Furthermore, we tested the stability of the devices without any encapsulation and found a 10.9% decrease after 1,394 h of storage in the dark with air at RT ( SI Appendix, Fig. S6A ). The observed stability is better than that of the state-of-the-art P3HT-based cells with an initial PCE of 7.7%, which show a T 80 lifetime of 1,200 h under the same storage condition ( 25 ). We also examined the thermal stability of our device in air without encapsulation. PCE retentions of 99.6% and 94.9% were observed for devices heated at 120 °C for 1 and 2 h, respectively ( SI Appendix, Fig. S6B ). Such good stability in ambient air is advantageous for the processing of OPVs by many other techniques, such as all-solution processing of OPVs ( 31 ) and integration with other devices on the same substrate ( 1 ). In sharp contrast to the good stability of the PBDTTT-OFT/IEICO-4F OPVs, faster degradation is observed for both PBDTTT-OFT/PC 71 BM blend and PBDTTT-EFT(PCE10)/IEICO-4F blend ( SI Appendix, Figs. S7 and S8 ). Therefore, we conclude that both stable donor polymers and nonfullerene acceptors are essential for achieving good environmental stability.

We then tested the environmental stability of the glass-encapsulated OPVs after postannealing at 150 °C for 5 min. When stored in the dark with air at RT the devices did not degrade within 2,000 h and only lost 4.8% of the initial efficiency after 4,736 h ( Fig. 1D ) with an estimated T 80 lifetime (80% of the initial PCE) of over 20,750 h ( SI Appendix, Fig. S3 ), which will potentially lead to an ultralong shelf lifetime of more than 11.4 y. Furthermore, devices stored in an environmental test chamber at 85 °C and 30% RH showed a T 80 lifetime of 1,050 h ( Fig. 1E ). The degradation mainly arises from the reduced FF ( SI Appendix, Fig. S4 ). Although it is difficult to compare the stability of our device with those reported previously for other devices owing to the variations in the test conditions, the stability of our OPVs is outstanding for a device with an initial PCE of >12% ( SI Appendix, Table S1 and Fig. S5 ). In contrast, under the same test conditions, in the absence of postannealing treatment at 150 °C, our device lost 31.1% of the initial efficiency after 169 h ( Fig. 1E ). These results clearly demonstrate that the thermal stability of our OPVs is improved significantly after the postannealing treatment at 150 °C, and the related mechanism will be discussed later.

Characterization of nonfullerene OPVs with high efficiency and good environmental stability. (A) Stack structure of the OPV. (B) Comparison of current density–voltage curves and (C) EQE–light wavelength curves of OPVs before and after the single-step postannealing treatment at 90 and 150 °C for 5 min. (D) Evolution of PCE with time when the devices were stored in the dark with air at RT. (E) Comparison of the thermal stability of devices before and after the postannealing treatment at 150 °C before the stability test. The devices were stored in a commercial environment box at 85 °C and 30% RH.

Fig. 1A shows the OPV structure with rigid glass encapsulation. In this device, a D/A blend of PBDTTT-OFT/IEICO-4F ( 8 , 30 ) ( SI Appendix, Fig. S1 ) and MoO x are utilized as the active layer and HTL, respectively. After the device fabrication process, we annealed the devices in an N 2 -filled glove box for only 5 min. Interestingly, the OPVs exhibit an unusual change in their performance after annealing at different temperatures ( Fig. 1B ). Before annealing, the devices have an average PCE of 13.0% ( Table 1 ), short-circuit current density (J SC ) of 25.5 mA/cm 2 , open circuit voltage (V OC ) of 0.698 V, and fill factor (FF) of 73.4%. Although the J SC decreased to 18.2 mA/cm 2 after annealing at 90 °C for 5 min in a nitrogen atmosphere, it increased to 26.6 mA/cm 2 after annealing at 150 °C for 5 min ( Table 1 ). This change in J SC was confirmed by both external and internal quantum efficiency (EQE and IQE, respectively) measurements ( Fig. 1C and SI Appendix, Fig. S2 ). The quantum efficiencies, EQE and IQE, changed uniformly in the whole range of the responsive wavelengths after annealing. The current densities calculated from EQE measurements are 23.2, 16.3, and 24.1 mA/cm 2 for the as-fabricated device and devices annealed at 90 and 150 °C for 5 min, respectively, with ∼9% difference in J SC in J–V curves for all conditions. In sharp contrast to J SC , both V OC and FF remained constant after such postannealing treatments. Although the device performance decreased dramatically after annealing at 90 °C, its PCE could be recovered to 13.0% again after annealing at 150 °C.

Mechanism of the Postannealing Effect on OPVs’ Performance.

To investigate the unique effects of the postannealing treatment on the device performance we heated the same device successively at different temperatures between 50 and 200 °C under nitrogen. Since we annealed the same device at each temperature for 5 min, this process is only for investigation of the abnormal performance. As shown in Fig. 2 A–D, the OPVs degraded rapidly upon heating at temperatures between 50 and 90 °C, mainly due to the decrease in J SC (SI Appendix, Table S2). Surprisingly, the device performance improved dramatically when heated between 90 and 140 °C, and the J SC increased relative to the initial value. Owing to the slightly decreased FF, the PCE only recovered to ∼90% of the initial value. The same device degraded slowly when heated at temperatures higher than 150 °C. Next, different devices were heated at a given temperature for different durations up to 23 h (SI Appendix, Fig. S10). Both the devices annealed at 85 and 130 °C first showed a decrease in performance, followed by an increase. The performance recovered to 84.1 and 96.8% of the initial value after annealing at 85 °C for 23 h and at 130 °C for 30 min, respectively. These results indicate that the device performance changes significantly with the postannealing condition.

Fig. 2. Fabrication of nonfullerene OPVs via effective postannealing treatment. (A–D) Evolution of normalized PCE, J SC , FF, and V OC with postannealing temperature. Note that the same device was heated from 50 to 200 °C; it was heated at each temperature for 5 min in a nitrogen atmosphere. (E) Comparison of the series resistance and shunt resistance during the postannealing treatment. (F) Comparison of the short circuit current as a function of light intensity during the postannealing treatment. (G) Comparison of transient photocurrent decay curves acquired at 0 V during the postannealing treatment.

To unravel the mechanism of such variation under the three conditions shown in Fig. 1, we considered the differences in the series resistance (R s ) and shunt resistance (R sh ), charge recombination, and charge extraction of the devices (Fig. 2 E–G). We extracted R s and R sh from the J–V curves (32). The initial R s of the device (1.7 Ω∙cm2) increased to 3.5 Ω∙cm2 after annealing at 90 °C and then decreased to 2.8 Ω∙cm2 after annealing at 150 °C, while the initial R sh (822 Ω∙cm2) decreased to 300 Ω∙cm2 after annealing at 90 °C and then increased to 571 Ω∙cm2 after annealing at 150 °C (Fig. 2E). Such variation in resistance is consistent with the variation in the device performance. We investigated the charge-recombination behaviors qualitatively by evaluating the dependence of J SC with light intensity (Fig. 2F). Under all three conditions, the devices showed a linear power-law dependence of J SC with light intensity, I ( J SC ∝ I α ) (33), with α values of 0.998, 0.987, and 0.996 for the untreated device and devices annealed at 90 and 150 °C, respectively. The α values of the untreated device and that annealed at 150 °C are closer to 1, suggesting negligible bimolecular recombination at the short circuit condition in these devices (33). The differences in charge-extraction dynamics under all three conditions were studied by transient photocurrent (34) measurements (Fig. 2G). The photocurrent decay time increased after annealing at 90 °C and then decreased after annealing at 150 °C as compared to the initial value. The results indicate that the charge extraction in OPVs became less efficient after annealing at 90 °C and much more efficient compared to that of the initial device after annealing at 150 °C.

To elucidate the origin of the thermal annealing-induced change in device performance, we investigated the roles of different materials by systematically changing only one layer or one material in the device. Upon only replacing the ZnO/PEIE bilayer with a ZnO single layer (SI Appendix, Fig. S11 and Table S3), or replacing PBDTTT-OFT/IEICO-4F with PBDTTT-EFT/IEICO-4F (SI Appendix, Fig. S12 and Table S4) or PBDTTT-OFT/PC 71 BM (SI Appendix, Fig. S13 and Table S5), a similar anomalous change in the efficiency was observed, yet with the highest recovered efficiencies of 11.6, 10.0, and 5.0%, respectively. Finally, we replaced MoO x with PEDOT:PSS and achieved 11.7% PCE with 66% FF (SI Appendix, Table S6). Interestingly, the decrease and recovery phenomena were not observed during the postannealing treatments (SI Appendix, Fig. S14). These results indicate that the anomalous change in the performance could be related to the changes in the MoO x layer during thermal annealing. Furthermore, we tested the reproducibility of such abnormal phenomena in OPVs with the MoO x HTL by changing the thickness and source of MoO x (SI Appendix, Fig. S15).

To investigate the change related to MoO x , we first characterized its composition, work function, and bandgap on ITO glass by X-ray photoelectron spectroscopy (XPS), UV photoelectron spectroscopy (UPS), and UV-visible (UV-vis) spectroscopy (SI Appendix, Fig. S16). Differing from previous reports (35, 36), we found no difference in the properties of MoO x after annealing, indicating that MoO x itself was stable after annealing.

We then characterized the interfacial property of the ITO/active layer/MoO x structure. We compared the UV-vis absorption spectra of the active layer and active layer/MoO x bilayer before and after annealing (Fig. 3A). For the active layer/MoO x bilayer, the light absorption in the range of 600 to 900 nm decreased after annealing at 150 °C, while that of the pure active layer changed negligibly. The light absorption at ∼726 nm changed negligibly for the active layer bilayer after the annealing process, while that of the active layer/MoO x bilayer decreased with a blue shift after annealing at 150 °C. Further, there was a very small new absorption peak at ∼1,100 nm after annealing at 90 and 150 °C, which indicates that polaronic or bipolaronic structures formed at the interface between the active layer and MoO x layers after the postannealing treatment (37). The above phenomenon indicates a change in the interfacial properties during the postannealing process. Fig. 3 B and C show the core-level XPS spectra of S and Mo for the active layer/MoO x bilayer. Compared with that in the initial state, after annealing at 90 and 150 °C, the S2p peak gradually shifted to a higher binding energy and the Mo+5 content increased with a concomitant decrease in the Mo+6 content (36). The UPS results show that the work function of the active layer/MoO x bilayer shifted from 5.9 eV to 5.6 and 5.5 eV after annealing at 90 and 150 °C for 5 min, respectively (SI Appendix, Fig. S17). We also measured the contact resistance of the devices using the ohmic contact region in the current density voltage curve. For the devices with the film of the pure PBDTTT-OTF, the contact resistance decreased from the initial 16.8 Ω∙cm2 to 1.4 and 0.1 Ω∙cm2 after annealing at 90 and 150 °C for 5 min (SI Appendix, Fig. S18). A similar phenomenon could be observed for the devices with the film of the PBDTTT-OTF/IEICO-4F.The above phenomenon indicates that the doping process happens between MoO x and components containing an S atom. Charge transfer occurs between the lone-pair electrons of S and the Lewis acid Mo during the postannealing process (SI Appendix, Fig. S19), leading to the doping process (38). This doping process can also happen to the S atom on other sites of PBDTTT-OFT and even on small-molecule acceptors. The doped interface between the active layer and MoO x at 150 °C can be more stable than the initial interface, and it contributes to the good device stability. This doping-induced stable interface approach has been utilized for fabricating stable organic light-emitting devices (39) and perovskite solar cells (40⇓–42).

Fig. 3. Mechanism of the high-temperature tolerance of nonfullerene blends. (A) Comparison of the UV-vis absorption spectra of nonfullerene blends (Top) and nonfullerene blend/MoO x bilayers (Bottom) before and after the postannealing treatment. It should be noted that the thickness of the active layer is 15 nm for the detection of the new absorption peak at ∼1,100 nm. (B and C) Evolution of XPS spectra of S (B) and Mo (C) at the interface between nonfullerene blends and MoO x film before and after the postannealing treatment. (D) Comparison of the pc-AFM photocurrent images and topography images of nonfullerene blends captured at 0 V before and after the postannealing treatment. (E and F) Comparison of 1D OOP (E) and IP (F) GIWAXS line profiles of the nonfullerene blends before and after the postannealing treatment.

We speculate that multiple factors cause the abnormal change in the device performance at different temperatures. According to our results, the doping levels between MoO x and the active layer at 90 and 150 °C are different, which should lead to different levels of enhancement of the device performance (19, 38, 43). The decreased device performance after annealing at 90 °C has to be caused by other factors, such as by a change in the interfacial morphology due to the doping process. Further characterization is needed in the future, which is not within the scope of this work.