This study investigated the feasibility of using either 23Na‐MRI or 1H‐MRI to calibrate the 13C excitation.Calibration of hyperpolarized 13C‐coils were used to estimate the transmit gain and center frequency for 13C and 23Na resonances. Simulations of the transmit B1 profile of a Helmholtz loop were performed. Noise correlation was measured for both nuclei. A retrospective analysis of human data assessing the use of the 1H resonance to predict [1‐13C]pyruvate center frequency was also performed. In vivo experiments were undertaken in the lower limbs of 6 pigs following injection of hyperpolarized 13C‐pyruvate.Commercial 23Na and [1‐13C]pyruvate was reproducible, with a mean scale factor of 1.05179 ± 0.00001 and 10.4 ± 0.2 dB, respectively. Utilizing the 1H water peak, it was possible to retrospectively predict the 13C‐pyruvate center frequency with a standard deviation of only 11 Hz sufficient for spectral–spatial excitation‐based studies.The difference in center frequencies and transmit gain between tissue 23Na and 1H resonances to calibrate the 13C transmit B1 using commercially available 13C‐coils. The method provides a simple approach for in vivo calibration and could improve clinical workflow.We demonstrate the feasibility of using the  Calibration is usually undertaken prior to the detection of the transient hyperpolarized 13C‐pyruvate signal in vivo. A number of solutions to this problem have been proposed, including the use of a phantom containing a high concentration of 13C‐enriched molecules for calibration, as well as phantom acquisition prior to patient scanning.B1 field over the whole FOV, which is particularly relevant because variations in the delivered flip angle are expected throughout the patient. Furthermore, the position and composition of the phantom increase the uncertainty of centre frequency estimation.Hyperpolarized equency f0, and es13C]pyruvate. Deviations in the flip angle across the FOV can introduce significant uncertainty in the kinetic parameters that are derived using this spectral–spatial imaging approach.13C renders these challenges difficult to solve without the use of a phantom. An interesting calibrationless approach has been reported that uses the pyruvate signal to estimate the sensitivity map for in vivo studies.A selective excitation approach is frequently employed to acquire high signal to noise ratio images of the metabolic products of pyruvate from this frequency measured in vivo was tabulated, and the mean offset determined from all the subjects was applied in each individual.Water Spectra were analyzed in MatLab to measure the actual frequency of both 2.723Na and 1H prescan was performed at site A. Six female, Danish domestic pigs weighing ∼30 kg and fasted overnight were imaged. The pigs received intravenous propofol  and intravenous fentanyl (8 μg/kg/h) and were mechanically ventilated. Catheterization was performed through the femoral veins and arteries for the administration of hyperpolarized [1‐13C]pyruvate and measurement of arterial blood pressure, respectively.A prospective study assessing the use of a 13C‐receive loop was positioned centrally on the biceps femoris muscle of the right lower limb. The receive coil and pig were placed in the central transmit field of the 13C‐transmit coil. Bloch‐Siegert experiments were performed prior to the 13C‐pyruvate injection for both the 23Na calibration acquisitions, with 1H center frequency taken from the previous series, and the 13C prescan acquired with 13C‐lactate phantom. Frequency scaling between 1H/23Na and 1H/13C pyruvate was calculated from Equations 13C‐pyruvate was injected in the left femoral vein over 10 s, followed by a 20 mL saline flush. 13C magnetic resonance spectroscopy was performed at the commencement of the 13C‐pyruvate injection using the following parameters: partially self‐refocusing sinc excitation,23 RF bandwidth = 2000 Hz, spectral acquisition bandwidth = 5000 Hz, number of samples = 2048, repetition time = 1 s, time points = 128, flip angle = 12°, and slice thickness = 40 mm.The pigs were imaged in a supine position, and the 2.8kPL) were calculated by solving the differential form of the modified Bloch equations in the time and frequency domains.13C]lactate:[1‐13C]pyruvate and [13C]bicarbonate:[1‐13C]pyruvate ratios were calculated. Ratiometric and kinetic data were averaged over all subjects.Hyperpolarized spectra were fit using a matching pursuit algorithm, and apparent kinetic rate constants for the exchange of pyruvate to lactate kPL were c2.9P < .05. The mean center frequency difference between 13C‐urea and 23Na/1H at site A was determined by assessing the difference in center frequency between each paired 13C and 23Na and the 1H measurement and averaging the results. Using in vivo data from site A, the ratiometric frequency difference between hyperpolarized 13C‐pyruvate and 1H/23Na across all porcine experiments was averaged.Statistical significance was defined as Differences in RF gain between different saline loading states were assessed by fitting a linear model to the RF gain data.Due to the small number of subjects used in this study, differences between time and frequency domain‐derived kinetic parameters for the in vivo data were assessed using a Mann‐Whitney U test.33.123Na (NaCl) and 13C (13C‐urea) f0 at site A was 1664497 ± 12 Hz (mean ± standard deviation [SD]) with a mean scaling factor of 1.05180 ± 0.00001, as derived from Equation 23Na and 13C RF gain was 10.4 ± 0.6 dB, as derived from Equation 2 < 0.1 and P > .05 in all cases; loop and paddle coil results shown in Figure 13C and 23Na per coil experiment. Inspection of the calibration files for each coil demonstrated that there was a 1.35 dB difference in the additional requested system RF gain for the 13C loop and 8‐channel paddles. The average RF gain for each coil setup is shown in Table The difference in 3.213C and 23Na B0 maps showed a mean shift of +1 ± 18 Hz and −6 ± 23 Hz over the 6 × 6 CSI grid, respectively. 13C and 23Na B1 maps showed a similar RF distribution over the FOV . The absolute average difference between the B1 maps , as demonstrated in Figure Noise correlation between the 8 channels in the paddle coil revealed good linearity between coils, with little correlation between channels from the off‐diagonal elements Figure . However3.53.5.1The site B frequencies of both pyruvate in vivo and urea in the phantom drifted slowly downward over the 31‐month timescale of the experiments Figure . The rat3.5.213C‐pyruvate in vivo and 13C‐urea in the phantom varied over a wide range (186 ± 43 Hz). The frequency separation was not correlated with date of acquisition (R2 = 0.05), suggesting it was unlikely to be due to phantom degradation.The difference between the frequencies of 3.5.31H frequency of water scaled to the 13C frequency by the respective gyromagnetic ratios, as described above, was on average 2620 ± 11 Hz higher than the pyruvate frequency in vivo.The 1H frequency performed better, with a SD of only 11 Hz, and there were no subjects in whom spectral–spatial excitation would have failed. All results are shown in Figure Frequency prediction using the temporal drift or by referencing to the external phantom both performed poorly Figure : the SD 3.623Na and 1H prescan, hyperpolarized 13C‐magnetic resonance spectroscopy was successfully performed from the musculature of 6 pigs following the injection of hyperpolarized 13C‐pyruvate signal from 13C‐pyruvate, 13C‐lactate, and 13C‐bicarbonate observed in all subjects. The frequency scaling between 23Na/13C and 13C/1H calculated using Equations 13C and 23Na RF gain was 156 ± 2 and 260 ± 3, respectively. Kinetic analysis revealed a mean kPL of 0.007 ± 0.002 and 0.007 ± 0.002 s−1 in the frequency and time domains, respectively, with no significant difference between methods (P > .05). The correlation (R2) between kPL in the frequency domain and lactate:pyruvate ratio was 0.9 with P = .01.Prospectively utilizing the 13C]lactate:[1‐13C]pyruvate (LAC/PYR) and [13C]bicarbonate:[1‐13C]pyruvate (BIC/PYR) ratios were 0.15 ± 0.04 and 0.006 ± 0.004, respectively.The mean [1‐413C‐MRI is a powerful technique to noninvasively probe tissue metabolism in many normal and diseased tissues, such as brain, heart, kidneys, and cancer.13C‐lactate in prostate cancer as well as in brain and breast tumors. Clinical hyperpolarized 13C studies face a number of technical challenges, including accurate frequency calibration and the requirement for system inhomogeneity correction. This is particularly problematic due to the low natural abundance of 13C in tissue, which results in poor signal from endogenous 13C nuclei, and which in turn renders the determination of the center frequency and TG difficult. Therefore, new approaches are needed to increase confidence in 13C prescan results, estimation of RF gain, transmission of B1 distribution, and other coil sensitivity parameters.Hyperpolarized 13C frequency can be accurately estimated using the water 1H frequency, which performed better than 2 alternative approaches that we explored: frequency drift and using an external 13C‐urea phantom. In addition, we have recently shown the feasibility of using a 13C‐tuned coil to image the distribution of 23Na in tissue due to the high natural abundance of the latter and the small frequency difference between the 2 nuclei.23Na signal could also be used to improve 13C‐MRI acquisition. The results show that the RF gain and f0 for 13C can be estimated using the endogenous 23Na resonance acquired using dedicated, commercially available and clinical 13C‐coils.Here, we show that the center 23Na resonance to calibrate the 13C experiments: the short T2 and T1 of the nucleus allow rapid data averaging to increase signal to noise ratio and thus increase the confidence in the calibration parameters. Conversely, the long T1 of 13C‐labeled molecules  can lead to lengthy acquisition times for calibration, which further compound the problems of the low natural abundance of the nucleus.There are additional benefits in utilizing the 1H and 23Na resonances to provide a robust prescan for hyperpolarized 13C experiments, with center frequency and RF gain, for single and multichannel clinical coils. Simulations showed consistent overflipping at 13C frequency by the clamshell, which was also experimentally observed. The simulated field 23Na frequency was highly homogeneous, which was also experimentally observed in the B1 maps. Therefore, this provides a commercially available solution to the prescan and calibration required for 13C‐MRI and could assist in simplifying hyperpolarized 13C‐MRI acquisition. In the future, multislice B1 maps with ultrashort echo time sequences could be acquired to complement current human imaging protocols, providing a slice‐by‐slice correction for transmit B1 for further quantitation. However, this will require further experimental optimization  before clinical implementation is achieved.1H and 13C‐pyruvate, it is advisable to utilize the 23Na or 1H prescan before each experiment rather than assuming a static center frequency because selective excitation strategies may otherwise fail.Here, we have demonstrated the potential for using the 13C coils, which significantly increased at the 23Na frequency. This could represent a problem for estimation of the sensitivity maps using array coils, but future 13C coil design could ensure that this effect is minimized. Future studies could assess this concept in human experiments, including a comparison of kPL values before and after correction with a sodium B1 map. Uncertainty in transmit B1 dramatically increases the error in kinetic modeling, which could have significant implications for the interpretation of the results as part of clinical trials.23Na frequency, leading to differences in TG required for a 90‐degree excitation. Reductions in transmit efficiency are expected when operating away from the intended frequency due to degradation in the quality of the tune and match of the coil, although the relatively small difference in TG needed here suggests this degradation is not severe. In the future, the TG could also be used to indirectly inform on the quality of the tune and match. Although we have demonstrated here that 13C coils can be successfully used to measure the 23Na resonance, this approach should be applied with caution by assessing the coil response at both frequencies on the bench and also by undertaking studies in phantoms prior to performing these studies in humans.A challenge to our proposed method is the coupling between multichannel 523Na and 1H resonances to calibrate the 13C transmit B1 using commercially available 13C‐coils. The method provides a simple approach for in vivo calibration and could improve clinical workflow.We have demonstrated the feasibility of using the Dr Rolf Schulte is an employee of GE Healthcare.FIGURE S1 Phantom setup. Saline filled buckets  with the 8‐channel paddle coils (A) or the single loop coil (B) inside the clamshell transmit coil. The black arrow points to the 13C enriched urea phantomFIGURE S2 Individual 1H (A) and x‐nuclei flip angle images for 23Na  and 13C  acquired from an 8 L phantom. 23Na and 13C images are normalized to the largest signal in their 80‐degree imagesClick here for additional data file.


II‐based single‐ion magnets. Using multi‐frequency EPR, susceptibility and magnetization measurements we found that these complexes show some of the highest zero field splittings reported for five‐coordinate CoII complexes to date. The predictable coordination behaviour of the interlocked ligands allowed the magnetic properties of their CoII complexes to be evaluated computationally a priori and our combined experimental and theoretical approach enabled us to rationalize the observed trends. The predictable magnetic behaviour of the rotaxane CoII complexes demonstrates that interlocked ligands offer a new strategy to design metal complexes with interesting functionality.Mechanically chelating ligands have untapped potential for the engineering of metal ion properties. Here we demonstrate this principle in the context of Co II single‐ion magnets illustrates this principle. The predictable coordination behaviour of the rotaxane enables remarkably accurate calculation of their zero‐field splittings, which are in agreement with experimental observations, and can be rationalized in terms of metal–ligand bonding.Mechanically interlocked ligands enable predictable metal ion coordination, offering a new strategy to engineer their properties. A series of field‐induced Co Interlocked molecules(ClO4)2 ](ClO4)2, suggesting that high spin (HS) and low spin (LS) configurations co‐exist in the solid state. The coordination spheres of the de novo HS and LS models of [Co(1)]2+ agree remarkably well, both in terms of geometry and bond lengths , with one of the structures observed by SCXRD  and models of [Co(1)]2+ derived from the corresponding SCXRD geometries , although it should be noted that predicting accurate energy gaps of multiconfigurational complexes is a challenge for DFT modelling. Consistent with this, the EPR spectra of polycrystalline [Co(1)](ClO4)2 show that the complex exhibits both HS and LS configurations in the solid state (vide infra), in line with the EPR data previously obtained on frozen solutions.Two independent structures with different bond lengths  were observed in the asymmetric unit of [Co(1)](ClO4)2, we modelled complexes based on interlocked ligands containing other readily available macrocycle components,2) is more rigid and the other (3) contains a potentially weakly coordinating ether unit near to the bipyridine ligand ]2+ was predicted to display an all‐neutral N, 5‐coordinate environment similar to that of [Co(1)](ClO4)2 but in this case, π‐Co and π‐π interactions result in distortion of the sbpy geometry and a twisting of the macrocycle relative to the axle ]2+, the weakly coordinating ether O was predicted to bind to the metal ion in lowest energy model obtained, resulting in a distorted‐octahedral geometry, although once again the N donors are arranged in a sbpy geometry ]e Figure . In the y Figure .2 and 3 were synthesized from the respective macrocylesII complexes prepared (see the Supporting Information). The asymmetric units of the solid‐state structures obtained by SCXRD of [Co(2)](ClO4)2 and [Co(3)](ClO4)2 contain two very similar structures . The SCXRD sbpy2)](ClO4)2 agrees well with the HS de novo model , with similar π‐Co and π‐π interactions observed in both, confirming that the modelling captures these inter‐component interactions with reasonable accuracy. Similarly, although the calculated and observed octahedral3)](ClO4)2 differ to a greater degree , perhaps in part due to the more disordered nature of the SCXRD structure ]2+ and [Co(3)]2+  and models constructed from the SCXRD coordinates  suggest a larger HS‐LS gap than in the case of [Co(1)]2+, consistent with a single geometry being observed in their SCXRD structures.Rotaxane‐based ligands 1–3)]2+ and their SCXRD‐derived structures confirms that the predictable nature of the coordination environment provided by a mechanically chelating ligand, combined with the ability of simple computational models to accurately capture weak, geometry distorting interactions, allows the structure of such complexes to be predicted with reasonable precision. The computationally predicted values of D obtained for the de novo HS models of [Co(1–3)](ClO4)2 agree remarkably well with models derived from the SCXRD geometries ](ClO4)2.The agreement between de novo models of [Co–3]2+ and1–3)](ClO4)2 were investigated to evaluate their spin configurations and provide an estimate of D. X‐band EPR spectra of polycrystalline samples at 10 K , is visible  in all complexes, the high‐field signals are dependent on the nature of the ligand. The signal at geff=7.9 must arise from a transition within the Ms=±3/2 doublet, characterized by gz >2.5,D is negative (see the Supporting Information). EPR spectra of frozen solutions  were similar to those of polycrystalline samples, showing that no coupling occurs in the solid state, presumably due to the steric bulk of the ligands preventing close approach of the CoII ions.The LS species exhibit similar e Figure  span a vD, high‐field EPR (HFEPR) measurements were carried out on pelletized samples at 4.5 K ](ClO4)2, Figure S24 for [Co(1)](ClO4)2 and [Co(3)](ClO4)2). The absence of additional EPR lines in the HFEPR spectra shows that no inter‐doublet transitions occur up to 375 GHz, setting |D|≥15 cm−1 but preventing accurate determination of D. The large (negative) D value is also responsible for the absence of Ms=±1/2 intra‐doublet signals. The assignment of the geff=7.9 signal to an intra‐Kramer transition is confirmed by the linear trend of the peak positions in the entire experimental frequency range .To determine the magnitude of 1)](ClO4)2 ≫ [Co(3)](ClO4)2 > [Co(2)](ClO4)2, as predicted by modelling, but that it affects the coordination environment of the HS state, in line with the SCXRD data. Importantly, the EPR data demonstrate experimentally that the HS configuration in complexes [Co(1–3)](ClO4)2 exhibits a large negative zero‐field splitting (D<−15 cm−1), in line with predictions for the de novo models.The EPR data show that the ligand environment not only influences the population of the LS state, with [Co(1–3)](ClO4)2 were assessed by direct current (DC) susceptibility and magnetization measurements on the pelletized samples previously measured by HFEPR ](ClO4)2, respectively, are higher than the value (χT=1.88 cm3 K mol−1) expected for S=3/2 systems with g=2 in the spin‐only approximation, but lie in the typical range for five‐coordinated CoII complexes with second‐order SOC.[χT ](ClO4)2 confirms the HS state as the dominant state of these complexes, and the slightly smaller χT measured for 3 compared to 2 is consistent with the higher amount of LS present in this sample 2 e Figure .1)](ClO4)2, the χT vs. T data point to an incomplete spin‐crossover (SCO) behaviour, with a transition from a mixture of states at low T to a progressive conversion to the HS state at T>100 K, suggesting that the LS configuration is lower in energy. Although the modelling suggests that the HS state is favored, it should be noted that the calculated energy gap HS‐LS is small (2.7 kJ mol−1) and was obtained in the gas phase and is thus inconclusive. Although many 4‐ and 6‐coordinate CoII complexes are known to undergo a SCO transition, including reversible switching between SCO and SIM,II SCO compounds are rare, especially in conjunction with SIM behaviour,1)](ClO4)2 only the second SCO complex characterized by a neutral CoN5 coordination.For (ClO4)2, the χT vs. T plots  were simultaneously fitted with PHI.gaverage of 2.33, 2.31 and 2.33 and D values of −78, −59 and −95 cm−1 for [Co(1–3)](ClO4)2, respectively , considerably larger than values obtained for other pentacoordinate mononuclear CoII complexes[1)](ClO4)2, only the fit of the magnetization was used to estimate the parameters (i.e. the χT vs. T plot was not fitted for [Co(1)](ClO4)2 in Figure 1–3)](ClO4)2, respectively (Table 3 K mol−1 for [Co(1–3)](ClO4)2, respectively, see the Supporting Information), all indicate that no J coupling needs to be considered. In contrast to EPR, the inclusion of the rhombic zero‐field parameter E in the magnetic data did not lead to a better fit, (due to the lower sensitivity of SQUID measurements). The large negative D revealed by DC measurements for [Co(1–3)](ClO4)2 2,–(ClO4)2,1–3)](ClO4)2 was measured in the temperature range 1.8–10 K under an oscillating (1–1500 Hz) field of 1.55 Oe (0.155 mT). In the absence of an external static field, none of the complexes displayed maxima in the out‐of‐phase susceptibility. Application of a static field reduces the quantum tunneling of magnetization (QTM) effect by removing the degeneracy between microstates and thus out‐of‐phase maxima for all complexes are observed . These shift to low frequency with the increase in the magnetic field strength, reaching a maximum shift under an applied field of 100–250 mT, considered to be the optimal field under which to observe the slow magnetic relaxation. Strongly temperature and frequency dependent ac susceptibility signals characteristic of single ion magnet (SIM) behaviour were observed in all cases . This characteristic field‐induced SIM behaviour, displayed by all complexes, is exemplified for [Co(2)](ClO4)2 in Figure The frequency dependence of the in‐phase (χ′) and the out‐of‐phase (χ′′) magnetic susceptibility for [Co and their distribution (α) at each temperature for [Co(1–3)](ClO4)2 (Tables S5–7). Reliable fits of the AC magnetic data were obtained for the temperature range between 1.8 and 4.8 K. The α values range between 0.16–0.30, 0.07–0.25 and 0.07–0.25 for [Co(1–3)](ClO4)2, respectively, showing a narrow distribution of the relaxation times. The curvature of the plot of ln(τ) vs. T−1 .where the first term designates the direct process, the second the Raman process and the third the Orbach process; QTM was set to zero given its dependence on D using the experimentally obtained D (224 K (156 cm−1), 170 K (118 cm−1) and 273 K (190 cm−1) for [Co(1–3)](ClO4)2, respectively; all parameters obtained from the fitting are summarized in Table D‐values derived from dc magnetic measurements are quite reliable, resulting in reasonable fits of the data, with minimal contribution from the Orbach process . The direct process was considered in the fitting, due to the dependence of the direct coupling between the |+3/2⟩ and |−3/2⟩ states to H4.(ClO4)2, calculations using the experimental geometries from SCXRD data yielded the splitting of the d‐orbital energy levels ](ClO4)2. We further simplified the model to reflect a purely conformational constraint; the rotaxane framework was replaced by NH3 ligands (and one molecule of H2O in the case of 3), with the ligand‐metal distances and related angles constrained to those obtained by SCXRD. Unsurprisingly, given the equivalent symmetry of their primary coordination sphere ](ClO4)2 > [Co(1)](ClO4)2 > [Co(2)](ClO4)2) and giving values in close to quantitative agreement with the experimental data for the complexes derived from 2 and 3. Even the de novo models gave predicted values of D in good qualitative agreement with those observed  depend on the matrix elements of the angular momentum operator, the effective SOCs ξand the energy gap between the first excited state (Q1) and the ground state (Q0) with Dkl∝-ξ2E(Q1)-E(Q0).(ClO4)2, respectively, it is evident that the latter largely determines D. Further analysis illustrates how the geometrical restrictions enforced by the rotaxane scaffold influences the electronic properties of the metal ion and so modulate D.Computational modelling allows the differences observed experimentally between the complexes to be rationalized. Large negative values of 0 is highly multiconfigurational (see the Supporting Information), the electronic configuration shown in Figure xzd)2(yzd)2(xyd)1(dx2-y2)1(dz2)1, contributes only 50 % to Q0 for [Co(1–2)](ClO4)2. In contrast for [Co(3)](ClO4)2, the electronic configuration (xzd)2(yzd)1(xyd)2(dx2-y2)1(dz2)1 represents 52 % of Q0. Q1 is dominated by the electronic configuration (xzd)1(yzd)2(xyd)2(dx2-y2)1(dz2)1 with contributions of 69 %, 76 % and 60 % for [Co(1–3)](ClO4)2. The xyd orbital thus has a significant population in the excited state and mainly controls the Q1‐Q0 gap. The stability of the dxy orbitals is determined by the intramolecular interactions in the xy plane ](ClO4)2, one of the d lobes is projected along a Co−N bond, increasing the energy of the dxy orbital. Both complexes display similar angles and the energies of their yzd orbitals are similar (2.5 eV). The smaller Q1‐Q0 gap in [Co(1)](ClO4)2 in comparison with complex [Co(2)](ClO4)2 is due to a relatively large contribution of 20 % from the lowest energy configuration to Q1 for the former in contrast with only 7 % for the latter. For complex [Co(3)](ClO4)2, the N‐Co‐N angles are closer to 90° and the dxy orbital does not significantly overlap with any of the Co−N bonds ](ClO4)2 resulting in a significantly smaller Q1‐Q0 gap and a more negative D term.For the three complexes, QD value. For rotaxanes with similar SOCs, imposing geometrical constrains to control the stability of the dxy orbital could be an effective strategy to tune the value of D.Our calculations reveal the interplay between the geometry of the frameworks and the electronic structure in tuning the 1–3)](ClO4)2 are the first examples to show field‐induced SIM behaviour. Our combined EPR and magnetic measurement approach enabled us to determine the magnetic parameters, including the rhombic anisotropy (E/D=0.10–0.22) and HS:LS populations, and the magnitude of D, with confidence. [Co(1–3)](ClO4)2 exhibit remarkably large negative D values compared to other penta‐coordinated cobalt complexes reported (summarized in Table S4) that exhibit SIM behaviour, such as bis(imino)pyridine pincer cobalt complexesD=−28 cm−1) and [Co(phen)(DMSO)Cl2]D≈−17 cm−1). Although [Co(1–3)](ClO4)2 all exhibited slow relaxation of the magnetization in the presence of a magnetic field, there is clearly room for improvement in tuning the structures to achieve SIM behaviour at higher temperature and in the absence of a magnetic field.The results presented support the proposal that interlocked molecules can provide a ligand platform for the development of SIMs and that computational modelling can be used to direct this process. Although hybrid organic‐inorganic rotaxanes have been proposed as qubits,D may be tuned through structural changes in [CoII(tbta)N3]+ complexes by varying the Lewis basicity of the axial ligand on the N3− site.II rotaxane‐based complexes presented here, our calculations suggest a new strategy to tune D, based on the control of the intramolecular angles that determine the stability and population of (in this case) the dxy orbital. We have shown that the HS‐LS energy gap may be estimated even using simple DFT de novo models, although it should be noted that higher levels of theory are required to determine D in these complexes exhibiting multiconfigurational ground states. Using CASSCF, good agreement between calculated and experimentally obtained D values could be reached with the SCXRD‐derived truncated models, demonstrating that the enforced coordination geometry is more important than the exact chemical structure of the ligands. As we have recently demonstrated,We have shown that the zero‐field splittings may be predicted computationally with reasonable accuracy even using de novo models, thanks to the predictable coordination environment provided by the mechanical bond. Magnetostructural correlations with computational methods previously revealed that the size of The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.SupplementaryClick here for additional data file.


Nevertheless, the members of stoichiometric 2DLMs are relatively limited. This renders them incompetent to fulfill the multitudinous scenarios across the breadth of electronic and optoelectronic applications since the characteristics exhibited by a specific material are relatively monotonous and limited. Inspiringly, alloying of 2DLMs can markedly broaden the 2D family through composition modulation and it has ushered a whole new research domain: 2DLM alloy nano‐electronics and nano‐optoelectronics. This review begins with a comprehensive survey on synthetic technologies for the production of 2DLM alloys, which include chemical vapor transport, chemical vapor deposition, pulsed‐laser deposition, and molecular beam epitaxy, spanning their development, as well as, advantages and disadvantages. Then, the up‐to‐date advances of 2DLM alloys in electronic devices are summarized. Subsequently, the up‐to‐date advances of 2DLM alloys in optoelectronic devices are summarized. In the end, the ongoing challenges of this emerging field are highlighted and the future opportunities are envisioned, which aim to navigate the coming exploration and fully exert the pivotal role of 2DLMs toward the next generation of electronic and optoelectronic devices. A comprehensive overview on the synthesis of 2D layered material alloys, as well as, the corresponding electronic and optoelectronic devices is provided. In addition, the ongoing challenges are highlighted and the future opportunities are envisioned, which aim to navigate the coming exploration and fully exert the pivotal role of 2D materials toward the next generation of electronic and optoelectronic devices. In broad terms, 2DLMs represent the emerging atomically thin materials where intramolecular interactions are covalent/ionic forces, whereas the intermolecular interactions are van der Waals (vdW) forces. Inspiringly, the scientific communities have by far identified and produced a broad catalog of 2DLMs beyond graphene, represented by group 15 semiconductors, nitrides, transition metal dichalcogenides (TMDCs), post‐transition metal chalcogenides, halides, multi‐elemental compounds, just to name a few. Benefiting from the prominent electrostatic tunability, dangling‐bond‐free top/bottom interface, intrinsic exemption to short‐channel effects, spatially confined structures, reasonably high carrier mobility, outstanding mechanical strength, and strong interactions with light, these fascinating 2D building blocks have come under the limelight of the nano‐electronics and nano‐optoelectronics realms. Thus far, numerous accomplishments beyond the established norms have been effectuated. For example, in 2019, Zhang et al. demonstrated a field‐effect transistor (FET) with an effective gate length down to 4 nm based on a 2D MoTe2 channel. This device manifests prominent gate‐induced switching characteristic with an on/off ratio of ≈105 and a subthreshold swing of 73 mV dec−1, revealing a broad prospect in the sub‐10 nm complimentary metal‐oxide‐semiconductor (CMOS) circuits. Recently, Maiti et al. have achieved a MoTe2 photodetector with a high responsivity of 0.5 A W−1 in the telecommunication C‐band (1550 nm) by leveraging strain engineering. Specifically, the introduction of ≈4% tensile strain in MoTe2 narrows its bandgap from 1.04 to 0.8 eV, which thus markedly broadens its effective wavelength range.The landmark cleavage of bulk graphite toward one‐atomic‐thick graphene in 2004,1 for  has been identified as a powerful strategy to tailor the electric and optoelectronic attributes in a wide range without breaking the structural integrity. To begin with, alloying semiconductors with distinct charge‐carrier polarities enables the modulation of electric transport properties. In addition, alloying represents a flexible route to tune the energy band structure of semiconductors, which can thus be exploited to tailor the optical properties. For example, as the As content increases from 0% to 100%, the peak position of the photoluminescence (PL) spectrum of the InAsxP1−x alloy continuously evolves from 860 to 3070 nm. Similarly, mercury cadmium telluride alloy with composition‐dependent direct bandgap has been developed for mid‐wave infrared to long‐wave infrared (3–15 µm) photodetection. Such wide‐range modulation of bandgap far exceeds other strategies such as quantum confinement, strain engineering, and regulation of dielectric environment. In this regard, alloying facilitates on‐demand design of exotic semiconductors according to the down‐to‐earth electronic and optoelectronic applications.Despite having been extensively investigated from both experimental and theoretical aspects, the category of stoichiometric 2D layered materials (S2DLMs), namely, 2DLMs with a fixed atomic composition and structure, is yet far less than that of conventional semiconductors. This renders them incompetent to fully satisfy the multitudinous scenarios across the entire breadth of electronic and optoelectronic applications, and it's thereby challenging for 2DLMs to take the place of traditional semiconductors. By virtue of the added degree of freedom in composition, alloying, a strategy widely exploited in conventional semiconductors,44 has1−xWxSe2 alloy with a low W content. In 2019, Zhang et al. predicted that the energy band alignment and energy band offset of the 2D Mo1−xWxY2  alloy based heterojunction can be readily modulated by tailoring the composition of the alloy. In particular, a remarkable photoelectric power conversion efficiency of up to 23.4% can be achieved based on the optimally designed heterostructure. It is noteworthy that a number of 2DLMs share a common stoichiometric formula  and analogous crystal structures. Such high similarities promise a high miscibility, that is, the complete alloying across the whole composition range without compromising the crystalline quality. In this regard, theoretically innumerable 2DLM alloys are available to tie in with variable scenarios in the integrated electronic and optoelectronic applications.Inspiringly, numerous theoretical investigations have prefigured that alloying can markedly expand the library of 2DLMs as well and it will bring a variety of exotic physical properties beyond the fundamental limits of S2DLMs. For example, in 2015, Huang et al. predicted that the deep‐level defect states can be markedly suppressed in 2D Mo2xSe2(1−x) alloy with complete composition tunability (0 ≤ x ≤ 1). In the wake of the prodigious accomplishments in material synthesis, a series of distinctive electronic and optoelectronic attributes of 2DLM alloys have been unveiled. In 2017, Rhodes et al. revealed a fascinating semiconductor to semi‐metal transition in 2D MoSe2xTe2(1−x) alloy with increasing Te content. In another work, Ernandes et al. found that WS0.4Se1.6 alloy exhibited a strong immunity to PL quenching with increasing thickness as compared to the binary WS2 and WSe2. Most recently, Guo et al. have achieved a high‐performance phototransistor based on a MoSe2xTe2−2x thin layer. This device demonstrates a pronounced photoswitching characteristic with an on/off ratio of ≈105. Furthermore, 2DLM alloys have also been exploited for vdW heterostructures. As the Se content increases from 0% to 100%, the band structure of a PbI2/WS2(1−x)Se2x heterojunction will transit from straddling type‐I alignment to staggered type‐II alignment. As a consequence, the PL of WS2(1−x)Se2x can be dramatically modulated from apparently enhanced to greatly quenched as compared to an independent monolayer WS2(1−x)Se2x counterpart. Beyond that, 2DLM alloys have also proven as compelling building blocks toward fundamental physics and novel discoveries. For example, it is found that the valley splitting of CoxMo1−xS2 can be markedly boosted by more than 100% as compared to pristine MoS2. On the whole, alloying can break the intrinsic limitations of S2DLMs and it has opened up unexpected and abundant opportunities to the design and implementation of exotic material properties and device functionalities.Thus far, researchers have made considerable progress in the synthesis of 2DLM alloys as well. For example, Li et al. have reported on a facile one‐step chemical vapor deposition (CVD) growth of 2D MoS To our knowledge, a timely and comprehensive overview on both the production of 2DLM alloys as well as their application in electronic and optoelectronic nanodevices is yet to be provided, despite it is probably a critical step toward navigating further exploration on this nascent field. To this end, this review begins with a comprehensive summary on various methods for the production of 2DLM alloys, including chemical vapor transport (CVT), CVD, pulsed‐laser deposition (PLD), and molecular beam epitaxy (MBE). Then, the application of 2DLM alloys in electronic devices is summarized. Subsequently, the application of 2DLM alloys in optoelectronic devices is summarized. In the end, the ongoing challenges of this rapidly progressing domain are highlighted and the potential schemes to address them are envisioned, which aim to navigate and motivate the future research and fully exert the pivotal role of 2DLMs toward the next generation of electronic and optoelectronic devices.Despite tremendous advancements, the studies on 2DLM alloy electronic and optoelectronic devices are scattered, and the underlying fundamentals are elusive. In addition, the published reviews on 2D devices have been mostly dedicated to S2DLMs.97 To 2 there are few natural alloyed counterparts. Therefore, the production of high‐quality 2DLM alloys is of fundamental importance for the in‐depth exploration of the fundamental physics and device demonstration. However, compared to S2DLMs, the growth of 2DLM alloys is much more technically challenging since the incorporation of heteroatoms may induce lattice distortions, and it will make the growth kinetics much more complicated. Thus far, extensive research enthusiasms have been dedicated to the production of 2DLM alloys for both electronic and optoelectronic applications. Inspiringly, significant accomplishments have been achieved based on a series of growth techniques, including CVT, CVD, PLD, and MBE. Each of these techniques manifests its own advantages and disadvantages, and they will be critically introduced in detail below. It is noteworthy that a series of solution‐processed synthesis methods, such as, hydrothermal reaction, thermolysis, Schlenk line synthesis, and liquid exfoliation, suffer from insurmountable contamination issues , low crystalline quality of products and poor compatibility with the state‐of‐the‐art micro‐fabrication technology, which severely hamper their application in the field of electronic and optoelectronic devices. Therefore, these approaches have not been included in this review.In contrast to the multiple stoichiometric layered materials naturally occurring in minerals,98 the2.1 A typical experimental configuration of CVT is presented in Figure 2(1−x)S2x alloy for example, Sn, S, and Se powders with the nominal atomic molar ratio are used as source materials, and a small amount of I2 powder is used as transport agent. Specifically, mixed Sn, S, Se, and I2 powders are loaded into a quartz ampoule, after which the ampoule is evacuated down to a low pressure and sealed to serve as a closed reaction chamber. Then, the ampoule is horizontally placed in a furnace with a temperature gradient between its two ends, where the source materials are posited at the hot end and the colder end is the deposition zone. In the growth process, the source materials react with the transport agent to form volatile intermediates through a reversible endothermic reaction. Then, these species transport to the cold end, where they decompose and crystallize into monocrystalline SnSe2(1−x)S2x alloy.CVT, a mature and powerful production technique for the crystallization of variable compositions, represents a category of reactions where source materials are generally evaporated/sublimated from a hot side and finally crystallize at a cold side via a sublimation‐transport‐solidification process.106 A 1−xWxS2, Mo1−xNbxSe2, NbxRe1−xSe2, MoSe2xTe2(1−x), WSe2(1−x)Te2x, ZrSxSe2−x, HfSxSe2−x, HfS2(1−x)Te2x, TaSe2−xSx, Sn(SxSe1−x)2, InxSn1−xS2, VxSn1−xS2, PbxSn1−xSe2, Ga1−xInxSe, MnPS3−xSex and Nb1−xTixS3, just to name a few, demonstrating the broad generalizability of this approach. In addition, the composition of the CVT‐derived alloys can be facilely customized in a wide range, since the source materials are sealed in an enclosed environment, which can effectively prevent the unintentional volatilization. For example, Chen et al. have produced a series of monolayer Mo1−xWxS2 alloys with the W content ranging from 0% to 100% from the corresponding CVT‐grown bulk crystals  under high‐temperature conditions, which is adverse to the direct back‐end‐of‐line integration regarding the strict thermal restriction of substrates. As a consequence, the CVT‐derived 2DLM alloys have been largely limited to fundamental studies in the current stage.Nevertheless, CVT has still been stuck with some deficiencies that preclude its widespread application in the field of 2DLM alloy electronic and optoelectronic devices. First, the alloys produced by CVT are generally bulk crystals. Therefore, an additional micromechanical exfoliation procedure, which is not scalable, is required to produce the ultrathin 2D counterparts. It is rather labor‐intensive and time‐consuming, making CVT unavailable for large‐scale application. Importantly, by reducing the amount of metal sources and transport agent whilst increasing the temperature gradient between the reaction and deposition sides, the direct growth of 2DLMs onto targeted substrates has been achieved Figure .[1272.2 Generally, for the synthesis of 2DLM alloys, multiple precursors containing various components are required to be exploited simultaneously. Taking the synthesis of 2D MoS2(1−x)Se2x alloy for example, S and Se powders are used as the chalcogen sources, while MoO3 powder is used as the metal source Se2x monolayers with a lateral size up to 200 µm through a salt‐assisted confined‐space CVD approach.On the whole, CVD also manifests broad generalizability for the growth of 2DLM alloys. For example, Zhou et al. have synthesized 13 kinds of 2DLM alloys by exploiting a single molten‐salt‐assisted CVD approach.134 Ins Figure .[133 This makes it suffer from a considerably large thermal budget and low degree of compatibility with the commercially available flexible substrates such as polymide, polydimethylsiloxane and polyethylene terephthalate, which is the precondition for top‐surface integration. This severely hampers the development of CVD for the construction of wearable devices. Fourth, it is noteworthy that there are a series of parameters, including but not limited to the precursors, substrates, gas atmosphere, gas flow rate, growth temperature, heating/cooling rate, location of precursors/substrates, and heating duration, that play a critical role in the morphology, orientation, composition, thickness, and crystalline quality of products. The parameter space for the CVD growth of a specific 2DLM alloy is rather narrow. Restricted by the rough controllability in the growth dynamics and the nonuniform supply of precursors, the repeatability of CVD is relatively poor and the products are sometimes messy . For example, in the CVD growth of 2D Mo1−xWxS2 alloy by Wang et al., monolayer and multilayer Mo1−xWxS2 nanosheets are randomly distributed throughout the substrates. In addition, contaminations and crystal defects will be inevitably induced by the precursors. For example, the use of metal oxide precursors for the growth of TMDCs may result in the formation of O‐passivated chalcogen vacancies and O interstitials.Third, CVD growth is generally performed under high temperature, which is normally higher than 900 K.148 Th2.3 Basically, a focused pulsed laser beam is directed onto the surface of a target, which bombards the surface layer and ejects plasma plumes containing a myriad of species stoichiometrically proportional to the target in an ultra‐short time scale . These highly active species then transport downstream and adhere to the surface of the targeted substrates, which are normally set to be opposite to the plasma plumes. Finally, they recrystallize, usually in a 2D thin film form, through a complicated process of diffusion, nucleation, crystallization and growth. Since the focused pulsed laser beam possesses high energy flux, the composition of targets can be non‐selectively and synchronously bombarded, which warrants stable and highly controlled growth conditions. In this regard, compared to other growth techniques, PLD holds the intrinsic benefits for the production of multi‐component compounds with complicated phase diagrams, making it a powerful growth technique for synthesizing 2DLM alloys. Actually, 2D Mo0.5W0.5S2, Mo0.5W0.5Se2, Sn1−xAsxSe, Bi2Te2.7Se0.3, and BiInSe3 alloys have been experimentally synthesized through PLD.PLD is a high‐vacuum physical vapor deposition technique that reproduces the composition of targets onto the intended substrates with the assisted of pulsed laser.151 Ba where the lateral size of products is generally limited by the size of substrates. Inspiringly, a series of centimeter‐scale, continuous and uniform 2DLM alloys have been produced via PLD. For example, in 2016, Yao et al. synthesized highly c‐axis oriented multilayer Mo0.5W0.5S2 alloy with uniform composition distribution over a 1.5 × 1.5 cm2 SiO2/Si substrate  and post‐growth treatments are in need of development to address these issues. Recently, Xie et al. have demonstrated that annealing of the PLD‐grown chalcogenides in a chalcogen‐rich atmosphere is a promising route to eliminate the oxides and chalcogen vacancies. Theoretically, this strategy can be generally applicable across various 2DLM platforms, making it a versatile scheme to ameliorate the crystalline quality of 2DLM alloys in the future.In spite of numerous advantages, PLD suffers from relatively high production cost as compared to CVD since a high‐vacuum growth environment is demanded. In addition, PLD‐derived 2DLMs are generally polycrystalline, and there are thus abundant crystallites/grain boundaries.168 Mo2.4−9 mbar), where the highly pure reaction precursors are fed through the thermal sublimation or electron beam evaporation of the corresponding source materials  is sublimated from a standard Knudsen cell. These precursors then transport downstream and attach to the surface of targeted substrates, which are set to be opposite to the sources, finally crystallizing into 2D Mo1−xWxSe2 alloy.MBE is a sophisticated layer‐by‐layer vapor deposition technique proceeded in an ultra‐high vacuum environment 2Te3, (Bi1−xSbx)2Se3, GaTexSe1−x, WSe2−xTex, MoxW1−xSe2, VxMo1−xSe2, and AsxTe1−x alloys.Compared with the above growth techniques, MBE manifests a series of peculiar advantages. To begin with, MBE is easily scalable and versatile for the growth of various 2DLM alloys. Thus far, it has been successfully exploited to synthesize 2D Bi1−xSbx21−xSbx)2Se3 alloy. This makes MBE a promising technique for the fundamental research on the exotic physics of 2DLM alloys. For example, the Fermi level of the above MBE‐grown (Bi1−xSbx)2Se3 topological insulator can be well tailored by precisely tuning the Sb content  of charge carriers, which results in the modulation of the channel conductance. Prior to further discussion, the important performance metrics of FETs are introduced. The on current and off current  represent the drain current under the on and off states, respectively. The on/off ratio refers to the ratio of on current to off current. The field‐effect mobility (μEF) represents the carrier mobility extracted from the transfer curve. It is calculated byL and W are the length and width of the channel, ε0 is the vacuum dielectric constant, εr is the dielectric constant of the dielectric layer, d is the thickness of the dielectric layer, Vds is the source‐drain voltage, Ids is the source‐drain current, and Vg is the back gate voltage. Experimentally, 2DLMs inevitably suffer from crystal defects, which result in unintentional doping. This may degrade the properties of the corresponding 2DLM FETs. Fortunately, alloying is a promising route to tailor the charge carrier concentration by incorporating heterogeneous atoms with different valance state to those of the host 2DLM matrix. For example, in 2018, Liu et al. revealed that the device performances of a SnSe2 FET can be markedly improved by cation substitution with Pb atoms from the same periodic element group to Sn  related with low on/off ratio and high energy consumption. As shown in Figure 0.036Sn0.964Se2 FET reaches ≈106, which is two orders of magnitude higher than that of a pristine SnSe2 FET. Such dramatic improvement is ascribed to the alloying induced modulation of carrier concentration. Due to the lower valence state of Pb as compared to Sn, Pb atoms act as acceptors in the PbxSn1−xSe2 alloy, partially compensating the excessively predominant electrons and thus leading to a decrease of electron concentration with respect to the intrinsically n‐type SnSe2. Specifically, electric transport measurements indicate that the substitution of ≈3.6% Sn atoms with Pb atoms results in a dramatic reduction of electron concentration by 6 times (from 2.08 × 1012 cm−2 to 3.31 × 1011 cm−2). This is also indirectly validated by a conspicuous positive shift of the threshold voltage. As a consequence, the standby current of the Pb0.036Sn0.964Se2 FET is markedly suppressed by two orders of magnitude (≈10−11–10−10 A) and the gate electric field induced switching character thus becomes much more pronounced. In particular, the field effect electron mobility of the SnSe2 and Pb0.036Sn0.964Se2 FETs is comparable (19.4 cm2 V−1 s−1 for the former and 15.7 cm2 V−1 s−1 for the latter). This is related to an intricate tradeoff between the impurity scattering and the electron‐electron scattering. That is, the incorporation of heterogeneous Pb atoms intensifies the impurity scattering, whereas the reduction of electron concentration suppresses the electron‐electron scattering, which conjointly result in roughly unchanged carrier mobility. Therefore, alloying is a powerful means to ameliorate the on/off ratio and energy consumption of 2DLM FETs whilst not compromising the on current (≈8.8 × 10−5 A for the SnSe2 FET and ≈5.9 × 10−5 A for the Pb0.036Sn0.964Se2 FET at a source‐drain voltage of 1 V).FETs represent devices with the output channel current that can be controlled by the electrostatic gating.177 Ba2 is a natively n‐type semiconductor, presumably because of the omnipresent electron‐donating sulfur vacancies, while WSe2 is a natively p‐type semiconductor. As the composition of 2DLM alloys can be markedly tuned over a wide range by virtue of the high miscibility of the host materials with a general stoichiometric formula and analogous crystal structure, it is plausible that the native propensity of electric transport characteristics of 2DLMs can be significantly modulated through alloy engineering.On the other hand, alloying can also regulate the electric transport properties of 2DLMs via counteraction by deliberately coupling host materials with distinct charge‐carrier polarities. Typically, pristine S2DLMs have relatively steadfast electric transport properties. That is, normally only one type of charge polarity tends to be formed for a specific S2DLM. For example, MoSacancies,178 wh2xSe2−2x alloys with variable compositions through a salt‐assisted CVD method and systematically explored their electric transport characteristics. As shown in Figure 2xSe2−2x alloy gradually evolves from an n‐type semiconductor, and then to an ambipolar semiconductor, and finally to a p‐type semiconductor. The dramatic modulation of the electric transport property validates that alloying provides a flexible and facile route for the regulation of the charge‐carrier polarity of 2DLMs. In another study, a transition from n‐type conductance to p‐type conductance is achieved by incorporating Te atoms into MoS2. Inspiringly, a remarkable room‐temperature hole mobility of 182 cm2 V−1 s−1 is realized. Most recently, a bipolar to p‐type transition has been revealed in a VxW1−xSe2 FET as the V content increases from 0% to 1.9%. In addition, this alloying strategy outperforms conventional approaches, such as, surface decoration, in terms of stability, as the heteroatoms in alloys are firmly secured by strong covalent bonding. This is of prodigious significance for the implementation of complicated CMOS circuits based on an individual material system, and it may substantially simplify the preparation of materials toward next‐generation microelectronic circuits.In 2020, Sun et al. synthesized a series of uniform monolayer WS3.2 Basically, a high current flow is generated when the external applied voltage bears the opposite direction to the built‐in electric field, as it can significantly lower the interfacial potential barrier. On the contrary, the current flow is markedly suppressed when the external applied voltage bears the same direction to the built‐in electric field, as it can significantly elevate the interfacial potential barrier. An important performance metric of diodes is the rectification ratio (RR), which is defined as the ratio of forward current to reverse current. It can be expressed asIV and I−V are the current flow under forward voltage and reverse voltage, respectively. By virtue of the dangling‐bond‐free top/bottom surface, 2DLMs can be discretionarily stacked to fabricate various heterostructures without strict lattice matching constrains as those of conventional heterojunctions. In addition, as discussed in the previous section, alloying can modulate the carrier concentration in a substantially broad range or even reverse the charge‐carrier polarity of semiconducting 2DLMs without compromising the layered crystal structures. For example, Liu et al. have demonstrated a p‐type to n‐type transition of the NbxRe1−xSe2 alloy as the Nb content increases from 0% to 12.5%. Therefore, alloying provides a viable option for the preparation of various building blocks toward constructing diodes.Diodes refer to devices with unidirectional electric conduction capability, which is generally originated from the built‐in electric field of heterostructures.187 Ba0.24Ga0.76Se heterojunction ‐type to n(p)‐type transition, semiconductor‐to‐metal transition has also been achieved in a host of 2DLM alloys such as MoSe2xTe2(1−x), WS2(1−x)Te2x, WSe2(1−x)Te2x, and Mo1−xRexSe2, just to name a few. Therefore, these 2D alloys provide prospective material platforms for the construction of Schottky diodes in the future. On the other hand, this strategy is applicable to 2D/3D heterojunctions as well. For example, in a recent work by Chanhan et al., it is found that the rectification ratio of a RexSn1−xSe2/p‐Si diode is ≈3 times higher than that of a pristine SnSe2/p‐Si diode. The ameliorated device property is ascribed to the alloying induced upshift of the Fermi level to conduction band, which renders elevated interfacial build‐in electric field.Most recently, Yu et al. have developed a high‐performance diode based on an InSe/Inn Figure 7a.[192r Figure , thus pr1 Figure . These s Basically, pristine MoS2 is a typical n‐type semiconductor , which accounts for the p‐type character of MoS2−xOx  refers to the generated photocurrent upon illumination with per unit of energy. It is calculated byILight and IDark are the channel current under illumination and in dark, P is incident light's power density, and A is the effective photosensitive area. External quantum efficiency (EQE) reflects the efficiency of incident photons to generate effective photocarriers. It is equal to the number of collected charge carriers by external circuits upon excitation of a single incident photon. It is calculated byh is the Planck constant, c is the velocity of light, λ is incident light's wavelength, and e is the elementary charge. Detectivity (D*) is employed to evaluate a PD's capability to distinguish the weak light signal. It is calculated byPhotodetectors are devices converting optical signals into electric signals, which are critical function units of extensive applications spanning LiFi communication,210 taSince both the optical and electric properties of 2DLMs can be markedly modulated by adjusting their composition, alloy engineering represents a powerful route to tailor the performance of 2DLM photodetectors.4.1 For example, Qiu et al. have identified the existence of numerous S vacancies (≈3.6 × 1013 cm−2) even in the mechanically exfoliated monocrystalline 2D MoS2 nanosheets by exploiting the aberration‐corrected transmission electron microscopy measurements. In addition, these vacancies may accelerate the oxidation of 2DLMs due to the unsaturated bond induced high chemical activity, which introduces a myriad of deep‐level defect states inside the bandgap. These mid‐gap states are scattering centers for free carriers and high‐efficiency recombination centers for photoexcited electron–hole pairs. They are detrimental to the electric and optical properties of 2DLMs. Furthermore, the large surface‐to‐volume ratio makes 2DLMs much more susceptible to these defects as compared to the bulk counterparts. Theory investigation has indicated that the defect state assisted annihilation of exciton in 2D TMDCs is 100–1000 times faster than that without defects. Most recently, first‐principles calculations have predicted a dramatic degradation of carrier mobility by two orders of magnitude as the vacancy concentration of 2D MoS2 increases from 0.1% to 3%. Both of these factors are inconducive to the accumulation and transport of photocarriers.Generally, 2DLMs, especially those synthesized under high reaction temperatures or those having been deposited in ambient environment for a substantially long period, will inevitably suffer from numerous crystal defects induced by thermal damage and air aging. Among various defects, the anion vacancies are the most likely to form regarding the lowest formation energy.218 Fo laser irradiation, and superacid treatment, have by far been developed for defect passivation. However, these treatments are complicated, and the efficacies are limited. In addition, the reliability of these approaches remains another severe issue impeding their practical application. For example, the healing of S vacancies of MoS2 by oxygen molecules can be easily invalidated by either vacuum or light excitation. Furthermore, it is unrealistic to completely eliminate all the crystal defects even with thorough‐paced post‐processing.A series of post‐synthesis approaches, such as adsorption of sulfur adatom clusters,223 la0.5W0.5S2 alloy via PLD and systematically investigated its photoresponse properties. Inspiringly, the 2D Mo0.5W0.5S2 alloy showcases a high Hall electron mobility of 35 cm2 V−1 s−1, which is ≈2 orders of magnitude higher than that of the reported PLD‐grown pristine WS2 and it is even 2 times higher than that of the annealed PLD‐grown WS2. As a consequence, the Mo0.5W0.5S2 photodetector achieves a high responsivity of 5.8 A W−1 as well as, an excellent EQE of 1135%, and the response/recovery time (τrise/τdecay) is less than 150 ms/150 ms  and the CVD‐derived single domain ones . The ameliorated device performances are mainly ascribed to the suppression of deep‐level defect states through alloy engineering, as is explained in detail below  of the defect states VS–W (S vacancy in vicinity to W atoms) of Mo0.5W0.5S2 is much lower than those of WS2, while the electron ionization energy of defect states VS–Mo (S vacancy in vicinity to Mo atoms) of Mo0.5W0.5S2 is close to those of MoS2. It is to be emphasized that the former kind of defect is dominant in the Mo0.5W0.5S2 alloy due to its relatively low formation energy. That is, most defects in Mo0.5W0.5S2 are shallow level defect states. As a consequence, the defect induced scattering and recombination of photocarriers are markedly suppressed, which thus gives rise to improved response rate and photosensitivity. The suppression of deep‐level defect states has also been validated by the increased quantum yield of PL of Mo0.5W0.5S2 as compared to those of pristine MoS2 and WS2 in another independent study by Bogaert and his co‐authors . In another work, Mo et al. found that the response time of a Mo1−xSnxS2 photodetector can be markedly shortened with increasing Sn content. Specifically, the response/recovery times of a Mo0.74Sn0.26S2 photodetector (20/23 ms) are 2–3 orders of magnitude shorter than those of a pristine MoS2 photodetector (2.12/0.62 s). Most recently, Kang et al. have demonstrated that the photoresponse of ReS2xSe2(1−x) alloy is also markedly modulated as compared to that of pristine ReS2 and ReSe2. Under 532 nm illumination, the optimized responsivity of a ReS0.92Se1.08 photodetector reaches 0.25 A W−1, which is several orders of magnitude higher than that of a pristine ReS2 photodetector (3 mA W−1), and it is also slightly higher than that of a pristine ReSe2 photodetector (0.14 A W−1). In addition, as the content of Se increases, the response time first decreases and then slightly increases, whereas the recovery time first decreases and then remains roughly unchanged. Specifically, at a Se content of 54% , the device simultaneously achieves the shortest response time and recovery time (15 m/15 ms), which are superior to that of pristine ReS2 (60 ms/350 ms) and ReS2 (35 ms/34 ms) photodetectors. Thus far, ameliorated device performances have been extensively demonstrated in numerous 2DLM alloys such as InxSn1−xS2, SnS2−xSex, and MoS2(1−x)Se2x, revealing broad expandability of this approach.Fortunately, alloying sheds light by leveraging energy band engineering. Through the rational design of alloys, the deep‐level defect states of 2D semiconductors can be significantly tailored by tuning their relative energy levels in the bandgap. To this end, in 2016, Yao et al. synthesized centimeter‐scale 2D Mooperties.155 Inw Figure . Accordis Figure .[230y Figure . These p By virtue of the suppression of deep‐level defect states, 2D SnS0.27Se0.75 exhibits the highest carrier mobility among the SnS1−xSex alloys with different S content from 0% to 100%, making it the most compelling candidate for 2D/3D heterojunction photodetectors. Under a self‐powered working mode, the fabricated SnS0.25Se0.75/Si photodetector achieves high responsivity, detectivity and on/off ratio of 377 mA W−1, 1011 Jones (1 Jones = 1 cm Hz1/2 W−1) and 450, respectively. In addition, the response/recovery times are 4.7/3.9 ms, which are several orders of magnitude shorter than those of pristine SnS and SnSe photodetectors. In addition, alloying can also be applied to the heterojunction based MSM photodetectors. Most recently, Hao et al. have fabricated a series of graphene/InSe1−xSx photodetectors  and systematically investigated their photoresponse. Taking advantages of the strong light–matter interactions of InSe1−xSx and high‐efficiency carrier transport of graphene, these devices demonstrate remarkable photosensitivity. In particular, the graphene/InSe0.9S0.1 photodetector achieves an outstanding responsivity of ≈4.9 × 106 A W−1 and a remarkable EQE of 8.7 × 108%, which are much higher than those of the graphene/InSe  and graphene/InSe0.8S0.2  photodetectors , where only incident light with photon energy larger than the bandgap of semiconductor can excite electron–hole pairs. Therefore, the cutoff wavelength (λ) is generally dictated by Broadband photodetectors are of pivotal importance for improving the versatility and integration level of modern optoelectronic systems by virtue of their broad applicability for a wide breadth of working environment.243 Th In recent years, a variety of strategies such as strain engineering, defect engineering, coupling of ferroelectric dielectric layers with 2DLMs, and heterogeneous systems have been developed to broaden the effective wavelength range of 2DLM photodetectors. However, these strategies are either complicated in processing or at the compromise of other device performances such as responsivity, response rate and on/off ratio. For example, Xie et al. recently reported the success of room‐temperature detection of 2.52 THz light illumination based on a Mo‐vacancy‐rich 2D MoS2.19 photosensitive channel. However, this device's responsivity, response/recovery time and on/off ratio are only 10 mA W−1, 5.12 s/6.33 s and ≈1/100, which are far inferior to those of the reported defect‐free MoS2 photodetectors.Since the extensively explored 2D TMDCs and post‐transition metal chalcogenides normally possess sizable bandgaps in the range of 1–2 eV, the effective wavelength range of the corresponding photodetectors is largely limited to the visible to near‐infrared range.245 InxRe1−xS2 alloy photodetector. As shown in Figure xRe1−xS2 alloy as x is below 0.9. As shown in Figure 2 and MoS2 photodetectors exhibit negligible photoresponse to the 980 nm and 1550 nm illumination with photon energy below their bandgaps. In contrast, distinct photoswitching is demonstrated in a MoxRe1−xS2 photodetector upon both 980 and 1550 nm periodic illumination. It is worth emphasizing that in contrast to the defect engineered broadband 2DLM photodetectors with sluggish response rate, the response time of this MoxRe1−xS2 alloy device suffers no degradation as compared to those of the pristine ReS2 and MoS2 photodetectors. Similarly, Ye et al. have revealed the bandgap narrowing phenomenon of HfS2 by Te substitution. Specifically, the bandgap of HfS2(1−x)Te2x alloy decreases prominently from 1.7 to 0.88 eV as the Te content increases from 0 to 0.095. Consequently, the absorption spectrum of HfS1.81Te0.19 is markedly broadened from visible (730 nm) of binary HfS2 to short‐wavelength infrared (≈1400 nm), which covers the free‐space laser communications wavelength and the second near‐infrared region in medicine. Importantly, the corresponding HfS1.81Te0.19 photodetector demonstrates broadband photoresponse from 520 to 1310 nm, far beyond the limitation of pristine HfS2 photodetectors (<1000 nm). In addition, this device also achieves a high responsivity of 2 A W−1. This value is significantly higher than many of the previously reported pristine HfS2 photodetectors (10−2–26.5 mA W−1). Furthermore, a satisfactory response/recovery time of 8.8 ms/75 ms is demonstrated. These values are in the range of the reported results (τrise ≈ 130 µs–55 ms/τdecay ≈ 155 µs/230 ms), and they can be further improved by strategies such as substrate passivation which is conducive to suppressing the interfacial defects.Incorporation of heterogeneous atoms into the pristine 2DLM matrix is a powerful route to introduce mid‐gap states. These intermediate states can participate in the excitation of carriers, which can thus reduce the lowest photon energy required for triggering the transition of electrons. As a result, these mid‐gap states can extend the effective wavelength range of 2DLMs without compromising the electric transport properties, provided little lattice defects, such as strain and distortion, are introduced. Therefore, in theory, alloy engineering represents a potential scheme to implement high‐performance broadband photodetection. In 2020, Deng et al. demonstrated broadband photodetection from 532 to 1550 nm based on a 2D Mo0.83P0.17) as the photosensitive channels. As shown in Figure 0.83P0.17 showcases broadband light absorption from near‐infrared to mid‐infrared with the absorption edge at ≈8.27 µm. Inspiringly, such wide light absorption range is far beyond the limit of both intrinsic black phosphorus and intrinsic black arisen, which is ascribed to the alloying induced bandgap modulation. As a consequence, the b‐As0.83P0.17 photodetector demonstrates room‐temperature broadband photoresponse from 405 nm to 8.05 µm, which is far beyond the limitation of pristine black phosphorus photodetectors (λ < 4.14 µm). Both responsivity and EQE decrease monotonically with increasing wavelength of incident light in the infrared range  and gate‐defined b‐P (0.35 mA W−1). However, they still far lag behind those of commercial devices. Therefore, prior to practical application, strategies need to be developed to further improve the device performance.In addition to TMDCs, alloy engineering has been applied to the mono‐elemental 2DLMs toward broadband photodetection as well. For example, Long et al. have achieved room‐temperature broadband photodetectors by exploiting black arisen phosphorus alloy  phase transition. For example, as the S content decreases from 100% to 0%, the WS2xSe2−2x alloy will gradually evolve from a highly doped p‐type semiconductor, then to a lightly doped p‐type semiconductor, then to a lightly doped n‐type semiconductor, and finally to a highly doped n‐type semiconductor. In particular, the synthesis of 2DLM alloys with different atomic molar ratios can be commonly effectuated with an individual growth technique by simply modulating the growth conditions such as the substrate temperature, precursor temperature, mass ratio of precursors, as well as, positions of precursors and substrates, etc. Furthermore, the heterojunctions of 2DLMs can effectively suppress the generation‐recombination noise by virtue of the dangling‐bond‐free heterointerface. Therefore, alloy engineering provides a fascinating and flexible avenue for the preparation of high‐performance 2DLM based self‐powered heterojunction photodetectors.Self‐powered photodetectors bear the advantages of low standby energy dissipation, high on/off ratio, excellent detectivity, and fast response rate,270 wh0.82Te0.18 heterojunction and systematically evaluated its potency for self‐powered photodetection. As shown in Figure 1−xTex alloy is strongly dependent on the content of Te. As the Te content increases from 0% to 18%, the InSe1−xTex alloy gradually evolves from an n‐type semiconductor  to a p‐type semiconductor . It is noteworthy that the InSe1−xTex alloy with variable compositions can be produced by an individual growth technique, which significantly simplifies the device construction. In addition, in this composition range (0 ≤ x ≤ 0.18), the layered crystal structure of the InSe1−xTex alloy is well maintained. Therefore, the stack of 2D InSe and InSe0.82Te0.18 flacks results in the formation of a high‐quality dangling‐bond‐free p‐n junction . The moderate performance is probably associated with the low interfacial built‐in electric field due to the far from optimized band alignment, which has also been validated by the inapparent rectification characteristic. In addition, this device not only displays a stable and reproducible photoresponse , GaTe/InSe (13.8 mA W−1), InSe/MoS2, and Se/InSe (110 mA W−1). Importantly, the corresponding detectivity reaches 1.55 × 1012 Jones, which is also comparable to those of the above InSe based devices (1.08 × 1010−2.2 × 1012 Jones).To this end, in 2020, Yu et al. developed a dry transfer method to fabricate a 2D p‐n diode based on an InSe/InSeetection.276 Asn Figure . As showe Figure , but alse Figure . Followi3 Figure . In addie Figure . Under a5 Figure . In addiOn the whole, the success of these studies has undoubtedly established that 2DLM alloys based heterostructures can provide abundant platforms for constructing high‐performance self‐powered photodetectors. In spite of the somewhat mediocre device performances in the current stage, which are probably ascribed to the far from optimized energy band structures, these pioneering studies present a novel, facile, and flexible gateway for the realization of self‐powered photodetectors through alloy engineering. In the future, the band alignments of the 2DLM alloys based heterostructures, including the CBM offset, VBM offset, and Fermi level offset, can be further optimized by simply tuning the components of the alloys, which promises a broad degree of freedom for the implementation of advanced self‐driven photoelectric devices.4.4IP is the photocurrent, Δn is the concentration of photocarriers, μ is the carrier mobility, and τ is the carrier lifetime. Thus far, it has been extensively demonstrated that coupling semiconductor quantum dots with 2DLM channels can markedly prolong the lifetime of photocarriers by spatially separating the photogenerated electron–hole pairs, which results in remarkably high photogain. However, this approach suffers fatal challenges, such as slow response rate due to the sluggish detrapping process of photocarriers seized by the quantum dots. By sweeping the photogenerated electrons and holes in opposite directions, the built‐in electric field of heterojunctions and homojunctions, including Schottky junctions and p‐n junctions, can contribute to high‐efficiency separation of photogenerated electron–hole pairs as well. Nevertheless, it also has its own shortcomings. For example, only the photocarriers in vicinity of the heterointerface can be effectively swept and separated. A large amount of photocarriers will recombine before being extracted to the external circuit, especially for the in‐plane heterojunction and homojunction photodetectors. This leads to relatively low responsivity. For example, the responsivity of a lateral MoS2/WS2 photodetector is as low as 4.36 mA W−1. To address this issue, Lu et al. introduced periodic built‐in electric field across the whole 2D photosensitive channel by deliberately coupling a 2D In2S3 flake with a SiO2 nanograting array, where strain generates in the unsupported In2S3. As a result, the In2S3/SiO2 photodetector achieves an excellent responsivity of 1810 A W−1 and a remarkable detectivity of 2.09 × 1015 Jones. Nevertheless, a complicated and high‐cost focused ion beam micromachining process is entailed to fabricate the nanograting, which makes this strategy impractical for large‐scale application.As for the photoconductive type photodetectors, the photoresponse is proportional to the lifetime of the photocarriers, which is dictated by 1−xWxS2 alloys with graded composition by a one‐step CVD approach and systematically explored its photoelectric characteristics  within a single domain is achieved by optimizing the positions of the precursors. As shown in Figure 1−xWxS2 alloy photodetector covers 532 to 1064 nm. Under 532, 671, and 1064 nm illumination, the optimal responsivity reaches 298.4, 224.4, and 28.7 A W−1. These performance metrics are much superior to those of the homogeneous Mo0.5W0.5S2 alloy photodetector (R ≈ 5.8 A W−1) and the compositionally abrupt MoS2/WS2 heterojunction photodetectors (R ≈ 68 mA W−1–2.3 A W−1). In addition, the corresponding detectivity reaches 7.6 × 1011, 5.9 × 1011, and 1.0 × 1011 Jones, respectively. The excellent device performances are mainly attributed to the high‐efficiency separation of photocarriers across the whole photosensitive channel, which is enabled by the graded composition distribution. Specifically, MoS2 and WS2 bear a staggered type‐II band alignment, where the CBM and VBM of WS2 are both higher than those of MoS2, respectively. As a result, the graded element distribution of the Mo1−xWxS2 alloy results in a gradient band offset, where the more the W atoms included, the higher the CBM and VBM are, as shown in Figure  In this study, graded MoS2(1−x)Se2x alloy with Se content increasing along the normal direction is employed as the photosensitive channel. Conforming to the gradient composition induced immanent out‐of‐plane built‐in electric field, which efficiently drives electrons and holes to transport to opposite directions, this device is operable under a self‐powered working mode, with an outstanding responsivity of 311 mA W−1. This value is significantly higher than those of the reported self‐powered photodetectors based on lateral MoS2/MoSe2 heterojunctions (R ≈ 18 mA W−1). In addition, this device achieves a satisfactory detectivity of ≈1011 Jones. Under a source‐drain bias of −0.5 V, the responsivity can be further boosted to 23.2, 191.5, and 26.2 A W−1 when excited by 405, 650, and 808 nm lasers, and the detectivity is at the magnitude of ≈1012 Jones. Moreover, the response/recovery time of the MoS2(1−x)Se2x photodetector is as short as 51 ms/51 ms. On the whole, these values also outperform those of pristine MoS2 , MoSe2  photodetectors, the mosaic MoS2/MoSe2 heterojunction photodetector  and comparable to those of the in‐situ fabricated vertical MoS2/MoSe2 heterojunction photodetector . In general, these experimental findings have established that the composition gradient provides an additional degree of freedom to tailor the photoresponse of 2DLMs‐based photodetectors, without structure abruption as in the case of heterojunctions. Such strategy can ameliorate the photosensitivity of 2DLM photodetectors whilst circumventing the complicated device structures such as the split‐gate configuration as well as the labor‐intensive device fabrication process. By optimizing the bandgap offset and span, the device performances can be further improved.To address these issues, in 2019, Yang et al. synthesized monolayer Mo As shown in Figure 1−xWxS2 alloy that span nearly the full composition range along a single domain is controllably synthesized through a diffusion‐based CVD method, that is, synthesis of defective WS2 skeleton followed by Mo incorporation in the subsequent step.Importantly, Bogaert et al. recently found that point defects played a vital role in the alloying of 2DLMs.230 As2 templates. Despite that only graded Mo1−xWxS2 alloy is synthesized in this work, this strategy theoretically can be generally applicable across the entire breadth of 2DLM platforms. Thus far, strategies such as focused ion beam irradiation, e‐beam irradiation, plasma bombardment, and vacuum annealing have been developed to deliberately create defects in 2DLMs, which can be powerful tools for engineering the defect distribution toward on‐demand design of graded 2DLM alloys. Therefore, persistent efforts are expected to be devoted to the precise customization of the composition profiles of 2DLM alloys for advanced photodetectors in the future.As shown in Figure 5 In addition, through rational design of energy band structure, alloying can effectively suppress deep‐level defect states, which gives rise to both ameliorated photosensitivity and expedited response rate of the corresponding 2DLM photodetectors. In a word, 2DLM alloys have markedly enriched the library of 2D materials and these fascinating members enable substantial property modulation toward on‐demand device design and applications. In this regard, 2DLM alloys promise fertile building blocks potentially complementing the shortcomings of the stoichiometric host materials. However, compared with S2DLMs, the research on 2DLM alloys is still in its infancy. In the wake of the deepening of 2DLM alloy research, a series of new issues head for researchers. Below, we highlight the ongoing challenges in this fast‐evolving domain, accompanied with the future opportunities.As a whole, alloying can markedly expand the technological landscape of 2DLMs by introducing fresh blood with various exceptional features, distinguished from the corresponding host materials, via composition engineering whilst inheriting the communal advantages of vdW materials. In the past several years, 2DLM alloys have witnessed a prosperous development in both fundamental science and device applications. For example, alloys of 2DLMs with opposite charge‐carrier polarities can be exploited to implement ambipolar FETs.180 In5.12−xTex alloy. As a consequence, numerous interspersed semiconducting domains and metallic domains co‐exist in the MBE‐derived 2D WSe2−xTex alloy. The phase segregation and non‐uniform composition distribution seriously preclude the exploration and harness of the intrinsic physical properties of the nominal 2DLM alloys. In this regard, they can render considerable devicetodevice variation, which is inconducive to practical application. Generally, the difficulty of producing high‐quality 2DLM alloys lies in the complexity and inadequate understanding of the phase diagrams of multicomponent layered materials, as well as, the inadequate reactions among multiple compositions. As for the former, theory calculations are efficient and powerful means to acquire the phase diagram and screen the thermodynamically stable 2DLM alloys. Most recently, machine learning has been identified as another effective route for predicting the thermodynamic stability of 2DLMs as well. Therefore, machine learning‐driven design of 2DLM alloys represents another potentially feasible direction of this domain. As for the latter, increasing the synthesis temperature is an effective strategy to improve the composition uniformity as it can provide sufficient energy for interdiffusion and to surmount the reaction barriers. However, it is to be concerned that this strategy may bring additional dilemmas such as increased thermal defect density.Unintentional and unmanageable phase segregation as well as non‐uniform composition distribution are still in the rank of the most pivotal challenges plaguing the preparation and practical application of 2DLM alloys, especially for those consisted of host materials with significantly distinct structural and physical properties. For example, there is a miscibility gap (from 14% to 70% Te concentration) where phase separation easily occurs for MBE growth of 2D WSe5.22 TMDC alloys, while the experimental realization of alloys based on other 2DLMs are relatively rare. This significantly limits the scope of property modulation enabled by alloy engineering. For example, there is still no report on the alloys of TDMCs and transition metal halides, although many of these members share a common stoichiometric formula  and can be plausibly completely alloyed. In this regard, the diversity of 2DLMs and the variation of stoichiometric ratio of the constituent elements theoretically provide infinite schemes for the design and implementation of novel 2DLM alloys. Theoretical analysis is a potent thrust to develop and screen mysterious 2DLM alloys with prospective physical properties, which help to circumvent the unavailing material synthesis and characterization to a certain extent. For example, recently, Hemmat et al. have excavated 25 kinds of 2DLM alloys through the equilibrium temperature‐composition phase diagrams acquired by first‐principles calculations, and they have synthesized 12 of them by employing a single CVT approach.Thus far, the research enthusiasms on 2DLM alloys have been mostly focused on the MX5.3 This is a serious oversight as the underneath physical origins of exceptional findings are fundamental to the design and implementation of future electronic and optoelectronic devices based on 2DLM alloys.Thus far, only a portion of studies on 2DLM alloys based electronic and optoelectronic devices have been explicitly elucidated. Many experimental advancements are still in need of in‐depth analysis.307 Th5.4 Therefore, through rational alloying, photosensitive channels that exactly coincide with all wavelengths from ultraviolet to near‐infrared can be available. On the other hand, non‐layered 2D materials have been successfully synthesized through a series of approaches represented by vdW epitaxy, 2D template assisted topotactic transformation, space‐confined growth method, self‐limited growth method, etc. The incorporation of non‐layered 2D materials for alloying prominently enriches the building blocks toward the construction of 2D semiconductor alloys. For example, layered Bi2Te3, non‐layered In2S3, and In2Te3 share a common stoichiometric formula (M2X3) but distinct crystal structures and physical properties, where Bi2Te3 is a topological insulator whereas In2S3 and In2Te3 are semiconductors. Therefore, the alloys based on them are appealing subjects for the research efforts in the future.Compared with binary and ternary alloys, 2DLM alloys with four or more elements bear more degrees of freedom for property modulation, enabling them theoretically hold more abundant physical properties. For example, the bandgap of ternary phosphorus chalcogenides lies in the range of 1.3–3.5 eV.308 Th5.5 Due to the favorable bandgap, outstanding electrostatic tunability and excellent mechanical properties, 2DLMs have demonstrated grand potential in LEDs as well. Nevertheless, there are still several issues to be addressed prior to the extensive utilization of 2DLMs in LED. The first one is that the active light‐emitting materials should hold a direct bandgap, which enables phonon‐free electron transition, thereby realizing high‐efficiency light emission. In addition, the active materials are anticipated to bear continuous bandgap values to meet the light‐emission requirements of different wavelengths. Moreover, there need to be 2D building blocks with various charge‐carrier polarities to enable the construction of p‐n junctions. Unfortunately, a substantial portion of 2DLMs are indirect bandgap semiconductors. For example, the most commonly studied group VIB TMDCs exhibit direct bandgap only in the monolayer form. Few‐layer and multilayer group VIB TMDCs are indirect bandgap semiconductors, which hinders their application in LEDs. In addition, the bandgap distribution of S2DLMs is relatively discrete, and they thus cannot fully meet the requirements of light emission with diverse wavelengths. Furthermore, most 2DLMs exhibit n‐type conductance, which hinders the design and construction of p‐n junctions. Fortunately, the latest investigations have suggested that alloying can effectively regulate the band structure of 2DLMs. For example, Ernandes et al. found that WS0.4Se1.6 alloy exhibited a strong immunity to PL quenching with increasing thickness as compared to binary WS2 and WSe2. Sainbileg et al. have revealed that Ni substitution on the Sn sites of SnS2 generates new band lines at the bottom of the conduction band due to strong orbital hybridization with states of its nearby sulfur atoms, which results in intriguing indirect‐to‐direct bandgap transition. On the other hand, as demonstrated in previous studies, both bandgap value and charge‐carrier polarity of 2DLMs can be markedly modulated by alloying. Therefore, 2DLM alloys hold enormous potential in LED applications, and further exploration is urgent need in the coming future.Light‐emitting diodes (LEDs) represent devices that can enable photon emission by driving electrons and holes to recombine with the assist of an external electric field.313 DuAs a concluding remark, alloying has served as a variable and flexible avenue to tailor the attributes of 2DLMs with a high degree of tunability, bringing a variety of exotic physical properties and applications beyond the fundamental limits of S2DLMs. Therefore, 2DLM alloys have ushered the new frontier of materials science and they are important contenders for constructing electronic and optoelectronic devices. It is envisioned that after overcoming the above‐mentioned predicaments, 2DLM alloys will be an indispensable complement to the existing S2DLMs and play a vital role in the forthcoming breakthroughs of semiconductor technology nodes.The authors declare no conflict of interest.