Quantum confinement and photoresponsivity of β-In₂Se₃ nanosheets grown by physical vapour transport

The integration of layered semiconductors with graphene to form heterostructure devices offers new routes to the fabrication of optoelectronic devices such as fast, ultrasensitive photodetectors [1–3]. Metal dichalcogenides (MoS₂ and WSe₂), III–VI semiconductors (InSe, GaSe and In₂Se₃), and elemental semiconductors (black phosphorus), are currently being studied as potential candidate materials for these applications [1–9]. The III–VI materials exhibit an interesting dependence of the bandgap on the layer thickness which has been identified for exfoliated flakes of InSe, GaSe, and GaTe, and has been attributed to the formation of a two-dimensional quantum well with potential barriers formed by the physical boundaries of the exfoliated flakes [10–12]. This effect offers the prospect of band-gap engineering in the III–VI system through the tailoring of layer thickness to control optoelectronic properties, notably the absorption edge and photoluminescence energy of emitted photons. Quantum confinement has so far been demonstrated clearly only for exfoliated crystals and has not been reported to date for thin films grown on a substrate. This raises questions as to whether thin film growth techniques provide material and interfaces of sufficient quality for the observation of such quantum effects.

In this paper we demonstrate quantum confinement in the In₂Se₃ material system, both for exfoliated flakes, and for thin films grown by physical vapour transport. Among the III–VI crystals, In₂Se₃ can occur as several different crystal structures and phases (α, β, γ, δ, and κ) [13, 14]. The α- and β-In₂Se₃ are van der Waals layered crystals and have attracted particular interest. For both phases, the primitive unit cell contains three layers, each consisting of five closely-packed, covalently bonded, monoatomic sheets in the sequence Se–In–Se–In–Se (figure 1(a)). As proposed in [15], in α-In₂Se₃ the outer Se-atoms in each layer are aligned, whereas in β-In₂Se₃ they are located in the interstitial sites of the Se-atoms in the neighbouring layers, thus leading to a smaller volume in β-In₂Se₃, but without changing the crystal symmetry [15]. The growth of thin films of α-In₂Se₃ with excellent structural properties has been reported recently [8, 16, 17], but the growth of two-dimensional β-In₂Se₃ remains largely unexplored although this phase exhibits significant differences in its electrical properties, including a stronger sensitivity of the electronic band structure to externally applied electric fields [18, 19].

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As we show below β-In₂Se₃ layers can be grown on several different substrates (SiO₂/Si, mica and graphite) and quantum confinement can indeed be realized for thin films grown on SiO₂/Si. Specifically, using γ-InSe as source material we demonstrate the growth of β-In₂Se₃ islands with typical widths 1–15 μm and thickness ranging from 100s of nanometers down to 2.8 nm. The presence of β-In₂Se₃ is confirmed using a combination of Raman spectroscopy, x-ray diffraction (XRD), elemental analysis and photoluminescence. Our results are unexpected since there have been several other studies which have found that a different phase, α-In₂Se₃, is commonly observed under somewhat similar growth conditions [8, 16, 17]. In addition, the stability of β-In₂Se₃ has been questioned and some interconversion between different phases has been reported [14, 18, 20]. However, we found that our grown films display excellent stability under a long period of storage in ambient conditions. Moreover, we observe a shift in the photoluminescence peak due to quantum confinement which is controlled by the physical thickness of the β-In₂Se₃ islands; this indicates that the effects previously reported for exfoliated layers can also be realized in films grown on a substrate, offering the prospect of band-gap engineering in large-area devices. Furthermore, our work implies that a high quality heterojunction is formed between the substrate and the β-In₂Se₃ island since this interface provides one of the barriers which leads to carrier confinement. The grown material may be readily incorporated into simple devices for high sensitivity photodetection.

2.1. Synthesis of In₂Se₃ layers

2.2. Characterisation

The β-In₂Se₃ layers were grown by a physical vapour transport method, as illustrated in figure 1(b). A ground powder of Bridgman-grown γ-InSe crystals was used as source material [21]. We found that growth on SiO₂/Si substrates produces near-circular, slightly facetted films with lateral size between 1 and 15 μm (figure 1(c)), while highly facetted hexagonal films with lateral size ∼100 μm grew on mica under similar conditions, see figure 1(d). On graphite, films with arbitrary shape grew (figure 1(e)) with preferential growth at step edge bunches on the substrate. Figure 2 shows optical micrographs (figure 2(a)), topographic atomic force microscopy (AFM) images (figure 2(b)) and corresponding height profiles (bottom of figure 2(b)) of typical films grown on a SiO₂/Si substrate. The thickness L of the In₂Se₃ layers was determined by AFM. The thinnest measured films have a thickness of L ≈ 2.8 nm, which corresponds to the size of a unit cell along the c-axis for β-In₂Se₃ (see figure 3(b)). Additionally, single monolayer steps (≈1 nm) were observed on the surface of some of the investigated films.

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Identifying the crystalline phase of the as-grown In₂Se₃ layers has proven challenging due to multiple phases (α, β, γ, δ, and κ) existing at different temperatures [13, 14]. While the crystal structures of γ- and δ-In₂Se₃ are hexagonal and trigonal, respectively [22], the rhombohedral crystal structures of α- and β-In₂Se₃ phases are quite similar and different symmetries (with space group R$3$m and R$\bar{3}$m) have been suggested by some reports [19, 22, 23], while others have argued that the two phases have the same symmetry [15]. In previous studies of bulk In₂Se₃ it has been shown that the α-phase, which is stable at RT, reversibly transforms into the β-phase at T = 200 °C by thermal annealing [14, 23]. However, a recent report has shown that the β-phase can persist in thin layers (∼4–5 nm) at RT [18]. To discriminate between different phases in our experiments we use a combination of Raman spectroscopy and XRD.

RT (T = 300 K) micro-Raman spectra of representative films grown on SiO₂/Si substrate are shown in figure 3(a). The Raman peaks are centred at ∼110, 175, and 205 cm⁻¹ and correspond to the phonon modes of β-In₂Se₃ [15, 18, 20, 24]. The main Raman line at ∼110 cm⁻¹ corresponds to the A₁(LO+TO) phonon mode and the weaker peaks at ∼175 and 205 cm⁻¹ are attributed to the A₁(TO) mode and A₁(LO) mode, respectively. In comparison, the A₁(LO+TO), A₁(TO), and A₁(LO) phonon modes in exfoliated flakes of α-In₂Se₃ crystals [24] grown by the Bridgman method are at ∼104, 181, and 200 cm⁻¹, respectively [15, 18, 20, 24], see bottom of figure 3(a) (further details can be found in supporting information figure S1). Therefore, our Raman spectra indicate that the grown layers are β-In₂Se₃ and that they are stable at RT consistent with a recent report [18].

The inset of figure 3(a) shows the Raman shift versus layer thickness of the A₁(LO+TO), A₁(TO), and A₁(LO) peaks of as-grown β-In₂Se₃. The A₁(LO+TO) mode at ∼110 cm⁻¹ does not depend on the layer thickness, although, the A₁(TO), and A₁(LO) phonon modes exhibit a small shift to lower wavenumbers with decreasing L. This could arise from the smaller vibration coherence length along the c-axis due to the weak van der Waals force along this direction [26]. When the layer thickness L decreases below ∼10 nm, the full width at half maximum (W) of the 110 cm⁻¹ Raman line tends to broaden by a factor of ∼1.4, possibly due to the roughness of the film and its interface with the SiO₂ substrate [27].

XRD and x-ray photoelectron spectroscopy (XPS) measurements are also consistent with the formation of the β-In₂Se₃ phase (figures 1(a) and 3(b)). Within each plane, atoms form hexagons with lattice parameter a = 4.00 Å; along the c-axis, the lattice parameter is c = 28.33 Å [22, 23]. The equivalent lattice parameters for α-In₂Se₃ are a = 4.025 Å and c = 28.762 Å [22, 23]. Figure 3(c) shows a typical XRD spectrum of as-grown films on a SiO₂/Si substrate at T = 300 K which confirms that the In₂Se₃ layers are highly crystalline with a lattice constant along the c-axis, c = 28.31 ± 0.05 Å, and along the hexagonal axes, a = 3.99 ± 0.02 Å (see supporting information figure S2 which shows the XRD measurements of the $(\bar{1}20)$ and (1$\bar{2}$0) reflections from which the value for a is derived). Our XRD values of the lattice constants (a and c) are in excellent agreement with the reported values of β-In₂Se₃ in [22, 23] and rule out the presence of α-In₂Se₃. Additionally, our XPS spectra show that the stoichiometric composition of the layers is [In] ≈ 42 ± 3 atomic % and [Se] ≈ 58 ± 3 atomic %, close to [In]:[Se] ≈ 2:3, see supporting information figure S3.

Room temperature (T = 300 K) μPL spectra of representative films grown on SiO₂/Si substrate are shown in figure 4(a). The PL peak occurs at a photon energy hν = 1.43 eV for films with L > 50 nm. With decreasing L, the PL emission exhibits a blue-shift by up to 160 meV (figures 4(a) and (b)), consistent with quantum confinement of photo-excited carriers by the external surfaces of the films. In addition, when L decreases below ∼30 nm, the PL intensity decreases by a factor of two or more partly due to the weaker absorption of light (α ≈ 1 × 10⁷ m⁻¹ at λ = 532 nm [28], where α is the absorption coefficient of In₂Se₃) and the full width at half maximum (W) of the PL emission increases by a factor of two possibly due to the roughness of the layers and/or crystal defects. The PL emission of these films persists for several months when they are stored under ambient conditions, confirming their chemical stability. No PL signal was detected at T = 300 K from the thinnest In₂Se₃ films (L = 2.8 nm) identified by AFM.

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Figure 4(b) shows the results of PL and AFM measurements of several as-grown β-In₂Se₃ layers (black dots) and our exfoliated Bridgman-grown α-In₂Se₃ [25] flakes (blue stars) on the SiO₂/Si substrate. They reveal the dependence of the band-to-band transition energy, E_2D, on the layer thickness L. RT PL spectra of representative exfoliated α-In₂Se₃ flakes are shown in the supporting information figure S4. Figure 4(b) also shows the RT PL peak energy of α-In₂Se₃ thin layers (red stars) grown by physical vapour transport as reported in [17].

We model the dependence of E_2D on L using a square quantum well potential of infinite height, i.e.

$$\begin{eqnarray}&&{E}_{2{\rm{D}}}=\;{E}_{{\rm{g}}}-\;{E}_{{\rm{b}}}+\;{\pi }^{2}{\hslash }^{2}/2{L}^{2}{\mu }_{\parallel c},\end{eqnarray} \tag{ 1 }$$

where ${E}_{{\rm{g}}}$ is the optical band gap energy at T = 300 K, ${E}_{{\rm{b}}}$ is the exciton binding energy, ${\mu }_{\parallel c}=\;{\left(\tfrac{1}{{m}_{\parallel c}^{e}}+\tfrac{1}{{m}_{\parallel c}^{h}}\right)}^{-1}$ is the exciton reduced mass for motion along the c-axis in In₂Se₃. The values of ${{E}}_{{\rm{g}}}-\;{{E}}_{{\rm{b}}}$ and ${\mu }_{\parallel c}$ are obtained from the best fit to the measured values of the PL peak energy (at T = 300 K) versus layer thickness. The value of the exciton reduced mass along the c-axis of bulk β-In₂Se₃, ${\mu }_{\parallel c}=0.030$ m_e, is smaller than the value for α-In₂Se₃ (${\mu }_{\parallel c}=0.080$ m_e) and for γ-InSe (${\mu }_{\parallel c}=0.054$ m_e) [10]; here m_e is the electron mass in vacuum. This indicates a stronger quantum confinement effect in β-In₂Se₃ compared to α-In₂Se₃ and γ-InSe. The scatter in the individual data points around the modelled curves (continuous lines) suggest that carrier confinement is influenced by the roughness of the layers and its interface with the substrate and/or crystal defects. In recent theoretical work it has been argued that bulk β-In₂Se₃ has an indirect band gap and remains indirect when the layer thickness is reduced to a single layer with a shift of the bandgap energy by 590 meV [19]. Our experimental data are in qualitative agreement with this theoretical study. However, the higher energy PL peak position for bulk β-In₂Se₃ compared to α-In₂Se₃ is opposite to that predicted in [19] and requires further investigation and modelling of the crystal structure for different atomic arrangements. RT PL emission has also been detected from our β-In₂Se₃ films (with L > 50 nm) grown on mica and graphite, shown in the supporting information figures S5 and S6, respectively.

In order to study the photoconductivity of the as-grown β-In₂Se₃ layers, we have fabricated photodetectors based on representative films with thickness down to L = 20 nm. For the electrodes, Ti/Au (10 and 100 nm, respectively) contacts were deposited using a combination of evaporation and electron beam lithography. The photocurrent spectra indicate a systematic blue-shift of the absorption edge with decreasing layer thickness and do not show a clear excitonic absorption band, suggesting that the β-In₂Se₃ has an indirect band gap, as predicted in [19], see figure 5.

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Figure 6(a) shows the current–voltage, I–V, characteristics of a β-In₂Se₃ photodetector (L = 76 nm) measured under dark and illuminated conditions (λ = 633 nm, P = 170 μW) at T = 300 K. Under illumination by a focused laser beam, the current increases, particularly for V > 0.5 V. The dependence of the photocurrent, ΔI, on the applied bias under different laser powers is shown in figure 6(b). A spatially resolved photocurrent map obtained by scanning a focused laser beam (λ = 633 nm and P = 17 nW) across the plane of the In₂Se₃ layer shows that photocurrent generation occurs primarily in the In₂Se₃ region of the film between the two Ti/Au electrodes (see inset of figure 6(a)).

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Under an applied bias V, photo-excited electrons and holes in In₂Se₃ are swept by the electric field in opposite directions, thus generating a photocurrent ΔI = [eLαP/hν](τ_l/τ_t), where α is the absorption coefficient of In₂Se₃ at the photon energy hν, P is the incident power, e is the electronic charge, and τ_l/τ_t is the ratio of the carrier lifetime (τ_l) and transit time (τ_t) of electrons in In₂Se₃ [5]. Thus the photoresponsivity R of our device can be described approximately by the relation R = ΔI/P = [eLατ_l /hντ_t]. Moreover, we can express the external and internal quantum efficiencies as EQE = Rhν/e = Lατ_l/τ_t and IQE = τ_l/τ_t, respectively.

The In₂Se₃ photodetectors exhibit a stable and reproducible photoresponsivity, R = ΔI/P, (figure 6(c)) with values of R of up to ∼1720 A W⁻¹ at V = 1 V, λ = 633 nm and low incident power P = 1.7 pW, which corresponds to an IQE = τ_l/τ_t ≈ 5.5 × 10³ for α ≈ 8 × 10⁶ m⁻¹ at hν = 1.96 eV (λ = 633 nm) [28] and L = 76 nm. From the maximum value of R (for L = 76 nm), we estimate an external quantum efficiency EQE ≈ 3370 and a specific detectivity D* = R(A/2eI)^1/2 ≈ 7 × 10¹⁰ m W^–1 s^–1/2, where A ≈ 20 μm² is the area of the film and I = 36 nA is the dark current at V = 1 V. A power law relation of the form R ∝ P⁻ⁿ provides a good empirical fit to the values of R versus P for devices with L = 30 and 76 nm, with n = 0.75 and 0.84, respectively.

The photodetector response time was measured by focusing a mechanically modulated laser beam of λ = 633 nm and P = 17 nW on the device (L = 76 nm). Figure 6(d) shows the photocurrent waveform in response to a series of cycles with the laser beam alternately on and off. Our In₂Se₃ photodetector exhibits a repeatable and stable response to the incident light. The measured rise (τ_r) and decay (τ_d) times are 0.6 and 2.5 ms, respectively. Our values of R and the rise/decay times for photodetectors based on as-grown nanosheets are comparable or superior to those reported previously for α-In₂Se₃ [29–31]. The direct growth of β-In₂Se₃ nanosheets on different substrates also offers flexibility for 2D electronic and optoelectronics.

In summary, we have grown β-In₂Se₃ layers by a physical vapour transport method. The β-In₂Se₃ layers are chemically stable and optically active at RT over periods of several months. Due to the smaller exciton mass along the c-axis, the 2D quantum confinement effects which we observe in β-In₂Se₃ are stronger than those previously reported in other III–VI van der Waals crystals, e.g., GaSe and GaTe [11, 12]. The thickness of the film can be used to tune the absorption and emission in the technologically relevant-midinfrared spectral range between 1.43 and 1.58 eV. The β-In₂Se₃ photodetectors showed excellent photoresponsivity and relatively fast response to light. These properties confirm that the β-In₂Se₃ layers are promising candidate materials for optoelectronic applications.

This work was supported by the Engineering and Physical Sciences Research Council (EPSRC) [under grants EP/M012700/1 and EP/K005138/1], the EU FP7 Graphene Flagship Project 604391, the University of Nottingham, and the Ukrainian Academy of Sciences.