Nonequilibrium Rock-Salt-Type Pb-Doped SnSe with High Carrier

Article pubs.acs.org/cm

Nonequilibrium Rock-Salt-Type Pb-Doped SnSe with High Carrier Mobilities ≈ 300 cm2/(Vs) Takeshi Inoue,† Hidenori Hiramatsu,*,†,‡ Hideo Hosono,†,‡ and Toshio Kamiya*,†,‡ †

Materials and Structures Laboratory, Tokyo Institute of Technology, Mailbox R3-4, 4259 Nagatsuta-cho, Midori-ku, Yokohama 226-8503, Japan ‡ Materials Research Center for Element Strategy, Tokyo Institute of Technology, Mailbox SE-6, 4259 Nagatsuta-cho, Midori-ku, Yokohama 226-8503, Japan S Supporting Information *

ABSTRACT: We synthesized nonequilibrium cubic rock-salt (RS)-type (Sn,Pb)Se and investigated their optoelectronic properties with the expectation that the RS-type structure would exhibit better carrier transport properties than the equilibrium GeS-type layered crystal structure in SnSe because the RS-type structure has the three-dimensional network of (PbSe6) octahedra, while the GeS-type structure has the two-dimensional network of (SnSe3), leading to larger band dispersions and smaller carrier effective masses. To stabilize the nonequilibrium phase, epitaxial thin films were grown by a unique method, a reactive solid-phase epitaxy with a thin RS-type PbSe epitaxial template layer. Additionally, a rapid quenching process from 600 °C to room temperature was also effective for stabilizing the nonequilibrium RS-type (Sn,Pb)Se epitaxial films. We succeeded in controlling Pb concentration continuously from 0 to 100% in the (Sn,Pb)Se solid-solution films. The minimum Pb concentration to stabilize the RS-type SnSe was extended from the previously reported value of 63% to 50%. A band gap bowing effect was observed with the smallest estimated band gap of ∼0.14 eV for the RS-type (Sn0.35Pb0.65)Se. The GeS-type to RS-type structural change increased hole mobility drastically from 60 for SnSe to 290 cm2/(Vs) for 58% Pb-doped RS-type (Sn,Pb)Se film as expected. It was found that p-type to n-type conversion occurs by further higher Pb concentrations ≥61%, and the highest electron mobility of 340 cm2/(Vs) was observed.

1. INTRODUCTION Simple tin-based binary chalcogenides such as SnS and SnSe are good candidates for optoelectronics applications because of their band gaps (∼1.1 eV for SnS1−5 and 0.8−1 eV for SnSe2,6−8) appropriate for solar cells and other optoelectronic devices. Among them, tin monoslenide SnSe is a native p-type semiconductor and has a great potential due to its high hole mobilities, i.e., 90−200 cm2/(Vs) for single crystals at room temperature (RT).6,9 Thus, the tin-based simple binary selenide is expected as an active layer of electronic devices. From this point of view, we recently obtained high-quality epitaxial films of pure SnSe with a hole mobility of 60 cm2/(Vs) and a hole density of 3 × 1016 cm−3 by pulsed laser deposition (PLD).8 This hole mobility is higher than that of SnS epitaxial films (∼40 cm2/(Vs) at RT)4 but still lower than those of the SnSe single crystals. As seen in the side-view in the left panel of Figure 1(a), SnSe has the layered crystal structure composed of an alternate stack of the (Sn2+Se2−)2 bimolecular layers along the a-axis [Figure 1(a) is drawn based on the space group Pnma (No. 62)]. Because each layer is charge neutral, the adjacent layers are bound by van der Waals interaction. On the other hand, the GeS-type structure can also be regarded as a distorted rock salt (RS)-type structure. As seen in the top-view shown in the right panel of Figure 1(a), the distorted face-centered structure is found in the √2 × √2 supercell unit (shown by © 2016 American Chemical Society

the red square); and if an adjacent layer is shifted slightly along the c-axis (indicated by the red arrow in the side-view shown in the left panel in Figure 1(a)), then the GeS-type structure is transformed to the RS-type structure. However, it has been reported that the RS-type SnSe is a metastable phase; thus there have been only a few reports on syntheses of RS-type SnSe so far, to our knowledge.10,11 Distortion of the orthorhombic structure reduces the number of Sn−Se direct bonds from six in the high symmetric RS-type structure to three and forms strong in-plane Se−Se bonds only along the b-axis, stabilizing the orthorhombic GeS-type structure at RT but deteriorating carrier transport properties because of larger Se− Se and Sn−Sn separation (i.e., due to the reduced coordination numbers). In contrast, a simple binary selenide PbSe has the RS-type structure, as shown in Figure 1(b). Different from SnSe, the RS-type structure is thermodynamically stable for PbSe at RT. Comparing these crystal structures, we expect that smaller hole effective mass and higher hole mobility would be realized if the crystal structure of SnSe is changed from the thermally equilibrium GeS-type one to the RS-type one because the Received: January 23, 2016 Revised: March 23, 2016 Published: March 29, 2016 2278

DOI: 10.1021/acs.chemmater.6b00307 Chem. Mater. 2016, 28, 2278−2286

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Chemistry of Materials

nonequilibrium RS-type (Sn,Pb)Se phase. We employed a unique growth method, reactive solid-phase epitaxy, in which a thin RS-type PbSe epitaxial template layer works as a sacrificial layer. Additionally, a quenching process from 600 °C to RT was also effective for stabilizing the nonequilibrium RS-type epitaxial (Sn,Pb)Se. Using this technique, we succeeded in varying the Pb concentration from 0 to 100% continuously. The minimum Pb concentration to stabilize the RS-type SnSe was 50%, which is the minimum Pb content ever reported. Structure transition from GeS-type to RS-type increased hole mobility drastically from 60 for SnSe to 290 cm2/(Vs) for 58% Pb-doped RS-type film. P-type to n-type conversion is also observed by further higher Pb doping up to 100%. The maximum electron mobility of 340 cm2/(Vs) was achieved by 61% Pb doping. Figure 1. Crystal structures of (a) orthorhombic GeS-type SnSe and (b) cubic rock-salt (RS)-type PbSe. GeS-type structure is drawn based on the space group Pnma. The rectangular boxes in the left panels (side view) and the black rectangle/square in the right panels (top view) show the unit cell. The red arrow in (a) shows that if an adjacent layer is shifted by (0, 0, 0.38), the GeS-type structure is transformed to the RS-type structure. The red square in (a) shows the unit cell of the quasi RS-type cell formed by the √2 × √2 supercell unit of the GeStype unit cell. The octahedral polyhedron (PbSe6) in the RS-type PbSe is shown by purple in (b). The top-view in the right panel of (b) shows the edge-sharing network of the (PbSe6) octahedra along {011} directions.

2. EXPERIMENTAL SECTION 2.1. Synthesis of Polycrystalline SnSe1.2, PbSe1.1, and (SnSe1.2+PbSe1.1) Disks for PLD Targets. Polycrystalline disks of the Se-rich SnSe1.2, PbSe1.1, and (SnSe1.2+PbSe1.1) mixture were synthesized by solid-state reactions for laser ablation because our previous study revealed the Se rich target was effective for growing high-quality stoichiometric SnSe epitaxial films.8 A tin rod (purity: 99.999%) was employed as a starting reagent to prevent surface oxidization of tin. To obtain high-purity tin powders, we ground the rod with a metal file in a glovebox filled with a dry inert Ar gas (dew point < −90 °C, oxygen concentration 600 °C, the as-deposited bilayer films completely disappeared from the substrates after thermal annealing. Therefore, we employed the maximum Ta = 600 °C without re-evaporation during thermal annealing. Then the ampule was subjected to rapid quenching in water from 600 °C to RT in order to stabilize the high-temperature RS-type (Sn,Pb)Se phase. In this method, formation of the PbSe epitaxial template layer is a key to achieving the epitaxial growth, and the annealing process converted the SnSe/PbSe bilayer to a uniform epitaxial film.

3. RESULTS AND DISCUSSION 3.1. Polycrystalline SnSe1.2, PbSe1.1, and (SnSe1.2+PbSe1.1) Disks for PLD Targets. Figure 3 shows XRD patterns of the SnSe1.2, PbSe1.1, and (SnSe1.2+PbSe1.1) disks. As seen in Figure 3(a), the SnSe1.2 can be assignable to almost the single phase of orthorhombic GeS-type SnSe, while a small amount of SnSe2 was detected at 2θ = 14.4° (see the inset of Figure 3(a)), which originates from the intentionally added excess Se. The obtained SnSe1.2 disk exhibited strong aaxis orientation because the relative intensities of the 111 (2θ = 30.4°, the highest peak in the simulated pattern) and the 400 (31.1°) diffractions were opposite to the simulated powder XRD pattern shown at bottom. This result implies that SnSe is easily oriented to its a-axis even in bulk samples because of the anisotropic layered structure. In contrast, for a PbSe1.1 disk, we cannot observe any detectable impurity (e.g., the highest peak position of tetragonal PbSe2 is 2θ = 30.4°) irrespective of the 10% excess Se addition. For the (SnSe1.2+PbSe1.1) disk, the major phase was the PbSe structure, but a small amount of the SnSe phase was included along with segregated SnSe2 as shown by a vertical arrow at 2θ = 14.4° (Figure 3(b)). The lattice parameter of the PbSe phase in the (SnSe1.2+PbSe1.1) disk became slightly smaller (0.6073 nm) than that of the PbSe1.1 (0.6143 nm) shown at the top in Figure 3(a), implying that some Sn substitute the Pb sites in the (SnSe1.2+PbSe1.1) disk. This phase separation is consistent with 2280

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Figure 3. XRD patterns of SnSe1.2, PbSe1.1, and (SnSe1.2+PbSe1.1) disks for PLD targets. The blue and red vertical bars indicate the positions of the Bragg diffraction peaks of the orthorhombic GeS-type SnSe and the cubic RS-type PbSe, respectively. (a) Observed patterns of SnSe1.2 and PbSe1.1 disks (red curves) with simulated patterns (black curves) for comparison. The inset shows the enlarged pattern of the SnSe1.2 sample in the 2θ region indicated by the dotted rectangle, showing the existence of impurity SnSe2 phase. (b) Observed patterns of a (SnSe1.2+PbSe1.1) disk (red curve) with simulated SnSe and PbSe patterns (black curves) for comparison. The vertical arrow shows the SnSe1.2 impurity phase.

lattice parameters unchanged. Consequently, we concluded that it was very difficult to obtain the RS-type (Sn,Pb)Se films directly by a simple PLD method because when we employed higher Ts ≥ 550 °C, no film was deposited on the substrate (see the top XRD pattern in Figure 4). This result means that reevaporation is more dominant rather than deposition at Ts ≥ 550 °C and such high-Ts growth process cannot be employed. Thus, we employed the R-SPE technique with water-quenching of RT-deposited SnSe/PbSe template on the MgO substrate. 3.3. Heteroepitaxial Growth of (Sn,Pb)Se Films by RSPE Combined with Quenching Process by Water. 3.3.1. Optimization of Ts for Epitaxial PbSe Template Layer by PLD. Next, we optimized the Ts for the PbSe template layer in the R-SPE process between 400 and 600 °C using the PbSe1.1 disk as a PLD target (Figure 5). At Ts = 400 °C, the obtained film was a-axis oriented (Figure 5(a)), while the preferential orientation was very weak because fwhm of the observed XRC was broad (1.8°, Figure 5(b)). With an increase in Ts to 500 °C, the film quality was much improved; i.e., a-axis orientation became stronger with a sharp fwhm of the XRC (0.5°). However, Pb metal began segregating at Ts = 500 °C because cubic Pb 220 diffraction was detected at 2θ = ∼52°. With further increasing Ts to 600 °C, no film was deposited; i.e., only the diffraction from the MgO substrate was observed. These results indicate that re-evaporation of Se already began at Ts = 500 °C although we employed the 10% Se-rich PbSe1.1 target. The in-plane orientation (Figure 5(c)) was also totally different between the Ts = 400 and 500 °C samples. At Ts = 400 °C, the ϕ scan of the PbSe 022 diffraction indicates that the film consists of two kinds of in-plane domains rotated by 45° with each other, while the 500 °C sample has a single domain structure. These results indicate that the Ts = 500 °C sample has the best quality although a small amount of Pb is segregated. Thus, we concluded that the optimum Ts is 500 °C for the epitaxial PbSe template layer between the top RTdeposited SnSe layer and the MgO substrate.

the reported SnSe−PbSe phase diagram at the Sn:Pb ratio = 1:1.12 3.2. Direct Film-Growth Using the Mixed Phase (SnSe1.2+PbSe1.1) Disk by PLD. At first, we deposited films using the (SnSe1.2+PbSe1.1) polycrystalline disk as a PLD target. Figure 4 shows XRD patterns of the films grown at Ts = RT −

Figure 4. Out-of-plane HRXRD patterns of the films grown at Ts = RT (bottom) − 550 °C (top) directly by PLD using the (SnSe1.2+PbSe1.1) disk as a target disk. At Ts = 550 °C, only MgO 200 diffraction is observed, indicating that no film is deposited. These patterns indicate that SnSe and PbSe phases are completely separated in the films.

550 °C. At Ts = RT, a mixture phase of [100]-oriented SnSe and [100]-oriented PbSe was obtained. With an increase in Ts to 400 °C, PbSe 111 diffraction started being observed at 2θ = 25.3°. With a further increase in Ts, the intensity of the diffractions from the GeS-type SnSe phase decreased but did not completely disappear; while that from the RS-type PbSe phase increased especially for 111 diffraction with keeping the 2281

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Figure 5. HRXRD patterns of the cubic RS-type PbSe epitaxial template layer grown at Ts = 400−600 °C by PLD using the PbSe1.1 disk as a target to optimize Ts. (a) Out-of-plane diffraction pattern. At Ts = 600 °C, no film was deposited. (b) Normalized out-of-plane XRC of the PbSe 200 diffraction of the films grown at Ts = 400 and 500 °C. (c) In-plane 2θχ-fixed ϕ scans of the PbSe 022 diffraction of the films grown at Ts = 400 and 500 °C.

with a facet structure began to be observed. In the finally obtained RS film, the facet structure (average size 130 nm) was observed more clearly. This facet structure seems to grow on the bottom granular layer; this may be the origin of the two peak-components observed in the XRC (Figure 6(b)). Next, we synthesized RS-type (Sn,Pb)Se epitaxial films with further higher Pb concentrations by the same R-SPE and rapidquenching method where the thickness ratio of the PbSe to the SnSe layer was increased. Figure 7(a) summarizes the out-ofplane HRXRD patterns of the films. The PbSe film (i.e., x = 1) grown by PLD is also shown for comparison. We observed clear peak shifts to lower 2θ with increasing Pb concentration, indicating the increase in lattice parameters (more detail will be discussed in Figure 8). All the XRC of the films (Figure 7(b)) exhibited two peaks, narrow and broad ones, similar to Figure 6(b) (see Figure S2 in Supporting Information for all the XRC in this study). The fwhm values are almost independent from x in the films; while volume fraction of both peaks slightly changed. The sharp fwhm component decreases with increasing x due probably to an increase in the in-plane lattice mismatch; i.e., √2a of MgO = 0.595 nm, b of (Sn0.5Pb0.5)Se film = 0.607 nm (see Figure 6(e)), and consequently the lattice mismatch increases with increasing x. These results indicate that our process performed in this study can continuously control the Pb concentration of the RS-type (Sn1−xPbx)Se epitaxial films in the region x ≥ 0.5. 3.3.3. In-Plane Lattice Parameters and Orthorhombic to Cubic Phase Transition. Figure 8 summarizes the lattice parameters of the (Sn1−xPbx)Se epitaxial films. For x = 0.2 (orthorhombic) and 0.3 (orthorhombic + cubic) films, we could not detect clear in-plane orientation of the orthorhombic phase (see Figure S3 in the Supporting Information for the inplane HRXRD). Therefore, we cannot show in-plane lattice parameters of these two samples in this figure. The orthorhombic GeS-type phase region is estimated roughly to be 0 ≤ x < ∼2.5 because we confirmed the 0 ≤ x ≤ 0.2 films are the single orthorhombic phase based on the out-of-plane HRXRD patterns (see Figure 6(a)). The orthorhombic and cubic mixed phase region is in 0.3 ≤ x < 0.5 because we observed single cubic phase at x = 0.5. After transition from the mixed phase to the single cubic phase, the lattice parameters linearly increased with increasing x. The lattice parameter of the end member x = 1 (PbSe) film (0.6128 nm) is almost the same as that of the bulk PbSe (0.613 nm, see Figure 2(b)), implying that the films grew heteroepitaxially on MgO single crystals with the crystalline lattice fully relaxed.

3.3.2. Growth of RS-Type (Sn,Pb)Se by R-SPE Combined with a Rapid Quenching Method. Figure 6(a) summarizes out-of-plane HRXRD patterns of the films annealed at Ta = 600 °C, which were quenched rapidly in water from 600 °C to RT. When the thickness ratio of the as-grown bilayer film was SnSe:PbSe = 3:1, only diffractions from the orthorhombic phase were observed in the quenched film. With an increase in the PbSe thickness ratio (see SnSe:PbSe = 1:1), the RS-phase 200 diffraction was observed at 2θ = 29°; and then we further increased the PbSe thickness (see SnSe:PbSe = 1:1.2). It should be noted that we succeeded in observing the single-phase diffraction pattern originating from the RS-type crystal structure. The chemical composition, i.e., Sn:Pb, is 1:1.0, determined by EPMA. The RS phase is thermodynamically unstable for this chemical composition at RT because the RS phase is stable at temperatures higher than 600−650 °C according to the SnSe−PbSe phase diagram.12 Thus, these results indicate that we succeeded in stabilizing its hightemperature nonequilibrium phase in the epitaxial films fabricated by the R-SPE and rapid-quenching combined method. This Pb concentration of 50% is 13% lower than the minimum Pb concentration in RS-type SnSe (50%) in previous reports,13,14 indicating the present technique is effective in stabilizing the high-temperature RS-type phase. The fwhm of the RS-type epitaxial film of the XRC 200 diffraction is composed of two components with different FWHMs (Figure 6(b)), i.e., sharp (fwhm = 0.05°) and very broad ones (0.6°), indicating that two kinds of crystallites exist in the film although in-plane ϕ scan in the inset of Figure 6(c) shows a singledomain 4-fold symmetry structure due to the cubic lattice. This in-plane HRXRD pattern provides us in-plane orientation relationship between the RS-type (Sn0.5Pb0.5)Se film and the MgO single crystal, which is illustrated in Figure 6(e). When the film has the orthorhombic GeS-type structure, the in-plane orientation is explained by a quasi cube-on-cube relationship as illustrated in Figure 6(d).8 On the other hand, after transition from the orthorhombic GeS-type to the cubic RS-type structure, the unit cell becomes larger and is rotated by 45°; consequently the epitaxial relationship is explained by 45°rotated epitaxy. Figure 6(f) summarizes the surface morphology of the obtained epitaxial films. The single-phase (but Pbsubstituted) orthorhombic film has a granular structure with an average grain size of 87 nm and root means square roughness (Rrms) of 1.6 nm. At the surface of the mixed phase film (i.e., orthorhombic and cubic phases), the average grain size (87 nm) and Rrms (1.6 nm) are the same as those of the orthorhombic film; but a grain (∼100−150 nm in lateral size) 2282

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(a) and in the IR region (b). As seen in Figure 9(a), the band gap of the x = 0 film (SnSe) was clearly observed with Eg ∼ 0.8 eV,8 while Eg was decreased sharply by Pb addition as low as 30%; i.e., all the Pb doped films did not show optical absorption edges in the visible−near IR region down to 0.5 eV. Therefore, we performed FT-IR measurements of the Pb doped films (Figure 9(b)). As observed here, the Eg of Pb-doped films are smaller than that of the end members SnSe (0.8 eV) and PbSe; i.e., a clear band gap bowing is observed as reported in refs 14 and 25. The Eg of PbSe estimated at α = 104 cm−1 is 0.2 eV, which is slightly smaller than those of reported values (0.25− 0.3 eV) because of large amounts of subgap states at hν < 0.2 eV. It is also very difficult to determine the Eg of the Pb-doped samples because strong tail states are also observed in the region at hν < 0.2 eV. However, we can see that the absorption spectrum of the x = 0.65 film is shifted to the lowest energy, suggesting that the band gap of (Sn1−xPbx)Se takes a minimum in the RS-type structure with x = 0.65 (Eg at α = 104 cm−1 is estimated to be ∼0.14 eV) and being consistent with previous reports.14,25 Figure 10 summarizes the relationship between carrier transport properties of the (Sn1−xPbx)Se films and x in (Sn1−xPbx)Se (i.e., structure types). Up to x = 0.2, the crystal structure keeps the orthorhombic one, and μHall slightly decreases from 60 (x = 0) to 50 cm2/(Vs) (x = 0.2), whose low mobilities are explained by the poorer crystallinity (see XRC of Figures S2(a) and (b) in the Supporting Information). With a further increase in x to 0.3, the crystal structure changed to the orthorhombic and cubic mixture phase, and then μHall jumped up from 50 to 272 cm2/(Vs), which would be attributed to the better carrier transport through the edgesharing network structure of [(Sn,Pb)Se6] octahedra as discussed in the Introduction. At x = 0.5, the structure transformed to the single cubic phase, its high hole mobility was kept (270 cm2/(Vs)) with increasing Nh to 3.2 × 1019 cm−3. The maximum hole mobility (290 cm2/(Vs)) was achieved at x = 0.58. From x = 0 to 0.58, all the films exhibited p-type characteristics. It should be noted that carrier type conversion to n-type occurred at x = 0.61, where the high electron mobilities of 340 (x = 0.61) − 183 (x = 1) are obtained with Ne = (1−2) × 1019 cm−3. The electron mobilities are higher for the solid-solution RS films than for the end member, PbSe (x = 1), whose result appears to contradict with the fact that the effective mass of PbSe is much smaller than that of SnSe9 and thus expected to be smaller than those of the Sn-rich solid solutions. This result might be affected by the randomness and more disordered local structures of the solid solutions that reduce the carrier relaxation time and consequently deteriorate the carrier mobility. However, we like to point that the band gap bowing effect14,25 observed in Figure 9 may have an essential effect because a narrower band gap forms more dispersed bands, resulting in smaller carrier effective masses and higher mobilities. In addition, we should also consider carrier transport deterioration due to the microstructures of the epitaxial films. Actually, the reported mobilities of a PbSe single crystal (830−930 and 1200−1300 cm2/(Vs) for hole and electron, respectively)9 are much higher than the values obtained in this study, whose behavior is very commonly observed in epitaxial films because domain boundaries and dislocations work as carrier scattering centers and carrier transport potential barriers. As for the origin of the n-type conduction, it would be attributed to the Se vacancy because in the case of PbSe it is known that a small amount of

Figure 6. R-SPE growth of the RS-type (Sn,Pb)Se epitaxial films. (a) Out-of-plane HRXRD patterns of the top SnSe/bottom PbSe template bilayer on MgO annealed thermally at Ta = 600 °C, followed by rapid water-quenching to RT. Thickness ratio of SnSe and PbSe layers before thermal annealing is shown on right. The bottom pattern is shown as an example of an as-grown film for comparison. (b) XRC of the cubic 200 diffraction of the (Sn0.5Pb0.5)Se epitaxial film. Circles and lines show the observed pattern and peak deconvolution results, respectively. (c) In-plane HRXRD patterns of the cubic (Sn0.5Pb0.5)Se epitaxial film. Inset shows the in-plane ϕ scan of the cubic 022 diffraction of the (Sn0.5Pb0.5)Se epitaxial film. (d,e) In-plane epitaxial relationship between the films and MgO(100) single-crystal substrates. The film unit cells are drawn by the blue rectangles. Crystallographic axis directions are drawn based on the films. The lattice parameters of the films and the substrates at RT are given as well. (d) Orthorhombic GeS-type film on MgO (100), showing the quasi cube-on-cube structure, i.e. with the in-plane relationship of orthorhombic SnSe [010] [001] || MgO [001]. (e) Cubic RS-type film on MgO (100), showing the 45°-rotated epitaxy from the cube-on-cube epitaxial relationship. (f) AFM images of (Sn,Pb)Se films; (i) orthorhombic GeS-type (Sn,Pb)Se film grown using the thickness ratio of SnSe:PbSe = 3:1, (ii) mixed phase film (orthorhombic and cubic) grown using a thickness ratio of SnSe:PbSe = 1:1, and (iii) cubic RS-type (Sn,Pb)Se film grown using a thickness ratio of SnSe:PbSe = 1:1.2, i.e., (Sn0.5Pb0.5)Se film.

3.3.4. Narrow Band Gap and High Carrier Mobility of RSType (Sn1−xPbx)Se Films. Figure 9 shows the optical absorption spectra of the (Sn1−xPbx)Se films in the visible−near IR region 2283

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Figure 7. R-SPE growth with quenching in water for RS-type (Sn1−xPbx)Se epitaxial films with x ≥ 0.5. (a) Out-of-plane HRXRD patterns. Right values indicate the Pb concentrations in the films (x) determined by EPMA. The pattern of the PbSe (x = 1) film given in Figure 5(a) is shown for comparison. Vertical bars indicate the peak positions of the (Sn1−xPbx)Se 600 diffractions. (b) FWHMs of out-of-plane XRC of the 200 diffractions as a function of x. Peaks 1 and 2 indicate broad and narrow peaks, respectively.

Figure 8. In-plane lattice parameters of (Sn1−xPbx)Se epitaxial films. End members, SnSe and PbSe, were grown directly by PLD. Solid solutions, i.e. 0 < x < 1, were grown by R-SPE and rapid quenching. There are the following three areas. (i) Orthorhombic GeS-type phase area in 0 ≤ x < ∼2.5. Data of the x = 0 film is taken from ref 8. (ii) Mixed phase area of orthorhombic and cubic phases in 0.3 ≤ x < 0.5. (iii) Cubic RS-type phase area in 0.5 ≤ x ≤ 1.0.

Figure 10. Carrier transport properties of the (Sn1−xPbx)Se films as a function of x. σ, Nh,e, and μHall show electrical conductivity, carrier concentration of hole or electron, and Hall mobility, respectively. Carrier types are indicated by the colors of the respective regions; the samples in red and blue regions are p-type and n-type, respectively. Crystalline phase for each Pb composition and boundary are shown on top. Data of the x = 0 film is taken from ref 8.

before thermal annealing in the high x region; i.e., a small amount of Pb metal is observed in the PbSe template layer as seen in Figure 5(a). These results indicate that the RS structure is effective to attaining high hole mobility in SnSe by nonequilibrium Pb doping. Further, carrier type conversion is also possible by tuning Pb concentration in (Sn1−xPbx)Se films.

4. CONCLUSIONS Isovalent Pb doping to the orthorhombic GeS-type SnSe was examined in order to stabilize the nonequilibrium RS-type (Sn,Pb)Se and realize high hole mobility in the nonequilibrium phase. A simple conventional PLD growth resulted in phaseseparated films composed of the SnSe and PbSe phases. The method combined the R-SPE with rapid-quenching successfully grew RS-type (Sn,Pb)Se heteroepitaxial films. The minimum Pb concentration to stabilize the RS-type SnSe was 50%, which

Figure 9. Optical absorption spectra of the (Sn1−xPbx)Se epitaxial films in (a) the visible−near IR region obtained with a conventional spectrometer and (b) the IR region obtained with a FT-IR spectrometer. Data of the x = 0 film is taken from ref 8.

Se vacancy contributes to its n-type character.9 Actually our EPMA results show that the Se concentration in the n-type samples is ∼0.91(1); i.e., ∼10% Se deficient. The Se deficiency originates from the thicker PbSe template layer in bilayer films 2284

DOI: 10.1021/acs.chemmater.6b00307 Chem. Mater. 2016, 28, 2278−2286

Article

Chemistry of Materials

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is 13% widened from the previous reports. Structure transition from GeS-type to RS-type increased hole mobility drastically from 60 for SnSe to 290 cm2/(Vs), which is explained by the edge-sharing network having shorter Sn−Sn and Se−Se chemical bonds in the RS-type structure. It should be noted that n-type conversion is also observed by further Pb doping ≥61% with the highest electron mobility of 340 cm2/(Vs). These characteristics open ways to utilize the RS-type SnSe as an active layer in electronic devices and also to develop a new topological insulator with a higher topological transition temperature in a nonequilibrium phase.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.6b00307. XRR spectra, XRC, and an in-plane HRXRD pattern (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Ministry of Education, Culture, Sports, Science and Technology (MEXT) through Element Strategy Initiative to Form Core Research Center and also by the New Energy and Industrial Technology Development Organization (NEDO) under the Ministry of Economy, Trade and Industry (METI). H. Hiramatsu was also supported by the Japan Society for the Promotion of Science (JSPS) through a Grant-in-Aid for Scientific Research on Innovative Areas “Nano Informatics” Grant Number 25106007 and Support for Tokyotech Advanced Research (STAR). Crystal structures of Figures 1, 8, and 10 were drawn by the VESTA program.26 XRD pattern simulation in Figure 3 was performed using the RIETAN-FP code.27

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REFERENCES

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DOI: 10.1021/acs.chemmater.6b00307 Chem. Mater. 2016, 28, 2278−2286

Article

Chemistry of Materials Semiconductors to Fe-Based Superconductors. Physica C 2009, 469, 657−666. (25) Strauss, A. J. Inversion of Conduction and Valence Bands in Pb1−xSnxSe Alloys. Phys. Rev. 1967, 157, 608−611. (26) Momma, K.; Izumi, F. VESTA 3 for Three-Dimensional Visualization of Crystal, Volumetric and Morphology Data. J. Appl. Crystallogr. 2011, 44, 1272−1276. (27) Izumi, F.; Momma, K. Three-Dimensional Visualization in Powder Diffraction. Solid State Phenom. 2007, 130, 15−20.

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DOI: 10.1021/acs.chemmater.6b00307 Chem. Mater. 2016, 28, 2278−2286