This review paper aims to provide a basic understanding of THP as a semiconductor by highlighting the physical structure, energy band structure, electrical properties, and doping mechanisms. Additionally, the key chemical instability issues of THPs are discussed, which are identified as the potential bottleneck for further device development. Based on the understanding of the THPs properties, the key recent progress of THP‐based solar cells and FETs is briefly discussed. To conclude, current challenges and perspective opportunities are highlighted.Because of its less toxicity and electronic structure analogous to that of lead, tin halide perovskite (THP) is currently one of the most favorable candidates as an active layer for optoelectronic and electric devices such as solar cells, photodiodes, and field‐effect transistors (FETs). Promising photovoltaics and FETs performances have been recently demonstrated because of their desirable electrical and optical properties. Nevertheless, THP's easy oxidation from Sn Tin halide perovskite (THP) is a promising material for active layer in optoelectronic and electric devices, due to its less toxicity and analogous electronic structure to that of lead. This review provides a basic understanding of the material, including lattice structure, energy band, doping and chemical instability, and presents key progress and perspectives of THP‐based high performance solar cells and transistors. Such prompt progress of MHPs in a broad spectrum of the scientific community is due to not only its impressive optoelectronic properties but also promising cost‐effective, high‐throughput, and low‐temperature manufacturing processes. The bulky three‐dimensional (3D) form of MHPs comprises of a corner‐sharing BX6 octahedral network with the chemical stoichiometry ABX3, where A is an organic or inorganic monovalent cation, B is a divalent metal ion, and X is a halide anion. The B‐X‐B configuration in a BX6 octahedron has a bonding angle of 180°, inducing a large dispersion of both the conduction and valence bands originating from super degeneracy. This electronic structure implies that MHPs possess intrinsically outstanding and desirable electrical and optoelectronic properties such as high charge‐carrier mobility, tunable bandgaps, large optical absorption coefficient, long exciton diffusion length, and ambipolar charge transport. Remarkable performances from MHP‐based FETs and MHP‐based solar cells (PSCs) (power conversion efficiency (PCE) > 25%) imply the potential scopes for developing future functional optoelectronic devices.To date, metal halide perovskite (MHP) can be considered as a promising candidate for various next‐generation optoelectronics and electronics, such as solar cells,1 fiel3 structure, the toxicity issue of Pb remains a long‐lasting debate. The concentration of immediately dangerous to life or health of Pb is 100 mg m−3, which means there is a risk of human exposure to Pb in the fabrication and handling of perovskite devices. Moreover, our biological systems might be in great jeopardy due to the water‐soluble properties of Pb. Thus, to prevent any perturbing issues related to Pb, the European Union has publicly announced that exposures of up to 1000 ppm are allowed for all electronic devices. Systematic encapsulation protocol and numerous encapsulation approaches have been given serious consideration to inhibit any possible leakage of Pb; however, it is difficult to ensure how safe it is when it comes to device fabrication as well as long‐term storage. Even with a sophisticated encapsulation procedure, there is no guarantee that potential Pb leakage can be prevented. Therefore, controlling, handling, and securing safety during the manufacturing of any related Pb‐based MHPs devices remain a huge challenge.Despite these exciting properties and advances, most MHP devices today are manufactured using Pb as the divalent metal in the ABX devices.15 Mor Bi, and Ge have been proposed to replace Pb, not only to reduce its toxicity but also to maintain the unique optoelectronic properties of perovskite. Among these materials, eco‐friendly material Sn has been widely utilized in various promising optoelectronic devices, including solar cells and FETs. Similar to that of Pb, Sn has an inactive outer shell s orbital, which is vital to realizing the distinctive electrical and optical properties of MHP. The effects of Sn substances may vary, and it is relatively safe for a human when the hydrogen bonds grow longer in Sn substances. Still, the influences and their adverse effect on human health have raised some eyebrows, there is a big question regarding its use.Apart from the encapsulation method, recent attention has been devoted to replacing Pb with low toxicity and non‐toxicity cations. Several potential low‐toxic and chemically compatible materials such as Sn,16 Bi, followed by a systematic study from Donaldson's research group. Later, the first hybrid THP was proposed by Yamada et al. in which Mitzi and colleagues expanded the work and conducted a study in the context of dimensional reduction, and demonstrated the first structural properties under pressure. These initial discoveries stimulate researchers to devote huge attention to THP‐based devices until now. In 2014, Kanatzidis et al. reported the first Sn‐based perovskite solar cells (TPSCs). Subsequently, numerous studies have utilized Sn instead of Pb for various type of optoelectronics devices; for instance, the latest breakthrough in TPSCs have witnessed the PCE of 14.8%, whereas the highest hole and electron mobilities of 55 and 2.1 cm2 V−1 s−1 for THP‐based FETs have been demonstrated. The rapid development in THP‐based optoelectronics devices makes it urgently necessary to review and assess their potential for several optoelectronic and electronic applications. In this regard, earlier published reviews have highlighted the progress and potential of THP's applications with solar cells, photodetectors, light‐emitting devices, and radiation detectors. With intense research for various applications, THP‐based electronic devices show encouraging performances. However, the stability of the THP‐based electronic devices is still poor as compared to the Pb‐based ones. In this article, we aim to provide a comprehensive understanding of THP and highlight the impressive latest results for their functional device applications. First, we discuss the structure, chemistry, and thin‐film formation of THP. Second, we introduce the structure‐electrical property relationship of THP with energy band structure and electrical doping through intrinsic and extrinsic doping. Third, we review the recent progress of TPSCs. In this section, we also discuss the various chemical bonding and chemical structure of THP and their effect on the fabrication of TPSCs. Fourth, we discuss recent progress in THP‐based FETs, including two‐dimensional (2D) perovskite, hybrid 2D/3D perovskite, and 3D perovskite. Finally, we review the prospects for THP‐based PSCs and FETs, aiming to motivate the scientific community and better understand the development of future optoelectronic devices  was successfully developed by Fisher et al. in 1974,21 fol22.13 stoichiometry  and X is a halide: I−, Br−, or Cl−. By permutation of the A‐site cations and anions, a huge number of THPs are possible. Besides the widely used Cs and Rb, 13 protonated amines can be incorporated into the perovskite framework, as A‐site cations and eight anions, to obtain different types of perovskites. Additional structures can also be achieved through doping with various compounds. Nevertheless, to realize stable THPs, much theoretical understanding and fundamental calculation are necessary. At room temperature, the amine cations , formamidinium (FA)) in the perovskite cavity are generally disordered, and the systems exhibit a remarkable spectrum of ferroelectric and multiferroic characteristics when cooled. However, to realize an ideal perovskite structure, the cations must be well‐matched with the anions. The Goldschmidt model provides an ideal method to determine the stable perovskite structure.aR, bR, and xR are the ionic radii of the A, B, and X sites, respectively. Hence, the tolerance factor (t) assesses whether the A‐cation can fit within the framework of corner‐sharing octahedra. Experimental data suggest that t values have to be in the range between 0.8 and 1.0 to maintain a 3D perovskite structure. If the t values are within 0.9 to 1.0, an ideal cubic phase is expected. For the largest value of bR and xR  within the optimum t limit (0.8 < t < 1), the ionic radius for A‐site cation to form perovskite framework was calculated at 1.60–2.60 Å. Hence, MA (2.17 Å) and FA (2.53 Å) are the most suitable organic cations for developing effective perovskite compounds. Incorporation of A‐cation beyond the above size limit could reduce the dimensionality of perovskite from 3D to 2D and 1D. Interestingly, when divalent B‐cation is replaced from large Pb to small Sn cation, the values of t increase and approach toward 1. One of the remarkable aspects making the crystal structures of perovskites versatile is the structural flexibility of organic groups  with a smaller cation (Cs+) too. The doping of FASnI3 with Cs could tune the t downward from 1.04 to 1.02. THPs have high electrical conductivity (σ) because of the marginal crossing of Sn 5s and Sn 5p bands near the Brillouin zone boundary (due to the large dispersion of Sn 5s hybridized band) point . Relativistic effects  are stronger in Pb than Sn, which leads to stabilization of 6s2 orbital of Pb relative to 5s2 of Sn.THPs consist of ABXy Figure 1a. To mapossible.35 Bess Figure . A t val2.1.13 perovskites were studied intensively because of their similarity with MAPbI3. MASnI3 shows a cubic (Pm3m) phase at 295 K; however, a phase transition from cubic to lower symmetry tetragonal (I4/mcm) occurs at 275 K. Because the SnI6 octahedra tilt about the vertical axis, the tetragonal structure corresponds to a √2a × √2a × √2a supercell extension of the cubic lattice. By lowering the temperature to 108–114 K, another phase transition from tetragonal to orthorhombic (Pbn21 or Pbnm) can be obtained. In the adjacent planes, the SnI6 octahedra tilt in the same direction around the c–axis , the CsSnI3 perovskite converts to a black cubic (B‐α) phase   phase is stable at room temperature. In this aspect, the B‐γ phase of CsSnI3 perovskite is most suitable for device applications. Nonetheless, studies suggest that exposure to ambient air or organic solvents in this phase tends to transform into a double‐chain structure as shown in Figure 2SnI6, which has poor optoelectric and electric properties making it unsuitable for photovoltaic and FET applications. Kubicki et al. performed a detailed solid‐state nuclear magnetic resonance (ssNMR) study to compare the structures of MASnBr3, FASnBr3, CsSnBr3, MASnI3, FASnI3, and CsSnI3 perovskites. For every I/Br and Br/Cl mixed ratio, the ssNMR results showed that iodide–bromide and bromide–chloride mixtures can form a stable phase. Iodide–chloride mixes, conversely, produced phase‐segregated mixtures of phases, despite being somewhat miscible.When compared with MAe Figure 4a. Upon ) Figure  phase. Aperature.57 In 2.2s orbital lone‐pairs . The bandgap of THP is ≈1.4 eV, which is smaller than the bandgap of LHP (≈1.5 eV). The bandgap of THP approaches Shockley–Queisser optimal bandgap for single‐junction solar cells, and under ideal fabrication conditions, TPSCs are expected to reach a theoretical PCE up to 33.4%. Additionally, the lone‐pair state of Sn 5s is more reactive than Pb 6s because the strong lanthanide shrinkage effect of Pb makes them an inert electron pair. Because of the absence of this effect, Sn becomes an easily oxidized element with two active Sn 5s electrons  is generated from the antibonding of Sn 5p and I 5p orbitals with dominant contribution by Sn 5p orbital  method in comparison with its Pb counterpart. The band structure of MASnI3 showed a direct bandgap of 1.1 eV (1.67 eV for MAPbI3 at the same point), comparable with the bandgap of ≈1.3 eV obtained experimentally  is also smaller than that of MAPbI3 (0.25 m0). In the case of CBM, the Sn 5p orbital is shallower and less dispersive than Pb 6p orbitals due to much weaker SOC, and therefore, the CBM of MASnI3 is less dispersive and positions at 0.2 eV higher than that of MAPbI3. Thus, the calculated average electron effective mass of MASnI3 (0.28 m0) is larger than that of MAPbI3 (0.19 m0).The energy band structure is the key to analyzing the optoelectronic and electronic properties of semiconductor materials. Through combined theoretical and experimental approaches, the band structures of THP can be examined. The valence band maximum (VBM) of THP is composed of the antibonding between Sn 5l Figure .‐Phenyl C61 butyric acid methyl ester (PCBM) (CB = −3.90 eV)‐, C60 (CB = −4.50 eV)‐, and ICBA (−3.70 eV)‐based TPSCs have shown reproducible performances with impressive PCEs  in the MASnI3 framework with a hypothesis on hollow MASnI3 perovskite. These {en}MASnI3 films showed a much lower electron‐hole recombination ratio because of the high surface coverage on top of the mesoporous TiO2 layer and the respective TPSCs attained a PCE of 6.63%. A‐cation exchange approach through a two‐step deposition process was proposed to reduce the background carrier density of the MASnI3 films. In the first step, a hydrazinium tin iodide (N2H5SnI3) layer was deposited by a simple spin‐coating technique, which was followed by a transformation into MASnI3 in a MA gas atmosphere through an organic cation displacement approach. The two‐step deposition method reduced the electron‐hole recombination of the MASnI3 film and resulted in dense and uniform MASnI3 film with micrometer‐sized grains and high crystallization and resulted in PSCs with PCE > 7%. Besides the solution‐processed mesoporous TPSCs, vapor‐assisted PSCs were evaluated by a low‐temperature vapor‐assisted solution process but showed poor photovoltaic performance.The early days of TPSCs applications were conducted using MASnI of 6.4%.68 TheVOC, ion exchange/insertion reactions between solid‐state SnF2 and gaseous methylammonium iodide was adopted for the preparation of MASnI3 films. Contrary to the traditional approach, the p‐i‐n was adopted where PEDOT:PSS was applied as the HTL. The respective MASnI3 films showed a highly uniform, pinhole‐free coverage with excess SnF2, and the respective film yielded a low content of Sn4+. This high‐quality perovskite film enables the realization of a PCE of 7.78%. Despite numerous efforts, the PCE of MASnI3‐based TPSCs is still low compared with other TPSCs. Table 3‐based PSCs  outperforms the n‐i‐p structured TPSCs (PCE = 4.34%) due to favorable energetic matching with their respective charge transport layers (ETLs and HTLs). Since then, almost all FASnI3‐based TPSCs studies have been performed utilizing the p‐i‐n structure. Hence, we only discuss here the development of FASnI3‐based TPSCs for p‐i‐n structure.Significant progress was achieved for TPSCs when MA cation is replaced by FA. FA cation can stabilize the crystal structure more than the MA and lead to higher photovoltaic performances.79 Fur2+ oxidation when introducing coadditives with SnF2 as they tend to form a complex within the perovskite system, which can encapsulate the perovskite grains and modify the crystal structure. In the coadditive engineering approach, SnF2 is often used as the main additive and “other” coadditives were added. For simplicity of discussion, only “other” coadditives are discussed where mentioned. To retard the oxidization of Sn2+, a few reducing agents such as Catechin, ethylenediammonium dihypophosphite, and hypo‐phosphoric acid have been evaluated, but they failed to show a high PCE. Hydrazinium chloride showed a remarkable performance to reduce the Sn2+ oxidation by 20% , and phenylhydrazine hydrochloride (PHCl) were evaluated for TPSCs. Trihydrazine dihydriodide coadditive engineered FASnI3 TPSCs showed improved PCE of 8.48%. PHCl as a coadditive showed a robust capacity to reduce the Sn2+ oxidation and shifted the VBM by 0.984 eV  of gallium acid (GA) and an excess of SnCl2 can lead to the formation of an amorphous GA–SnCl2 complex and result in the self‐encapsulated phase of the FASnI3 perovskite, which increases the stability of TPSCs. The large bandgap of GA‐SnCl2 based FASnI3 prohibits the transfer of both charge carriers from FASnI3 to the adjacent charge transport layers in TPSCs. The FASnI3‐based TPSCs fabricated by this complex showed a PCE of 9.03% with stable photovoltaic behavior without any degradation for 1500 h in N2 environment. Consecutively, Trimethyl thiourea (3T), with structural features of both a Lewis base and H‐bond donor can influence the fabrication of high‐quality FASnI3 films. The unique bifunctional properties of 3T as a co‐additive can successfully enhance the morphology and texture of FASnI3 films by spreading and joining individual crystal grains. With such attributes, the charge carrier lifetime of FASnI3 perovskite was increased up to 123 ns leading to low VOC loss  in the respective TPSCs. Due to the high VOC of 0.92 V, JSC of 20.4 mA cm−2 and FF of 0.76, the 3T‐FASnI3‐based TPSCs showed a high PCE of 14.05%. It is noticeable that even with staggering VOC, the JSC remained low, which was attributed to the low perovskite film thickness rather than the charge‐carrier diffusion length. However, even increasing the film thickness did not result in higher JSC, rather resulted in lower VOC.Coadditive engineering has shown promising aspects to retard the SnCatechin,150 etV Figure , which iV Figure . The pre2+ in FASnI3‐based perovskite due to the rapid crystallization of the perovskite layer. Even if the precursor solution is heat treated at 100 °C in a glovebox filled with PPM level oxygen before spin‐coating and antisolvent treatment, the organocation FA+ and iodide ion (I−) are easily volatilized and are not properly integrated into the perovskite structure. This phenomenon induces rapid oxidation and results in high crystal defects of the FASnI3 film. To decelerate the crystallization process, Jokar et al. highlighted ethylenediammonium diiodide (EDAI2) in trace amounts can lower the crystallization rate and promote the growth of highly crystalline FASnI3 films more slowly. During the slow crystallization process, EDAI2 can retard the VSn inside the FASnI3 crystal by decelerating crystal growth through a kinetic balance between nucleation and crystal growth. The 1 mol% EDAI2‐FASnI3‐based TPSCs can achieve a PCE of 8.5% and exhibit stable performance for 1400 h in an N2 environment. The π‐conjugated Lewis‐base molecule 2‐cyano‐3‐[5‐[4‐(diphenylamino)phenyl]‐2‐thienyl]‐propenoic acid (CDTA) shows electron‐accepting properties to reduce the oxidation and binding to the SnI2 at the precursor stage. This strong electron‐accepting behavior contributes to a more stable Lewis adduct during the nucleation process. The CDTA‐based FASnI3 films suppressed the in‐plane and out‐of‐plane rotation of [SnI6]4− and reduced the carrier recombination of the FASnI3 films. Consequently, the corresponding TPSCs showed a high PCE of 10.32%, which is attributed to a 0.13 V increase in VOC and a 10% increase in FF compared to the pristine TPSC. Controlling the crystallization via electron‐accepting small molecules of CDTA attained the first time observation of light soaking stability of the TPSCs at maximum power point tracking (MPPT) for 1000 h. Another molecule poly (PVA) with a high density of hydroxyl groups can make a strong O—H—I− hydrogen bond with the FASnI3 system leading to nucleation sites which induce lowered crystallization rate. As a result, the FASnI3 perovskite crystals are more oriented and can inhibit the migration of the iodide ions to the adjacent charge transport layers. With these attributes, the VOC value is increased from 0.55 V (for pristine TPSC) to 0.63 V (PVA based TPSC) and resulted in long‐term stable photovoltaic performance with a PCE of 8.9%. By taking the advantage of the slower nucleation of the H—I bond and inducing highly crystalline films, a bifunctional compound, hydroxylamine hydrochloride (HaHc) has recently been applied to reduce electronic defects of the FASnI3 perovskites. The hydroxyl group in the HaHc formed a hydrogen bond with iodide ion in the FASnI3 perovskite system to slow the crystallization, and the Cl− ion coordinated with the under coordinated Sn2+ ions of the FASnI3 structure. With these bifunctional aspects of the HaHc, the corresponding HaHc‐FASnI3 based THP films showed lower trap assisted recombination, bimolecular recombination, and trap density as compared to the pristine FASnI3. With such beneficial improved electronic properties, the VSn was reduced and improved VOC up to 0.676 V for the HaHc‐FASnI3 based TPSCs was observed with 500 h light soaking stability. In terms of VSn reduction, the H—I bond inclusion seems to be an effective approach, although the resultant TPSCs performance is still lower when compared to coadditive engineering approaches. Meng et al. proposed that the crystallization of FASnI3 perovskite is more vulnerable to the solution air surface. By using a tailormade fluorinated organic cation, pentafluorophen‐oxy‐ethylammonium iodide (FOEI) in the FASnI3 precursor solution, the surface energy of the solution/air surface was reduced, which simultaneously served as a template and resulted in highly crystalline films. When the FOEI‐FASnI3 films were measured using depth‐dependent grazing incident X‐ray diffraction measurements at different incident angles  was evaluated as potential to further stabilize the structure of the FASnI3 perovskite by Dai et al. When compared to the conventionalFASnI3:xSnF2 precursor, the MAI·3MA: FASnI3:xSnF2 precursor can induce chemical stability owing to due to the extra donor electron groups in the amine complex. Due to the complex formation, the crystallization rate from precursor to THP film is significantly reduced resulting in highly crystalline THP films with low Sn2+ oxidation and non‐radiative recombination centers. As a result, the VSn is reduced and the corresponding TPSC can result in VOC of 0.65 V, leading to a PCE of 9.53% with a light soaking stability of 1000 h. However, in a recent breakthrough study, Jiang et al. highlighted the importance of SnI2 complex for regulating the crystal growth of FASnI3 perovskite. Typically SnI2 complexes are formed using two step synthesis process (TSS). However, SnI2 complexes prepared by TSS method result in SnI2 segregation. The segregation causes unsolvated edge‐sharing SnI2 clusters which limit the crystal modulation of the THP growth. To promote smooth, compact FASnI3 layer, and to retard the SnI2 segregation, a one‐step synthesis (OSS) of SnI2. (DMSO)x colloidal complexes were developed, where in situ, tin metal and I2‐DMSO in solution complex are obtained. The SnI2.(DMSO)x synthesized by the OSS method can be well dispersed to obtain SnI2 precursor without the formation of any edge‐sharing clusters. With these beneficial aspects, the stable SnI2. (DMSO)x colloidal complexes can assist to bind with the cations of FAI and phenylethylamonium bromide (PEABr) and promote the growth of highly crystalline and smooth THP film. As a result, the electron diffusion length of the THP film was improved to 290 ± 20 nm (versus pristine = 210 ± 20 nm). As a result, the corresponding TPSCs show a high VOC of 0.91 V, JSC of 20.6 mA·cm−2, and a FF of 0.77; leading to a PCE of 14.63%. It must be noted that by this approach TPSCs showed first time photovoltaic performance toward 15% PCE.Amine complex CH2+ oxidation. In recent reports, template‐assisted seed growth has shown more success to fabricate high‐quality FASnI3 film for TPSCs application. In a seeded growth (SG) approach proposed by Cao et al., a layer of perovskite film is first deposited by an antisolvent method, and the same step is repeated. In this method, the first spin‐coated FASnI3 layer acts as a template to promote the growth of uniform THP layer with low grain boundaries and highly crystalline perovskite films with compact and large grains and a longer carrier lifetime. TPSCs fabricated by the SG method showed a VOC of 0.49 V, which is far below other conventional reported methods. To facilitate template assisted growth, poly(ethylene‐co‐vinyl acetate) (EVA) self‐sealing polymer with the ability of complex formation with SnI2 was tested and verified. The C=O groups of EVA act as a powerful Lewis acid–base and forms a SnI2 complexation with uncoordinated tin atoms in perovskite grains, which influences the improvement of the grain size, optimized grain orientation, and decrease the surface defects of FASnI3 films. Furthermore, The FASnI3.EVA based THPs show a self‐encapsulation effect which can repeal the moisture and oxygen ingression and result in humidity (60%) stable TPSCs.Throughout the THPs progress the solvent to film conversion rate has been highlighted as a potential source for Sn3 surface is more prone to Sn2+ oxidation and leads to extreme surface defects, the post‐treatment method was considered a possible solution. Chowdhury et al. introduced a postdeposition vapor annealing method (PDVA), where after the antisolvent treatment, the FASnI3 films were treated in MACl enclosed vapor system. The PDVA method induced the growth of high‐quality, defect‐free FASnI3 film and showed a PCE of 8.48% when implemented in TPSCs. A solution approach for post‐treatment by bidentate amine‐edamine was evaluated by Kamaruddin et al. Before the antisolvent treatment, edamine treatment was introduced on top of the FASnI3 films. The addition of edamine which has a free electron pair acts as an electron donating group and facilitates the co‐ordination with the under‐coordinated Sn2+ of the FASnI3 system  with 1000 h light soaking stability under MPPT conditions by the pretreatment method of FASnI3 film  as IL to induce Ostwald ripening effect of FASnI3 films. During the thermal annealing of perovskite films, BMIBr IL domains generated by a lower melting point act as Ostwald ripening agents for dissolving the perovskite grains. The FASnI3 films fabricated by inclusion of BMIBr showed larger grain size with fewer grain boundaries which eventually reduces the defect states. The successful suppression of defect states reduces the VSn of the FASnI3 films and the corresponding TPSCs showed VOC up to 0.68 V, with JSC of 19.63 mA cm−2 and FF of 0.72 (PCE = 10.09%). In an alternate approach, IL formamidine acetate (FAAc) improved the FASnI3‐based TPSCs performance up to 9.96%. The FAAc contains CH3COO− a group, which can form an intermediate phase and slow the nucleation of the FASnI3 film. Even after the thermal annealing treatment, the FAAc remained within the perovskite system and retarded the cationic vacancies of the respective THP film, helping to retain 82% of the initial PCE during the 1500 h light soaking stability test.n‐butylammonium acetate (BAAc) IL showed coordination with specific O…Sn chelating bonds and N‐H…X hydrogen bonds in the FASnI3 system retarded the Sn2+ oxidation in the precursor solution. Contrary to the other reports, the BAAc can suppress precursor Sn2+ oxidation even when heated at 100°C for 2 h and can form highly crystalline FASnI3 films with fewer surface defects. The BAAc‐FASnI3‐based TPSCs exhibited a PCE of 10.4% and maintained 80% of their initial photovoltaic performance for up to 100 h at 85°C. Note that this is the first report on TPSCs showing stable photovoltaic performance at elevated temperatures.Ionic liquids (ILs) are well known for their very low vapor pressure, making them ideal candidates for slower crystallization. When a small amount is added to the perovskite precursor, it tends to remain in the resulting perovskite film after all the solvent has been evaporated during the annealing process. With controlled crystal growth, the surface defects at the perovskite grain boundaries can be reduced effectively. Additionally, ILs can form a dipole at the perovskite/charge transport layer interface, which can improve the energy level alignment due to the reduced interfacial energy barriers.170 LiVOC of the TPSCs is often low. Abdoul‐Sahkour et al. highlighted the importance of the charge collection loss of the FASnI3/ETL interface and introduced diaminomaleonitrile (DAMN) in the FASnI3 system to collect the photogenerated charges more effectively. The DAMN with two cyano groups which perform as an electron‐withdrawing group can successfully extract electrons effectively from the FASnI3 layer and transfer them to the adjacent C60 ETL, which was confirmed by a 42% increment of electron mobility. Consequently, the VOC of the DAMN‐FASnI3‐based TPSC improved by 0.5 V compared with the pristine FASnI3. Due to the better charge management in the FASnI3 system, the DAMN‐FASnI3‐based TPSC showed stable photovoltaic behavior at MPPT conditions for 300 h. In a recent report, Chen et al. highlighted the importance of the charge transport at the HTL/FASn0.9Ge0.1I3 interface. To promote efficient charge transport at the HTL/FASnI3, the FASnI3 perovskite was doped with 1 mol% GeI2. Unlike the usual dopants or additives which contribute to the surface of the perovskite layer, the GeI2 induces in situ formation of a thin (≈3 nm) amorphous interfacial GeO2 layer at the NiOX/FASn0.9Ge0.1I3 interface. The GeO2 layer prevents the Sn2+ oxidation at the interface, provides more robust mechanical bonding, and promotes effective charge transportation in the TPSC. As a result, the corresponding TPSCs showed a high PCE of 10.43%.Due to the less charge extraction ability at the perovskite/ETL interface, the 3 can tune the bandgap, push the VBM to a more shallower region, increase the light‐harvesting capability of the THP films and provide a basis for high‐performance TPSC fabrication. Compositional engineering by cationic displacement for FASnI3‐based TPSCs was first evaluated by Zhao et al. where MA cation was partially introduced in the FASnI3 system. The increasing content of MA cation in the FASnI3 perovskite system shifts the VBM level and induced a favorable VBM level with the adjacent PEDOT:PSS HTL 0.25SnI3‐based TPSCs showed PCE of 8.12%, which were higher than the MASnI3 (PCE = 4.29%)‐ and FASnI3 (PCE = 6.60%)‐based TPSCs. The improved photovoltaic response was attributed to the improved VOC by 0.15 V compared with the MASnI3 or FASnI3. Liu et al. discussed the importance of the appropriate antisolvent for the fabrication of FAXMAXSnI3 films and revealed that chlorobenzene (CB) is the more suitable antisolvent. When compared with the other antisolvents of diethyl ether and toluene, CB as an antisolvent could assist the formation of a dense, compact FA0.75MA0.25SnI3 layer and the respective TPSCs showed 9.06% PCE.Cationic or anionic displacement of FASnIrication.177 Co0.75MA0.25SnI3 perovskite films to suppress the nonradiative recombination, losses were reported by Wan et al. Before the antisolvent treatment of the pre‐annealed FA0.75MA0.25SnI3 perovskite films were placed under high vacuum of 5 × 10−4 Pa. The introduction of vacuum treatment before the annealing in order to reduce the trap density reduced the nonradiative recombination and promoted interfacial charge transfer in the respective TPSCs leading to a VOC of 0.631 V and FF of 0.75 (PCE of 10.3%). Inclusion of Diethylammonium iodide (DEAI) as a cation in the (DEAxFA1–x)0.98EDA0.01SnI3 perovskite system can shift the VBM to a much deeper level (from −4.97 to −5.15 eV) and reduce the deep trap states of the respective THP film with low defect densities. With such favorable VBM level, the hole transportation between the PEDOT:PSS and the (DEAxFA1–x)0.98EDA0.01SnI3 layer is more efficient and the VOC loss is reduced in the respective TPSC. With an optimal 10% DEA in this mixed cation‐based TPSC a high VOC of 0.67 V (versus pristine = 0.59 V) was attained. By introducing a trace amount of ethylammonium (EA) as the cation to the FASnI3 systems, a noteworthy improvement in TPSCs was observed by Nishimura et al. The FA0.98EA0.01SnI3 based THP showed a shift of 0.30 eV in VBM and 0.15 eV shift in CBM. The favorable shift of both ends facilitated obtaining a high VOC of 0.62 V while maintaining a high JSC of 23.05 mA cm−2 leading to a PCE of 10.80%. The inclusion of EA as a cation could successfully repress the trap states of the THP layer which increased charge carrier mobility. Further surface treatment of the TPSCs with EDAI elevated the VOC up to 0.84 V and reached a PCE of 13.24%. However, after EDAI surface treatment, the JSC decreased to 20.32 from 23.05 mA cm−2, which highlights the double‐sided effect of the surface treatment process for mixed cation THP perovskites. To overcome the drawbacks of the reduced JSC, the same research group recently introduced trimethylsilyl bromide (Me3SiBr) as a surface passivation agent instead of EDAI. Upon surface treatment with 0.15 mm Me3SiBr on top of the FA0.98EA0.01SnI3 perovskite, the surface morphology was improved, which showed lower defect densities, larger grains, and higher carrier density. As a result, the JSC of 24.11 mA cm−2 can be maintained even after surface treatment. However, the VOC of the respective TPSC was reduced to 0.70 V, leading to a PCE of 12.22%. Additionally, the Me3Si+ remains as the counter cation on the surface and forms a hydrophobic protective layer which protects the FA0.98EA0.01SnI3 film from oxidation which showed stable photovoltaic performance in a N2 glovebox for 92 days of storage in N2 filled glovebox by maintaining 80% of its initial efficiency.In an alternate approach by vacuum‐assisted self‐healing of defects in FA0.75FA0.15PEA0.1Sn(Br0.33 I77)3 can increase the bandgap from 1.29 to 1.50 eV. Although a significant improvement in TPSCs performance was expected, their fabricated TPSCs showed a low PCE of 4.63%. The bandgap of FASnI3 films was further enhanced to 1.61 eV by the addition of GuBr by Chen et al. with an optimized composition of GA0.06(FA0.8Cs0.2)0.94SnI2Br, which showed matched energy level alignment with the adjacent charge transport layer and reduced the trap density (both surface and bulk). The best performing TPSC showed OCV of 0.64 V, JSC of 15.16 mA cm−2, and FF of 0.72 after aging. In a recent report, Br mixed perovskite system of EDA0.01(GA0.06(FA0.8Cs0.2)0.94)0.98SnI2Br, the PCE was further enhanced up to 8.66%. However, the improved photovoltaic performance of these wide bandgap TPSCs was attributed to the influence of replacement of PEDOT:PSS by a monolayer of 2PACZ rather than the influence of the wide bandgap THP layer  revealed that SnCl2 is the most suitable compound to obtain smooth surface morphology and maintain the B‐γ phase of CsSnI3 for TPSCs applications with PCE of 3.56%. However, the partial replacement of Cl and I was not observed; rather, an ultrathin layer of SnCl2 layer formed on top of the perovskite layer and provided assistance to achieve stability.Because of the absence of hydrophilic organic cations of MA and FA, CsSnXith LHPs.190 Ho2 additives in the CsSnI3 framework induced some photovoltaic response but the PCE was significantly lower than 5%. The SnX2 additives failed to address the critical Sn2+ oxidation even in an N2 atmosphere with a trace amount of oxygen (1 ppm). To increase the photovoltaic response, taking for example from FASnI3 perovskites, the coadditive strategy was adopted for CsSnI3 perovskite films. The addition of 2‐aminopyrazine with 20% SnF2 remarkably suppressed the Sn2+ oxidation, which was attributed to the electron‐donating capability of the amino group in the aminopyrazine and the formation of the aminopyrazine–SnF2 complex. Upon spin‐coating, the precursor followed by antisolvent and thermal annealing at 100°C resulted in smooth CsSnI3 perovskite films with compact grains. Although the VSn of the respective TPSC was not reduced which can be observed from the low VOC of 0.40 V. A higher PCE of 7.50% was attained by utilizing the lone electron pairs of −NH and −CO functional groups of methylenebis(acrylamide) (MBAA)  was used as a coadditive with SnF2, the electron density of CsSnI3 was successfully passivated, and smooth pinhole‐free surface morphology was observed  substrate and immediately transferred to the vacuum chamber to remove the solvent. The introduction of such a fast solvent removal technique improved the crystallization and surface properties but attained only a PCE of 3.8%. Yin et al. proposed a sequential vapor deposition strategy, which combined the advantages of both the vapor deposition and the surface passivation. Thiosemicarbazide (TSC) with S‐C‐N functional groups was introduced as a passivator in addition to the SnI2 and CsI evaporation source. The vapor deposition was performed sequentially by depositing SnI2, TSC, and CsI sequentially and annealing at 120°C to form a (B‐γ) CsSnI3 with large grain size and uniform coverage. The fabrication of extraordinary (B‐γ) CsSnI3 films prepared using PASVD was attributed to the strong electrostatic attachment and coordination interaction, which leads to electron cloud density and reduced the VSn. When TPSCs were fabricated in a p‐i‐n configuration, the VOC enhancement from 0.47 to 0.63 V was observed which is ascribed to the reduced deep level trap‐state density of the successfully passivated THP films. Interestingly, by adoption of this approach, a PCE > 8.0% was noted for CsSnI3‐based TPSCs with stable photovoltaic performance up to 500 h  method and the vacuum flash‐assisted solution processing (VASP) method have been evaluated for the fabrication of CsSnIof 2.23%.198 Th4.3.43 film quality, the conventional 2D/3D approach was examined. Perovskites with a Goldschmidt tolerance factor (t) between 0.8 and 1.0 show a photoactive black phase, which can induce high‐quality MASnI3 film. Ji et al. adopted EA cation into MASnI3 to induce a preferential orientation perpendicular to the substrate resulting in a 2D/3D MASnI3‐based PSC with PCE of 9.24%. The (EAXMA(0.98−X)2D/3D)SnI3 films could successfully passivate the hole traps, induce preferential crystallization rate, produce a uniform and smooth surface, and suppress the trap‐assisted nonradiative recombination. Shao et al. introduced a lower concentration of PEAI (0.16 M) to FASnI3 to obtain 2D/3D structures and verified that the crystallinity of 3D THPs is significantly better when compared with the 3D FASnI3‐based THPs, which increased the VOC of the respective TPSCs by 0.67 V when compared with the 3D FASnI3‐based TPSC and showed PCE of 9%. In follow‐up work, the authors evaluated the importance of stoichiometry of 2D/3D PEA0.08FAxSnI3 (where x refers to the concentration of the FA cation in the case of fixed PEA concentration). It was observed the concentration of FA cation significantly impacts the crystallinity of the 2D/3D films. A hierarchical 2D‐quasi‐2D–3D structure using NH4SCN as an additive was proposed to regulate the crystal growth of PEAxFA1‐xSnI3 THP films by Wang et al. With regulated crystal growth, the quasi‐2D/3D THP films showed higher carrier mobility and reduced the carrier density resulting in a PCE of 9.41% in the respective TPSCs. They hypothesize that a thin layer of 2D layer of PEA2SnI4 on top of the 3D FASnI3 film can provide resilience against ambient air and moisture. Because of this, the respective TPSCs can maintain 90% of their initial PCE even after 600 h storage in air. Chen et al. added a bulky divalent organic cation, 4‐(aminomethyl)‐piperidinium (4‐AMP) to FASnI3 system. Interestingly, the addition of 15 mol% of 4‐AMP with FASnI3 did not influence 2D/3D formation but rather remained at the surface and grain boundaries. Interestingly, this concentration was not enough to create detectable low‐dimensional phases. To further elevate the PCE, Jiang et al. proposed that for the development of 2D/3D aside from the THP layer the adjacent charge carriers play a pivotal role. They replaced the ETL (PCBM) with ICBA and HTL (NiOX) with PEDOT:PSS in the PEAxFA1‐xSnI3+NH4SCN‐based TPSCs and increased the VOC up to 0.94 V, leading to a remarkable PCE of 12.4%. Such high VOC was attributed to the matched energy level of the HTL/perovskite and perovskite/ETL interface. By using the same HTL and ETLs, the best performing TPSCs with 14.81% were reported with 4‐fluoro‐phenethylammonium cations (FPEABr) in FASnI3 system. It is well known that the fluorinated cations can create stronger Van der Waals interaction or hydrogen bonding reducing the Sn2+ oxidation and creating a 2D microstructure to influence the 3D phase growth. Additionally, fluorinated compounds provide the unique characteristics of hydrophobicity due to the presence of fluorine atoms, which resulted in the high PCE‐based TPSCs with longer stability. Despite the significant efforts by numerous research groups, the VOC loss of TPSCs was minimized up to 0.45 V. Wang et al. proposed a vacuum treatment method to establish a 2D/3D stratified heterojunction which can enable charge separation at the interface. Prior to the 2D/3D stratified heterojunction formation, guanidinium thiocyanate (GuaSCN) was introduced as a co‐additive to improve the crystallinity of the 2D perovskite phase. By the synergetic effects of the vacuum treatment and GuaSCN, the respective THP films showed high hole mobility, passivated the surface traps, and prolonged carrier lifetime up to 140 ns (the highest carrier lifetime for TPSCs). The energy level diagram of the pure quasi‐2D perovskite (n = 2) film shows a similar valence band level EV (−5.08 eV) to 3D FASnI3 film (−5.10 eV) while its conduction band level EC (−3.26 eV) is much higher than that of FASnI3 (−3.71 eV), which can prohibit electron transfer across the quasi‐2D/3D mixed phase. With such beneficial aspects, the TPSCs fabricated by this method with PEA2FAn− 1SnnI3n+ 1(n = 10) showed a high VOC of 1.01 V, which minimizes the VOC loss of any type of TPSCs to 0.39 V  was tested for the fabrication of cm2‐sized TPSC by Ran et al. The addition of PPA to FASnI3 could promote the grain size, passivate the grains, reduce trap density, extract photogenerated charges, and induce preferential orientation of the perovskite film. Consequently, a PCE of 7.08% with photovoltaic parameters of VOC of 0.56 V, JSC of 17.57 mA cm−2, and FF of 0.72 . However, it shows a high driving voltage because of the low dielectric constant, when the thickness is reduced, the charge carriers penetrate thin layers of SiO2, resulting in a high gate leakage current, which is a critical fault of device downscaling. Hence, materials with high dielectric constant may replace SiO2 in the future for the downscaling of FETs. The active layer is where charge‐carrier transport takes place and determines the field‐effect conductivity of the devices. For the n‐type active layer, amorphous metal oxides, especially indium gallium zinc oxide (IGZO), have succeeded in commercialization for backplane TFT of the large‐area organic light‐emitting display. For the p‐type active layer, various materials such as silicon (Si) and other inorganic III‐V semiconductors have been developed until now but suffer from the high cost and complexity of fabrication. Although the search for alternative semiconductors beyond inorganics has taken its place in organic materials with low cost, convenience, and scalability of fabrication, its low carrier transport mobility still limits the performance. Recently, MHP stood out as a promising semiconductor candidate for FET active layer due to suitable features, such as high charge‐carrier transport mobility and solution processability  creates a potential difference between gate and source, which is a driving electric force for charge carriers to accumulate at the channel between source and gate electrode in enhancement mode devices. Drain voltage generates a potential difference between source and drain that induces the accumulated charge carriers at the channel to transport from source to drain. Gate and drain voltage are applied simultaneously and form a controllable current between source and drain. Thus, the magnitude of applied gate voltage can change the carrier density participating in transport through the channel, as well as the total current. The magnitude of applied source–drain voltage affects the charge‐carrier concentration gradient at the channel.In a typical FET, a bias voltage is applied to the gate and drain electrodes, whereas the source electrode is grounded. Gate voltage Vg createsVg drives electrons(holes) at the insulator/active layer interface. Some of the induced charge carriers fill in the trap states during the transport, the others are mobile carriers. The carrier density at the channel is related to the magnitude of Vg and capacitance (C) of the insulator. When there are sufficient mobile carriers at the active channel, Vg should be higher than VTh (Vg − VTh > 0). If only Vg is applied without source–drain voltage, the charge carriers at the channel are uniformly distributed. iii) When a small source–drain voltage (VDS << Vg) is applied, a linear gradient of carrier concentration is formed from source to drain electrode. VDS will be increased to the point of VDS = Vg − VTh, where a pinch‐off point exists at the transport channel adjacent to the drain (source) electrode. This is also called the depletion region, at which the difference between Vg and VDS matches VTh. At this point, the electric field is not sufficient to induce additional mobile carriers to accumulate at the channel near the drain electrode. iv) When VDS is further increased (VDS > Vg − VTh), the same effect is enhanced, whereas the transport channel length is reduced. This is also called the saturation region, where the carriers can be driven across the narrow depletion zone and form a space‐charge‐limited current. Further boost of VDS expands the depletion region, which cannot substantially increase source–drain current IDS and saturates at a certain level called ID,sat.The operation of FET with polycrystalline or amorphous semiconducting films can be divided into four main regions, as illustrated in Figure µFE), on/off current ratio, threshold voltage (VTh), and subthreshold swing (SS). µFE can explain the carrier transport capacity and can be extracted from the linear region of the transfer curve. On/off current ratio characterizes the ability to control channel current with Vg and can be calculated by the ratio of IDS at on‐ and off‐state. VTh describes the interface charge traps and SS shows the ability to switch on and off the device.Some of the important parameters for quantifying the FET performance include field‐effect mobility µFE, on/of These attractive properties enabled the development of high‐performance, highly reliable THP‐based FETs over the last two decades, which will be discussed in the next section.As discussed above, FETs are building blocks for modern electronic technologies. Through the discovery of promising semiconductor materials, innovative device structures, and electric circuit applications, FETs have successfully obtained a highly mature level of fabrication. Among the incumbent and emerging semiconductor materials, THPs offer excellent optoelectronic properties, including low effective mass, high carrier mobility, long charge‐carrier lifetime, and defect tolerance character.239 Th5.25.2.1 Because of these advantages, the first‐ever made perovskite transistor was based on 2D THP than any other 3D perovskites or LHPs. Almost two decades ago, Kagan et al. used (PEA)2SnI4 as the channel layer for BGBC FET, as illustrated in Figure  The device showed pure p‐channel characteristics with mobility of µh = 0.62 cm2 V−1 s−1 and a reasonable on/off current ratio of 104 at room temperature 2SnI4 accelerated only in the recent few years encouraged by remarkable advances in solar cells and LEDs. Recently, the performance of 2D THP‐based FETs remarkably developed through precise engineering of film morphology, device structure/interface engineering, crystal formation, and organic spacer synthesis.2D layered THPs are highly stable with long‐chain organic bulky ligand in between the inorganic slabs to suppress ion migration and moisture penetration. Additionally, this family of materials is especially attractive for allowing versatile structure and tunable optoelectronic properties through compositional engineering of organic cations.242 Bee Figure . This woe Figure . Additio2SnI4 by vacuum deposition on top of octadecyltricholorosilane (OTS)‐treated Si/SiO2 substrates. By optimizing the growth rate and temperature of OTS on substrates, the film morphology was much improved and the fabricated FETs exhibited hole mobility of µh = 0.78 cm2 V−1 s−1 and an on/off current ratio of 105 at room temperature. The same group proposed a significant development in solution‐processed (PEA)2SnI4 FETs by applying a series of optimization processes, including a self‐assembled monolayer (SAM)‐treatment on substrates and the addition of MoOx as a hole‐injection layer between the channel and the electrode 2SnI4‐based FET by increasing L to reduce the contact resistance (Rc) between the perovskite semiconductor and source/drain electrodes while using the same TCTG device configuration. With a negligible contribution of Rc relative to the total resistance, the hole mobility increased and saturated to a value of 26 cm2 V−1 s−1.Matsushima et al. first fabricated (PEA)bstrates.245 Bye Figure .thiophen2‐yl)‐‐5yl) ethan‐1‐ammonium, also known as TT Figure .[2596 Figure . The pere Figure . Their v5.2.3p‐doping. Until the works of Shao et al., there were no reports that used 3D THPs as the semiconducting layer for FETs. The authors successfully reduced the hole density of 3D FASnI3 by incorporating a small amount of 2D (PEA)2SnI4. The addition of (PEA)2SnI4 greatly enhanced the crystallinity and orientation of the 3D perovskite phase, which was critical for FET application 2SnI4.97 Thee Figure . Through4 Figure .2SnI4 into 3D FASnI3 from an alternative perspective. The authors presented a new concept of 2D–3D core–shell structure, where the 3D core is fully isolated by the 2D counterpart. The suggested structure provided two major benefits: independent control of VTh by 2D component and significantly improved grain boundary resistance by the 2D/3D interface. The 2D (PEA)2SnI4 FETs suffer from abnormally saturated on‐current caused by large series resistance at the grain boundaries, whereas 3D THP‐based FET exhibits the excessive carrier density that only permits on‐state 2SnI4‐FASnI3‐based FET showed hole mobility of 25 cm2 V−1 s−1 and an on/off current ratio of 106 2CsSn2I7 crystal was applied as a channel layer to BCBG FET, which exhibited hole mobility of 34 cm2 V−1 s−1 at a low temperature of 77 K.Recently, Kim et al. also reported a study on the incorporation of 2D (PEA)spective.261 The Figure . This 2D6 Figure . Additio3 itself has high hole concentration and hole mobility due to self p‐doping nature of 3D THP perovskite, which is outstanding to be used as a phototransistor. Recently, Liu et al. developed a high‐performance FASnI3‐based phototransistor with a high responsivity of ≈105 A W−1 at a low operating voltage. The authors also incorporated FASnI3/PEDOT:PSS heterojunction due to the enhanced photogate effect from the heterojunction, reaching a maximum responsivity of 2.6 × 106 A W−1.In another perspective, FASnI voltage.263 Th5.2.4p‐type tin perovskite‐based transistors was recently accomplished by our group through developing a 3D CsSnI3‐based channel layer. The fabricated champion device achieved a record performance with hole mobility over 50 cm2 V−1 s−1 and on/off current ratios exceeding 108  anion engineering on the improvement of film morphology and suppression of tin/iodide vacancy, the authors developed high‐performance hysteresis‐free TFTs with a high hole mobility of 20 cm2 V−1 s−1 and on/off ratio exceeding 107 2SnI4 TFT to develop the first all‐perovskite complementary inverter with a high‐gain voltage over 38 63, FASnX3, and CsSnX3) and presented their current issues and perspective opportunities. Thus far, despite the high device performance of the potential of THPs, there has been no report showing the ambient stability for the commercialization level. Therefore, to further improve our standing on device performance and its stability, we must indeed pay attention to overcoming several challenges including i) pure and non‐oxidized initial chemicals and non‐oxidizing solvents, ii) fully understanding of film formation dynamics, iii) precise control of cation and anion defects in the film, iv) suppression of Sn2+ oxidation, v) commercial‐grade device operation stability in the air ambient, and vi) better energy band alignment. The vital difference between THPs and its Pb‐based counterparts originated from the instability of Sn2+ in 3D structure and the tendency to realize VSn. Thus, to aid our discussion, we hence also review the mechanism governing Sn oxidation, as well as the efficient route to reduce Sn4+ and the encapsulation approach to alleviating the poor stability, especially in ambient air. It is understood that the under‐performance of THP‐based devices is due to the formation of VSn; thus, a greater understanding of the Sn oxidation would lead to better‐performing THP‐based devices. To date, numerous techniques have been introduced to reduce such effects, namely, the use of additives to act as reducing agents, control crystallization, partial ion substitution, and reduced dimensionality, and we indeed believe that an important breakthrough will be reported in the near future. Henceforward, finding new, stable, and effective solvents for THP‐based devices, having a deep understanding of the Sn chemistry via film formation, interfacial engineering to better match the energy bands of THPs, and stabilizing Sn2+ through device fabrication procedures are greatly important. Alternatively, the implementation of high‐performing devices using vacuum‐deposited THPs can be another breakthrough to fundamentally escape from the issues of various solubility of precursors and additives or stable and effective solvents. Last, although huge progress has recently been made in THP‐based optoelectronics devices, a deep understanding of the photo‐physics and photo‐chemistry processes is hardly obtained. Many approaches have been introduced to improve the performance of those devices including employing reducing agents, morphological control, compositional engineering, device engineering as well as interfacial engineering, but the impact of these approaches on the structural properties and device performance requires more systematic investigation. This will ultimately improve not only device performance but also device stability, which is critical to commercialization.To conclude, we reviewed the structural property relationship, device performance, and stability of several THPs (MASnIThe authors declare no conflict of interest.


Here we measure MR relaxation times at 50 mT and compare results with commonly used models based on both in vivo and ex vivo measurements.Low‐field  were 327 ± 10 ms, 275 ± 5 ms, and 3695 ± 287 ms, respectively. Corresponding values of T2 were 102 ± 6 ms, 102 ± 6 ms, and 1584 ± 124 ms. T1 times in the calf muscle were measured to be 171 ± 11 ms and were 130 ± 5 ms in subcutaneous and bone marrow lipid. Corresponding T2 times were 39 ± 2 ms in muscle and 90 ± 13 ms in lipid.The average T1 times are much shorter than reported at higher fields and generally lie within the range of different models in the literature. As expected, T2 times are similar to those seen at typical clinical field strengths. Analysis of the relaxation maps indicates that segmentation of white and gray matter based purely on T1 or T2 will be quite challenging at low field given the relatively small difference in relaxation times.For tissues except for CSF, the measured T The general relaxation model has 2 components—free and bound water—undergoing exchange that is rapid compared to the MR measurement time scale. The fast exchange 2‐state model shows that the measured relaxation rate is the weighted average of bound and free water. As discussed in Korb and Bryant,1 value than tissue water.Some of this loss in SNR may, however, be recovered because of favorable relaxation times and fewer limitations on RF power due to the lower specific absorption rate.2 times show a much weaker field strength dependence,2 values at lower fields.2 < T1, the lower the frequency dependence.T1 relaxation times for various tissues as a function of magnetic field strength, albeit with quite a large standard deviation in the predicted values. Rooney et al1 relaxation, albeit with different coefficients than those of Bottomley. Fischer et alWhereas in vivo relaxation times have extensively been studied in both healthy subjects and in many pathologies at clinical field strengths, the available literature at low field strength is much sparser. Bottomley1 and T2 relaxation time maps of the brain and lower leg in healthy subjects on a custom‐built 50 mT permanent magnet‐based MRI scanner. We compare the measured values with those predicted by the models mentioned in the previous paragraph, and also with selected measurements at similar fields. We also discuss the implications for image segmentation at low field.In this work, we acquire in vivo T22.10 orientation intrinsic to cylindrical Halbach arrays.All data were acquired on a custom‐built 50 mT (2.15 MHz) Halbach‐based MRI scanner that was previously described in detail.0 shimming was performed using the standard automatic shimming algorithm included with the spectrometer, which maximizes the peak of the spectrum of a nonselective FID by applying pseudo‐random gradient offsets to 3 linear gradient coils. Line widths  were less than 1 kHz on all volunteers after shimming. The center frequency of the sample was measured prior to every scan.Brain imaging was performed using an elliptical 20‐cm wide, 25‐cm tall, 15‐cm deep spiral‐solenoid head coil2.21 and T2, respectively. Total imaging times were relatively long and could potentially be shortened considerably by use of more efficient sequences such as Look‐Locker,1 time in particular results in a long imaging session. Our imaging protocol is limited to 1 h maximum; thus, separate measurement sessions were required for each relaxation time.Human studies were conducted with the approval of the institutional review board. Conventional inversion‐recovery and multiple spin‐echo sequences were used to map TData were corrected for any frequency drift induced by heating of the magnet during the acquisition by applying appropriate phase shifts of the k‐space data. The center frequency was measured immediately before and after the scan, and the drift is assumed to be spatially homogeneous over the imaging region and linear over time. Values were typically 500 Hz over the ~45 min imaging time. All data were acquired on healthy volunteers aged 27 to 62 years . Informed consent was obtained from all subjects prior to scanning.2.31 maps of the gray matter (GM) and white matter (WM) were reconstructed from 6 whole‐brain 3D inversion recovery scans with a turbo spin echo readout acquired with the following scan parameters: resolution: 2.5 × 2.5 × 5 mm3; TR/TE/TEeff: 1250 ms/13 ms/13 ms (center‐out Cartesian k‐space trajectory); no signal averaging; echo train length: 6; and acquisition bandwidth: 20 kHz. Scans were acquired with inversion times of 50, 100, 150, 200, 300, and 500 ms; total scan duration was around 36 min.T2 maps of the GM and WM were reconstructed from a 3D Carr‐Purcell‐Meiboom‐Gill (CPMG) imaging sequence acquired with the following scan parameters: resolution 2.5 × 2.5 × 5 mm3; TR: 1250 ms; no signal averaging; and acquisition bandwidth: 20 kHz. A total of 10 echoes were acquired with a constant echo spacing of 20 ms; images were reconstructed separately for each of the 10 echo times (TEs).T1 and T2 of CSF due to their much longer values than for WM/GM. For T1: resolution: 2.5 × 2.5 × 5 mm3; TR/TE/TEeff = 12000 ms/11 ms/500 ms (low‐to‐high Cartesian k‐space trajectory); echo train length: 90; and inversion times: 500 ms, 1500 ms, 2000 ms, 2500 ms, 3000 ms, and 4000 ms. Total scan duration for the T1 mapping sequence was around 36 min. Given the relatively coarse spatial resolution, the influence of partial voluming of voxels with CSF and WM/GM could be problematic, which is why a long effective TE using low‐to‐high k‐space coverage (approximately 5 times the T2 of WM/GM) was used to suppress the signal from the WM/GM. T2 maps of the CSF were reconstructed from 7 different turbo spin echo scans: TR = 10,000 ms; echo train length = 60; and TE/TEeff = 15/450, 30/900, 45/1350, 60/1800, 75/2250, 90/2700, 105/3150 ms (low to high Cartesian k‐space coverage). Total acquisition time was around 42 min.Separate scans were used to more accurately map the T2.41 maps: spatial resolution: 2.5 × 2.5 × 5 mm3; TR/TE/TEeff = 850 ms/10 ms/10 ms; echo train length: 4; no signal averaging; imaging bandwidth: 20 kHz. Scans were acquired with 6 different inversion times: 25, 50, 75, 100, 150, and 400 ms. Total scan duration for the T1 mapping sequence was around 36 min. T2 maps: spatial resolution: 2 × 2 × 6 mm3; TR: 800 ms; TE: 10 different TEs 11 to 110 ms in 11 ms steps; and no signal averaging. Scan duration was approximately 28 min.T2.51 mapping the data were fitted to:All images were reconstructed using the Numpy fast Fourier transform implementation, with k‐space filtered using a sine‐bell‐squared filter. Parameter maps were reconstructed from a central slice of the 3D reconstruction on a voxel‐by‐voxel basis using the least squares function in SciPy 1.6.0 running in Python 3.7.3. For Tρ and T1 provided to the least squares function are given by:The initial guesses for 2 mapping, the data were fitted to:For Tρ and T2 are given by:The initial guesses for Gaussian curves were fit to the histogram, as in Ref.  using th31 map acquired from 1 of the healthy volunteers showing T1 contrast between GM and WM as well as lipid. It should be noted that the T1 values obtained in the CSF are inaccurate due to saturation effects, resulting from incomplete longitudinal relaxation due to the short repetition time (1250 ms) relative to the T1 time (around 4 s), which is why a separate scan was performed for accurate quantification. The measured T1 times for each individual are reported in Table 1 times across all 3 subjects are 327 ± 10 ms in the GM and 275 ± 5 ms in the WM.Figure 2 map of the brain of a different volunteer is shown in Figure 2 measured in both the GM and WM across the 3 subjects is 102 ± 6 ms. Note that the T2 of times of CSF are not accurately represented in this data set due to T1 saturation effects.A T1 and T2 times in the CSF using a very long repetition time to allow for full longitudinal relaxation. The measured T1 and T2 times in the CSF for each subject are reported in Table 1 and T2 maps were not acquired in the same subjects). The mean T1 in the CSF is measured to be 3695 ± 287 ms, and the mean T2 is 1584 ± 124 ms.Figure 1 map of 1 of the subjects reconstructed from 6 inversion recovery acquisitions, which are also shown. The mean T1 across all 3 subjects is 171 ± 11 ms in the muscle and 130 ± 5 ms in the lipid. Two areas of much higher T1 values are located at the anterior and posterior tibial arteries and veins and are likely due to flow effects and the relatively short repetition time relative to the expected T1 of blood of around 400 msFor muscle and lipid measurements in the lower leg, Figure 2 map, together with individual images from a subset of the different TEs. The mean T2 across all 3 subjects for muscle is 39 ± 2 ms and 90 ± 13 ms for lipid. There is no measurable difference in the T1 and T2 between subcutaneous and bone marrow lipid. The measured T1 and T2 times in the muscle and lipid for each individual are reported in Table Figure 41 and T2 maps of the brain  and lower leg  at 50 mT. As expected, T1 times are lower at 50 mT than at clinical field strengths; gray and white matter T1s were measured to be 327 ± 10 ms and 275 ± 5 ms, respectively, compared to 1200 ms and 650 ms at 1.5 T1 times of 171 ± 11 ms for muscle and 130 ± 5 ms for lipid were measured. Again, these are much shorter than at high field, with T1 of the muscle reported as 1130 ms and 1420 ms for 1.5T, and 3T and a T1 of 250 ms and a range 380‐450 ms for subcutaneous lipid at 1.5T2 times closely match data reported at different field strengths: this is also to be expected given the relative insensitivity of T2 with respect to B0.In this work, we acquired in vivo T1 relaxation times was performed by Koenig and Brown1 values of ~320 ± 55 ms for GM, ~240 ± 40 ms for WM, ~210 ± 38 ms for skeletal muscle, and ~145 ± 40 ms for lipid. Rooney1 = 0.583(B0)0.382 (WM) and T1 = 0.857(B0)0.376 (GM). These equations correspond to values of 186 ms (WM) and 286 ms (GM).  ; D is the baseline; A is the height of the dispersion step; fc is the inflection frequency; and β′ is the steepness of the dispersion step. Using a sample of 13 WM and 10 GM samples from 4 different brains, they empirically derive values of D , A , fc , and β′ . This gives values of 230 ms for WM and 395 ms for GM.In terms of previous work on the field dependence of relaxation times, fundmental work on T1(GM) of 327 ± 10 ms is right at the center of the range from Bottomley, well above the value from Rooney and below the value from Fischer; T1(WM) of 275 ± 5 ms is at the upper end of the range from Bottomley, well above the value from Rooney and above that from Fischer; T1 of 171 ± 11 ms is at the lower end of the Bottomley range; and T1(lipid) of 130 ± 5 ms is in the center of the Bottomley range.Placing our experimental results in the context of these different models: T1 values of 200‐220 ms for WM and 185‐200 ms for GM: these values are lower than our measured ones, in line with the lower magnetic field used. Dean et al1 values was 96‐133 ms, which is shorter than those we measured, as expected.Comparing to previous in vivo low field studies, Agartz et al2 values are generally assumed to be independent of frequency, with Bottomley reporting values of 101 ±13 ms for GM, 92 ± 20 ms for WM, 47 ± 6 ms for skeletal muscle, and 84 ± 36 ms for lipid. Our measured values in vivo agree well with these predictions. T2 values of 92‐98 ms for WM and 81‐87 ms for GM were also reported in the study by Agartz et al at 20 mT, which are very similar to our results. T2 values for lipid in breast at 20 mT were 57‐74 ms in the work by Dean, which is slightly lower than our values.T1 values1 time for CSF shows no significant B0 dependence, with a value of around 4400 ms from 0.15T to 7T.1 of CSF (measured at 1.5 and 3T) was significantly greater than that of pure water at room temperature. They ascribed this difference to be mainly due to the higher temperature in the body. Tsukiashi1 of water increases from approximately 3.2 s at 25°C to 4 s at body temperature, whereas the T2 is much less temperature‐dependent, rising from 2.1 to 2.2 s. Qin2 of CSF, with values at 3T of greater than 2000 ms for ventricle, but lower values (~1600 ms) for frontal cortex and ~1500 ms for both temporal cortex and occipital cortex. The paper attributed the lower T2 values measured in the subarachnoid spaces to be caused by a higher partial pressure of oxygen, which would correlate with the lower T1 values measured by Zaharchuk et al.2 values of ~1400 s were measured by Spijkerman et al at 7T.1 and 1760 for the T2 at 0.15T. Placing our results in context, the average T1 value was ~3700 ms and T2 ~1550 ms, in line with the various literature references.The relaxation times of CSF are notoriously difficult to measure in vivo, and a wide variety of values have been reported in the literature, in particular for the T1 and T2 relaxation times. In all the data, there was some evidence of multi‐exponential behavior, although we believe that the source of this is the partial volume effect rather than the tissue itself.The major challenges in accurate relaxation time mapping at low field are partial volume effects resulting from the relatively coarse spatial resolution and the low intrinsic SNR. In this work, we used very conventional and somewhat time‐inefficient inversion recovery and multiple spin‐echo sequences. We were only able to measure 1 parameter per imaging session of ~45 min, with separate measurements also necessary for CSF due to its very long T510 inhomogeneity is known to cause an underestimation of the T2 times due to diffusion effects,Practical relaxation time mapping for image segmentation purposes in longitudinal studies, or characterization of pathologies in patients, obviously will require much faster scanning. Reconstruction techniques such as compressive sensing from undersampled k‐space data have already been shown to be applicable to low‐field imaging.1 and T2 times measured in this work show that generating inter‐tissue contrast based purely on relaxation times will be more challenging at low field than at higher fields. The absolute difference in T1 relaxation times between tissues becomes smaller at low field, which means conventional methods for generating T1 weighting in images, such as inversion recovery‐based sequences and imaging with a short repetition time, may be suboptimal. It may be that other contrast‐generating techniques such as magnetization transfer1 relaxation times , and the very small chemical shift of less than 10 Hz makes spectrally selective‐based approaches such as spectral presaturation with inversion recovery (SPIR) and DIXON impractical.The T1 and T2 times are also known to correlate with disease progressionFinally, in this preliminary work, we did not differentiate between age and gender, both of which have been shown to affect relaxation times.


Susceptibility maps are usually derived from local magnetic field estimations by minimizing a functional composed of a data consistency term and a regularization term. The data‐consistency term measures the difference between the desired solution and the measured data using typically the L2‐norm. It has been proposed to replace this L2‐norm with the L1‐norm, due to its robustness to outliers and reduction of streaking artifacts arising from highly noisy or strongly perturbed regions. However, in regions with high SNR, the L1‐norm yields a suboptimal denoising performance. In this work, we present a hybrid data fidelity approach that uses the L1‐norm and subsequently the L2‐norm to exploit the strengths of both norms.We developed a hybrid data fidelity term approach for QSM (HD‐QSM) based on linear susceptibility inversion methods, with total variation regularization. Each functional is solved with ADMM. The HD‐QSM approach is a two‐stage method that first finds a fast solution of the L1‐norm functional and then uses this solution to initialize the L2‐norm functional. In both norms we included spatially variable weights that improve the quality of the reconstructions.The HD‐QSM approach produced good quantitative reconstructions in terms of structural definition, noise reduction, and avoiding streaking artifacts comparable with nonlinear methods, but with higher computational efficiency. Reconstructions performed with this method achieved first place at the lowest RMS error category in stage 1 of the 2019 QSM Reconstruction Challenge.The proposed method allows robust and accurate QSM reconstructions, obtaining superior performance to state‐of‐the‐art methods. Optimization models minimize a functional, usually composed of two terms: a data consistency and a regularization term. The regularization term is used to include prior information about the solution, which promotes desired characteristics, such as smoothness or continuous solutions . The squared L2‐norm is a mathematically and computationally efficient function that also defines a convex penalty function that gives a unique solution. For QSM, this approach performs relatively well with data that have been corrupted by moderate amount of noise. However, the squared L2‐norm heavily penalizes large discrepancies produced by strong noise or other sources of discrepancy such as preprocessing artifacts. This high penalty tends to produce susceptibility maps with streaking artifacts, especially in low‐SNR areas. This behavior might be explained from a Bayesian point of view. Finding the solution that minimizes the L2‐norm with noise‐corrupted measured data is equivalent to finding the maximum likelihood estimate in a maximum a posteriori probability problem, but this only happens when the noise source has a Gaussian distribution. Indeed, whereas the noise distribution in phase MRI signals with high SNR can be approximated as Gaussian, this approximation is no longer valid for low SNR.To address this problem, Liu et alIn this paper we present a hybrid data fidelity term approach for QSM (HD‐QSM). This dipole inversion algorithm sequentially uses linear L1‐norms and L2‐norms for data consistency. The resulting algorithm successfully combines the strengths of both norms. The HD‐QSM approach participated at the 2019 QSM Reconstruction Challenge ,22.1‖⋅‖1 is the L1‐norm; F is the Fourier transform with its inverse FH; D=γH0TE13−kz2k2 is the dipole kernel; ϕ is the local phase map; χ1 is the susceptibility distribution obtained in this first stage; TV(·) is the total variation regularizerλ1L1 is the regularization weight used in the first stage. The value of w is a region‐of‐interest binary mask or a magnitude‐based weight, as follows:N is the number of echoes; Magi is the magnitude of echo i; and TEi is the ith TE. The idea behind the weight is to improve the SNR to equal that of the local field, using magnitude‐weighted least‐squares magnitude fitting similar to echo combination.The proposed method HD‐QSM consists of two stages. The first stage finds a suitable initial solution that is robust to streaking artifacts.χ1) as the initialization:‖⋅‖2 is the L2‐norm. The value of W is a spatially variable weight modulated by the voxel‐wise phase discrepancy factor between the solution obtained in the first stage (convolved by the dipole kernel) and the acquired local phase, as follows:The second stage solves the following linear functional, using the solution of the previous stage  to generate strong phase inconsistencies. These phase simulations were used as input for the QSM reconstructions of L1, L2, and L1L2 methods to compare their performances. Optimal reconstructions were obtained for each method by optimizing the normalized RMS error (NRMSE). We performed a sensitivity analysis evaluating the quality of the reconstructions around the optimal regularization parameter λ* within a range defined by 0.1·λ*,10·λ*, sampled at λi=λ*·10i30, with i=.We used the COSMOS 3.2× 256 × 256 to prevent large‐scale aliasing and other artifacts. All reconstructions were stopped when they reached 300 iterations, and the reconstruction parameters were optimized to minimize NRMSE. For each optimal NRMSE reconstructions, we computed the error metrics used in RC2, namelyIn the context of the 2019 QSM challenge,3.31 = 4.92 ms, ΔTE = 4.92 ms, TR = 35 ms, flip angle = 15°, matrix = 232 × 288 × 64 with 0.8 × 0.8 × 2 mm3 voxel size, and acquisition time = 4:51 min. Phase unwrapping was performed with SEGUE ,We performed an in vivo acquisition on a Siemens 3T scanner  with a 12‐channel phased‐array head coil. We used a gradient‐echo sequence with six echoes of a patient showing extensive brain hemorrhage with the following sequence parameters: TEMR images from patients with brain hemorrhage were selected retrospectively. The study was approved by the IRB and the subjects gave informed consent.44.1The L1L2 reconstruction obtained the best performance for medium and low SNRs Figure , whereasAlthough L2 reconstructions yielded streaking artifacts in the presence of phase inconsistencies, both L1 and L1L2 could successfully suppress those artifacts  reconstruction achieved the best performance in most of the analyzed metrics, especially when considering RMS error–based metrics (Table i1) was 40 and 280 for L1L2 and L1L2wH, respectively. For different i1, the optimized reconstructions of L1L2 and L1L2wH showed differences less than 1% in NRMSE  were smaller than those respectively obtained for the L1 and L2 methods . Initialization based on solutions that do not require parameter tuning can also be explored , in which case this model would serve as a refinement step.6The HD‐QSM method combines the beneficial features of the L1‐norm and L2‐norm to obtain high‐quality QSM reconstructions , while maintaining the computational complexity of a linear method. We also proposed a simple and effective heuristic that reduces fine‐tuning to only one parameter to achieve optimal performance. The HD‐QSM method demonstrated exquisite numerical performance in the QSM challenge 2.0 and in pathological MRI data sets with structural abnormalities and conspicuous features.Table S1. Local measurements  of evaluation areas for the COSMOS –based phantomFigure S1. Normalized RMS error (NRMSE) evolution of Sim1 optimal reconstructions for different i1. The error difference obtained between the free parameter method and the proposed heuristic is less than 1 point for all i1. The difference between the best and the worst reconstruction is less than 1 point, indicating that the number of iterations of L1‐norm is not an extremely determinant factor, which confirms the hypothesis that the stage of an L1‐norm solution is a better starting point than 0Figure S2. The NRMSE evolution of Sim2 optimal reconstructions for different i1Figure S3. Solutions at the end of stage 1 and stage 2 for Sim1 and Sim2. The solutions at the end of stage 1 show a noisy appearance with streaking artifacts , but with good structural definition. The final solutions maintain the structural details but do not show the noise and streaking artifactsFigure S4. The first column presents the solution at the end of stage 1; the second column shows the discrepancy factor, which weights the data‐consistency weight; the third column shows the solution of stage 2 using the adjustment factor; and the fourth column shows the solution without using the adjustment factor with the same parametersFigure S5. The NRMSE‐optimized solutions of the hybrid data fidelity term approach for QSM (HD‐QSM) without the discrepancy factor. The first column presents the solution at the end of stage 1, and the second column shows the final solutionFigure S6. Optimal NRMSE reconstructions of Sim 1. For the search of the optimum of L1L2wH, a search was performed in a vector space of 5 × 5 × 5 × 5 λ1L1,μ1L1,λ1L2,μ1L2; once the optimum of this space was located, a second search was performed in a space of the same size in the vicinity of the optimum. In total, three search processes were performed for each simulation, which equates to 3750 reconstructions, whereas for L1L2 only 50 reconstructions were necessaryFigure S7. Optimal NRMSE reconstructions of Sim2. To find the optimum of L1L2wH, a search was performed in a vector space of 5 × 5 × 5 × 5 λ1L1,μ1L1,λ1L2,μ1L2. Once the optimum of this space was located, a second search was performed in a space of the same size in the vicinity of the optimum. A total of three search processes were performed for each simulation, which equates to 3750 reconstructions, whereas for L1L2, only 50 reconstructions were necessaryFigure S8. Additional in vivo reconstruction of a healthy patient. The red arrows indicate a cortical vein. L1L2 and nlL1 were able to correctly reconstruct the veins of the cortexFigure S9. Additional in vivo reconstruction of a patient with brain calcifications. Red arrows indicate the origin of a streaking artifact. The L1L2 reconstruction succeeds in mitigating this spread.Click here for additional data file.


Pancreatic cells, especially the endocrine cells, share similar developmental paths, gene expression profiles, and epigenetic landscapes.The prevalence of diabetes has increased dramatically, with over 537 million adults affected worldwide.1in vivo, and in primary human islets in vitro.Following physiological or pathological stress, β cells lose their mature, healthy characteristics and revert to progenitor‐like cells, which is called cell dedifferentiation.2Pax4,Nkx6.1,Pdx1 and MafAArxPDX1 and MAFA also show conversion of α cells to β cells and a glucose‐lowering effect in the diabetic recipient mice after the transplantation.Pancreatic α cells are the second most abundant cells in islets (about 18% in mice and 38% in humans); the α cell mass is preserved or even increased in diabetes,Ins1, Ins2, Pdx1, Mafa, Ucn3, and Nkx6.1) to the normal β cells, suggesting that the trans‐differentiated β cells are at least partly functional.Notably, defective α cells (particularly in terms of their glucose‐sensing abilities) in diabetes might cause the functional deficits of trans‐differentiated β cells. The previous studies mostly used toxins  to induce β cell loss, and transcriptomic profiles of α cells in these conditions did not seem to be different from normoglycemic control mice.3Pancreatic δ cells, which mainly secrete somatostatin, are located in the islet mantle similar to the α cells. Some somatostatin and insulin co‐expressing cells can be found in normal individuals and patients with type 2 diabetes,Recently, Gribben et al.4PDX1 and MAFA can be reprogrammed to insulin‐producing cells in vitro.Pancreatic PP cells and β cells are thought to be derived from a common precursor.5Collectively, pancreatic β cells can regenerate from cell phenotype conversion among the pancreatic endocrine cells. As for the conversion path, re‐differentiation of the dedifferentiated β cells takes the spotlight because of its accessibility to mature β cells. Transdifferentiation of α cells to β cells attracts lots of attention because of larger cell mass and compensatory hyperplasia in α cells under diabetic conditions.in vivo.There are several concerns that need to be addressed. First, different lineage‐tracing models may be associated with varying levels of accuracy and efficiency of the transgenic animals, thereby leading to variable results.The authors declare no conflict of interest. Professor Rui Wei and Tianpei Hong are members of Chronic Diseases and Translational Medicine editorial board and are not involved in the peer review process of this article.


Supplementary Material. Video files 5 to 8 with MOV file extension were updated to 1 to 4 with AVI file extension. The MOV files 5 to 8 are now removed from the Supplementary Material.In the published article, there was an error in Materials and Methods” section, sub-section “ACE2-S1-RBD Molecular Dynamics Simulation Trajectory Analysis”, paragraph 1. This sentence previously stated:In the published article, there was also an error in the ““Representative trajectory movies of the 10 ns simulations were prepared from 5,000 trajectory snapshots (50 snapshots/ns) generated using VMD Movie Maker Tool  and compiled using Fiji distribution of ImageJ software  at a frame rate of 300 fps.”The corrected sentence appears below:“Representative trajectory movies of the 100 ns simulations were prepared from 500 trajectory snapshots (5 snapshots/ns) generated using VMD Movie Maker Tool  and compiled using Fiji distribution of ImageJ software  at a frame rate of 60 fps.”The authors apologize for these errors and state that this does not change the scientific conclusions of the article in any way. The original article has been updated.