Phase‐Restructured SnS₂ Nanodots: Unlocking ORR Activity through Thermal Stress‐Mediated Fracturing

ABSTRACT

The production of electrocatalytically active materials from inert precursors through intricate engineering represents a frontier in sustainable energy research. The present study unveils a transformative thermal stress‐mediated fracturing strategy to activate the oxygen reduction reaction (ORR) in SnS_2, a material traditionally considered unsuitable for such electrocatalytic applications. By subjecting bulk SnS₂ hexagonal plates to extreme thermal shock via rapid heating followed by instant quenching in liquid nitrogen, a simultaneous reduction in dimensionality and a critical 1T‐to‐1H phase transition has been achieved, yielding ultra‐small 1H@1T‐SnS₂ nanodots (2–5 nm). High‐resolution aberration‐corrected STEM imaging reveals a restructured atomic lattice where the emergence of the 1H phase, coupled with plentiful edge sites, fundamentally reconfigures the electronic landscape. While the parent 1T‐SnS₂ follows an inefficient 2‐electron pathway, the phase‐restructured nanodots exhibit a robust 4‐electron oxygen reduction mechanism. Theoretical insights from Density Functional Theory (DFT) confirm that the synergy between the 1H‐phase transition and edge‐site enrichment lowers the Gibbs free energy for intermediate adsorption, effectively “unlocking” the catalytic potential of SnS₂. This work provides an erudite blueprint for the phase‐and‐size engineering of non‐precious metal dichalcogenides, transmuting mundane semiconductors into high‐performance electrocatalysts for the next generation of fuel cells and metal‐air batteries.

Thermal shock generated phase and size engineering of SnS₂ toward electrocatalysis, unlocking the oxygen reduction reaction performance.

1. Introduction

With the discovery of graphene, the wondrous journey of the layered materials has indeed been perpetuated in various forms like double hydroxides, Mxenes, layered halide perovskites, transition metal dichalcogenides (TMDs), and single‐atom layers. Strikingly, of all these layered materials, TMDs have sprung up to exceptional predominance owing to the uniquely rich blend of electronic, physicochemical, and catalytic properties they exhibit, with applications in various fields [1, 2, 3, 4, 5]. Inevitably, this unprecedented magnificence exhibited by TMDs, especially MoS₂, provoked a greater urge for finding more and more layered functional materials. Consequently, a ubiquitous question has inevitably popped up as to what happens when the transition metal in the TMDs is replaced by any other metal. The intensive research work published so far, indeed, reveals the emergence of amazingly rich, exotic, and new physico‐chemical properties in the ultrathin limit when a group III (Ga, In) or a group IV (Sn) or a group V (Sb, Bi) metal replaced the transition metal. Eventually, a new class of layered metal dichalcogenides (LMDs) emerged, endowed with immense possibilities of leveraging intense superiorities in future applications [6, 7, 8, 9].

Of the numerous LMDs explored so far, tin disulfide (SnS₂), with structure and properties akin to MoS₂, has been able to garner special interest in them. SnS₂, occurring naturally as berndtite, is an n‐type semiconductor with a band gap of 2.2 eV‐2.4 eV [10]. Contrary to most TMDs, SnS₂ doesn't show any indirect to direct crossover in bandgap as it is trimmed down to the monolayer limit, rather its value increases to 2.41 eV [11]. Owing to its low cost, natural abundance, low toxicity, structural peculiarities, and manifold physicochemical virtuosities [12, 13, 14], SnS₂ has found widespread applications in electronics, optoelectronics, sensors, air purification, wastewater treatment, N₂ fixation, CO₂ reduction, energy storage, etc [15, 16, 17, 18, 19, 20, 21, 22, 23]. However, in an era manifested by rational craving to develop sustainable and clean energy, the use of SnS₂ in energy generation has been distinctly sparse and sporadic, with preponderance being entirely on photocatalytic hydrogen generation [24, 25]. Interestingly, earlier studies suggest the unsuitability of SnS₂ as an oxygen reduction reaction (ORR) or oxygen evolution reaction (OER) catalyst, but identified a possibility of hydrogen evolution reaction (HER) [26] in it. Recently, a few reports on electrocatalytic HER have emerged to raise a reassured prospect of finding a good electrocatalyst in SnS₂ [27, 28, 29].

Interestingly, the conduction and valence band energies in a single crystal SnS₂ bridge the oxidation and reduction potential of water [14]. In addition, SnS₂ has carrier concentrations of 10¹³ cm⁻³ to 10¹⁷ cm⁻³ [30], high carrier mobility (18 to cm²V⁻¹s⁻¹) [10, 13], high electron affinity of 4.16 eV [14], and work function (lower than Pt) [14]. These traits of SnS₂ are seemingly conducive to achieving superior electrocatalytic performance in ORR. Yet, surprisingly, SnS₂ remains inapposite to ORR with no reports being published so far in the domain of ORR research. However, previous reports suggest that in SnS_2, the edges are more reactive than the basal plane [31] and can be effectively manipulated to induce desired changes in properties through different edge terminations at various S‐coverages and orientations [31]. In fact, Liu et al. had recently reported efficient and stable conversion of CO₂ in edge‐site‐enriched SnS₂ nanoplates [32]. The inherent structural peculiarities and anisotropies in SnS₂ may bestow propitious opportunities to foster newer functionalities in it through scrupulous engineering. Most interestingly, SnS₂, which resembles the CdI₂ crystal structure [10] with the central Sn atom bonded to six other S‐atoms occupying the octahedral sites to form a closely packed hexagonal unit cell, can have a large number of stacking‐polytypes. Theoretically, SnS₂ can have as many as 70 stacking polytypes with diversified features through different stacking sequences [14, 33]. However, phase restructuring has been very rarely looked upon as a means to stimulate ORR activity in 2D materials. Interestingly, in recent times, there have been two reports showing that the incorporation of metallic 1H phase improved the CO₂ reduction properties of SnS₂ [16, 17]. So far, phase engineering in 2D layers has been achieved mostly through laser or e‐beam irradiation, Ar‐plasma exposure, intercalation methods, chemical functionalization, and extremely high temperatures and pressures [34, 35]. Even though widely used, these techniques are harsh, complex, low‐throughput, uneconomical, and lack stability and purity of the phases achieved. Whilst size reduction/edge tailoring/ defect engineering /phase restructuring has been independently operated as individual techniques deemed to improve catalytic activities, a detailed, comprehensive study that could holistically expound the possibility of enormously improved performance when these mechanisms work in tandem is by far lacking. Therefore, devising a simplistic, sustainable, non‐hazardous, cost‐effective technique in which all of size reduction, edge tailoring, defect engineering, and phase restructuring can be achieved concomitantly in a single procedure is imperative and a clamant demand in order to effectuate an impressively augmented catalytic performance.

Herein, for the first time, we devise a facile and cost‐effective technique to reduce the dimensions of the layers as well as to cut down the layer numbers through a thermal stress‐mediated fracturing process in which phase restructuring together with drastic modification in the defect states and edges can be achieved concomitantly in a single process as a natural sequel to the art of making. In fact, in this work, we report the achievement of stable and scalable conversion of 1T to 1H phase SnS₂, as 1H‐phase rich 1H@1T‐SnS₂ nanodots were prepared from bulkier 1T‐SnS₂ hexagonal plates through our newly contrived and unique thermal stress‐assisted fracturing method. In contrast to the previous three works [16, 17, 36] which reported the achievement of the metastable 1H‐phase of SnS₂ through harsh, complicated methods like e‐beam irradiation, hydrogen assisted calcination and CVD grown techniques [16, 17, 36], we have achieved a stable 1T‐phase to 1H‐phase conversion through facile, non‐hazardous, purely mechanical, cost‐effective, surfactant and solvent free technologically viable technique with high yield. Even though SnS₂ has been used as an archetype for the proposed thermal stress‐mediated fracturing process, this technique can be very well extended to other VdW materials to achieve simultaneous control over dimensionality and crystallographic phase restructuring. It may be predicted that a precise tuning of the parameters of control, like temperature gradient, time of exposure, the number of cycles of repetitions, etc., could effectively manoeuvre the craftsmanship towards achieving the intended dimensionality, the extent of phase restructuring, and defect incorporation. The phase restructured nanodots exhibited a drastically dropped band gap energy of 0.856 eV, along with copiously abundant ${\mathrm{S}}_2^{2-}$ species, suggestive of a possible rearrangement in its electronic configuration and engineered edge sites, which could, in a way, greatly aid their applications in electrocatalytic activities. The phase and size engineered SnS₂ nanodots show remarkably enhanced electrocatalytic ORR activities in 0.1 m KOH. Unlike bulkier 1T‐SnS₂ plates, which showed poorly performed ORR activities through a 2‐electron transfer pathway in alkaline medium, the1H@1T‐SnS₂ nanodots adopted the strongly preferred 4‐electron transfer mediated mechanism with largely enhanced current densities and high selectivity. Our experimental studies and DFT calculations for the first time comprehensively revealed that in SnS₂, the 1H‐phase indeed acts as an effective promoter towards ORR activity. This process of making nanodots with a restructured phase is a unique technique endowed with immense futuristic prospects in the large‐scale production of phase‐restructured 0D functional materials.

2. Results and Discussion

Thick hexagonal plate‐like structures of SnS₂, henceforth referred to as SnS₂ plates, were synthesized by facile hydrothermal method (Section S1.2). The scanning electron microscopy (SEM) study on these plates revealed that the plates were very thick with a lateral size far greater than 200 nm (Figure S1). These so obtained SnS₂ plates served as the precursor to our newly devised thermal stress mediated fracturing process in which we adopt a ‘tear and wear’ strategy and intend to fragment these plates to their 0D versions. Indeed, these forms, hereafter referred to as SnS₂ nanodots, are produced from the plates enormously truncated along their lateral dimensions in tandem with coinciding curtailment on their layer numbers. The details of the procedure of making of SnS₂ nanodots are given in Section S1.2 and schematically illustrated in Figure 1a. The rapid heating of SnS₂ plates in the first step, causes the layers to swell up as a result of massive thermal stress built up within them. The local tensile forces those come into play as an aftermath, would inevitably overpower the interlayer van der Waal's force and as a matter of course, thinner layers of SnS₂ are peeled off from their bulky parent. In the subsequent step, when very hot SnS₂ is instantly quenched in liquid N₂, small cracks would emerge in the enlarged layers as they try to contract abruptly in reaction to massive thermal shock they had to encounter. In the succeeding cycles of repetition, these cracks would proliferate, widen up and propagate leading to eventual segmentation of the plates into smaller fragments. The impact of these sudden and ultrafast cycles of successive expansions and contractions is very likely to impart cleavages into the intralayer bonds in addition to concurrent lateral constrictions. The process leads to eventual disintegration of the thick SnS₂ plates into their nanodot version. Apart from this, the mechanism of action intrinsic to this procedure of fracturing seemingly enhances the likelihood of achieving a stable phase transition in the sample along with a vastly altered landscape of inherently existent defects. In view of the large and asymmetric local stresses, originating from the huge thermal shock imparted, there would be veritable alterations in the atomic bonds lengths which would cause denser concentration of atoms at local acute angle corners. This would eventually trigger the gliding of the Sn‐ and S‐ atomic planes towards areas with lesser concentration of the atoms so as to release the massive stress instilled in them. The consecutively repetitive gliding of Sn‐, S‐, Sn+S atomic planes would possibly give rise to an entirely different co‐ordination environment concomitant to phase restructuring. Apart from the simplistic gliding of atomic planes, the rapturing process would potentially involve numerous atomic displacements as well. This preconceived proposition of an act of displacement of atoms, in turn assertively posits the accomplishment of an entirely different electronic co‐ordination intrinsic to phase restructuring.

FIGURE 1.

(a) Schematic representation of the thermal stress mediated rupturing procedure. The components of the schematic were generated by copilot software and remastered to build up the schematic representation. (b) X‐ray diffraction patterns of SnS₂ plates and SnS₂ dots (JCPDS No. 023–0667). (c) The Raman spectra for SnS₂ plates and SnS₂ nanodots. The inset shows the magnified spectra within the range of 182 cm⁻¹ to 220 cm⁻¹ for A_1g Raman mode.

The powder X‐ray diffraction technique (XRD) was employed to decipher the crystal structure existent in the formed SnS₂ samples as well as ascertain the phase purity. As seen from Figure 1b, the primary diffraction peaks of SnS₂ plates can be well indexed to hexagonal SnS₂ (JCPDS: 023–0677) confirming that pure and well crystallized SnS₂ has been obtained [9, 37]. In SnS₂ nanodots, however, the intensity of the diffraction peak due to the (001) plane, as evident from Figure 1b, had immensely diminished with a significant increase in FWHM from 0.231° to 0.366°. This can very well imply an appreciably large reduction in overall size of the nanodots alongside a great inhibition of growth in c‐direction thereby restricting growth to only a few layers [38, 39, 40], as preconceived and desired in our approach. The interplanar spacing d corresponding to the (001) planes as calculated by SE16 show an increase to 0.603 nm in SnS₂ nanodots as compared to 0.591 nm in the SnS₂ plates indicating existence of lattice strain. Furthermore, the intensities of the remaining diffraction peaks, other than the one due to the (001) facet of SnS₂, had also greatly reduced in the SnS₂ nanodots, impelling them to almost virtually disappear from the XRD spectra which in consequence indicate an extensive diminution in the sizes of the nanodots [38, 40].

Raman spectroscopy can further provide a comprehensive insight into the occurrence of phases in SnS₂ plates and nanodots. As seen from Figure 1c, a very intense and sharp peak located at $\sim 311.29$ cm⁻¹ along with a very little hump near 203.13 cm⁻¹ is observed in SnS₂ plates. According to earlier reports, the peak at $\sim 311.28$ cm⁻¹ can be assigned to the out of plane A_1g mode of vibration for thick samples of SnS₂ in 1T‐phase while that at 203.13 cm⁻¹ is due to the in‐plane vibration of E_g mode³⁷. Any presence of SnS₂ in 4H phase can be right away negated as we don't see the occurrence of a doublet near 200 cm⁻¹ and 214 cm⁻¹ in SnS₂ plates [8]. Likewise, the position of the A_1g mode for SnS₂ plates at value far lower than the reported values $\sim 314$ cm⁻¹ for 2H phase in SnS₂ also invalidates the likelihood of presence of 2H‐phase in traces [41]. Therefore, it may be presumed that the SnS₂ plates are purely in 1T phase. As can be seen from Figure 1c, there is not much of difference in position of the A_1g Raman peak for the SnS₂ nanodots in comparison to that of SnS₂ plates with just an inappreciably slight upshifting of $\sim 0.48$ cm⁻¹. However, as can be seen from the inset of Figure 1c the intensity of the little hump corresponding to the E_g mode in 1T phase, had intensely reduced which in accordance to earlier reports can be attributed to decrease in number of layers [8, 37]. Moreover, the intensity of the A_1g peak has also greatly reduced in case of the nanodots with the corresponding difference (ΔK) between the A_1g and E_2g peaks being reduced from 108.15 cm⁻¹ (in case of SnS₂ plates) to 106.92 cm⁻¹ (in case of SnS₂ nanodots) which further reaffirms a decrease in the number of layers in line with previous reports [37]. Therefore, in congruence with X‐ray diffraction studies, these findings can further reinstate our claims of thinning of the SnS₂ nanodots to almost 1–2 layers. In order to uncover the structural features, transmission electron microscopy (TEM) was done on both SnS₂ plates and nanodots. The TEM images (Figure S2a) and SAED pattern (Figure S3a) of the SnS₂ plates on the other hand reaffirmed that they acquire hexagonal plate like structures with high crystallinity. The HRTEM image of the SnS₂ reveal a lattice fringe separation of $\sim d=0.329$ nm which can be ascribed to the (100) crystal planes in 1T‐SnS₂ in congruence to previous reports [37, 42]. The low‐resolution TEM images (Figure 2a) and high‐resolution TEM (HRTEM) images (Figure 2b,c) of the SnS₂ nanodots, on the other hand, revealed nearly circular shapes of the nanodots with visible imperfections in the boundaries (Figure 2c). However, the SnS₂ nanodots displayed some degrees of polydispersity in their lateral dimensions as the diameters of the nanodots ranged from 2 nm to 5 nm (Figure 2a). The HRTEM image of SnS2 nanodots as shown in Figure 2c recognizes a lattice fringe spacing of $d\sim 0.338$ nm. The corresponding SAED pattern, indexed with respect to [010] zone axis using CrystBox software [43] reveal the high crystallinity of the samples with visible resemblance to the hexagonal crystal structure in the nanodots. The energy dispersive spectroscopy (EDS) and elemental analysis reveal Sn:S atomic ratios of Sn and S as 1.75 and 1.67 for SnS₂ plates and nanodots respectively (Figures S4 and S6). The corresponding EDS mapping images are shown in Figures respectively (Figures S3b–d and S5a–d). The UV–vis spectra and the PL spectra are shown in Figure S9. As can be seen from the Uv–vis spectra, there is a distinctive shift of the absorption edge of SnS₂ nanodots to longer wavelength as compared to the SnS₂ plates. Further, the PL‐emission peak located at 325.41 nm for SnS₂‐nanodots becomes completely extinct in the SnS₂ plates. This, in seriatim unequivocally foresee a possible reorganization in the electronic arrangements which require further in‐depth and detailed investigation for precise interpretation. Therefore, we carried out Scanning tunnelling microscopy (STM) and scanning tunnelling spectroscopy (STS) experiments on the SnS₂ nanodots in order to figure out the true nature of shapes of the nanodots, their thickness and their electronic configuration. Figure 2e shows STM image of a single SnS₂ nanodot with a lateral size of approximately 4 nm with a thickness of  ≈ 0.225 nm. The height profile details of this are given in Figure S7a,e. Further, Figure S7b–h shows STM images with relevant height profiles where we captured groups of 3–5 SnS₂ nanodots of 2 to 3 nm in lateral dimension and $\sim$1.5 nm in the vertical direction. Therefore, it is worthy to note here that the lateral fracturing restricts the lateral length of SnS₂ nanodots to 2–3 nm and the vertical height indicates that typically the nanodots are comprised of 2‐ 3 SnS₂ layers [44]. These findings are well in harmony with our results from TEM micrographs. With intent to probe the electronic structures of the SnS₂ nanodots, we have carried out I–V measurements on the selected particle (Figure 2e). The data is the average of several measurements taken from different positions of the nanodot. Calculated $dI/dV$ vs V variation for the SnS₂ nanodots (Figure 2f) clearly indicate the bandgap energy of the present SnS₂ structures as 0.856 V which is well below the reported values for layered 1T‐SnS₂ so far [14]. But there are previous reports which show a variation in band gap energy of SnS₂ with phase transition. Therefore, in congruence with the trends predicted by UV–vis and PL spectroscopy, the lower bandgap energy in case of the SnS₂ nanodots is indicative of a possible change in electronic configuration demanding an in‐depth investigation into the crystal structure of the sample and comprehensively deciphers the anomalous behaviour.

FIGURE 2.

(a) Low resolution TEM image of SnS₂ nanodots. The inset shows the particle size distribution histograms obtained from TEM micrograph(a) and the fitted normal distribution curve. (b,c) HRTEM images of SnS₂ nanodots. The lower inset in (c) is the magnified and processed portion of the nanodot marked 2 for improved clarity and analysis. The upper inset in (c) is the intensity profile corresponding to the blue line drawn in lower inset of (c) for determining the interplanar distance d in the nanodots. (d) SAED pattern of SnS₂ nanodots with the inset showing the region of acquisition marked in the white rectangle. (e) STM image of SnS₂ nanodot deposited on HOPG. The inserted color bar is a guide to the eyes for the height profile reference of the nanodot. Steady state STM imaging was performed in the constant current mode with a set point of 950 pA and a bias voltage of 0.8 V. The inset in (e) shows the I–V profile for the nanodot obtained from STS spectroscopy. (f)The corresponding averaged $dI/dV$ profile for the nanodot from STS spectroscopy.

In order to have a clear view of the crystal arrangements in both samples at the atomic level, aberration corrected high‐angle annular dark field scanning transmission electron microscopy (HAADF‐STEM) were carried out on both of them. The experimental HAADF‐STEM images of SnS₂ plates and the nanodots as evident from Figure 3a–e distinctly show two different patterns of atomic arrangements in them. The as obtained experimental HAADF‐STEM image and enlarged view of a portion of the same image of the SnS₂ plates (Figure 3a) indexed to [001] zone axis of SnS₂ suggest that they are purely in 1T‐phase in harmony to earlier studies [16, 36] and aptly correlate to the atomic model depicted in Figure 3i,l The atomic intensity ratios corresponding to the light golden line drawn in Figure 2a confirms the same. On the other hand, enlarged portions from the as obtained experimental HAADF‐STEM image (Figure 3c) of the SnS₂ nanodots delineate two distinctive patterns of atomic assemblage. The line profile corresponding to the blue line in Figure 3c perfectly matches to the atomic model as illustrated in Figure 3j,m and establishes the formation of pure 1H‐phase SnS₂ in majority of the sample space [16, 36]. However, in some smaller sections we observe big and significant differences in atomic intensity ratios between the peaks corresponding to the line profile of the light green line in Figure 3c. This in all probability can be attributed to a mixed phase of 1H and 1T which is residual to an incomplete phase transition in that part [17]. The atomic model corroborative to this sort of crystal arrangement in a mixed 1H@1T phase is shown in Figure 3k,n. In order to elucidate the fact that we have actually obtained 1T to 1H phase transition in the SnS₂ nanodots, other HAADF‐STEM images from different random regions of the sample with different magnifications along with line profiles are provided in the supporting information (Figure S10a–f). The fast Fourier transformation (FFT) of the atomic images (Figure 3a,c] exhibit almost similar kind of hexagonal symmetry without much difference and is shown in Figure 3p,q), respectively. This happens because both of 1H‐SnS₂ and 1T‐SnS₂ are made up of the same unit cells and have same long‐range ordering [36]. The 1T phase of SnS₂ resembles a CdI₂‐type structure that belongs to P3m1 space group, and the 1H replicate a MoS₂‐like structure and belongs to P63mc space group. In the 1T phase, Sn atoms are octahedrally co‐ordinated by six S atoms as shown in Figure S14c, with adjacent S‐Sn‐S tri‐layers arranged in an ABC stacking sequence as shown in Figure S14a. By contrast, in the 1H phase, Sn atoms are coordinated in a trigonal prismatic environment as shown in Figure S14f, and the tri‐layers follow an ABA stacking sequence as shown in Figure S14d. In both cases, strong covalent Sn─S bonds hold each tri‐layer together, while adjacent tri‐layers are coupled only through weak van der Waals interactions along the crystallographic c‐axis. In view of the striking resemblances, HRTEM, XRD or SAED couldn't have brought out the subtle changes in the atomic structure of 1H and 1T SnS₂ and it could have been only precisely differentiated through atomic resolution HAADF‐STEM images. The incorporation of 1H phase in the SnS₂ nanodots gives rise to lattice strain as indicated by our XRD analysis.

FIGURE 3.

The as obtained experimental HAADF‐STEM images of (a,b) SnS₂ plates and (c–e) SnS₂ nanodots. (b)Zoomed region enclosing the golden line from (a) for clearer view of atomic arrangements to the eyes. (d,e) Magnified view of two different regions in (c), one surrounding the cyan line and the other circumambient to the light green light in order to demarcate two distinct patterns of atomic arrangements. The white scale bar in (b,d,e) represents a length of 1 nm. (f,g,h) Intensity profiles corresponding to the golden, cyan and light green lines in (a and c). The colour scheme of the plot profile in (f,g,h) corresponds to the individual line colours. Side views of models for atomic arrangements in (i) 1T‐SnS₂, (j) 1‐ SnS₂ and (k) Mixed phase of 1T and 1H SnS₂. Top views of atomic models of (l) 1T‐SnS₂, (m) 1H‐ SnS₂ and (n) Mixed phase of 1T and 1H SnS₂ (p) FFT patterns of SnS₂ plates and (q) FFT patterns of SnS₂ nanodots corresponding to panels (a) and (c) respectively. (o) represents colour scheme for Sn and S atoms for 1T and 1H phase of SnS₂ respectively.

To acquire comprehensive understanding of the surface elemental composition and the chemical states of the synthesized samples X‐ray photoelectron spectroscopy (XPS) studies were done on both of them. Strikingly, the high resolution XPS spectra for the SnS₂ nanodots (Figure 4a,c) reveal that the peaks corresponding to both Sn 3d and S 2p regions have been visibly shifted downwards lower binding energies in comparison to the peaks of SnS₂ plates (Figure 4b,d). This downshifting can in no way be attributed to formation of larger number of S‐vacancies that would have caused the Sn 3d and S 2p peaks to move in opposite directions corresponding to differently oxidized/reduced S and Sn centres [45, 46] which does not occur here. Further, any likelihood of conversion of SnS₂ to SnS can be right away eliminated as that would have made the peaks shift by energies greater than 1 eV [47]. In addition, the separation between the Sn 3d_3/2 and Sn 3d_5/2 peaks is ≈8.4 eV indicating that Sn in both samples exist predominantly in the Sn⁴⁺ state, consistent with earlier reports on XPS analyses of pure SnS₂ [48]. The possibility of formation of the Sn oxides is very much unlikely to have occurred as we don't see any Sn─O bond in the high‐resolution XPS spectra of the O 1s region for both the samples (Figure S13). Therefore, in congruence with our HAADF‐STEM findings, this down‐shifting of the XPS peaks in case of the SnS₂ nanodots can be indubitably assigned to the 1T to 1H phase transition of SnS₂ in the nanodots. The deconvolution of the Sn 3d (Figure 4b) and S 2p regions (Figure 4d) identifies that the SnS₂ plates are purely in 1T phase with the Sn 3d_5/2 and Sn 3d_3/2 peaks located at 486.67 eV and 495.063 eV respectively [37]. On the other hand, the S 2p_3/2 and the S 2p_1/2 peaks were located at 161.454 eV and 162.648 eV, respectively. The little hump at 163.24 eV can be associated with bridging ${\mathrm{S}}_2^{2-}$ species in accordance to earlier works [49, 50, 51]. The bridging ${\mathrm{S}}_2^{2-}$ entities possess higher electron binding energies as compared to other types of S‐active atoms and are prominently known for being superbly active catalytic sites [52]. In SnS₂ nanodots, however, the high resolution XPS peaks corresponding to both Sn 3d (Figure 4a) and S 2p (Figure 4c) orbit signals are significantly broadened when compared to those in the SnS₂ plates and therefore must be accommodating additional peak. Deconvoluting the Sn 3d peaks for SnS₂ nanodots yield a pair of doublets with one doublet being positioned at 494.729 eV and 486.293 eV and the other doublet stood at 495.316 eV and 486.91 eV respectively. In congruence to earlier reports, the doublet at 494.729 eV and 486.293 eV originates because of the emergence of 1H‐phase of SnS₂ in the nanodots whilst the later doublet located at energies higher by ≈ 0.6 eV is due to the 1T‐phase [16, 17]. Likewise, the S 2p region for the nanodots can be resolved with a pair of doublets located at 161.07 eV and 162.265 eV due to 1H‐phase and 161.76 eV and 162.96 eV for 1T‐phase respectively [16, 17]. The hump corresponding to ${\mathrm{S}}_2^{2-}$ state blue shifts to163.53 eV in the nanodots. The integrated area ratio of Sn 3d and S 2p peaks, normalised by relative sensitivity factors (R.S.F.) is utilized to obtain atomic ratios of Sn to S as 1.796 and 1.759 in SnS₂ plates and nanodots respectively. Therefore, we find that there is no drastic variation in content of S‐vacancies in the two samples as the number of S‐vacancies rise from 10.2% in the SnS₂ plates to 12.05% in SnS₂ nanodots respectively. But interestingly, there is an overwhelming increase by168.59% in the integrated area ratios of ${\mathrm{S}}_2^{2-}$ peak normalised by R.S.F. as we move from SnS₂ plates to the SnS₂ nanodots. This considerably large increase of the unsaturated ${\mathrm{S}}_2^{2-}$ peak in addition to decrease in layer numbers might actually imply a substantiative rise in number of bridging and terminal ${\mathrm{S}}_2^{2-}$ ligands [51]. In consequence, this effectively can indicate presence of larger number of uncoordinated Sn and S‐atoms together with proliferation in chemically active sites at the edges [51]. Finally, the survey spectrum from XPS of SnS₂ plates and SnS₂ nanodots are shown in Figures S11 and S12, respectively. Overall, our XPS investigations confirm to the achievement of stable 1T to 1H phase conversion and increased presence of ${\mathrm{S}}_2^{2-}$ ligands in SnS₂ nanodots.

FIGURE 4.

High‐resolution XPS spectra corresponding to the Sn 3d region in (a) SnS₂ nanodots and (b) SnS₂ plates. High Resolution XPS spectra of S 2p region in (c) SnS₂ nanodot and, (d) SnS₂ plates.

Further, the synthesized phase restructured SnS₂ nanodots were employed to study electrocatalytic ORR in them in great detail. Cyclic voltammetry (CV) studies were carried out in separate solutions of 0.1 m KOH saturated with Ar and O₂ respectively within the voltage span of 0 V–− 0.8 V (vs Ag/AgCl) in order to primarily estimate the performance of the samples towards electrocatalytic reduction of O₂. The voltage sweep rate was kept fixed at 50 mV/sec in the measurements. It is evident from Figure 5a that the SnS₂ nanodots gave rise to a distinctly intense cathodic peak at −0.43 V (vs Ag/AgCl) in oxygen saturation, unveiling their competence in electrocatalytic ORR activity. In contrast, no such reductive peak was observed in argon saturated solution. Furthermore, to get an insight into the kinetics associated with the seen ORR activity, Linear sweep voltammetry (LSV) measurements were recorded using a rotating disk electrode (RDE) at different rotation speeds and scan rate fixed at 10 mV/sec in 0.1 m KOH solution saturated with O₂. The obtained current densities were normalised to the geometrical area of an electrode of 5 mm diameter. Figure 5b shows ORR polarization curves acquired at 1600 for SnS₂ plates and SnS₂ nanodots. A typical LSV curve for a Pt/C (benchmark electrocatalyst for ORR) is provided in the inset for reference. Clearly, the phase restructured SnS₂ nanodots exhibited superior ORR activity in comparison to SnS₂ nanoplates with a possible 4‐elctron transfer pathway, the more desired essentiality for energy related applications. 1T‐SnS₂ is seen to exhibit poor current density and was found to adopt a 2‐electron transfer mediated mechanism. However, the current density in case of the nanodots increased manifold compared to that of 1T‐SnS₂ plates. The increased current density in case of the 1H@1T‐SnS₂ nanodots may be attributed to the interplay between proliferation in number of catalytically active sites. The onset potential (defined as the potential at which the current density is −0.05 mA/cm²) [54] for the nanodots is −0.219 V (vs Ag/AgCl). Figure 5c represent RDE voltammograms of SnS₂ nanodots composed electrode for ORR at different rotating rates. The current density values are seen to show monotonic increase with increasing rotating speeds possibly due to reduction in diffusion lengths at higher rotation speeds [55]. The different kinetic parameters for 1H@1T‐SnS₂ nanodots, counting the electron transfer number (n) and kinetic current density (J_k ) were evaluated applying the Koutecky‐Levich (K‐L) equation. Figure 5d shows the K‐L plots for 1H@1T‐SnS₂ nanodots at potential window of −0.65 V to −0.85 V. The linearity of the Koutecky‐Levich plots and the near parallelism of the fitted lines indicate that the oxygen reduction by 1H@1T‐SnS₂ nanodots follow the first‐order kinetics and almost identical n values for different potentials.

FIGURE 5.

Electrocatalytic performance of SnS₂ nanodots towards ORR. (a) CV curves of SnS₂ nanodots in Ar and O₂ saturated 0.1 m KOH. (b) RDE polarisation curves for SnS₂ nanodots and SnS₂ plates at a rotational speed of 1600 rpm. The inset in (b) shows typical LSV curve of benchmark catalyst Pt/C under identical conditions [taken from ChemSusChe https://doi.org/10.1002/cssc.201800982 [53]. (c) Polarisation curves for SnS₂ nanodots in O₂‐saturated 0.1 m KOH at various rotation speed of the disk electrode. (d) Koutecky‐Levich plot for SnS₂ nanodots at different voltages and derived from (c). The inset shows the variation of electron transfer number n with voltage.

The electron transfer number of value $\sim 4$ for different potentials, calculated from the Levich slopes clearly indicate a direct one‐step 4‐electron transfer mediated oxygen reduction pathway adopted by the 1H@1T‐SnS₂ nanodots (Figure 5d). In order to ascertain the selectivity and further corroboration of previously mentioned results RRDE (rotating ring disk electrode) measurements were performed. As seen from Figure 6a, the ring current (I_r ) values are considerably insignificant when compared to the disk current (I_d ) values. The hydrogen peroxide (H₂O₂) yield and the corresponding electron transfer numbers computed from equations (SE3) and (SE4) are displayed in Figure 6b. The peroxide yield was well below 26% over the voltage range of −0.5–0.9 V. The electron transfer numbers calculated from equation (SE4) are in good agreement with those obtained from the RDE quantification derived K‐L plot, thereby confirming that the kinetics of ORR activity exhibited by the 1H@1T‐SnS₂ nanodots occur through the greatly intended 4‐electron transfer mediated pathway. Here, in case of the SnS₂ nanodots, the profusely occupying catalytically active sites, achieved through reduction in dimensions of SnS₂, can significantly boost the interfacial electron transfer between the surfaces of the catalyst and the intermediates. That effectively decreases the activation energy required to dissociate O₂ to the intermediates. This, together with incorporation of the 1H‐phase, in turn may help out to effectively foster a direct 4‐electron transfer mediated electrocatalytic ORR activity with exalted results. Intriguingly, the S‐vacancies, the most prolific intrinsic defect states in SnS₂, apparently doesn't seem to contribute overwhelmingly towards drastic transformation of ORR performance as there is only a small (≈ 2%) increase in their numbers as we move from the SnS₂ plates to nanodot configuration ‐a result expounded by our XPS analysis.

FIGURE 6.

(a) LSV of SnS₂ nanodots acquired on a RRDE at a rotational speed of 1000 rpm. (b) Corresponding variation of electron transfer number and the hydrogen peroxide yield with different potentials. (c) Chronoamperometric measurements of SnS₂ nanodots and commercially available Pt/C at ‐0.6 V (V vs Ag/AgCl). The insets show the methanol immunity of SnS₂ nanodots. (d) The derived Tafel plots for SnS₂ nanodots and commercially available Pt/C.

Stability of ORR catalysts in harsh conditions has been a long‐standing issue of massive concern so far. In order to estimate the durability of the synthesized nanodots catalysts, chronoamperometric studies were carried out under a constant potential of −0.6 V at a rotational pace of 1000 rpm in O₂ saturated electrolyte. Figure 6c, explicitly depicts that the 1H@1T‐SnS₂ nanodots had long‐term durability in 0.1 m KOH with meagre and tardy attenuation in the current values after hours. That way, as regards the stability, the1H@1T‐SnS₂ nanodots outperformed the commercially available Pt/C which showed comparatively larger decline in the current values within the same interval of time and under same experimental conditions. In addition, the SnS₂ nanodots were also checked for methanol crossover tolerance by chronoamperometric response aftermath of addition of 2% (v/v) methanol to the single compartment test cell. As a matter of fact, methanol crossover remains a critical issue hitherto as it effectively acts as a massive deterrent towards commercialization of the fuel cell systems. In methanol crossover, unreacted methanol permeates through the anode to the cathode to eventually get oxidised and induce a mixed sort of potential at the cathode. Therefore, instead of just promoting the desired ORR, a simultaneously competing oxidation reaction of methanol would be enforced. Consequently, the cathodic ORR performance gets critically decreased accompanied by an overall lowering in fuel cell efficiency. Therefore, developing electrocatalysts with remarkable ORR performance exhibiting simultaneous excellent methanol tolerance is imperative at this point of time. As apparent from inset of Figure 6c, the SnS₂ nanodots, suffered a little sink in the current values on methanol addition, but retrieved to virtually same current values and strikingly in no time. This very well suggest that no electrooxidation of methanol occurs in case of the SnS₂ nanodots. Pt/C, the benchmark ORR electrocatalyst, on the other hand, is seen to suffer an irrecoverable plunge in current values because of electrooxidation of methanol. In addition, the other inset in Figure 6c displaying the LSVs of SnS₂ nanodots recorded in presence of 3 m CH₃OH, show no rise in anodic current or emergence of distinct oxidation peak. This further reinstates that no electrooxidation of methanol occurs in case of SnS₂ nanodots. Moreover, the onset potential for ORR in case of SnS₂ nanodots virtually remain invariant under methanol dosage as evident from the LSVs. The stability of the SnS₂ nanodots was further evaluated by acquiring X‐ray diffraction and TEM patterns after completion of electrochemical ORR which show no significant change (Figures S14 and S15).

In order to comprehensively understand the kinetics involved in the extraordinary ORR activity demonstrated by the 1H@1T‐SnS₂ nanodots, the corresponding Tafel slopes were plotted and are shown in Figure 6d. In general, lower Tafel slope at a low overpotential domain specify efficient kinetic mechanisms involved in the catalytic process. Here, for the SnS₂ nanodots, in the region where the current density is low, the Tafel slope of 112 mV/decade indicates a Temkin kind of adsorption mechanism for the intermediates [56], very much like that of Pt/C catalyst. Interestingly, a Tafel slope of 120 mV/decade corresponds to the first‐electron reduction of oxygen as the rate determining step (RDS) which makes the succeeding reduction step and cleaving of O─O bonds much easier to accomplish [57]. The Tafel slope in the high current density region for the 1H@1T‐SnS₂ nanodots SnS₂ nanodots is 498 mV/decade which is very close to that of Pt/C. A change in the Tafel slope for SnS₂ nanodots, unambiguously indicate a switchover from Temkin to Langmuir kind of adsorption mechanism [56, 58], a phenomenon reported very often in earlier studies.

The density functional theory (DFT) calculations were further performed to delve into the electronic structures of the SnS₂ plates and nanodots so as to reveal the properties envisaged to have engendered such augmented ORR performance in the nanodots. The details of the DFT methodology are given in S1.8. We simulated finite nanodots of both 1T‐ and 1H‐SnS₂. Each nanodot consists of two S‐Sn‐S tri‐layers, containing a total of 32 Sn and 64 S atoms [59]. To prevent interactions with periodic replicas, a vacuum spacing of 20 Å was introduced in all three directions, as illustrated in Figure S14a,b,d,e. This configuration ensures that the nanodots behave as isolated clusters rather than as slabs or bulk‐like structures. The finite nanodot geometry results in under coordinated Sn and S atoms at the edges, which lead to the formation of chemically active edge sites‐a result already established through our experimental findings. This is evident from the sharp peaks in the density of states, shown in Figure 7b,d, indicating a high density of accessible electronic states for adsorbates. These dangling bonds are expected to dominate the catalytic behaviour, in contrast to the relatively inert basal planes. Given the number of atoms and the lateral dimensions of the parent SnS₂ lattice, the simulated nanodots have lateral sizes ranging from approximately 1.5 nm to 2.0 nm. These nanodots represent the ultrasmall limit of 2D SnS₂, where edge effects play a significant role in shaping both the electronic structure and catalytic properties. After simulating the nanodots, we performed ground‐state energy calculations for both the 1T and 1H phases and found that the 1T phase was more stable than the 1H phase, with an energy difference of 13 meV per atom. This indicates that, although both phases are energetically favorable, the 1T polymorph is thermodynamically preferred under ambient conditions. Next, we calculated the projected density of states (PDOS) for both the 1T and 1H phases, as shown in Figure 2a–d. The PDOS analysis reveals that the S‐p orbitals (depicted in blue) make the dominant contribution to the electronic states near the Fermi level for both the bulk and nanodot forms of SnS₂. In the bulk structures, the 1T phase exhibits a band gap of 1.63 eV, while the 1H phase shows a slightly smaller band gap of 1.25 eV. However, when transitioning to the nanodot geometry, the band gaps in both phases are significantly reduced due to the increased availability of localized electrons at the edges. Specifically, the band gap in the 1T nanodot is reduced to 0.2 eV, while the 1H nanodot becomes metallic with no band gap. While our theoretical calculations successfully capture the experimental trend of a narrowing band gap during the 1T to 1H phase transition in SnS₂, a significant disparity remains between our calculated values and the experimental data. The reason of the contradiction can be very well attributed to the fact that our nanodots are in mixed phases with varied layer numbers in accordance to previous studies which predict an increase in band gap for mixed states of 1T and 1H as the layer numbers are increased [36]. However, this transition from a semiconductor to a metallic state is primarily driven by the formation of edge states, which introduce new electronic states at the Fermi level. These edge states are particularly relevant for catalysis, as they provide active sites for adsorption and electron transfer processes, which are crucial for catalytic reactions like the oxygen reduction reaction (ORR).

FIGURE 7.

Panel (a), (c), and panel (b), (d) show the projected density of states for SnS₂ bulk and its nanoflake, respectively, for 1T and 1H phase.

The ORR on both 1T‐ and 1H‐SnS₂ proceeds via the four standard proton coupled electron transfer (PCET) steps in alkaline media is outlined in equations SE1, SE2, SE3 and SE4. The calculated free‐energy diagrams derived from 1H‐ SnS₂ (black lines) and 1T‐ SnS₂ (red lines) are shown in Figure 8. Figure 8a,c presents the free‐energy profiles for the 4e− ORR pathway at U = 0 V and U = 1.23 V, respectively. At U = 0 V (Figure 8a), 1H‐ SnS₂ exhibits a slightly uphill activation step associated with OOH^* formation which is thermodynamically demanding, while the subsequent reductions to O^∗ and OH^∗ are thermodynamically favorable. In contrast, 1T‐ SnS₂ displays downhill energetics for the first three steps, reflecting strong stabilization of oxygenated intermediates. However, upon increasing the potential to U = 1.23 V (Figure 8c), the final OH^∗ → OH⁻ step remains thermodynamically demanding for both phases, identifying this step as the rate‐limiting process. Notably, the residual free‐energy penalty is significantly smaller for 1H‐SnS₂, indicating more balanced adsorption of O^∗ and OH^∗, whereas the large penalty for 1T‐ SnS₂ signifies severe OH^∗ poisoning. T). The final conversion step requires an uphill free‐energy change of approximately 1.45 V (U = 0 V) and 2.68 V (U = 1.23 V) for the 1H phase, whereas the corresponding step in the 1T phase is associated with a larger energy penalty of approximately 2.23 V (U = 0 V) and 3.46 V (U = 1.23 V). Therefore, the 1H phase exhibits a lower theoretical overpotential. The lower overpotential of the 1H‐SnS₂ nanostructure originates from its more favorable interaction with ORR intermediates. The 1H phase binds the reaction intermediates strongly enough to activate them, but not so strongly that their subsequent removal becomes energetically difficult. The 1H phase achieves this balance more effectively than the 1T phase in the case of 4e‐ pathway. The 2e⁻ ORR pathway is illustrated in Figure 8b,d at U = 0 V and U = 1.23 V, respectively. At U = 0 V (Figure 8b), the initial O₂ → OOH^∗ step on 1H‐ SnS₂ is endothermic, making O₂ activation the potential‐determining step, whereas the subsequent reduction to HO₂‐ is strongly exothermic. In contrast, 1T‐ SnS₂ exhibits downhill energetics for both steps, suggesting a thermodynamically favourable pathway for HO₂‐ formation. This trend persists at U = 1.23 V (Figure 8d), where the free‐energy landscape for 1T‐ SnS₂ remains favorable for the 2e⁻ process, while 1H‐ SnS₂ retains an energetic penalty associated with OOH^∗ formation.

FIGURE 8.

Panels show the energy profile graph of 1T and 1H SnS₂ nanodot for U = 0 V in (a) and (b), and U = 1.23 V in (c) and (d). Panels (a) and (c) show the 4‐step electron transfer process via an associative mechanism, and panels (b) and (d) show the 2‐step electron transfer process. (^* stands for the free energy of the ligand).

Although neither phase exhibits ideal ORR energetics, the comparative analysis of Figure 8 reveals distinct phase‐dependent selectivity. The metastable 1H‐ SnS₂ phase favours the 4e⁻ ORR pathway due to its balanced adsorption of O^∗ and OH^∗ intermediates, despite being limited by OOH^∗ formation. Conversely, the metallic 1T‐ SnS₂ phase is constrained by strong OH^∗ binding in the 4e⁻ pathway but preferentially promotes the 2e⁻ ORR mechanism owing to weaker OOH^∗ stabilization and facilitated HO₂‐ formation. Theoretical studies in conjunction with our experimental findings suggest that the overpotentials for ORR can be lower in 1H‐SnS₂ compared to 1T. To identify the catalytically active sites, partial charge density analysis for O^∗ and OH^∗ adsorbed intermediates were performed on the 1H‐ and 1T‐SnS₂ nanostructures. The corresponding charge‐density distributions clearly reveal that the surface S‐atoms exhibit pronounced charge accumulation and strong electronic overlap with the adsorbates, as shown in Figure S17a,b. For both O^∗ and OH^∗ adsorption configurations, the charge density is localized around the surface S atoms, showing that these surface S‐atoms actively participate in charge transfer during the ORR process. This confirms that the exposed surface S‐sites serve as the primary catalytic active centres for oxygen reduction. The present phase‐and‐size‐engineered SnS₂ nanodots proliferate in chemically active S‐ edge sites and 1H‐phase supported favourable electronic structure which worked in synergy to effectuate an enhanced electrocatalytic ORR mechanism in SnS₂ nanodots. Therefore, this work successfully provides unprecedented insights into the ways of infusing ORR activity in materials which are otherwise ORR inactive and in conjunction opens up a new avenue to advantageously engineer low cost, earth abundant materials through a facile process of production towards futuristic energy applications.

3. Conclusions

In summary, ultrathin nearly circular 1H@1T SnS₂ nanodots were elegantly sculpted from bulk 1T‐phase‐SnS₂ plates through an improvised thermal stress‐mediated fracturing technique. While a subtle polydispersity was observed, with lateral diameters ranging from 2 nm to 5 nm, the true allure of these nanostructures lies in their anomalous electronic evolution. Contrary to the conventional expectations, a dramatic reduction in the bandgap to 0.856 eV was recorded by STS study. This quasi‐metallic transition was elucidated via HAADF‐STEM and DFT simulations, which traced the phenomenon to the emergence of edge states and the strategic transformation of the 1H phase within the 1T lattice. XPS studies revealed that reduction in size and layer numbers had introduced copiously abundant ${\mathrm{S}}_2^{2-}$ species, famously known for being catalytically active sites. The synthesized nanodots exhibited a masterful performance in the Oxygen Reduction Reaction (ORR) within 0.1 M KOH, facilitating a much desired 4‐electron transfer pathway in stark contrast to the sluggish 2‐electron mechanism of the bulk SnS₂ plates. A reduction in the Gibbs free energy for intermediate adsorption is theoretically confirmed via DFT. This exaltation in electro catalytic prowess is attributed to a synergistic harmony between the burgeoning of active sites, edge‐state modification, and phase hybridization. Ultimately, this fracturing methodology opens a novel horizon for sustainable materials engineering, enabling the design of SnS₂‐based electrocatalysts with unprecedented efficiency.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: smll74575‐sup‐0001‐SuppMat.docx.

Basu P., Saxena S., Chakraborty J., et al. “Phase‐Restructured SnS₂ Nanodots: Unlocking ORR Activity through Thermal Stress‐Mediated Fracturing.” Small 22, no. 49 (2026): e74575. 10.1002/smll.74575

Data Availability Statement

The data that support the findings of this study are available on request from the corresponding author. The data are not publicly available due to privacy or ethical restrictions.