 cm² V^− 1 s^− 1 for PBP-treated perovskite devices, an enhancement of 0.53 × 10^− 3 cm² V^− 1 s^− 1. For the hole-only devices (Supplementary Fig. 4b), the hole mobility of the PBP-treated perovskite film is 1.65 × 10^− 3 cm² V^− 1 s^− 1, much higher than that of the untreated film (9.14 × 10^− 4 cm² V^− 1 s^− 1), indicating faster carrier transport in PBP-treated devices. Additionally, the hole-to-electron mobility ratio in untreated and PBP-treated devices is 0.78 and 0.97, respectively, suggesting more balanced carrier transport in PBP-treated perovskites. This balance minimizes recombination and asymmetric transport losses in the devices.

Time-resolved photoluminescence (TRPL) spectroscopy was used to analyze the carrier dynamics and recombination kinetics in perovskite films. As depicted in Supplementary Fig. 5, a double exponential decay model was used to fit the TRPL decay curves, obtaining carrier lifetime parameters. Detailed fitting parameters are listed in Supplementary Table 1. For the untreated perovskite film, the fast recombination process τ₁ is 50.0 ns, while for the PBP-treated perovskite film, τ₁ increases to 102.3 ns, suggesting that the shallow traps are reduced and carriers can exist for a longer duration without undergoing fast recombination in the PBP-treated film. Additionally, Compared to the untreated film, which has a slow recombination process τ₂ of 204.3 ns, the PBP-treated film exhibits a significantly higher τ₂ of 563.0 ns, indicating a reduction in deep defects and more efficient carrier transport. Furthermore, the average decay lifetime of the untreated perovskite devices is 187.3 ns, increasing to 550.2 ns for the PBP-treated devices, which implies a reduced defect density and suppressed non-radiative recombination²⁶. This indicates that the introduction of PBP into perovskite films effectively improves charge separation efficiency and minimizes charge recombination losses.

Sequentially, electrochemical impedance spectroscopy (EIS) was performed to study the impact of PBP treatment on carrier transport and recombination mechanisms in PerSCs. Supplementary Fig. 6 shows the Nyquist plots obtained from the EIS measurements of untreated and PBP-treated PerSCs under dark conditions and zero bias. The equivalent circuit model used for fitting is provided as the inset of Supplementary Fig. 6. The Nyquist plots include a semicircle in the high-frequency region and an arc in the mid-frequency region. The semicircle in the high-frequency region reflects charge transfer resistance (R_ct), while the arc in the mid-frequency region reflects recombination resistance (R_rec). The fitting parameters are listed in Supplementary Table 2. With the reduction in R_ct from 9.6 × 10⁴ Ω in untreated perovskite solar cells to 6.3 × 10⁴ Ω after PBP treatment, the interfacial charge transport efficiency is improved, enabling carriers to traverse the interface more quickly and effectively. The increase in R_rec from 0.76 × 10⁶ Ω in untreated PerSCs to 1.12 × 10⁶ Ω in PBP-treated PerSCs indicates a more efficient charge separation process, resulting in reduced charge recombination²⁷. Supplementary Fig. 7 displays the current density-voltage (J-V) characteristics of PerSCs tested under dark conditions. The dark current density of PBP-treated PerSCs is significantly lower than that of untreated PerSCs, suggesting reduced leakage current paths and fewer recombination centers²⁸. Subsequently, to investigate the carrier dynamics in devices, we measured the J-V characteristics under varying light intensities from 1 to 100 mW/cm². The relationship between light intensity and V_OC can be used to assess the recombination mechanisms in PerSCs. Theoretically, the slope of the curve corresponds to the ideality factor (n), which can be calculated using the following formula²⁹:

n = \(\:\frac{e}{{K}_{B}T}\frac{d{V}_{OC}}{dlnI}\) (3)

where K_B is the Boltzmann constant, T is the absolute temperature, I is the light intensity. When n approaches 1, it indicates that bimolecular radiative recombination is the dominant recombination mechanism. When n approaches 2, it implies that Shockley-Read-Hall (SRH) recombination and non-radiative recombination caused by trap states are dominant. As shown in Supplementary Fig. 8a, the value of n decreases from 1.74 in untreated PerSCs to 1.61 in PBP-treated PerSCs, reflecting a significant suppression of charge recombination and a reduction in defect state density. Next, we investigated the trend of J_SC with varying light intensities to further verify the carrier recombination mechanisms. Generally, the relationship between J_SC and light intensity can be described by the following formula: J_SC ∝ \(\:{P}_{light}^{{\alpha\:}}\), where α is an exponent related to the recombination mechanism. An α value less than 1 indicates the presence of defect-assisted non-radiative recombination. As illustrated in Supplementary Fig. 8b, the α values for untreated and PBP-treated perovskite solar cells are 0.97 and 0.99, respectively, indicating that both devices exhibit defect state-induced SRH recombination.

On the basis of the above-mentioned results, we fabricated inverted WBG perovskite devices to study the impact of PBP treatment on device performance. The optimal concentration of PBP was found to be 0.5 mg/mL (Supplementary Fig. 9 and Supplementary Table 3). Figure 2c describes the J-V curves of the best-performing untreated and PBP-treated PerSCs. The corresponding optimal performance parameters are given in Supplementary Table 4. The untreated PerSC achieves a V_OC of 1.29 V, a J_SC of 17.59 mA cm^− 2, an FF of 77.6%, and a PCE of 17.73%. In comparison, the optimized PBP-treated PerSC exhibits a higher V_OC of 1.33 V, a J_SC of 17.72 mA cm^− 2, and an improved FF of 81.2%, yielding a high PCE of 19.08%. The external quantum efficiency (EQE) spectra of the PerSCs are shown in Fig. 2d. The integrated current density for the untreated PerSC is 16.73 mA cm^− 2, while that for the PBP-treated PerSC is 17.34 mA cm^− 2, which are consistent with the J_SC values obtained from the J-V curve. The bandgap of the target device is determined to be 1.79 eV by differentiating its EQE spectrum, which exhibited an inflection point at 691 nm (Supplementary Fig. 10). Figure 2e makes a clear comparison of energy loss. Energy losses are 0.5 and 0.46 eV for untreated and PBP-treated PerSCs, respectively. Supplementary Fig. 11 presents the statistical photovoltaic parameters of 15 individual PerSCs from different batches. The significant improvement in the photovoltaic parameters of PBP-treated PerSCs can be attributed to the passivation effect of PBP on perovskite defects, which results in more efficient charge extraction and reduced recombination losses. These results demonstrate the reliability of our method in improving the performance of PerSCs. Supplementary Fig. 12 shows the operational stability of the PerSCs. The PBP-treated devices exhibit significantly enhanced stability, retaining 80% of their initial efficiency after 160 hours of continuous LED illumination. Additionally, we also achieved improved thermal stability for the PBP-treated PerSC, which retained 90% of its initial PCE after heating at 65°C for 600 h in the N₂ atmosphere (Supplementary Fig. 13).

Design of ICL with high transmittance and conductivity

In PO-TSCs, the ICLs control charge collection and recombination from subcells and are critical in determining the performance of PO-TSCs. The design of high-quality ICLs aims to avoid or minimize reverse junction formation, parasitic light absorption, and junction resistance. Here, we designed a V₂O₅-based ICL structure with low optical/electrical loss and compared it with the traditional MoO₃-based ICLs. To investigate the differences in optical and electrical properties of the V₂O₅-based and MoO₃-based ICLs, we conducted a series of characterizations, including transmittance, conductivity, and hole mobility. First, we compared the UV-Vis spectra of the MoO₃-based and V₂O₅-based ICLs. As displayed in Fig. 3a, within the wavelength range of 550 to 1100 nm, the transmittance of the MoO₃-based ICLs is significantly lower than that of V₂O₅-based ICLs. The low optical loss of the V₂O₅-based ICLs in the NIR region indicates that the front junction perovskite subcell and the ICL have lower utilization efficiency for the NIR part of the solar spectrum, favoring the transmission of photons in the NIR region. This contributes to the enhancement of current in the organic bottom cells, which promotes current matching between the front and rear subcells in the PO-TSCs.

Regarding conductivity (σ), both MoO₃-based and V₂O₅-based ICLs show a linear relationship between current density and voltage, indicating the formation of good ohmic contacts (Fig. 3b)³⁰. The σ of the ICLs was calculated using the formula³¹: σ = Jd/V, where d represents the thickness of the ICL, J represents the current density, and V represents the voltage. The V₂O₅-based ICLs show a higher σ value of 2.23×10^− 3 mS cm^− 1 compared with MoO₃-based ICLs (1.84×10^− 3 mS cm^− 1). Obviously, the hole mobility of V₂O₅ is also superior to that of MoO₃ (Fig. 3c). These results indicate that V₂O₅ can effectively facilitate rapid and efficient charge transport, which enhances hole selectivity^{32, 33}.

The ultraviolet photoelectron spectroscopy (UPS) measurement of V₂O₅ shows the work function (WF) and the valence band maximum (VBM) of 5.10 eV and 7.87 eV (Supplementary Fig. 14). Then the conduction band minimum (CBM) of the V₂O₅ film is determined to be 4.72 eV based on its bandgap of 3.15 eV (Supplementary Fig. 15). Supplementary Fig. 16 illustrates the energy level diagram. The detailed energy diagrams following Fermi level alignment are depicted in Fig. 3d. A significant band bending occurs at the C60/MoO₃ interface due to the considerable work function mismatch. This bending is expected to induce a large Schottky barrier at the interface, which hampers efficient charge recombination within the ICL. Replacing MoO₃ with V₂O₅ shifts the energy levels upward, substantially lowering the barrier and facilitating barrier-free electron transport from the WBG perovskite front subcells into the ICL, where the electrons efficiently recombine with the holes from the organic rear subcells.

We selected the combination of PBDB-T-2F (PM6), BTP-eC9 and PC₇₁BM as the photoactive layer for the organic subcells^{34, 35}. The detailed chemical structures of the donor and acceptor materials are illustrated in Supplementary Fig. 17. UV-Vis absorption spectroscopy was then used to compare the absorbance and transmittance of MoO₃ and V₂O₅ thin films. As seen in Supplementary Fig. 18, within the absorption range of 600 to 1100 nm, the absorption intensity of the MoO₃ films is higher than that of V₂O₅ films. The higher absorption intensity of MoO₃ films in the NIR region reduces the effectiveness of the organic subcell in capturing low-energy photons, leading to current mismatch in the PO-TSCs. Correspondingly, the transmittance of the MoO₃ films is lower than that of V₂O₅ films. We then established an optical transfer matrix model to analyze the light field distribution within the absorber layers. The device structure used for the simulation is ITO/MoO₃ or V₂O₅/PM6:BTP-eC9/TPMA/Ag. As presented in Fig. 3e, the MoO₃-based and V₂O₅-based OSCs exhibit similar simulated J_SC within the 300–600 nm range owing to the comparable absorption capacity of both for high-energy photons. However, the V₂O₅-based OSCs improve the utilization of low-energy photons in the near-infrared range of 600 to 1000 nm, leading to a higher simulated J_SC, which is beneficial for the current balancing between the front and rear subcells in PO-TSCs.

Transient photovoltage (TPV) and transient photocurrent (TPC) experiments were performed to gain insight into the charge recombination and extraction kinetics process³⁶. As for TPC decay lifetimes (Supplementary Fig. 19), V₂O₅-based OSCs exhibit a shorter extraction time of 1.91 µs compared to 2.13 µs for MoO₃-based OSCs, suggesting enhanced charge transport capacity. Photogenerated carrier lifetimes obtained from TPV measurements are 5.26 and 8.27 µs for MoO₃-based and V₂O₅-based OSCs, respectively, indicating charge recombination has been suppressed in the V₂O₅-based OSCs. The charge generation properties were characterized by measuring their exciton dissociation probabilities (P_diss), which were estimated by surveying the photocurrent densities versus effective voltage curves. As indicated in Supplementary Fig. 20, the calculated P_diss values were 94.47% for MoO₃-based OSCs and 97.13% for V₂O₅-based OSCs, with the higher P_diss partly explaining the improvements in J_SC and FF³⁷.

Utilizing AFM, we studied the surface morphology of the films. The AFM images (Supplementary Fig. 21) reveal stripe-like structures for both MoO₃ and V₂O₅ films, likely caused by anisotropic growth during thermal evaporation³⁸. Additionally, the MoO₃ films show a calculated root mean square roughness (RMS) of 0.50 nm, while V₂O₅ films exhibit an RMS of 0.39 nm. The extremely low roughness and smooth surface morphology of V₂O₅ films facilitate the formation of well-ordered molecular arrangements and crystallization in the absorber layer. Subsequently, absorber layers were fabricated on ITO/V₂O₅ and ITO/MoO₃ films to validate the above conclusions. The smoother surface with a roughness of 2.25 nm is achieved for the V2O5-based absorber layer, while a rougher surface with a roughness of 3.53 nm is observed for the MoO₃-based absorber layer.

According to our previous work, the interface plays an important role in controlling the vertical phase distribution³⁹. We then studied the vertical phase distribution of the absorber layer. The vertical phase distribution is relevant to miscibility between the donor and acceptor. We used the Flory-Huggins interaction parameter (χ) to quantize the miscibility of the donor and acceptor, where larger χ values indicate poorer miscibility^{40, 41}. To calculate χ, we carried out the contact angle tests on water and diiodomethane, and the results are presented in Supplementary Fig. 22. Through the contact angles, we could obtain the surface tensions (γ) for donor and acceptor materials. The γ was calculated using the Owens-Wendt-Rabel-Kaelble (OWRK) method. The relevant calculation results are listed in Supplementary Table 5. The γ of the V₂O₅-based absorber layer (39.5 mN m^− 1) is closer to that of the acceptor on V₂O₅ films (41.3 mN m^− 1), suggesting that the acceptor phase tends to enrich on the surface of the absorber layer. This vertical phase distribution is conducive to more efficient charge transport. Then χ can be calculated using the formula:

χ = (\(\:\sqrt{{{\gamma\:}}_{D}}\)−\(\:\sqrt{{{\gamma\:}}_{A}}\))² (4)

where D represents the donor, and A represents the acceptor. It was found that χ values were 0.2 and 0.3 based on MoO₃ and V₂O₅, respectively. The decreased miscibility may avoid the excessive mixing of donor and acceptor, leading to the formation of the desired domain size and purity.

Based on the above analysis, we investigate the impact of V₂O₅ or MoO₃ on the device performance. The corresponding J-V curves are exhibited in Supplementary Fig. 23, and the photovoltaic parameters are summarized in Supplementary Table 6. Figure 3f displays the J-V curves of the best-performing MoO₃-based and V₂O₅-based OSCs. The MoO₃-based OSCs show a PCE of 17.34% with a V_OC of 0.844 V, a J_SC of 26.75 mA cm^− 2, and an FF of 76.8%. In comparison, the V₂O₅-based OSC exhibits a V_OC of 0.859 V, a J_SC of 27.29 mA cm^− 2, an FF of 78.4%, and a PCE of 18.39%. Supplementary Fig. 24 presents the statistical photovoltaic parameters of 15 individual OSCs from different batches, showing significant enhancement in V₂O₅-based OSCs compared to MoO₃-based OSCs. Moreover, we observed an enhanced EQE in V₂O₅-based OSCs from 300 to 850 nm (Fig. 3g). The integrated current density derived from the EQE curves increased from 25.67 mA cm^− 2 for the MoO₃-based device to 26.31 mA cm^− 2 for the V₂O₅-based device.

Photovoltaic performance of Monolithic PO-TSCs

Encouraged by the reduced energy loss of subcells and minimized optical/electrical loss of ICL, we fabricated PO-TSCs with the architecture of ITO/NiO_x/4PADCB/perovskite/EDADI/C60/PEI/ITO/V₂O₅ or MoO₃/PM6:BTP-eC9:PC₇₁BM/TPMA/Ag (Fig. 4a). We compared the performance of PO-TSCs prepared with different thicknesses of ICLs. The J-V curves of the PO-TSCs are depicted in Supplementary Fig. 25. Supplementary Table 7 summarizes the corresponding photovoltaics parameters. The control PO-TSCs yield a PCE of 23.20%, with a V_OC of 2.07 V, a J_SC of 14.41 mA cm^–2 and an FF of 77.7%. In contrast, the optimized PO-TSCs exhibited an FF of 81.1%, a J_SC of 14.68 mA/cm², a V_OC of 2.10 V, and a PCE of 25.07% (Fig. 4b). As shown in Fig. 4c, the optimized PO-TSCs show good current matching, where the J_SC values of 14.11 and 13.98 mA cm^− 2 are integrated from the EQE spectra of the perovskite and organic subcells, respectively. For the control PO-TSCs, the integrated current densities are 13.97 and 13.56 mA cm^− 2, respectively. Additionally, the steady-state power output of the optimized PO-TSCs was 25.02% at a bias of 1.77 V (Fig. 4d). The statistical photovoltaic parameters of 30 individual PO-TSCs from different batches confirmed the advantages of the PO-TSCs with V₂O₅-based ICL (Fig. 4e and Supplementary Fig. 26). The significant enhancement in the performance of PO-TSCs can be attributed to the reduced energy loss of perovskite subcells and the low electrical and optical loss of V₂O₅-based ICL. Figure 4f depicts the light field distribution simulated by the optical transfer matrix model. The simulated J_SC values of the PO-TSCs with V₂O₅-based ICL are higher than those with MoO₃-based ICL, which could be due to the higher light field intensity in the 650 to 850 nm wavelength range for the PO-TSCs with V₂O₅-based ICL.

Subsequently, we evaluated the stability of the PO-TSCs. The thermal stability of the unencapsulated PO-TSCs in a nitrogen atmosphere was verified. After 200 hours of storage in a nitrogen atmosphere at 65°C fixed temperature (Fig. 4g), the PCE of the optimized PO-TSCs decreased by 9%, while the PCE of the control PO-TSCs decreased by 31%. As depicted in Fig. 4h, we also assessed the operational stability of the PO-TSCs. The stability of the TSCs under continuous operation is mainly determined by that of the perovskite subcell. The optimized PO-TSCs still maintained 82% of their initial efficiency after 460 hours of continuous LED light illumination, while the efficiency of the control PO-TSCs decreased to 80% of their original value after 180 hours.

In summary, efficient PO-TSCs with a high PCE of 25.1% has been obtained via minimizing the optical and electrical energy losses simultanously. By passivating the defects in WBG perovskites, we have significantly reduced the E_loss, boosting the V_OC of perovskite subcells. We also demonstrated a designed ICL structure with low optical and electrical loss. Compared with the MoO₃-based ICL, the V₂O₅-based ICL exhibits higher transmittance in the NIR region. In addition, the V₂O₅-based ICL has better electrical conductivity, higher carrier mobility, and more compatible energy level alignment, which helps improve charge transport and extraction, reduce charge recombination, and minimize charge accumulation at the ICL, thereby balancing the current between the perovskite and organic subcell. These results emphasize the importance of simultaneously minimizing energy losses in subcells along with optical and electrical losses in ICLs, which is crucial for constructing high-performance PO-TSCs.

Materials

All of the materials were used without further purification. Caesium iodide (CsI), formamidinium iodide (FAI), methylammonium iodide (MAI), lead bromide (PbBr₂), EDADI, 4PADCB, fullerene (C60), were purchased from Xi'an Yuri Solar Co., Ltd. Lead iodide (PbI₂), NiO_x, and BCP were purchased from Advanced Election Technology. PM6, BTP-eC9 and PC₇₁BM were purchased from Solarmer Materials. PEI was purchased from Beijing Innochem Science & Technology Co., Ltd. MoO₃, V₂O₅, and Ag were purchased from ZhongNuo Advanced Material (Beijing) Technology Co., Ltd. TPMA were obtained from the previous literature. Dimethylformamide (DMF), dimethyl sulfoxide (DMSO), chlorobenzene (CB), chloroform (CF), and isopropanol (IPA) were purchased from Acros Organics. PBP was purchased from Alfa.

Perovskite subcell fabrication

The ITO glass substrates were sequentially cleaned by sonication with detergent, deionized water and ethanol for 15 min, respectively. The cleaned ITO glass substrates were then transferred into an oven at 100°C for 10 min and treated with ultraviolet ozone for 20 min before use. The NiO_x (10 mg ml^− 1) dispersed in deionized water was deposited on the ITO glass substrates by spin-coating at 2,000 rpm for 45 s and then annealed at 100°C for 15 min in air. Then, the samples were transferred into a nitrogen glovebox for deposition of the 4PADCB and perovskite layers. 4PADCB (0.5 mg ml^− 1) dissolved in IPA was deposited on the NiO_x substrate by spin-coating at 4,000 rpm for 40 s, followed by annealing at 100°C on a hot plate for 10 min. The precursor solution was prepared by mixing 0.1 M CsI, 0.8 M FAI, 0.1 M MAI, 0.6 M PbBr₂ and 0.9 M PbI₂ in DMF/DMSO (4:1) with 1 M solution. PBP with different concentrations were incorporated into the precursors for the target sample. The precursor solution was spin-coated onto the NiOx/4PADCB layer in the first step at 1000 rpm for 10 s and the second step at 4000 rpm for 40 s. During the spin-coating process, 150 µL CB antisolvent was dropped at 25 s of the second step before the end of the spin-coating program. Then, the films were annealed at 100°C for 10 min. After cooling to room temperature, 50 µL EDADI solution (0.5 mg ml^–1) in IPA was dripped on the perovskite surface and spin-coated at 5000 r.p.m. for 30 s followed by annealing at 100°C for 10 min. Finally, all the samples were transferred into a high-vacuum chamber to evaporate 25 nm C₆₀, 6 nm BCP and 100 nm Ag electrode.

Organic subcell fabrication

MoO₃ and V₂O₅ were thermally evaporated as hole transporting layers on the ITO glass substrates. The PM6:BTP-eC9:PC₇₁BM was dripped on the MoO₃ and V₂O₅ hole transporting layer at 4000 r.p.m. for 30 s, followed by annealing at 80°C for 10 min. After cooling, 50 µL TPMA solution (0.5 mg ml^–1) in methanol was spin-coated onto the absorber layer at 4000 r.p.m. for 30 s. Finally, all the samples were transferred into a high-vacuum chamber to evaporate 100 nm Ag electrode.

PO-TSC fabrication

After thermal evaporation of C₆₀ in wide-bandgap subcells, 50 µL PEI solution (1 mg ml^–1) in IPA was spin-coated onto the C₆₀ layer at 5000 r.p.m. for 30 s. Then ITO recombination layer was sputtered through a shadow mask by magnetron sputtering with a radiofrequency power of 100 W. Subsequently, MoO₃ or V₂O₅ was thermally evaporated on top of the ITO, to form an ICL structure for the TSCs. Next, the PM6:BTP-eC9:PC₇₁BM active layer was spin-coated on MoO₃ or V₂O₅. TPMA was spin-coated on the organic active layer at 4000 r.p.m. for 30 s. Finally, all the samples were transferred into a high-vacuum chamber to evaporate 100 nm Ag electrode.

Characterization of the solar cells

J − V characteristics of the devices were measured with a Keithley B2901A source meter under the simulated AM 1.5 G illumination (100 mW·cm^− 2). The devices were measured both in reserve and forward scan with a step of 0.02 V. Light intensity was calibrated by a standard silicon cell (SRC-2020, EnliTech). The EQE spectra were performed using the Enlitech QE-R system. For TSCs, when measuring the EQE values of the perovskite subcells, a bias illumination with an 800 nm optical filter was applied to saturate the corresponding organic subcells. Similarly, when measuring the EQE values of the organic subcells, a bias illumination with a 500 nm optical filter was applied to saturate the corresponding perovskite subcells. No bias voltage was applied during the EQE measurements of either subcell. The dark I–V characteristics curves were recorded using a Keithley B2901A source meter. The steady-state power output measurements were recorded by tracing the current density at a bias (voltage at maximum power point).

Characterization of the films

Absorption and transmittance spectra were measured using a Shimadzu UV-Vis-NIR 3600 spectrometer. PL and TRPL were measured by FLS980 Photoluminescence Spectrophotometer (Edinburgh Instruments Ltd.). Krüss DSA30 optical contact angle meter. XPS measurements were performed on a THERMO VG ESCALAB 250. UPS measurements were performed on an AXIS ULTRA DLD, Kratos (by using an He I source) under high vacuum (10^− 8 Torr). The top-view and cross-sectional morphology were analyzed by a SEM (SU8600, HITACHI). The AFM characterizations were conducted by Bruker DMFASTSCAN2-SYS. Ellipsometry was carried out by an ellipsometer (ES01, ELLITOP SCIENTIFIC CO., LTD.). The EIS was measured by the CHI660E electrochemical workstation.

Other characterizations

¹H NMR spectra were measured on Bruker AVANCE ΙΙΙ (400 MHz) spectrometers.

Characterization of the device stability

Operational stability tests were performed on the encapsulated solar cells. The tests were conducted using a white-light LED array under a continuous simulation intensity of 100 mW cm^− 2 (Suzhou D&R instruments, PVLT-G8001M-32B). For thermal stability measurements, the devices without encapsulation were kept at a fixed temperature in the dark except during the J–V measurement, and the J–V curves were recorded at specific time intervals.

Data availability

All data supporting the results of this study are provided in this article and its supplementary information. Any other information can be requested from the corresponding author.

Acknowledgments

This work was financially supported by the National Natural Science Foundation of China (52373169 and 22379011).

Author information

Authors and Affiliations

Beijing Advanced Innovation Center for Soft Matter Science and Engineering, Beijing University of Chemical Technology, Beijing, 100029, China

Zhangwei He, Runnan Yu, Yiman Dong, Ruyue Wang, Yuling Zhang, Zhan’ao Tan

Contributions

Z.T. directed and supervised the project. Z.H. proposed the idea and designed the experiment. Z.H., Y.D., R.W. and Y.Z. fabricated and characterized the solar cells. Z.H. conducted data analysis. The manuscript was prepared by Z.H. The manuscript was revised and finalized by R.Y. and Z.T. All authors participated in the discussion and commented on the manuscript.

Corresponding authors

Correspondence to Runnan Yu and Zhan’ao Tan.

Ethics declarations

Competing interests

The authors declare no competing interests.

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Minimized Optical/Electrical Energy Loss for 25.1% Monolithic Perovskite/Organic Tandem Solar Cells

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Abstract

Figures

Introduction

Results

Reducing energy loss in WBG perovskite subcells

Design of ICL with high transmittance and conductivity

Photovoltaic performance of Monolithic PO-TSCs

Conclusions

Methods

Materials

Perovskite subcell fabrication

Organic subcell fabrication

PO-TSC fabrication

Characterization of the solar cells

Characterization of the films

Other characterizations

Characterization of the device stability

Declarations

References

Additional Declarations

Supplementary Files

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