Insights into the multifunctional applications of strategically Co doped MoS₂ nanoflakes

Abstract

Simultaneous tuning of magnetic, transport and electrochemical properties through strategic doping of cobalt (Co) ions in hydrothermally treated multi-layered MoS₂ nanoflakes (NFs) without having a secondary phase has been regarded as cutting-edge research. In our study, we have successfully incorporated Co into MoS₂ NFs at various percentages (0%, 2%, 4%, 6%, 8%) with a significant presence of defects, strain and sulfur vacancies, enabling prompt transformation of the surrounding 2H-MoS₂ local lattice into a trigonal (1T-MoS₂) phase. Effective amplification of magnetic property (ferromagnetic coupling on the scale of p_eff ∼ 4.37μ_B) in 8% Co-doped MoS₂ NFs has been evidenced from VSM measurements. The key reasons are probably attributed to the doping induced 1T phase, the presence of zigzag edges well-established from TEM and Raman measurements, and exchange interactions between ferromagnetically ordered sulfur vacancies and Mo⁴⁺ and Co²⁺ ions. The experimental observations on magnetic measurements have been fitted well with the well-known density-functional theory (DFT) computation. Further, the effect of intentional doping on transport property has been evaluated by employing a four probe linear geometry setup. The increased carrier concentration and decreased resistance result in improved transport properties. Various transport models such as variable range hopping (VRH) and nearest neighbour hopping (NNH) of the Co-doped MoS₂ systems have been successfully fitted in different temperature regimes with a tunable temperature coefficient of resistance (TCR) ∼ 3.0 × 10⁻² K⁻¹. Additionally, electrochemical measurements revealed a significant increase in electrochemical activity with the highest proportion of Co doping (8% Co), which is likely due to increased defect levels and active surface area with expanded interlayer separation, as well as exposure of the electrochemically more active metallic (1T phase) Mo atoms in the edge planes. Therefore, our approach in achieving mixed-phase defect-rich (1T and 2H) Co-doped MoS₂ NFs exhibiting room-temperature ferromagnetism, high TCR and improved electrochemical performance makes them an excellent multifunctional candidate in spintronics, infrared (IR) detection and energy storage devices.

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1. Introduction

Molybdenum disulfide (MoS₂), as one typical candidate of graphene analogues and a member of the transition metal dichalcogenides (TMDs), has recently drawn tremendous attention in a wide range of novel applications such as field-effect transistors,¹ flexible devices,² digital electronics,³ and nano- and opto-electronics^1,4 as well as in basic research^5–7 due to its unique electrical, optical, mechanical and electronic properties. Based on the demanding application of diluted magnetic semiconductors (DMS), manipulating the magnetic properties of MoS₂ nanoflakes is crucial for expanding their applications in nanoelectronics and spintronics. However, pristine MoS₂ nanoflakes are intrinsically nonmagnetic. Therefore, the realization of stable magnetism in MoS₂ is highly desirable.

Motivated by the potential applications and the great demand for magnetic MoS₂ nanostructures, a variety of methods have been developed and explored in the past few years. However, there are many effective ways in which the magnetism of 2D materials can be realized such as the introduction of defects,^8,9 surface functionalization (i.e., hydrogenation),¹⁰ cutting 2D nanoflakes into one-dimensional (1D) nanoribbons (NRs),^11,12 and transition-metal (TM) doping.¹³ Moreover enormous ferromagnetism in MoS₂ nanoflakes requires the introduction of a large number of defects which in turn results in the reduction of carrier mobility because of the formation of scattering centres or charge-trapping sites in the MoS₂ sheets. Although surface functionalization is capable of producing stable magnetism, experimentally it faces a big challenge. For example, surface hydrogenation requires the implementation of external stress in MoS₂ nanoflakes. The introduction of ferromagnetism in MoS₂ NRs is very crucial and the edge type of the nanoribbon needs to be well-controlled as zigzag NRs are ferromagnetic whereas armchair NRs are nonmagnetic. Structure and vacancy-related ferromagnetism contributed by zigzag edges of MoS₂ nanoribbons and sulfur vacancies in MoS₂ flakes has been reported in several theoretical calculations.^11,14,15 In this regard, TM doping is a facile and conventional way to achieve stable magnetism in 2D semiconductor materials. Moreover, the magnitude of magnetism and the dopant concentration of MoS₂ nanoflakes can be controlled in the doping process. Ferromagnetism in layered TMDs via substitutional doping of magnetic 3d transition-metal atoms has been reported by Ramasubramaniam et al.¹⁶ and possible ferromagnetism in MoS₂ by single absorption of the Co atom has been claimed by Chen et al.¹³ for the first time. Evidence of room temperature ferromagnetism in Mn, Fe, Co and Ni-doped MoS₂ in monolayer and bulk form, initiated by substitutional defects and defect complexes, has been reported by Wang et al.¹⁷ Additionally, proton beam irradiation⁸ and doping with nonmetal elements¹⁸ are considered to produce ferromagnetism in few-layer non-magnetic MoS₂. It has been shown in previous studies that localized nonbonding 3d electrons of TM atoms produce charge polarization in TM substituted 2D systems.^19–21 The extra nonbonding valence electrons present in the TM metal ion compared to Mo contribute magnetic moment with substitutional doping in layered materials. For example Mn–MoS₂ = 1μ_B, Fe–MoS₂ = 2μ_B, Co–MoS₂ = 3μ_B, Ni–MoS₂ = 4μ_B, Cu–MoS₂ = 5μ_B moments possess 1, 2, 3, 4 and 5 extra electrons in nonbonding orbitals.²² It is worth noting that a maximum of only 5 electrons can be accommodated by the 3d orbital in the single-filled states. Hence, based on the above theory, the magnetism of the TM doped MoS₂ structure can be tuned successfully by the internal TM engineering which plays a crucial role in the functionalization of MoS₂ nanomaterials in possessing a defined magnetic moment. However, Co being a member of the transition metal ion group possesses a comparable ionic radius (0.70 Å) to molybdenum (0.68 Å). Doping of the Co-ion in MoS₂ NSs can effectively modulate the structural properties by introducing vacancies, defects and strain which are also responsible for the ferromagnetism. The magnetism in MoS₂ nanostructures can be much influenced by the zigzag edges, and the experimental and computational study has shown that the magnetism in Co-doped MoS₂ nanoflakes is sensitive to the edges as well.²³ Theoretically it has been shown that Co substitution at the Mo site leads to a spin-polarized state²⁴ contributing three electrons in the 3d orbital in the single-filled states. The partial replacement of Mo with the transition metals like Co creates magnetism in MoS₂ making it a suitable candidate as a new kind of diluted magnetic semiconductor.

Moreover, other effects such as coupling between the adjacent layers are known to stabilize the magnetism and can enhance the overall magnetization originating from the grain boundaries.²⁵ The ferromagnetic signal was attributed to localized electronic states at grain boundaries in other nonmagnetic layered systems such as graphite.²⁶ Even though the interlayer coupling might enhance the ferromagnetic signal, such coupling is not a necessary condition for the existence of ferromagnetism as the presence of zigzag edges is sufficient and magnetization measurements certainly confirms that the magnetization originating from zigzag edges.^11,27 Nevertheless, studies on the tuning of magnetic and transport properties of MoS₂ simultaneously by doping of transition metal (TM) atoms are extremely critical in experimental work.

In earlier reports, it has been shown that n-type MoS₂-based field-effect transistors (FETs) face limitations in certain applications due to having high contact resistance. However, to reduce contact resistance, metallic MoS₂ with a lithium-solution treatment has recently been introduced^28,29 which is dissimilar to the conventional methods, such as chemical doping and ion implantation, used for the reduction of contact resistance in other devices. Li-treatment of MoS₂ results in the coexistence of various MoS₂ phases and consequently develops significantly different electronic and magnetic properties.^30,31 Kim et al. reported the transport properties of Li-ion treated polymorphic MoS₂ with a large temperature coefficient of resistance (TCR), enabling efficient IR detection at room temperature.³² Meanwhile recent studies on transport suggest that some factors such as charged impurities^33–35 and localized states^34,36,37 play a crucial role in controlling the dominant scattering processes that limit the carrier mobility in mechanically exfoliated samples, and this remains unexplored. Recent studies on CVD-grown MoS₂ suggest that the contact resistance due to Schottky barriers needs to be carefully considered. The adverse impact of dominant scattering processes that limit the carrier mobility should be critically evaluated.^38,39 The advantage of incorporation of the 1T metallic phase in the semiconducting 2H phase through Li-ion treatment or TM ion doping is that it controls the junction barrier resistance in TMD-based devices. In this work, we investigate the nature of various dominating charge transport mechanisms which come into play in the high- and low-temperature region of multi-layered undoped (0% Co) and doped (with various amounts of cobalt, i.e., 2% Co, 4% Co, 6% Co and 8% Co) MoS₂ nanoflakes. Doping-induced sulfur vacancies (SVs) and cobalt dopants act as electron donors and induce localized states in the band gap leading to the narrowing of the band gap.⁴⁰ In undoped MoS₂ the electron transport is dominated by thermally excited nearest-neighbour hopping (NNH) throughout the entire temperature region (50–300 K) due to the lack of doping-related defect induced localized electrons. The localized electron transport in the cobalt doped MoS₂ sample is dominated by thermally excited NNH from moderate to high temperature and variable-range hopping (VRH) at sufficiently low temperatures. In this report, doping-induced ferromagnetism by cobalt ion doping in MoS₂ nanoflakes has been explored and verified both experimentally (VSM, EPR) and theoretically (DFT). Additionally, the presence of zigzag edges has been verified by TEM measurements corresponding to 8% Co-doped MoS₂ NFs which probably gives rise to the ferromagnetic behaviour synergistically in association with the doping effect. The ferromagnetic behaviour of 8% Co-doped MoS₂ NFs having saturation magnetization (M_S) ∼ 1.2 emu g⁻¹ and coercivity ∼70 Oe at 5 K has been observed. The role of the substituted Co dopant and the doping induced 1T phase in enhancing ferromagnetism has been explained through the bound magnetic polaron (BMP) model. According to the BMP model, only those nearby Mo⁴⁺ ions in the 1T phase and Co²⁺/Co³⁺ ions in octahedral geometry located within the effective radius of the same polaron (around a sulfur vacancy center) could interact with each other ferromagnetically giving rise to ferromagnetism in 1T@2H-MoS₂ nanoflakes. A detailed first principle DFT study on undoped and cobalt doped systems reveals the increasing nature of ferromagnetism assuming substitutional Co doping in a pure 2H MoS₂ matrix. Therefore, we have successfully incorporated zigzag edges, defects and vacancies in a controlled way by doping Co in MoS₂ NFs while improving their electrical properties.

Layered TMDs, such as MoS₂, have proven their efficacy in the electrochemical energy storage/conversion field, owing to their large surface area that allows maximum exposure of surface active sites, high mechanical strength, and flexibility in the atomic scale dimension.⁴¹ The heterogeneous electron transfer (HET) on the edge planes of MoS₂ is predicted to be higher than that on the basal planes, and thus the edge sites are catalytically and electrochemically more active than the basal planes.⁴² Hence, it is always encouraged to fabricate MoS₂-based electrodes with the edge-exposed surface rather than with the basal planes exposed, and of these the former improves the catalytic and electrochemical performances.^43,44 Moreover, the electrical conductivity of MoS₂ is another key factor that influences the catalytic and electrochemical performance of MoS₂-based electrodes. Our strategy of systematic Co-doping in MoS₂ increases the number of active sites in the edge planes, enhances the conductivity by introducing the metallic 1T phase and increases interlayer separation aiding in the intercalation of ions easily and defect states with modulation of the band gap. Such salient features facilitate the Co-doped MoS₂ to function as a better electrode material (∼1.7 times higher specific capacitance observed in 8% Co doped MoS₂ compared to the undoped one) having an improved charge storage mechanism. Our work provides a detailed investigation of the manipulation of magnetic property for achieving room-temperature ferromagnetism and of transport property by varying Co ion dopant content in MoS₂ NFs as well as by introducing the 1T phase, defects, and edges with simultaneous enhancement in electrochemical property for future application in energy storage devices.

2. Experimental details

2.1. Chemicals used

High purity (99.9% purity) chemicals such as sodium molybdate dihydrate (Na₂MoO₄·2H₂O), thiourea (CH₄N₂S), and cobalt acetate tetrahydrate (CO(CH₃COO)₂·4H₂O) were procured from Sigma-Aldrich. All the reagents were of analytical reagent grade and used without further purification.

2.2. Synthesis of ultra-thin Co-doped MoS₂ NFs

The undoped and Co-doped (with various percentages of Co) MoS₂ NFs were synthesized using the facile one-step hydrothermal technique using Na₂MoO₄·2H₂O, Co (CH₃COO)₂·4H₂O and CH₄N₂S as starting reagents. In a typical preparation process, 2 mmol of Na₂MoO₄·2H₂O and 5 mmol of CH₄N₂S were taken and dissolved in 40 mL of deionized water and the pH value of the solution was adjusted to 5.0 by adding 0.02 mol L⁻¹ of acetic acid (CH₃COOH). Then the mixed solution was transferred into a 50 mL Teflon-lined stainless-steel autoclave for hydrothermal treatment at 200 °C for 24 h and cooled down to room temperature naturally. MoS₂ in the form of a black powder was obtained by centrifugation at 3000 rpm for 10 min. Then the black powder was washed with deionized water and ethanol several times and dried in a vacuum oven at 60 °C for 12 h. For the synthesis of 2 at wt%, 4 at wt%, 6 at wt% and 8 at wt% cobalt doped MoS₂ NFs, the two precursors Na₂MoO₄·2H₂O and Co (CH₃COO)₂·4H₂O were taken in the ratio of 0.98 : 0.02, 0.96 : 0.04, 0.94 : 0.06 and 0.92 : 0.08 (in mmol) while using the same amount of CH₄N₂S in the aqueous solution and the synthesis process was repeated.

The MoS₂ NFs with various amounts of cobalt doping (0 at wt%, 2 at wt%, 4 at wt%, 6 at wt% and 8 at wt%) were labelled as undoped, 2% Co, 4% Co, 6% Co and 8% Co, respectively.

2.3. Characterization techniques

The instrumental specifications of X-ray diffraction (XRD), Raman measurement, scanning electron microscopy (SEM) and high-resolution transmission electron microscopy (HRTEM) techniques to understand the existing crystalline phase and different vibrational modes present in our samples have already been reported.⁴⁰ The room temperature EPR spectra of the samples were recorded using an X-band Bruker ELEXSYS 580 EPR spectrometer with 9.8492 GHz of X-band frequency. A 7400 series vibrating sample magnetometer (VSM) (Lake Shore Cryotronics, USA) was used to measure the magnetic property based on the M vs. H (applied field ∼ −2 T to +2 T) and M vs. T plot (10–300 K, 500 Oe). The characterization of the transport property of the samples was carried out using a Lakeshore four-probe linear geometry setup using a low-temperature closed-cycle helium refrigeration cryostat. The powder form MoS₂ samples with various amounts of Co dopants were pelletized and conducting silver contacts were made for transport property measurements.

2.4. Electrode preparation

Electrochemical measurements like cyclic voltammetry (CV), galvanostatic charge–discharge (GCD) and ac impedance were performed using a CHI600E Electrochemical Workstation (CH Instruments, USA) at room temperature containing a glassy carbon electrode loaded with undoped and Co-doped MoS₂ samples, Ag/AgCl (in standard 3 M KCl aqueous solution), and Pt electrode as the working electrode, reference electrode, and counter electrode, respectively, using 1 M Na₂SO₄ aqueous solution as the electrolyte. Before starting all the measurements, the bare glassy carbon electrode (GCE, dia. 3 mm) was polished mechanically with 1, 0.3 and 0.05 μm alumina slurry sequentially to obtain a mirror finish, and then washed with acetone, ethanol and DI water thoroughly. The slurry was a mixture consisting of the active material, activated carbon, and poly(vinylidene fluoride) binder in the weight ratio of 80 : 15 : 5 with N,N-dimethylformamide as the solvent. Then, a certain amount of the slurry was drop cast on the GCE surface carefully. Finally, the electrodes were dried at 80 °C for 12 h in vacuum.

2.5. Computational details

To elucidate the impact of Co dopants on the magnetic properties of MoS₂ nanoflakes, first-principles calculations were performed using the Vienna Ab initio Simulation Package (VASP),⁴⁵ within the generalized gradient approximation (GGA)⁴⁶ considering the Perdew–Burke–Ernzerhof (PBE)⁴⁶ parametrization to describe the exchange-correlation interaction. The cut-off energy for the plane-wave basis set was set at 450 eV in all the computations. The relaxation convergence of energy was taken to be less than 1.0 × 10⁻⁶ eV and all the atoms were allowed to relax using a conjugate gradient scheme until the inter-atomic forces fell below 0.01 eV Å⁻¹. All the structural optimizations and the scf calculations were done using a well-converged Monkhorst–Pack⁴⁷ 11 × 11 × 11 and 21 × 21 × 21 k-point grid, respectively. The electronic structure calculations were done with a much denser k-point mesh. In our calculation, we assumed that unit cells have 2H atomic arrangements as the 2H phase is seen to be dominant over the 1T phase in all cases. To account for van der Waals interaction between two layers of the unit cell, the DFT-D2⁴⁸ method as implemented within VASP was adopted in all calculations.

3. Results and discussion

3.1. XRD data analysis

The XRD patterns of the undoped and cobalt-doped (with various percentages of Co) MoS₂ nanostructures recorded over the 2θ range of 10°–70° have been shown in our previous research paper.⁴⁰ Here the XRD peaks of all undoped and Co-doped samples exactly match with hexagonal MoS₂ having identical peak positions. The XRD graphical plot of the samples with JCPDS card no. 00-037-1492 is shown in Fig. 1(a). The signature of the diffraction peak corresponding to the 1T phase (ESI,† Fig. S1) is more prominently observed in the XRD pattern of 8% Co-doped MoS₂ compared to other samples with low concentration Co doping. This peak shifts to a lower angle 10.6° to 10.5° when the Co dopant percentage is increased from 2% to 8%, possibly due to the appearance of tensile strain within the lattice. The slight shift of the main (002) diffraction peak to a lower angle and increased lattice parameters (a and c) confirm the existence of uniaxial lattice strain (tensile) in the lattice due to substitutional Co doping.⁴⁹ The XRD pattern of the undoped and Co-doped MoS₂ nanoflakes was successfully refined with a 2H-MoS₂ phase (P63/mmc, a = 3.14 Å, b = 3.14 Å, c = 12.53 Å, α = 90°, β = 90°, γ = 120°) having no extra crystalline impurity phase. The lattice parameters as obtained after refinement are found to be a = 3.158 Å, b = 3.158 Å, c = 12.815 Å for undoped and a = 3.174 Å, b = 3.174 Å, c = 12.879 Å for 8% Co-doped MoS₂ NFs. The refined lattice parameters agree well with previous literature.⁵⁰ The average crystallite size of all the MoS₂ NFs varies from 8 nm to 6 nm. It has been observed that the crystallite size as obtained by using the Scherrer formula decreased with the increase of the cobalt dopant amount. This can be attributed to the possible appearance of lattice strain and defects induced by doping. The microstrain and dislocation density increased significantly with increasing doping concentration as evidenced by increasing defects and strain reported earlier.⁴⁰

Fig. 1 (a) XRD pattern of MoS₂ NFs doped with various percentages of cobalt. Rietveld refinement of (b) undoped and (c) 8% Co doped MoS₂.

3.2. Raman spectroscopy study

Raman spectroscopy was employed to demonstrate the doping effect in MoS₂ NFs using a 532 nm excitation laser with the calibrated silicon Raman peak at 521 cm⁻¹. The recorded Raman spectra of MoS₂ and Co-doped MoS₂ are shown in Fig. 2. Furthermore, for the undoped MoS₂ nanoflakes, the Raman spectra display characteristic peaks at 281, 377, 404, and 454 cm⁻¹ arising from the E_1g, E¹_2g, A_1g and longitudinal acoustic phonon mode (A_2u) of 2H-MoS₂.^51,52 The two intense peaks are located at ∼377 cm⁻¹ (E¹_2g mode) corresponding to the in-plane vibration of S and Mo atoms and at ∼404 cm⁻¹ (A_1g) attributed to the relative vibration of S atoms in the out of plane direction.^53,54 In addition to these 2H vibrational Raman modes, two other weak vibrational modes at 232 cm⁻¹ (J₂) and 335 cm⁻¹ (J₃) appear in the undoped one assigned to the 1T phase. By contrast, for 8% Co-doped MoS₂ NFs, the intensity of J₂ and J₃ has increased significantly along with the presence of the J₁ phase attributed to the 1T-MoS₂ phase. In 8% Co-doped MoS₂ NFs the suppression of the E¹_2g and A_1g peak intensity has been observed. This suggests that a high amount of cobalt doping enables phase transformation of the 2H phase to 1T phase in 8% Co-doped MoS₂ NFs. In the case of 8% Co-doped MoS₂ NFs, redshifts of the two peaks (E¹_2g and A_1g) were detected from 377.72 cm⁻¹ to 374.78 cm⁻¹ and from 404.27 cm⁻¹ to 402.32 cm⁻¹, respectively. Specifically, the absence of sulfur atoms in few-layer MoS₂ would induce redshifts of both Raman peaks as reported by Kim et al.⁵⁵ Tensile strain appearing in MoS₂, directly related to the increase of the lattice parameters, redshifts both E¹_2g and A_1g peaks as observed by Yang et al.⁵⁶ and Chaudhury et al.⁵⁷ Thus, a possible explanation is that sulfur vacancies and tensile strain lead to bond softening and increased lattice parameters. Additionally, we found that the redshift of the E¹_2g mode contributed by the in-plane opposite vibration of two S atoms with respect to the Mo atom is always larger than that of the A_1g mode caused by the out-of-plane vibration of S atoms.⁵⁸ The absence of sulfur atoms would have a larger effect on the E¹_2g mode due to the in-plane chemical bonds between S and Mo atoms⁵⁹ showing a gradual increase in Δk (difference in k) between E¹_2g and A_1g modes. Table 1 shows that there is a gradual increase of Δk between E¹_2g and A_1g in 8% Co-doped MoS₂ NFs compared to the undoped ones. Additionally, chemical doping induced strain in the lattice changes the bond length and hence an increase in Δk occurs. The observed increase in the E¹_2g/A_1g intensity ratio and the increased FWHM of A_1g in 8% Co-doped MoS₂ compared to the undoped one are an indication of stacking of more number of layers along the in-plane direction and suppression of the out-of-plane movement related to the mode.⁶⁰ The significant increase in FWHM and the downshift of A_1g frequency as observed in the 8% Co-doped MoS₂ case support the presence of zigzag edges forming structural disorder.⁶¹ The shifts and differences between the vibrational Raman E¹_2g and A_1g peaks with Co doping are presented in Table 1.

Fig. 2 Different Raman modes showing the presence of 1T and 2H phases in undoped and 8% Co doped MoS₂.

Table 1 Variation of characteristic Raman modes of undoped and 8% Co doped MoS₂ NFs

Co doping percentage (%)	E^12g (cm^−1)	A1g (cm^−1)	Δk (cm^−1)	FWHM ratio (A1g/E^12g)
0	377.72	404.27	26.55	0.82
8	374.78	402.32	27.54	1.62

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3.3. Morphology and microstructure study

3.3.1. SEM image analysis. Fig. 3(a and b) show the FESEM images of the as-prepared 8% Co-doped MoS₂ nanoflakes indicating the formation of a 3D hierarchical fluffy petaloid architecture. Each flower consists of several interconnected nanopetals in the form of nanoflakes. The average diameter of the undoped and 8% Co-doped MoS₂ flowers is about ∼200 nm and ∼1.5 μm, lateral dimensions are ∼130 nm and ∼240 nm and the thickness lies within 10 nm. The lateral size of the MoS₂ NF increases with the incorporation of more Co dopants. The reason for the formation of larger nanoflowers with higher dopant content may be ascribed to the accumulation of dopant and defect-induced charge on the surface and edge of the lamellar structure contributing to the electrostatic attraction between them which can be validated by zeta potential measurements.⁴⁰Fig. 3(c and d) show the EDAX elemental data analysis and plot indicating the proper distribution and presence of Mo, S and Co in 8% Co-doped MoS₂ NFs. However the theoretically calculated stoichiometric ratio of Co, Mo, and S (0.03 : 0.31 : 0.66) for 8% Co-doped MoS₂ matches well with the experimentally found value with a slight variation (0.05 : 0.41 : 0.53) possibly due to the presence of S vacancies.

Fig. 3 SEM images of (a) undoped and (b) 8% Co doped MoS₂ nanostructures. The inset of (b) shows a zoomed view of the encircled area indicating the microstructure of nanoflakes after doping. (c and d) EDAX analysis of 8% Co doped MoS₂ indicating the percentage of Mo, Co and S elements present in the sample.

3.3.2. TEM image analysis. The detailed information about the microstructure of undoped and 8% Co-doped MoS₂ NFs can be further unfolded by TEM image analysis. The thin interlaced flakes of MoS₂ convert to a bigger nanostructure having sheet-like morphology with larger lateral dimensions which can be well understood in Fig. 4(a and d). The inset of Fig. 4(a) shows diffused rings in the SAED pattern of undoped MoS₂ whereas the inset of Fig. 4(d) shows distinct spots in SAED pattern diffraction planes indicating an increased crystalline nature upon doping in 8% Co-doped MoS₂. Fig. 4(b) clearly shows the magnified image of interlaced flakes of undoped MoS₂ where the interlayer distance is found to be ∼0.62 nm forming a 14–16-layer system. In the case of 8% Co-doped MoS₂, it has been observed that the number of layers lies within 4–8 layers having an extended interlayer distance ∼0.65 nm as shown in Fig. 4(e) due to the appearance of doping-induced lattice strain.³⁹Fig. 4(c) shows the prominent presence of 2H and 1T phases and defects in the vicinity of the 1T phase. The presence of edge-terminated zigzag orientation in 8% Co-doped MoS₂ has been indicated by an arrow as shown in Fig. 4(f). According to Tinico et al., such type of zigzag orientation in vertically oriented MoS₂ is responsible for metallic behaviour.⁶²

Fig. 4 TEM images showing (a and b) thin nanoflake like morphology and interplanar spacing (0.62 nm) of undoped MoS₂, (c) prominent presence of 2H and 1T phases and sulfur vacancies (V_s, encircled) in 8% Co doped MoS₂, (d) interlaced flakes forming larger nanoflakes of 8% Co MoS₂, and (e and f) the extended interplanar spacing (0.65 nm) and presence of zigzag edges in 8% Co doped MoS₂.

3.4. XPS study

The chemical composition and prominent presence of the 1T phase along with the 2H phase in undoped⁴⁰ and 8% Co-doped MoS₂ have been verified by X-ray photoelectron spectroscopic (XPS) measurement. The typical BE peaks of Mo 3d_5/2 and Mo 3d_3/2 corresponding to the 2H phase of 8% Co-doped MoS₂ appeared at 228.37 and 231.75 eV, as shown in Fig. 5(a). Two extra peaks corresponding to Mo 3d_5/2 and 3d_3/2 of the 1T characteristic phase were shifted to ∼0.5–1 eV. Similarly, the binding energy peaks corresponding to S 2p_3/2 and S 2p_1/2 of the 2H phase appeared at 161.63 and 162.79 eV in the case of 8% Co-doped MoS₂, as shown in Fig. 5(b). Two other extra peaks shifted about ∼0.5 eV from S 2p_3/2 and S 2p_1/2 of the 2H phase are assigned to the 1T phase. It has been observed that the doublets corresponding to Mo⁴⁺ and S²⁻ are shifted to lower binding energy with a gradual increase of Co dopant as observed in 8% Co-doped MoS₂ when compared with the undoped one as shown in Fig. 5(a and b), confirming the loss of sulfur atoms in the sample.⁵⁸ The difference between Mo 3d_3/2 and 3d_5/2 binding energy peaks corresponding to the 1T phase is 3.1 eV confirming the emergence of Mo⁵⁺ ions in 8% Co-doped MoS₂.⁶³ By analyzing the XPS spectra, the S : Mo peak ratio for 4% Co-doped MoS₂ is 37% smaller than that for 2% Co-doped MoS₂ NFs.⁶⁴ In addition, based on the analysis of the Mo 3d peak area, it is deduced that the content of Mo⁵⁺ in MoS₂ NFs is around 26%, demonstrating a defect-rich structure.⁶⁵ Four main peaks corresponding to Co 2p_3/2 at 780.34 eV (Co³⁺) and 785.27 eV (Co²⁺) and Co 2p_1/2 at 791.33 eV (Co³⁺) and 798.07 eV (Co²⁺) are found in 8% Co-doped MoS₂ as shown in Fig. 3(c) similar to the report in the case of Co species found in CoS₂.^66–69 Also, another two peaks arising at 790.73 eV and 803.24 eV signify the existence of satellite shake-up peaks. The binding energy peak of Co 2p_3/2 at 780.34 eV resembles the CoMo₂S₄ phase, confirming the uniform substitution of Mo atoms by Co²⁺ along the {002} or the S-edge planes.⁷⁰ The presence of the elements C, Mo, S, O and Co is revealed in the survey spectrum (wide scan) of undoped and 8% Co-doped MoS₂ as shown in Fig. 5(d). The BE peaks corresponding to different oxidation states of Mo, S, and Co elements present in our samples are listed in Table 2.

Fig. 5 XPS narrow scan spectra of undoped and 8% Co doped MoS₂ (a) Mo 3d, (b) S 2p and (c) Co 2p of 8% Co MoS₂. Survey spectra of Mo (3d, 3p, 3s), S 2p, C 1s, and Co 2p are shown in (d) corresponding to undoped and 8% Co doped MoS₂.

Table 2 Position of binding energy peaks of Mo, S and Co elements corresponding to the 1T and 2H phase of undoped and 8% Co-doped MoS₂

Doping amount (%)	1T		2H		1T		2H		Co^3+		Co^2+
	Mo 3d5/2 (eV)	Mo 3d3/2 (eV)	Mo 3d5/2 (eV)	Mo 3d3/2 (eV)	S 2p3/2 (eV)	S 2p1/2 (eV)	S 2p3/2 (eV)	S 2p1/2 (eV)	Co 2p3/2 (eV)	Co 2p1/2 (eV)	Co 2p3/2 (eV)	Co 2p1/2 (eV)
0	228.25^40	231.12^40	229^40	231.5^40	161.7^40	162.0^40	162.96^40	163.37^40	—		—
8	227.88	231.25	228.37	231.75	161.17	162.31	161.63	162.79	780.34	791.33	785.27	798.07

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3.5. Electronic transport property study

Fig. 6 shows the temperature-dependent resistivity curves for all undoped and Co-doped MoS₂, exhibiting typical semiconductor behaviour. The resistivity eventually decreases with an increase in doping concentration and it is clearly seen in low-temperature regimes. This supports the increase in carrier concentration with the incorporation of an increased proportion of the 1T phase with Co doping.⁷¹ Thermal excitation of electrons from donor levels to the conduction band takes place with the increase in temperature. Temperature-dependent resistance can be written in an Arrhenius form:

σ(T) = σ₀exp(−E_a/k_BT)

where σ is the conductivity, σ₀ is the pre-exponential factor, k_B is the Boltzmann constant, and E_a is the activation energy. Fig. 6(a–f) shows the plots of the conductivity in log scale as a function of 1000/T from 217 K to 250 K of undoped and 2%, 4%, 6%, and 8% Co-doped samples. Different thermal activation energy regions have been obtained beyond 200 K. The data are fitted to a line with different slopes in this temperature region. The activation energy values for undoped and 2%, 4%, 6% and 8% Co-doped samples are 85.13 meV, 65.73 meV, 47.38 meV, 33.9 meV and 85.3 meV respectively (above 200 K). The activation energy is observed to be much lower than the band gap energy of MoS₂ samples. Such low activation energy strongly suggests that native defects such as sulfur vacancies and impurities are present in the as-synthesized MoS₂ samples as already verified in the XPS result. The decrease of activation energy with an increase of doping concentration suggests that donor carrier density lifts the Fermi level position in the band gap which supports the result obtained from UV-visible absorption spectra.⁴⁰ It has been studied in a doped semiconductor that a major portion of the free electrons is recaptured by donor atoms themselves at low temperatures.⁷² The electrons suffer from lack of sufficient energy to jump from donor levels to the conduction band. At this stage, the electrons transport through hopping from one level to another in the impurity band and free-electron band conduction becomes less significant.⁷² Thermal activation is due to deep donor level ions contributing to the conduction band. For higher cobalt doped MoS₂ at temperature <200 K, the activation energy (E_a) drops and other transport mechanisms could play a role in addition to the nearest neighbour hopping. The weak temperature dependence of resistance at low temperatures can be explained by the mechanism of electron hopping through impurity levels formed near the conduction band edge of cobalt-doped MoS₂. In addition to this, the decrease of temperature is accompanied by a rapid rise of electrical resistivity as electrons hop from the conduction band to the localized impurity states. Therefore, the conduction mechanisms in these Co-doped MoS₂ samples can be explained by the combination of two types of conduction models such as variable range hopping (VRH), σ_VRH = σ_0VRH exp [−(T_0VRH/T)^1/4], and thermal excitation by nearest neighbour hopping (NNH), σ_th = σ_0th exp[(E_a/kT)], where σ_0VRH, T_0VRH, and σ_0th are constants. The first one describes carrier hopping in localized states at a lower temperature range whereas the second one describes carriers that have been thermally excited from the localized state to the conduction band. The parameters were evaluated from the linear fitting method by plotting ln σ ∼ T^−1/4(VRH) and ln σ ∼ 1000/T(NNH) in different temperature regions to investigate the involved temperature-dependent transport processes. These parameters are listed in Table 3. The resistivity rises at low temperature and passes the VRH test for 2%, 4%, and 8% Co-doped samples (50–70 K) and for 6% Co (50–110 K) at a low-temperature range as shown in Fig. 7(a, b, d and c), respectively. It suggests that the conduction in these pellet-formed powder samples is due to the thermally assisted tunneling of the charge carriers through the grain boundary barrier and transition from the donor level to the conduction band. It is observed that at low temperature, the resistivity of 8% Co-doped MoS₂ is very less compared to the undoped one due to incorporation of the metallic 1T phase in 2H MoS₂ with the increase of cobalt doping concentration. Increasing the Co-doping concentration will reduce the resistance, and in Table 3 it has been observed that the activation energy decreases from 65.73 meV to 33.59 meV as the concentration of Co atoms increases from 2% to 6%. It has been observed that in the case of heavy doping of the Co ion (8% Co) in MoS₂, the activation energy is increased. This is attributed to the increase in atomic defects which are electrically active at the grain boundary.⁷³

Fig. 6 Logarithmic plot of resistivity with temperature of (a) undoped and Co doped MoS₂. Linear variation of conductivity with inversion of temperature to test the (b) undoped and (c) 2%, (d) 4%, (e) 6% and (f) 8% Co doped MoS₂ which satisfy the thermal excitation process of transport.

Table 3 Various transport mechanisms in different temperature regimes observed in cobalt doped MoS₂

Samples	σ 0VRH (Ω m)^−1	T 0VRH (K)	σ 0EVRH (Ω m)^−1	T 0EVRH (K)	σ 0Th (Ω m)^−1	E A (meV)	Adj. R^2	Temperature coefficient of resistance (TCR) (K^−1)
MoS2	—	—	—	—	3.7397	85.13	0.9998	−3.74 × 10^−2
2% Co	3.932 × 10^9	1.67 × 10^8	352.7	30043.28	0.2686	65.73	0.99825	−3.12 × 10^−2
4% Co	5.92 × 10^11	1.08 × 10^7	6.5023	9583.31	0.4832	47.37	0.99828	−1.85 × 10^−2
6% Co	0.7780	2 931 986	7.5341	12 500	0.1267	33.59	0.99871	−1.29 × 10^−2
8% Co	5597.07	2.673 × 10^7	1.769	18770.31	0.00647	85.3	0.9997	−2.93 × 10^−2

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Fig. 7 Temperature dependent conductivity plot of 2%, 4%, 6% and 8% Co doped MoS₂ (a–d) and ln σ ∼ T^−1/4 plot of 2%, 4%, 6% and 8% Co doped MoS₂ to test the VRH mechanism.

3.6. EPR analysis

The room temperature electron paramagnetic resonance spectra (EPR) of MoS₂ nanostructures doped with various percentages of cobalt are shown in Fig. 8. The undoped MoS₂ shows a prominent resonant peak of low intensity at a magnetic field of 3446 G with g = 2.0005. This signal may come from the unpaired electrons generated by unsaturated coordination structures, e.g., vacancies and defects such as S vacancies, acting as paramagnetic centers.⁷⁴ The EPR signal in the higher magnetic field is indicated with an arrow in Fig. 8 at g = 1.89, which is very close to g values of oxygenated species of Mo⁵⁺ in amorphous or unsupported sulfides.⁷⁴ The maximum EPR signal has been observed in 2% Co-doped MoS₂ NFs and decreases gradually with an increase in Co dopant amount. The expression

g = hν/βH₀

is used to calculate the electron g factors, where ν = 9.65 GHz (frequency), h = 4.135 × 10⁻¹⁵ eV s⁻¹ (Planck constant), β = 5.788 × 10⁻⁵ eV T⁻¹ (Bohr magnetron) and H₀ is the resonance magnetic field. Co doping at different concentrations in MoS₂ does not induce any new peak or change in the EPR spectra. However, a peak shift towards a lower magnetic field has been observed in all the doped MoS₂ samples. In Fig. 8, the prominent resonance peaks are witnessed at the magnetic field of 3435 G (g = 2.007), 3433 G (g = 2.008), 3439 G (g = 2.005), and 3443 G (g = 2.002) for 2 at% Co, 4 at% Co, 6 at% Co and 8 at% Co-doped MoS₂ samples, respectively, which are attributed to the enhanced ferromagnetic coupling with the increase in Co concentration. Eventually, the g value reaches a maximum for 4 at% Co-doped MoS₂. The 1T regions tend to produce more Mo⁴⁺ spins by expanding in the less dense S vacancy regions. However, in the case of higher cobalt concentration (8 at%), the proportion of the 1T phase has increased up to 48%. In this case, S atoms are surrounded by a majority of the Mo atoms compared to S vacancies whereas a few of Mo atoms exist as isolated Mo⁴⁺ spins. This explains the origin of the ferromagnetic coupling with S atoms, resulting in a rise of ferromagnetism in MoS₂ nanoflakes macroscopically.³¹

Fig. 8 EPR spectra of MoS₂ NFs doped with various percentages of cobalt.

3.7. VSM data analysis

The magnetic properties (M–T and M–H at 10 K, 100 K and 300 K) of undoped and 2% Co, 6% Co and the highest Co (8% at wt.) doped MoS₂ samples were studied through VSM as shown in Fig. 9. Both undoped and 2% Co-doped MoS₂ samples show paramagnetic behaviour originating from free electrons contributed from defects and a small proportion of the 1T phase whereas 2% Co-doped MoS₂ samples show higher magnetic moment compared to the undoped one as shown in the inset of Fig. 9(a). As reported by Nair et al. the existence of point defects such as vacancies and the introduction of fluorine adatoms in graphene induce notable paramagnetism.⁷⁵ However, a paramagnetic to ferromagnetic transition has been distinctly found above 2% Co doping. The ZFC–FC feature in the M–T curve of 8% Co-doped MoS₂ NFs, as shown in Fig. 9(b), almost retraces the same path indicating intrinsic single-phase ferromagnetism. The highest saturation magnetization (M_S) at 10 K is 0.42 emu g⁻¹ and 1.2 emu g⁻¹, remanence magnetization (Mr) is 0.06 emu g⁻¹ and 0.045 emu g⁻¹ and coercivity is ∼150 Oe and 70 Oe corresponding to 6% Co- and 8% Co-doped MoS₂ nanoflakes respectively as shown in Fig. S2(c) (ESI†) and Fig. 9(d). Since our undoped MoS₂ possesses a feeble paramagnetic nature, it is fair to assume that the robust magnetism in 6% Co- and 8% Co-doped MoS₂ arises as a consequence of doping. The magnetism in multilayer MoS₂ and its doped analogues depends on the presence of zigzag edges, defect type, vacancies, lattice strain, and dopant concentration. Theoretical calculations by Wang et al.⁷⁶ revealed that low concentration Co doping at 4% and 6% in the Mo vacant sites of the basal planes results in stable magnetic moments at room temperature. Yun et al.⁷⁷ and Saab et al.⁷⁸ also reported tuning of electronic and magnetic properties due to doping of metal ions in the MoS₂ lattice. The high M_S observed in Co-doped MoS₂ suggests that the robust ferromagnetism here originates not from the strain in the layer but rather from the ferromagnetically favourable ordering of Co²⁺ ions.⁵⁰ The effective Bohr magneton number (p_eff(Co)) and the value of the Curie constant (C) were calculated using the following formulae:

(p_eff(Co))² = xg²S_(Co²⁺)(S_(Co²⁺) + 1) + (1 − x)g²S_(Co³⁺)(S_(Co³⁺) + 1) (1)

(p_eff(Co))² = 3Ck/Nμ_B² (2)

where g = 2 represents the Lande’ g-factor of the cobalt ions, S_(Co²⁺) and S_(Co³⁺) represent the total spin quantum numbers for Co²⁺ and Co³⁺ ions respectively, k is the Boltzmann constant, N is the Avogadro number and μ_B is the Bohr magneton. The value of p_eff(Co) obtained for MoS₂ NFs doped with various amounts of cobalt lies between the theoretical value of p_eff(Co²⁺) = 3.87 and p_eff(Co³⁺) = 4.89 due to the presence of the mixed-valence state of Co in Co-doped MoS₂ NFs. Again, the value of Curie temperature of 6% and 8% Co-doped MoS₂ NFs was obtained from the minima of the dm/dT vs. T plot. The magnitude of the Curie constant (C) for 6% and 8% Co-doped MoS₂ nanoflakes was calculated using eqn (2) by putting the value of p_eff(Co). The variation of saturation magnetization (M_s), Curie temperature (T_c), Curie constant (C), p_eff(Co) and the percentage amount of Co²⁺/Co³⁺ ions present in MoS₂ NFs doped with various amounts of cobalt are listed in Table 4.

Fig. 9 (a) M vs. T plot of undoped and 2%, 6% and 8% Co doped MoS₂ at an applied field of 500 Oe; the inset shows the M–T curve of undoped and 2% Co doped MoS₂. (b) Individual ZFC–FC curve of 8% Co doped MoS₂ indicating T_c ∼ 139 K. (c) dm/dT plot of 8% Co doped MoS₂ showing T_c ∼ 139 K. (d) Linearly fitted 1/χ ∼ T (T_c ∼ 139 K) corresponding to 8% Co doped MoS₂; the inset shows a magnified image of the M vs. H hysteresis loop of 8% Co doped MoS₂ at 10 K having 70 Oe coercive field. (e) Hysteresis loop at 10 K, 50 K, 100 K and 300 K. (f) M–H loop of 8% Co doped MoS₂ showing room-temperature ferromagnetism at 300 K at a magnified scale.

Table 4 Variation of magnetic parameters with various amounts of Co doping (0%, 2%, 6%, 8%) in Co-doped MoS₂ NFs

Doping percentage (%)	M s (emu g^−1)	T c (K)	C (emu K Oe^−1 mol^−1)	p eff(Co^2+, Co^3+) (calculated)	Variation of Co^2+/Co^3+ ions in percentage with Co doping
					Co^2+	Co^3+
0	—	—	—	—	—	—
2	—	—	2.24	4.26	65	35
6	0.42	131	2.32	4.32	59	41
8	1.2	139	2.45	4.37	54	46

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To study the origin and variation of the observed ferromagnetism in MoS₂ NFs with Co doping, it is necessary to point out the role of the 1T phase and cobalt ion incorporation in the 2H MoS₂ matrix. The fact that Co dopant atoms can effectively introduce ferromagnetism in MoS₂ nanoflakes has been confirmed by magnetic measurement results. Herein, we only consider the magnetic moment introduced in the MoS₂ NFs through the substitution of Mo atoms by Co atoms. According to crystal field theory, 4d and 3d orbitals corresponding to Co²⁺/Co³⁺ remain in different degenerate states in 2H and 1T coordination.³¹ The Mo⁴⁺ ions corresponding to 2H coordination split into three energy levels: (d_z²), (d_x²−y², d_xy), and (d_xz, d_yz) while the 1T has two levels: (d_xz, d_yz, d_xy) and (d_x²−y², d_z²) as shown in Fig. 10(a). As described in Fig. 10(a), two 4d electrons occupy the d_z² orbital of the Mo⁴⁺ ion in a spin-antiparallel manner, thus giving net zero magnetic moments in the 2H phase. In the case of the 1T phase of MoS₂, the lower energy levels (d_xz, d_yz, d_xy) are occupied by the two 4d electrons spin-parallelly contributing net magnet moments in the 1T phase of Mo atoms.⁷⁹ Therefore, it can be assumed that the presence of ferromagnetism is directly connected to the proportion of both 1T and 2H phases in MoS₂ nanoflakes. Seven and six 3d electrons corresponding to the Co²⁺ and Co³⁺ configuration contribute three and four unpaired electrons in the octahedral geometry of the Co²⁺ and Co³⁺ configuration, respectively. The observed magnetic moment varies monotonically with the proportion of the 1T phase present in the sample. In undoped MoS₂, the proportion of the 1T phase was 18% having the lowest magnetic moment as shown in EPR analysis. The proportion of the 1T phase increases gradually with the incorporation of the cobalt dopant. The proportion of the 1T phase in undoped and 2%, 4%, 6% and 8% Co-doped MoS₂ samples is found to be 18%, 26%, 39%, 42% and 48%, respectively. The magnetic moment value increases as the proportion of the 1T phase increases with an increase in the cobalt dopant as revealed by experimental results.

Fig. 10 (a) Electron orbital degenerate states of Mo atoms (Mo⁴⁺ ionic state) in 1T and 2H configuration and Co atoms (Co²⁺ and Co³⁺ ionic state). (b) Schematic representation of the bound magnetic polaron model mediated by vacancy and Co dopant in MoS₂ NFs. 1T phases are induced around vacancy and Co atom forming spin polarized states.

To further explain the origin of ferromagnetism in Co-doped MoS₂ systems, the bound magnetic polaron (BMPs) model is utilized.^80–82 Two main results are obtained from the XPS spectra: (i) S vacancies are typical defects in hydrothermally synthesized MoS₂ nanoflakes and (ii) the substituted Co atoms can contribute two holes. Thus the Co²⁺ dopants provide the local magnetic moments in addition to the carriers needed to couple these moments.⁸⁰ Under this condition, the electron distribution changes causing a strong interaction among the spin impurities. Unintentional sulfur vacancies are formed in MoS₂ nanoflakes during the hydrothermal synthesis process. As a result, BMPs could be realized with localized holes and the S vacancies surrounded by a large number of Mo⁴⁺ spins.⁸⁰ The Mo⁴⁺ spins and S vacancy spin are aligned parallel, leading to the formation of a BMP. Cai et al. observed that the 2H–1T phase transformation always occurs near the defects.⁵² Such ferromagnetism appears due to the overlapping of BMPs and their ferromagnetic coupling to each other as shown in Fig. 10(b).

Additionally, with the Mo⁴⁺ ions, the doped Co ions can be bounded by the S vacancies^82,83 as shown in Fig. 10(b). Only those Co ions arranged at suitable distances can interact ferromagnetically possibly through some indirect coupling mechanism.⁵⁰ In this manner, numerous BMPs can be created by the Co ions in the vicinity of the S vacancies. They are ordered ferromagnetically and establish long-range room-temperature ferromagnetism in Co-doped MoS₂ nanoflakes.⁸⁴

3.8. Theoretical result

To simulate 2%, 6% and 8% Co doping in MoS₂, 5 × 5 × 1, 3 × 3 × 1 and 3 × 2 × 1 bulk MoS₂ supercells were created respectively and then one Mo atom from the supercell was substituted with one Co atom, to incorporate the desired doping level, as shown in Fig. 11(a and b). The average Mo–Mo and Mo–S bond length obtained theoretically from our calculation agrees well with the experimental value⁸⁵ in the case of undoped and doped samples. In all the doped cases a small variation of the average Co–Mo bond length values and the average Mo–Mo bond length around the dopant site is observed from the Mo–Mo bond length value in the undoped cases. This suggests that the substitutional doping of Co at Mo does not change the symmetry around the dopant site. All the calculations are performed on the semiconducting 2H phase, since this phase is seen to be the dominant one over the metallic 1T phase, in all the cases.

Fig. 11 The atomic structure of (a) the undoped MoS₂ supercell and (b) the 3 × 3 × 1 supercell of 6% Co doped MoS₂ from top view with a Co atom substituting the Mo site. The big purple balls represent Mo atoms, small yellow balls represent sulphur atoms and the blue ball represents the cobalt atom. The projected density of states (PDOS) corresponding to Mo d, S p and Co d states of (c) the undoped MoS₂ supercell, (d) the supercell with 2% Co dopant, (e) the supercell with 6% Co dopant and (f) the supercell with 8% Co dopant (the vertical dashed line indicates the Fermi level).

The total and projected densities of states (TDOS and PDOS) of undoped MoS₂ and 2%, 6%, and 8% Co-doped MoS₂ are displayed in Fig. S3(a–d) (ESI†) and Fig. 11(a–d), respectively. The symmetry of the spin-up and spin-down DOS throughout the energy range, as shown in Fig. 11(a), validates the nonmagnetic nature of the undoped MoS₂ NFs. In the case of doped systems, it is seen that the doping of the Co atom in one layer does not affect the other layer. The reason might be the weak interlayer interaction and the large interlayer separation. The isolated Co atom possesses an electronic configuration of 3d⁷4s² having three additional valence electrons compared to the Mo atom, and it contributes to the magnetic moment in the Co-doped system. The overall magnetic moment of the supercell with one Co atom substituting one Mo atom and the local magnetic moment of the constituent atoms for 2%, 6%, and 8% Co-doped MoS₂ NFs are presented in Table 5. The Co atom induces spin polarization in the neighbouring 6 Mo atoms and S atoms. The spins of the six neighbouring Mo atoms are coupled to the Co atom with an average moment of 0.2μ_B in the 6% Co-doped case. An average magnetic moment of 0.04μ_B is contributed by spin-polarized p orbitals of the S atoms. More obviously, the Co dopant site remains in trigonal prismatic symmetry even after structural relaxation. As shown in Fig. 11(c–f), with the doping of the Co atom replacing a Mo atom, a distinct spin splitting appears near the Fermi level contributing to large ferromagnetism. This is attributed to the defect states associated with the doped Co atom, p states of the adjacent S atoms, and d states of the nearby Mo atoms, although mainly Co-d and Mo-d states. Due to the overlapping among the Co-d, Mo-d, and S-p orbitals the band edge structures change significantly in comparison with undoped MoS₂. For the C_3v symmetry around the dopant site, the doping produces five 3d states within the band gap which split into three sub-levels e₁, e₂, and a, among which one is mono degenerate and two doubly degenerate. The Co e₁ (d_xy, d_x²−y²) orbitals and e₂ (d_xz, d_yz) orbitals are strongly hybridized with S 3p orbitals and remain at relatively low energy levels in the spin-up channel, whereas in the spin-down channel they possess higher energy.¹³ The 3 extra non-bonding valence electrons occupy the lowest energy mono-degenerate a (d_z²) and the first doubly degenerate (d_xy, d_x²−y²) states in the spin-up channel. The spin-down channel remains vacant at the Fermi level, at least in the low doping regime. The results obtained from the calculation indicate an obvious asymmetry of the DOS around the Fermi level of the cobalt doped MoS₂, confirming that the substitution of Co at the Mo site gives rise to the spin-polarized state. Interestingly, the DOS indicates that Co-doped MoS₂ exhibits half-metallic behaviour in the 2% doping case, i.e., the DOS shows no gap in the spin-up channel and a semiconducting band gap in the spin-down channel which means being a semiconductor in one spin channel and a metal in another.^13,86,87 It is obvious from the calculation that, the asymmetry around the Fermi level has increased significantly with the increase in Co dopant concentrations, resulting in a net increase of total magnetic moment. Experimentally it has been demonstrated that Co doping up to a certain percentage can enhance ferromagnetism very well. The theoretical calculation for the Co doping in the dominant phase 2H supercell of MoS₂ shows the increasing nature of magnetic moment with an increase in doping concentration which supports the experimental evidence obtained from VSM measurements. This is probably because the small contribution of the 1T phase has been neglected in the calculation since the 2H phase has always been the dominant one over the 1T phase. However, the calculation presents a clear picture of the environment around the dopant site, the local and global magnetism and the induction of magnetism from the dopant to the nearby Mo and S atoms. To visualize the nature of the impurity state initiated by the S defect, Salehi et al.⁸⁸ discussed the influence of chalcogen defect on the electronic structure of the monolayer MoS₂ by assuming a single sulphur vacancy per supercell. Evidently, at high defect concentration, a midgap band is introduced by the sulphur vacancies near the VBM having a bandwidth of ∼0.6 eV. Such types of defect states tend to become more delocalized due to the interaction between S vacancies. Additionally, a flat band appears just below the MoS₂ CBM which does not change notably with the concentration changes. The Mo 4d orbitals around the vacancies play the main role in the creation of midgap states, while the S 3p and 3d orbitals do not contribute considerably to the defect states of the gap region. The presence of sulphur vacancies shifts the Fermi level to the bottom of the conduction band due to the unsaturated electrons in the Mo orbitals around the vacancy defect. In the case of multi-layered MoS₂, the effect of layers and their interactions come into the picture. In the case of multi-layered MoS₂ with S vacancy in each layer, an increase in the electron density of the unremoved S atoms has been observed which increases the repulsion of electrons between layers, increasing the interlayer distance as reported by Zhu et al.⁸⁹ Additionally, a direct band gap with Γ → Γ transition is observed in addition to a decrease in the band gap by introducing mid-gap defect in the forbidden region for the MoS₂ structures with S vacancies. This is quite different from the Γ → K transition as observed in bulk MoS₂ without defects. In the MoS₂ structures with S vacancies, the VBM is mainly composed of the 4d orbital of the Mo atom at the Γ point, while the doping band minimum is almost composed of the 4d orbital of the Mo atom and the 3p orbital of the S atom at the Γ point. This can be considered to be responsible for such direct nature of Γ → Γ transition. Such type of defect and 1T phase enriched Co-doped half-metallic MoS₂ possessing robust ferromagnetism may show greater potential for application in spintronic and detection-based devices.

Table 5 Local magnetic moments of the dopant Co and the six neighbouring Mo and total magnetic moments in MoS₂ multi-layered systems doped with various percentages of cobalt

Doping percentage (%)	Local magnetic moments (μB)		Total magnetic moments
			(μB)	(emu g^−1)
	Mo	Co
0	0	—	0	0
2	0.24	0.91	2.96	4.16
6	0.20	1.10	2.94	5.77
8	0.14	1.34	2.63	7.76

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3.9. Electrochemical measurement

To unravel the effect of Co-doping on MoS₂ NFs, the electrochemical performance of undoped MoS₂ and 2% Co, 4% Co and 8% Co-doped MoS₂ as electrode materials was characterized by cyclic voltammetry (CV), chronopotentiometry (CP) and electrochemical impedance spectroscopy (EIS) measurements. The measurement was performed with three-electrode equipment, which consists of Ag/AgCl as the reference electrode, Pt as the counter electrode and one glassy carbon electrode loaded with a sample on it as the working electrode. The prepared working electrode was analysed at different scan rates (10–100 mV s⁻¹) within a potential window of −0.6 V to 0.2 V (undoped, 2% Co-doped), 0 V to 0.6 V (4% Co) and −0.8 V to 0.6 V (8% Co-doped MoS₂) using a 1 M Na₂SO₄ electrolyte. The comparative CV representation of undoped and 2% Co, 4% Co and 8% Co-doped MoS₂ at a scan rate of 10 mV s⁻¹ is provided in Fig. 12(a). It is observed from Fig. 12(a) that 8% Co-doped MoS₂ shows the highest CV area at 10 mV s⁻¹ scan rate which indicates better electrochemical performance compared to other MoS₂ moieties. Therefore, the CV plots of 8% Co-doped MoS₂ at different scan rates (10, 20, 40, 60, 80, and 100 mV s⁻¹) are depicted in Fig. 12(b). The variation of specific capacitance (calculated from the CV profile) at different scan rates of undoped and 2% Co, 4% Co and 8% Co-doped MoS₂ is depicted in Fig. 12(c) and the corresponding numerical values are listed in Table S1 (ESI†). The greater capacitive behaviour of 8% Co-doped MoS₂ is attributed to the following factors: (1) the presence of a more metallic 1T phase resulting in high electrical conductivity and low internal resistance; (2) 1T phase being more electrochemically active in the edge planes;⁹⁰ and (3) increased defect levels,⁹¹ validated from XPS and steady-state UV-vis absorption spectra and BET surface area (for undoped and 8% Co-doped MoS₂ the BET surface area is 33.84 m² g⁻¹ and 43.47 m² g⁻¹ respectively), facilitating charge transport and adsorption of transported ions within 8% Co-doped MoS₂ leading to rapid redox reactions. The area under the CV curve shows a symmetric rectangular-like shape revealing the better choice for energy storage application as an ideal supercapacitor.⁹² The specific capacitance of the symmetrical supercapacitor was calculated from the CV profile following eqn (3):where is the area under the CV curve, m is the mass of the electrode materials, ν is the scan rate, and ΔV is the potential window taken. 8% Co-doped MoS₂ shows the highest specific capacitance of 217.9 F g⁻¹ at a voltage scan rate of 10 mV s⁻¹ which is approximately 1.7, 1.5 and 1.3 times greater than that of undoped (98.1 F g⁻¹), 2% Co, and 4% Co-MoS₂ electrode (123.7 F g⁻¹) materials, respectively. It has been observed that upon increasing the scan rate, the capacitive value decreases. Usually, at a higher scan rate, sufficient ion diffusion cannot take place within a constant time, and therefore, the capacitance value decreases.^92,93 Similarly, the galvanostatic charge–discharge (GCD) technique was carried out to further evaluate the capacitance behaviour of as-synthesized materials at different current densities (Fig. 12(d) and Fig. S4(b, d and f), ESI†). The GCD curve of 8% Co-doped MoS₂ shows the symmetric charging and discharging curve which reveals the good capacitive behaviour of the 8% Co-doped sample. The specific capacitance of the hybrid material was derived from the GCD curve following eqn (4):

C_sp = I × Δt/mΔV (4)

where I is the current (mA), Δt is the discharge time (s), m is the active electrode material (mg) and ΔV is the potential window (V). The specific capacitance values of undoped and 2% Co, 4% Co, and 8% Co doped MoS₂ electrodes with varying current densities obtained from the GCD plot are listed in Table S2 (ESI†). It has been observed that specific capacitance values obtained from CV and GCD measurements are nearly the same which indicates validation of our measurements and stability of the as-prepared electrode materials. To better understand the reason for the improved electrochemical performance of 8% Co-doped MoS₂, electrochemical impedance spectroscopy (EIS) of undoped and 2% Co, 4% Co and 8% Co-doped MoS₂ in the frequency range of 1 Hz–100 kHz was performed (Fig. 12(e)). The prepared MoS₂ nanostructures show the supercapacitor feature, i.e., a semicircle at higher frequencies, which is related to the surface area and electrical conductivity, and a straight line in the lower frequency region. The intercept on the real axis is ascribed to the equivalent series resistance (R_s), and the diameter of the semicircle corresponds to the charge transfer resistance (R_ct) in the electrode/electrolyte system. The R_s values for undoped and 2% Co, 4% Co and 8% Co-doped MoS₂ are 31.1 Ω, 24.8 Ω, 8.7 Ω and 18.9 Ω, respectively. With the increase in cobalt doping the R_s value decreases from low-doping (2% Co) to high-doping concentration (8% Co-doped MoS₂) which confirms that Co doping has enhanced the electrical conductivity of MoS₂ NFs. The 8% Co-doped MoS₂ electrode exhibits the smallest semicircle, indicating that the electrode has lower charge transfer resistance than the other two electrodes because the existing 1T phase and more available active sites decrease the charge transfer resistance. The 8% Co-doped MoS₂ electrode shows a nearly vertical shape because electrons could transfer fast as rapid ion diffusion occurs due to less resistance at the electrode–electrolyte interface. The low-frequency straight line shows the capacitive behaviour of MoS₂ nanomaterials (Fig. 12(e)). The highest slope of the straight line parallel to the imaginary axis is obtained for 8% Co-doped MoS₂, suggesting that 8% Co-doped MoS₂ has optimal performance among the prepared MoS₂ nanostructures. Table 6 presents the comparative electrochemical performance of MoS₂ (pure or in composite form when used as a single electrode material), as published in previous works, with the results obtained from our highest Co-doped (8% Co) MoS₂ NFs, indicating improved electrochemical behaviour of our material. Therefore, the doping-induced phase and defect engineering of MoS₂ NFs for enhanced electrochemical property in supercapacitors paves the way toward the production of more efficient electrode materials by designing Co-doped 1T phase and defect-rich MoS₂ ultrathin NF-based composites.

Fig. 12 (a) Comparative CV curves of undoped and 2% Co, 4% Co and 8% Co doped MoS₂ NFs, (b) CV scans of 8% Co doped MoS₂ at different scan rates, (c) plot of specific capacitance at different scan rates for undoped and 2% Co, 4% Co and 8% Co doped MoS₂, (d) GCD plot of 8% Co-doped MoS₂ at different current densities, and (e) Nyquist plots of undoped and 2% Co, 4% Co and 8% Co doped MoS₂ electrodes.

Table 6 Electrochemical performance of MoS₂ and its derived electrodes in supercapacitors

Nanostructured electrode material	Electrolyte	Specific capacitance (single electrode)	Ref.
MoS2 nanosheets	1 M KCl	148 F g^−1 at 1 A g^−1	94
MoS2/PEI–GO	2 M Na2SO4	153.9 F g^−1 at 1 A g^−1	95
MoS2/graphene	1 M Na2SO4	270 F g^−1 at 0.1 A g^−1	96
MoS2/graphene	6 M KOH	258 F g^−1 at 2 A g^−1	97
MoS2/a-C	6 M NaOH	163.7 F g^−1 at 0.5 A g^−1	98
MoS2/CNT	1 M Na2SO4	74.05 F g^−1 at 2 A g^−1	99
Carbon-MoS2 yolk−shell microspheres	1 M Na2SO4	122.6 F g^−1 at 1 A g^−1	100
8% Co doped MoS2	1 M Na2SO4	201.4 F g^−1 at 1.5 A g^−1	This work

---

4. Conclusion

The well-matched 1T phase incorporation in 2H-phase MoS₂ NFs using Co as a dopant at different weight percentages (0%, 2%, 4%, 6%, 8%) has been successfully carried out via a facile one-step hydrothermal technique. The evolution of the 1T phase, defects and vacancies and introduction of zigzag edges with Co doping have been verified using Raman, TEM and XPS measurement techniques. The robust ferromagnetism with Co doping was shown to have originated from ferromagnetic exchange interaction of Mo⁴⁺ and Co²⁺ with sulfur vacancies via the BMP model and magnetic moment contributed from zigzag edges. The asymmetry nature of the DOS and spin splitting near the Fermi level of Co-doped MoS₂ NFs contributed by d states of Mo and Co atoms and p states of neighbouring S atoms confirms that substitutional doping of the Co atom at the Mo site can effectively induce a spin-polarized state corresponding to ferromagnetism. The prominent feature of half-metallicity and formation of the defect level within the MoS₂ band gap has been evidenced from DOS results. In addition, various types of transport mechanisms existing in the different temperature regime have been explained and fitted well in the temperature variation resistivity curves. Moreover, Co-doping in MoS₂ shows better electrochemical performances due to the presence of more electrochemically active metallic 1T sites, edge defects, increased BET surface area and enhanced interlayer separation due to intercalated Co²⁺/Co³⁺ ions. This work demonstrates that Co-doped ferromagnetic 1T-MoS₂ NFs are promising for constructing Co-doped 1T-MoS₂-based heterojunctions for application in spintronic and detection-based devices as well as energy storage devices in the future.

Conflicts of interest

The authors declare no conflict of interest.

Acknowledgements

The authors acknowledge the XPS facility, Department of Physics, IIT Kharagpur, and Central Research Facility (CRF), IIT Kharagpur, for providing the infrastructural support for the characterization of the samples. R. Rahman would like to acknowledge the INSPIRE research fellowship provided by DST, India, to carry out the present research work.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d2ma00738j

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