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Growth of FeSb2 thin films by magnetron sputtering
Thin Solid Films 519 (2011) 5397–5402
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Thin Solid Films j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / t s f
Growth of FeSb2 thin films by magnetron sputtering Y. Sun a, E. Zhang b, S. Johnsen a, M. Sillassen b, P. Sun c, F. Steglich c, J. Bøttiger b, B.B. Iversen a,⁎ a b c
Center for Energy Materials, Department of Chemistry and iNANO, Aarhus University, DK-8000 Aarhus C, Denmark Department of Physics and Astronomy and iNANO, Aarhus University, DK-8000 Aarhus C, Denmark Max Planck Institute for Chemical Physics of Solids, D-01187 Dresden, Germany
a r t i c l e
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Article history: Received 10 August 2010 Received in revised form 15 February 2011 Accepted 15 February 2011 Available online 24 February 2011 Keywords: Thermoelectrics Strongly correlated systems Sputtering FeSb2 Iron antimonide Scanning electron microscopy X-ray diffraction
a b s t r a c t The detailed growth of FeSb2 films formed on quartz (0001) substrates by magnetron sputtering is reported. FeSb2 films with different orientations and compositions can be produced by adjusting the Ar working gas pressure and the substrate temperature. By employing FeSb2 thin layers produced at different substrate temperatures as templates, b 101N-, b 120N- and b 002N-textured FeSb2 films were produced under identical growth conditions. The thermoelectric properties of film samples grown at different temperatures were measured and the effects of Sb and FeSb impurities were investigated. © 2011 Elsevier B.V. All rights reserved.
1. Introduction FeSb2 was studied extensively 40 years ago [1–4], but it has recently been rediscovered as a material with promising application potential due to its exciting physical properties [5–13]. FeSb2 has been characterized by various techniques as a strongly correlated semiconductor with a small hybridization gap at the Fermi level [5–7]. It was suggested that electron–electron correlations lead to a large electronic density of states at the band edges of the hybridization gap, which results in much enhanced absolute values of the Seebeck coefficient (thermopower, |S|) [14]. Correlated semiconductors potentially can reach enormous values of the power factor (P = S2σ, where σ is electrical conductivity) at low temperature by combining electron-correlation enhanced thermopower with a semiconductorlike conductance [15]. As an archetypical example of correlated semiconductors, single crystals of FeSb2 were discovered to exhibit colossal thermopower values up to ~ − 45,000 μV K− 1 and record high power factors up to ∼ 2300 μW K− 2 cm− 1 at 12 K [8]. The maximum value of P is 65 times larger than that of state-of-the-art Bi2Te3-based thermoelectric materials. Further investigation revealed that the significant thermoelectricity of FeSb2 is dominated by diffusive electrons of narrow d-bands rather than phonon drag effect [9–11]. FeSb2 therefore has promising application potential for thermoelectric cooling at cryogenic temperatures. Moreover, it was also reported that
⁎ Corresponding author. E-mail address: [email protected] (B.B. Iversen). 0040-6090/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2011.02.053
FeSb2 has giant carrier mobilities of up to ~ 105 cm2 V− 1 s− 1 and Codoped FeSb2 has colossal magnetoresistance. These results imply potential application of FeSb2 in high-speed electronic and spintronic devices [12,13]. Although FeSb2 single crystals present huge thermopower values at cryogenic temperatures, the thermoelectric performance is restricted due to a large lattice thermal conductivity (κL). Thin film thermoelectric materials (TMs) are known to have much reduced κL due to surface and grain-boundary scattering of phonons, and thus remarkably enhanced thermoelectric performance compared with bulk TMs [16–18]. The calculated mean free path of the dominant phonons of single crystal FeSb2 at 12 K is as large as ~ 100 μm [8], and therefore FeSb2 thin films are expected to have remarkably reduced lattice thermal conductivity at low temperature. Additionally, by combining the remarkable properties of FeSb2 with state-of-the-art semiconductor techniques, FeSb2 films could be integrated into electronic devices to fulfill their thermoelectric and electronic application potential [19]. In our previous work, the growth and thermoelectric properties of FeSb2 films prepared by magnetron sputtering were studied [20,21]. Here we report in-depth parametric studies of the growth of FeSb2 thin films. The effect of growth parameters on the structures and orientations of the film samples were explored. Furthermore, by employing different template layers, b101N-, b120N- and b002N-textured FeSb2 films can be produced under identical growth conditions. The anisotropic transport properties of FeSb2 have been reported by several groups. Both the resistivity and the thermopower of FeSb2 were found to be crystal axisdependent [3,8,22]. Therefore, controllable growth and, especially,
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orientation control of FeSb2 films could be a significant step for their property optimization. 2. Experimental details
3. Results and discussion Controllable growth of FeSb2 films was realized by adjusting the Ar working gas pressure and substrate temperature. By sputtering the Type A target (produced with Fe chips and Sb powder in stoichiometric ratio of 1:2), the optimized growth conditions for nearly phase-pure FeSb2 films grown on quartz were found at a substrate temperature of 350 °C and an Ar pressure of ~ 0.6 Pa [20]. An XRD pattern from such an FeSb2 film grown for 3 h is shown in Fig. 1. The FeSb2 (101) and (202) peaks dominate the whole pattern, indicating formation of highly b101N-textured FeSb2 films. Top view and crosssectional view SEM images of that FeSb2 film are presented in the inset of Fig. 1. These images confirm the formation of a uniform polycrystalline film with a thickness of ~ 600 nm. 3.1. Effect of Ar pressure At the substrate, the average kinetic energy and flux of the deposited species depend on the sputter-gas pressure through collisions with gas atoms. This may have a significant influence on the growth rate as well as the film structure. Furthermore, if the target contains heavier elements, there is a significant probability that the sputter-gas ions (Ar+) will be reflected as neutral atoms. By increasing the sputter-gas pressure, a higher flux of reflected neutrals will eventually hit the growing film and thereby affect the composition and film structure. Additionally, a higher flux of bombarding Ar+ ions from the plasma may
Fig. 1. XRD pattern from an FeSb2 film sample grown at 350 °C with Ar pressure of 0.6 Pa for 3 h by sputtering a Type A target (produced with Fe chips and Sb powder in stoichiometric ratio of 1:2). The inset shows the top and cross-sectional views of SEM images of the FeSb2 film. The scale bar applies to both SEM images.
arise as a result of increased sputter-gas pressure if the substrate is kept at a floating potential (~− 10 V) or negatively biased. Fig. 2 shows normalized XRD patterns of three FeSb2 films grown at 350 °C for 3 h but with different Ar pressures of 0.6 Pa, 1.6 Pa, and 1.8 Pa, respectively. The FeSb2 film grown with an Ar pressure of 0.6 Pa exhibits b101N-orientation, and the FeSb2 (101) and (202) peaks dominate the whole XRD pattern. Increasing Ar pressure leads to degradation of the b101N-orientation; and (120), (111) and (002) peaks are present in the XRD pattern of FeSb2 films grown with an Ar pressure of 1.6 Pa. A further increase in Ar pressure to 1.8 Pa results in the formation of b002N-textured FeSb2 films. Fig. 3 presents top-view and cross-sectional SEM images of the FeSb2 films shown in Fig. 2. The thicknesses of these three films were measured as (a) 600 nm, (b) 140 nm, and (c) 110 nm, respectively. This indicates that the Ar pressure has an important influence on the growth rates of the FeSb2 films. A higher Ar pressure results in a lower growth rate and thinner FeSb2 films due to fewer sputtered target atoms arriving at the substrate. The average crystallite sizes of the films are also observed to decrease with increasing Ar pressure,
(a) 0.6 Pa (101)
Normalized Intensity
FeSb2 films were deposited on quartz (0001) wafers by sputtering of specifically prepared compound targets. The 1-inch-targets were made by heating 99.99% Fe chips and 99.9999% Sb powder in a stoichiometric ratio of 1:2 (Type A target) or in a ratio of 1:4 (Type B target) in a corundum crucible for reaction, and subsequently the target was annealed in an induction furnace with ~ 1 MPa Ar atmosphere for 3 h. The growth chamber was equipped with an on-axis unbalanced magnetron source with a target-to-substrate distance of 10 cm. The chamber base pressure was approximately 1 × 10− 5 Pa. Ar (purity 99.9996%) was used as sputter-gas at a flow rate of 8 sccm. During the depositions, a constant power of 10 W was applied to the targets and the substrates were electrically floating. Different substrate temperatures (~ 25–400 °C), Ar pressures (0.6– 1.8 Pa) and growth times (10 min–3 h) were used to control the thin film growth. The as-deposited samples were characterized and analyzed by scanning electron microscopy (SEM, Nova600 NanoLab, FEI) with Energy Dispersive X-ray analysis (EDX), X-ray diffraction (XRD, D8 Discover, Bruker AXS) in θ–2θ geometry with CuKα radiation, and Rutherford Backscattering Spectrometry (RBS, using 2 MeV 4He+ and a scattering angle of 161°). The RBS spectra were simulated using the SIMNRA software [23]. Thermoelectric transport properties of the FeSb2 films were measured with a Quantum Design physical property measurement system. The resistivity (ρ) and S of the film samples were measured simultaneously in the thermal transport option using a standard fourpoint setup. Gold-coated copper wires were mounted on the film samples (8 mm × 2 mm) using conducting silver epoxy. The Hall resistivity (ρH) of the FeSb2 films was measured while sweeping the magnetic field (B) in two opposite directions in the AC transport option. Platinum wires were mounted on the film using conducting silver paste. After eliminating the resistive contributions to ρH by taking the average values in the two fields, the Hall coefficient (RH) was determined by ρH = RHB.
(202)
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2θ (°) Fig. 2. Normalized XRD patterns of three FeSb2 films grown at 350 °C for 3 h by sputtering a Type A target but with different Ar pressures of (a) 0.6 Pa, (b) 1.6 Pa and (c) 1.8 Pa, respectively. The peaks marked with asterisks are associated with the substrate.
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Fig. 3. Top-view and cross-sectional SEM images of the FeSb2 films grown at 350 °C for 3 h by sputtering a Type A target but with different Ar pressures of (a) 0.6 Pa, (b) 1.6 Pa and (c) 1.8 Pa, respectively. The scale bar applies to all the SEM images.
presumably due to that at higher Ar pressures a higher flux of energetic Ar+ ions (being accelerated from the plasma by the floating substrate potential) reaches the substrate, thereby creating more nucleation centers and hence smaller crystallites. The EDX results reveal that the Fe:Sb ratios of the three FeSb2 films shown in Fig. 2a, b and c are about 33:67, 38:62 and 39:61, respectively. We suggest that the lower Sb content at high Ar pressures originates from resputtering of weakly bound Sb atoms in the films due to a high flux of reflected Ar neutrals (from the target) reaching the substrate. In summary, with increase of the Ar pressure from 0.6 Pa to 1.8 Pa, the FeSb2 films grown at 350 °C evolve from stoichiometric FeSb2 films with b101N-orientation to Fe-rich FeSb2 films with b002N-orientation. Further study revealed that the b002N-orientation of the FeSb2 films does not depend on the composition of the films, since some Sb-rich FeSb2 films also have b002N-orientation (shown in Fig. 4a). Detailed studies on orientation control of FeSb2 films will be discussed below.
3.2. Effect of the substrate temperature The substrate temperature is one of the most important parameters for the growth of high quality FeSb2 films. Since the vapor pressure of
Fig. 4. Normalized XRD patterns of film samples produced with Ar pressure of 0.6 Pa for 15 min at (a) 250 °C, (b) 300 °C, (c) 350 °C, and (d) 400 °C. A Type A target was employed to produce these films. The peaks marked with asterisks are associated with the substrate. The insets show top view SEM images of the FeSb2 films. The scale bar applies to all the SEM images.
Fe is much lower than that of Sb, the composition of the films depends strongly on the substrate temperatures, yielding lower Sb concentrations at higher temperatures. In addition, our work reveals that the substrate temperature also influences the orientations of FeSb2 films. Fig. 4a, b, c and d shows normalized XRD patterns of FeSb2 films produced with Ar pressure of 0.6 Pa for 15 min at 250 °C, 300 °C, 350 °C and 400 °C, respectively. The FeSb2 (002) peak dominates the XRD pattern of the film grown at 250 °C. The existence of a small Sb peak in Fig. 4a indicates that the sample grown at 250 °C is antimony rich, which may be due to the lower Sb vapor pressure at 250 °C. In contrast, only FeSb peaks are observed in Fig. 4d, showing the higher evaporation rate of Sb at 400 °C. The FeSb2 films grown at 300 °C and 350 °C have clear b101N-orientation. Weak Sb and FeSb2 (002) peaks are observed in Fig. 4b, but both peaks can barely be identified in the b101N-textured film shown in Fig. 4c. The evolution of the XRD patterns of the film samples with substrate temperatures clearly reflects the effect of substrate temperature on the compositions and orientations of FeSb2 films. It is notable that the sample presented in Fig. 4a is an Sb-rich b002N-textured FeSb2 film. An Fe-rich b002N-textured FeSb2 film was also prepared at 350 °C with a high Ar pressure (shown in Fig. 2c) which suggests that the growth of the b002N-oriented FeSb2 films does not directly depend on the compositions of the films. b002Ntextured FeSb2 films were formed at 350 °C with Ar pressure of 1.8 Pa or at 250 °C with Ar pressure of 0.6 Pa; while a substrate temperature of 350 °C and an Ar pressure of 0.6 Pa lead to growth of b101Ntextured FeSb2 films. This indicates that either a reduced average energy of the reflected Ar atoms reaching the substrate (due to a high Ar pressure) or a low substrate temperature, or both, can be responsible for the formation of the b002N-textured FeSb2 films. In contrast, the (101) surface appears at low Ar pressure (corresponding to an increase in the energy of reflected Ar neutrals from the target while assuming that the contribution from a lower flux is less significant) and high temperature (350 °C). In both cases, contributions to increase the “effective” surface temperature are made. This suggests that the (101) surface has the lowest energy, but that the (002) surface appears because of kinetic restrictions at lower “effective” temperatures. Top view SEM images of the film samples grown at different substrate temperatures are shown in the insets of Fig. 4. With increasing substrate temperature from 250 °C to 350 °C, the morphology of the FeSb2 film samples evolves from high density islands to continuous films. The crystallite sizes in these FeSb2 thin films are in the range of 20– 100 nm, lower than the crystallite sizes of the FeSb film grown at 400 °C (~200 nm). The very thin FeSb2 or FeSb films contain a large number of uniform small crystallites, which could offer nucleation sites, and therefore could be used as templates to not only guide the orientations but also the crystallite sizes of the following growth of the FeSb2 films, as will be discussed below.
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Y. Sun et al. / Thin Solid Films 519 (2011) 5397–5402 on silica at 375°C for 3h (a) Grown (120)(101)
Normalized Intensity
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Grown on a template layer (200°C) at 375°C for 3h (002) (101) (120)
on a template layer (350°C) at 375°C for 3h (c) Grown(101)
(120) on a template layer (400°C) at 375°C for 3h (d) Grown(120)
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2θ (°) Fig. 5. Normalized XRD patterns of FeSb2 films grown at 375 °C for 3 h with an Ar pressure of 0.6 Pa by sputtering a Type B target (produced with Fe chips and Sb powder in stoichiometric ratio of 1:4) on (a) quartz substrate, (b) thin template layer formed at 200 °C, (c) thin template layer formed at 350 °C and (d) thin template layer formed at 400 °C.
3.3. Effect of template layers In our previous work, a pre-deposited FeSb2 thin layer grown at 200 °C for 10 min has been used as template to control the orientation of FeSb2 films grown at 350 °C. As a result, b002N-textured FeSb2 films on the template layer and b101N-textured FeSb2 films on quartz substrates were formed under identical growth conditions [21]. These two FeSb2 films have nearly the same composition but different orientations and crystallite sizes. The orientation control of FeSb2 films could be critical for the future work to realize epitaxial growth of FeSb2 films as well as the property optimization of FeSb2 films.
In the present work, sputtering of the Type A targets leads to b101N-textured stoichiometric FeSb2 films on quartz wafers at 350 °C with Ar pressure of 0.6 Pa. In order to produce stoichiometric FeSb2 films at a substrate temperature higher than 350 °C, Type B targets produced with Fe and Sb in atomic ratio of 1:4 were employed. As a result, nearly phase pure FeSb2 films were produced at 375 °C with an Ar pressure of 0.6 Pa. The normalized XRD pattern of a stoichiometric FeSb2 film grown on quartz at 375 °C is shown in Fig. 5a. The similar intensities of the FeSb2 (120) and FeSb2 (101) peaks in Fig. 5a indicate that the films directly grown on quartz wafers at 375 °C are not highly textured films. To produce highly textured FeSb2 films at 375 °C, three different pre-deposited layers grown at 200 °C, 350 °C and 400 °C with pressure of 0.6 Pa for 10 min were used as template layers. After deposition of a thin template layer on the quartz substrate, the substrate temperature was adjusted to 375 °C and kept at this temperature for 1 h, and then the film deposition was started again at 375 °C for 3 h. Normalized XRD patterns of the FeSb2 films grown on 200 °C, 350 °C and 400 °C template layers are shown in Fig. 5b, c and d, respectively. As was reported in our previous work, an FeSb2 thin layer deposited at 200 °C has b002N-orientation and a thin layer deposited at 350 °C has b101N-orientation [21]. Thus it is reasonable that the FeSb2 (002) peak dominates the XRD pattern in Fig. 5b and the FeSb2 (101) peak dominates the XRD pattern in Fig. 5c. It is worth noting that the XRD pattern presented in Fig. 5d indicates formation of a highly b120N-textured FeSb2 film on a template layer produced at 400 °C, which is practically an FeSb film. Because the lattice plane spacing of FeSb2 (120) is 0.2916 nm, which is very close to that of FeSb (101) (0.2900 nm), it is not surprising that the FeSb crystallites with b101N-orientation can guide the subsequent growth of b120N-textured FeSb2 films. Top view SEM images of the four FeSb2 films presented in Fig. 5 are shown in Fig. 6. Although these four films were produced under quite similar growth conditions and their film thicknesses are all about 400 nm, it can be seen that the average crystalline sizes of the highly textured films presented in Fig. 6b, c and d are larger than the FeSb2 film
Fig. 6. Top view SEM images of FeSb2 films grown at 375 °C for 3 h with Ar pressure of 0.6 Pa by sputtering a Type B target on (a) quartz substrate, (b) thin template layer formed at 200 °C, (c) thin template layer formed at 350 °C and (d) thin template layer formed at 400 °C. The scale bar applies to all the SEM images.
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grown directly on quartz substrates. In addition, some rectangular-like crystal faces can be observed in Fig. 6a, c and d, while absent in the SEM image of b002N-textured FeSb2 films (Fig. 6b). These SEM results confirm the formation of FeSb2 films with different orientations and the important influence of template layers on it. 3.4. Thermoelectric properties In our previous work, the thermoelectric properties of FeSb2 films as well as their anisotropy were reported [20,21]. It was shown that the thermoelectric properties of the FeSb2 films are dominated by the intrinsic properties of FeSb2 with, however, high charge carrier concentrations [20]. Anisotropy of FeSb2 films were demonstrated by comparing in-plane thermoelectric properties of b101N- and b002Ntextured FeSb2 films grown at 350 °C [21]. In this work, b002N-, b101Nand b120N-textured FeSb2 films were produced at 375 °C. Here, we report on the in-plane thermoelectric transport properties of three film samples formed with Ar pressure of 0.6 Pa for 3 h but at different substrate temperatures: an Sb-rich FeSb2 film grown at 250 °C, a stoichiometric FeSb2 film grown at 350 °C and an FeSb film grown at 400 °C. The data potentially reveal the effect of Sb and FeSb impurities on the thermoelectric properties of FeSb2 films, and they improve our understanding on thermoelectric properties of FeSb2. Further studies on anisotropy of FeSb2 films, especially the measurement on cross-plane thermoelectric properties of FeSb2 films are ongoing. Normalized XRD patterns and top view SEM images of these three samples are shown in Fig. 7. An Sb peak together with the FeSb2 (101) and (111) peaks dominate the XRD pattern of the Sb-rich FeSb2 film (shown in Fig. 7a). EDX reveals that the atomic ratio of Fe:Sb is about 27:73. Both the EDX and XRD results confirm the formation of an Sb-rich FeSb2 film at 250 °C. The SEM image in the inset of Fig. 7a shows that this Sb-rich FeSb2 film consists of a large number of small crystallites with sizes in the range of 20–200 nm. It is notable that the FeSb2 films grown at 250 °C for 15 min has b002N-orientation (shown in Fig. 4a), while the 3 h sample presented in Fig. 7a has lost its b002N-orientation. The excess amount of Sb may segregate to the grain-boundaries thereby limiting coarsening and local epitaxial growth of individual crystallites
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during coalescence and film growth. Thus, renucleation may lead to a film structure consisting of small crystallites with random orientations, as observed in Fig. 7a. The film sample grown at 350 °C is a highly b101N-textured, stoichiometric FeSb2 film. The FeSb2 (101) peak dominates the XRD pattern and an Fe:Sb atomic ratio of 33:67 measured by EDX is consistent with all other samples produced using this optimized growth condition. The crystallite sizes of this continuous film are in the range of 400–600 nm (shown in the inset of Fig. 7b). Fig. 7c shows the XRD pattern of the sample grown at 400 °C. The FeSb (110) peak has the strongest intensity in contrast to the very weak FeSb2 (101) peak. The Fe:Sb atomic ratio of this FeSb2 film is measured as 55:45; and its crystallite sizes are in the range of 400–1000 nm. In-plane S(T) and ρ(T) of these three film samples are shown in Fig. 8a and b, respectively. At T b 165 K, S of the stoichiometric FeSb2 film is negative and the maximum absolute value of S ~ 160 μV K− 1 occurs at ~ 50 K. At T N 165 K, S changes to positive value and increases with T to ~ 34 μV K− 1 at 300 K. The measured S(T) of the Sb-rich FeSb2 film grown at 250 °C present similar temperature dependence to that of the stoichiometric FeSb2 film. The maximum absolute value of S ~ 62 μV K− 1 occurs at ~ 60 K, and S changes to positive value at T N 200 K and increases to ~ 30 μV K− 1 at 300 K. However, |S| of the FeSb film is very small and almost within the measurement uncertainties in the whole temperature range of the measurements. This indicates that the FeSb film produced in this work have very poor thermoelectric performance. ρ(T) of the film samples was measured simultaneously with S(T). Fig. 8b shows the in-plane ρ(T) of the three film samples. ρ(T) of the stoichiometric FeSb2 film decrease with increasing temperature. A band gap energy of ~ 22.6 meV at 100–250 K is estimated based on the Arrhenius plot (presented in the inset of Fig. 8b), which is lower than previously published values of 26–36 meV on polycrystalline samples and single crystals of FeSb2 [11,22]. ρ(T) of the Sb-rich FeSb2 film is smaller than that of the stoichiometric FeSb2 film in the whole temperature range of the measurements. It increases with temperature up to ~ 50 K and then subsequently decreases. The estimated band gap energy of the Sb-rich FeSb2 film is ~ 18.2 meV at 100–250 K. ρ of the FeSb film presents similar value to that of the stoichiometric FeSb2 film at 300 K, but it only increases slightly with decreasing temperature, and
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T (K) Fig. 7. Normalized XRD patterns and top view SEM images of film samples grown for 3 h with an Ar pressure of 0.6 Pa by sputtering a Type A target at (a) 250 °C, (b) 350 °C, and (c) 400 °C. These three films are characterized as (a) an Sb-rich FeSb2 film, (b) a stoichiometric FeSb2 film, and (c) an FeSb film, respectively. The peak marked with an asterisk is associated with the substrate. The scale bar applies to all the SEM images.
Fig. 8. In-plane (a) S(T) and (b) ρ(T) of the Sb-rich FeSb2 film, the stoichiometric FeSb2 film, and the FeSb film presented in Fig. 7. The inset of b shows the Arrhenius plots of the ρ(T) curves of the Sb-rich and stoichiometric FeSb2 films. Their band gap energies were estimated as ~ 18.2 meV and ~ 22.6 meV at 100–250 K, respectively.
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ρ(20 K) of the FeSb film is about one-fourth of that of the stoichiometric FeSb2 film. Obviously, excess Sb in the Sb-rich films should account for the different thermoelectric properties of stoichiometric and Sb-rich FeSb2 films. S(T) and ρ(T) of thin polycrystalline antimony films have been reported by several groups [24,25]. It was shown that S(T) of the Sb films increases with temperature above 80 K. S(80 K) and S (300 K) of Sb films are reported to be ~ 5 and ~ 30 μV K− 1, respectively [24]. Since stoichiometric FeSb2 films present n-type thermoelectric properties at T b 165 K, the existence of p-type Sb impurities in the FeSb2 films should reduce the thermopower of the FeSb2 films at low temperature. While, Sb and FeSb2 films have similar positive values of S at room temperature, and the room temperature S of Sb-rich FeSb2 films should be similar to that of stoichiometric FeSb2 films. ρ(T) of antimony films was reported to increase with temperature at T N 80 K [24]. The small ρ of Sb (~ 1 × 10− 6 Ωm [24]) will result in the Sb-rich FeSb2 film having a smaller ρ(T) than that of stoichiometric FeSb2 film (~ 1–7 × 10− 5 Ωm). As a correlated semiconductor with a narrow band gap, the physical properties of FeSb2 itself, including the band structure, could be strongly influenced by slight impurities and defects, which may also lead to metallic-like conductivity of FeSb2 film at T b ~ 50 K. The thermopower of FeSb has not been reported previously. In order to confirm the thermoelectric properties of the FeSb films in this work, FeSb bulk samples were produced by heating 99.99% Fe chips and 99.9999% Sb powder in a stoichiometric ratio of 1:1 in an induction furnace followed by pressing by Spark Plasma Sintering. According to the XRD patterns FeSb dominate the bulk samples with trace amounts of FeSb2 impurity. |S| of the FeSb bulk samples was measured to be smaller than 5 μV K− 1 throughout the temperature range of measurement (T b 300 K), which is consistent with the results of the FeSb films. Based on the thermoelectric properties of FeSb films measured in this work, the effect of FeSb impurities on the thermoelectric properties of FeSb2 films has been revealed. Due to its very low thermopower of ~ 0, FeSb impurities will reduce the thermopower of FeSb2 films in the whole temperature range of the measurements. On the other hand, FeSb impurities are not supposed to strongly affect ρ(T) of FeSb2 films due to their similar ρ at T b 300 K. Since both Sb and FeSb impurities will reduce the thermoelectric performance of FeSb2 films at low temperature, it is essential to grow stoichiometric FeSb2 films with high quality and purity. Slight impurities and crystal defects of FeSb2 films will result in largely reduced thermopower at low temperature compared with FeSb2 single crystals. The realization of epitaxial growth of FeSb2 films and ultimately production of single crystalline FeSb2 films clearly merit further study. The current studies on the growth of FeSb2 films, especially orientation control of FeSb2 films should be important for growth of single crystalline FeSb2 films and their property optimization. 4. Conclusion FeSb2 films were formed on quartz (0001) substrate by means of magnetron sputtering. The influence of substrate temperature and Ar working gas pressure on the structures, compositions and orienta-
tions of the film samples were explored. The orientations of the FeSb2 films could be controlled by adjusting the substrate temperature, the Ar working gas pressure or employing thin template layers. The template layers grown at different temperature facilitate synthesis of b101N-, b120N- and b002N-textured FeSb2 films under identical growth conditions. These results are important for studies of anisotropy and property optimization of FeSb2 films. The thermoelectric properties of FeSb films, stoichiometric and Sb-rich FeSb2 films produced at different growth temperature were studied and compared. The effects of Sb and FeSb impurities on the thermoelectric properties of FeSb2 films were discussed. These findings should serve to strengthen interest in application of FeSb2 films in thermoelectrics and electronics. Acknowledgments The authors are grateful to the Danish Strategic Research Council (Center for Energy Materials) and The Danish National Research Foundation (Center for Materials Crystallography) for financial support. References [1] H. Holseth, A. Kjekshus, Acta Chem. Scand. 23 (1969) 3043. [2] H. Holseth, A. Kjekshus, A.F. Andresen, Acta Chem. Scand. 24 (1970) 3309. [3] A.K.L. Fan, G.H. Rosenthal, H.L. McKinzie, A. Wold, J. Solid State Chem. 5 (1972) 136. [4] J. Steger, E. Kostiner, J. Solid State Chem. 5 (1972) 131. [5] C. Petrovic, Y. Lee, T. Vogt, N.D. Lazarov, S.L. Bud'ko, P.C. Canfield, Phys. Rev. B 72 (2005) 045103. [6] A. Bentien, G.K.H. Madsen, S. Johnsen, B.B. Iversen, Phys. Rev. B 74 (2006) 205105. [7] A. Perucchi, L. Degiorgi, R.W. Hu, C. Petrovic, V.F. Mitrovic, Eur. Phys. J. B 54 (2006) 175. [8] A. Bentien, S. Johnsen, G.K.H. Madsen, B.B. Iversen, F. Steglich, Eur. Phys. Lett. 80 (2007) 17008. [9] P. Sun, N. Oeschler, S. Johnsen, B.B. Iversen, F. Steglich, Phys. Rev. B 79 (2009) 153308. [10] P. Sun, N. Oeschler, S. Johnsen, B.B. Iversen, F. Steglich, Appl. Phys. Express 2 (2009) 091102. [11] P. Sun, N. Oeschler, S. Johnsen, B.B. Iversen, F. Steglich, Dalton Trans. 39 (2010) 1012. [12] R. Hu, V.F. Mitrovic, C. Petrovic, Appl. Phys. Lett. 92 (2008) 182108. [13] R. Hu, K.J. Thomas, Y. Lee, T. Vogt, E.S. Choi, V.F. Mitrović, R.P. Hermann, F. Grandjean, P.C. Canfield, J.W. Kim, A.I. Goldman, C. Petrovic, Phys. Rev. B 77 (2008) 085212. [14] G.D. Mahan, Solid State Phys. 51 (1998) 81. [15] S. Paschen, Thermoelectric Handbook, in: D.M. Rowe (Ed.), CRC Press, Taylor & Francis Group, Boca Raton, 20068, Chap. 15. [16] X.F. Qiu, L.N. Austin, P.A. Muscarella, J.S. Dyck, C. Burda, Angew. Chem. Int. Ed. 45 (2006) 5656. [17] R. Venkatasubramanian, E. Siilvola, T. Colpitts, B. O'Quinn, Nature 413 (2001) 597. [18] M. Takashiri, K. Miyazaki, S. Tanaka, J. Kurosaki, D. Nagai, H. Tsukamoto, J. Appl. Phys. 104 (2008) 084302. [19] I. Chowdhury, R. Prasher, K. Lofgreen, G. Chrysler, S. Narasimhan, R. Mahajan, D. Koester, R. Alley, R. Venkatasubramanian, Nat. Nanotechnol. 4 (2008) 235. [20] Y. Sun, S. Johnsen, P. Eklund, M. Sillassen, J. Bøttiger, N. Oeschler, P. Sun, F. Steglich, B.B. Iversen, J. Appl. Phys. 106 (2009) 033710. [21] Y. Sun, E. Zhang, S. Johnsen, M. Sillassen, P. Sun, F. Steglich, J. Bøttiger, B.B. Iversen, J. Phys. D: Appl. Phys. 43 (2010) 205402. [22] C. Petrovic, J.W. Kim, S.L. Bud'ko, A.I. Goldman, P.C. Canfield, Phys. Rev. B 67 (2003) 155205. [23] M. Mayer, Nucl. Instrum. Meth. Phys. Res. B 194 (2002) 177. [24] F. Volklein, Thin Solid Films 191 (1990) 1. [25] A. Boyer, D. Deschacht, E. Groubert, Thin Solid Films 76 (1981) 119.
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Thin Solid Films j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / t s f
Growth of FeSb2 thin films by magnetron sputtering Y. Sun a, E. Zhang b, S. Johnsen a, M. Sillassen b, P. Sun c, F. Steglich c, J. Bøttiger b, B.B. Iversen a,⁎ a b c
Center for Energy Materials, Department of Chemistry and iNANO, Aarhus University, DK-8000 Aarhus C, Denmark Department of Physics and Astronomy and iNANO, Aarhus University, DK-8000 Aarhus C, Denmark Max Planck Institute for Chemical Physics of Solids, D-01187 Dresden, Germany
a r t i c l e
i n f o
Article history: Received 10 August 2010 Received in revised form 15 February 2011 Accepted 15 February 2011 Available online 24 February 2011 Keywords: Thermoelectrics Strongly correlated systems Sputtering FeSb2 Iron antimonide Scanning electron microscopy X-ray diffraction
a b s t r a c t The detailed growth of FeSb2 films formed on quartz (0001) substrates by magnetron sputtering is reported. FeSb2 films with different orientations and compositions can be produced by adjusting the Ar working gas pressure and the substrate temperature. By employing FeSb2 thin layers produced at different substrate temperatures as templates, b 101N-, b 120N- and b 002N-textured FeSb2 films were produced under identical growth conditions. The thermoelectric properties of film samples grown at different temperatures were measured and the effects of Sb and FeSb impurities were investigated. © 2011 Elsevier B.V. All rights reserved.
1. Introduction FeSb2 was studied extensively 40 years ago [1–4], but it has recently been rediscovered as a material with promising application potential due to its exciting physical properties [5–13]. FeSb2 has been characterized by various techniques as a strongly correlated semiconductor with a small hybridization gap at the Fermi level [5–7]. It was suggested that electron–electron correlations lead to a large electronic density of states at the band edges of the hybridization gap, which results in much enhanced absolute values of the Seebeck coefficient (thermopower, |S|) [14]. Correlated semiconductors potentially can reach enormous values of the power factor (P = S2σ, where σ is electrical conductivity) at low temperature by combining electron-correlation enhanced thermopower with a semiconductorlike conductance [15]. As an archetypical example of correlated semiconductors, single crystals of FeSb2 were discovered to exhibit colossal thermopower values up to ~ − 45,000 μV K− 1 and record high power factors up to ∼ 2300 μW K− 2 cm− 1 at 12 K [8]. The maximum value of P is 65 times larger than that of state-of-the-art Bi2Te3-based thermoelectric materials. Further investigation revealed that the significant thermoelectricity of FeSb2 is dominated by diffusive electrons of narrow d-bands rather than phonon drag effect [9–11]. FeSb2 therefore has promising application potential for thermoelectric cooling at cryogenic temperatures. Moreover, it was also reported that
⁎ Corresponding author. E-mail address: [email protected] (B.B. Iversen). 0040-6090/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2011.02.053
FeSb2 has giant carrier mobilities of up to ~ 105 cm2 V− 1 s− 1 and Codoped FeSb2 has colossal magnetoresistance. These results imply potential application of FeSb2 in high-speed electronic and spintronic devices [12,13]. Although FeSb2 single crystals present huge thermopower values at cryogenic temperatures, the thermoelectric performance is restricted due to a large lattice thermal conductivity (κL). Thin film thermoelectric materials (TMs) are known to have much reduced κL due to surface and grain-boundary scattering of phonons, and thus remarkably enhanced thermoelectric performance compared with bulk TMs [16–18]. The calculated mean free path of the dominant phonons of single crystal FeSb2 at 12 K is as large as ~ 100 μm [8], and therefore FeSb2 thin films are expected to have remarkably reduced lattice thermal conductivity at low temperature. Additionally, by combining the remarkable properties of FeSb2 with state-of-the-art semiconductor techniques, FeSb2 films could be integrated into electronic devices to fulfill their thermoelectric and electronic application potential [19]. In our previous work, the growth and thermoelectric properties of FeSb2 films prepared by magnetron sputtering were studied [20,21]. Here we report in-depth parametric studies of the growth of FeSb2 thin films. The effect of growth parameters on the structures and orientations of the film samples were explored. Furthermore, by employing different template layers, b101N-, b120N- and b002N-textured FeSb2 films can be produced under identical growth conditions. The anisotropic transport properties of FeSb2 have been reported by several groups. Both the resistivity and the thermopower of FeSb2 were found to be crystal axisdependent [3,8,22]. Therefore, controllable growth and, especially,
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orientation control of FeSb2 films could be a significant step for their property optimization. 2. Experimental details
3. Results and discussion Controllable growth of FeSb2 films was realized by adjusting the Ar working gas pressure and substrate temperature. By sputtering the Type A target (produced with Fe chips and Sb powder in stoichiometric ratio of 1:2), the optimized growth conditions for nearly phase-pure FeSb2 films grown on quartz were found at a substrate temperature of 350 °C and an Ar pressure of ~ 0.6 Pa [20]. An XRD pattern from such an FeSb2 film grown for 3 h is shown in Fig. 1. The FeSb2 (101) and (202) peaks dominate the whole pattern, indicating formation of highly b101N-textured FeSb2 films. Top view and crosssectional view SEM images of that FeSb2 film are presented in the inset of Fig. 1. These images confirm the formation of a uniform polycrystalline film with a thickness of ~ 600 nm. 3.1. Effect of Ar pressure At the substrate, the average kinetic energy and flux of the deposited species depend on the sputter-gas pressure through collisions with gas atoms. This may have a significant influence on the growth rate as well as the film structure. Furthermore, if the target contains heavier elements, there is a significant probability that the sputter-gas ions (Ar+) will be reflected as neutral atoms. By increasing the sputter-gas pressure, a higher flux of reflected neutrals will eventually hit the growing film and thereby affect the composition and film structure. Additionally, a higher flux of bombarding Ar+ ions from the plasma may
Fig. 1. XRD pattern from an FeSb2 film sample grown at 350 °C with Ar pressure of 0.6 Pa for 3 h by sputtering a Type A target (produced with Fe chips and Sb powder in stoichiometric ratio of 1:2). The inset shows the top and cross-sectional views of SEM images of the FeSb2 film. The scale bar applies to both SEM images.
arise as a result of increased sputter-gas pressure if the substrate is kept at a floating potential (~− 10 V) or negatively biased. Fig. 2 shows normalized XRD patterns of three FeSb2 films grown at 350 °C for 3 h but with different Ar pressures of 0.6 Pa, 1.6 Pa, and 1.8 Pa, respectively. The FeSb2 film grown with an Ar pressure of 0.6 Pa exhibits b101N-orientation, and the FeSb2 (101) and (202) peaks dominate the whole XRD pattern. Increasing Ar pressure leads to degradation of the b101N-orientation; and (120), (111) and (002) peaks are present in the XRD pattern of FeSb2 films grown with an Ar pressure of 1.6 Pa. A further increase in Ar pressure to 1.8 Pa results in the formation of b002N-textured FeSb2 films. Fig. 3 presents top-view and cross-sectional SEM images of the FeSb2 films shown in Fig. 2. The thicknesses of these three films were measured as (a) 600 nm, (b) 140 nm, and (c) 110 nm, respectively. This indicates that the Ar pressure has an important influence on the growth rates of the FeSb2 films. A higher Ar pressure results in a lower growth rate and thinner FeSb2 films due to fewer sputtered target atoms arriving at the substrate. The average crystallite sizes of the films are also observed to decrease with increasing Ar pressure,
(a) 0.6 Pa (101)
Normalized Intensity
FeSb2 films were deposited on quartz (0001) wafers by sputtering of specifically prepared compound targets. The 1-inch-targets were made by heating 99.99% Fe chips and 99.9999% Sb powder in a stoichiometric ratio of 1:2 (Type A target) or in a ratio of 1:4 (Type B target) in a corundum crucible for reaction, and subsequently the target was annealed in an induction furnace with ~ 1 MPa Ar atmosphere for 3 h. The growth chamber was equipped with an on-axis unbalanced magnetron source with a target-to-substrate distance of 10 cm. The chamber base pressure was approximately 1 × 10− 5 Pa. Ar (purity 99.9996%) was used as sputter-gas at a flow rate of 8 sccm. During the depositions, a constant power of 10 W was applied to the targets and the substrates were electrically floating. Different substrate temperatures (~ 25–400 °C), Ar pressures (0.6– 1.8 Pa) and growth times (10 min–3 h) were used to control the thin film growth. The as-deposited samples were characterized and analyzed by scanning electron microscopy (SEM, Nova600 NanoLab, FEI) with Energy Dispersive X-ray analysis (EDX), X-ray diffraction (XRD, D8 Discover, Bruker AXS) in θ–2θ geometry with CuKα radiation, and Rutherford Backscattering Spectrometry (RBS, using 2 MeV 4He+ and a scattering angle of 161°). The RBS spectra were simulated using the SIMNRA software [23]. Thermoelectric transport properties of the FeSb2 films were measured with a Quantum Design physical property measurement system. The resistivity (ρ) and S of the film samples were measured simultaneously in the thermal transport option using a standard fourpoint setup. Gold-coated copper wires were mounted on the film samples (8 mm × 2 mm) using conducting silver epoxy. The Hall resistivity (ρH) of the FeSb2 films was measured while sweeping the magnetic field (B) in two opposite directions in the AC transport option. Platinum wires were mounted on the film using conducting silver paste. After eliminating the resistive contributions to ρH by taking the average values in the two fields, the Hall coefficient (RH) was determined by ρH = RHB.
(202)
(b) 1.6 Pa (101)
(120)
(111)
(002)
*
(202)
(c) 1.8 Pa (002) (101) (111)
*
20
30
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50
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2θ (°) Fig. 2. Normalized XRD patterns of three FeSb2 films grown at 350 °C for 3 h by sputtering a Type A target but with different Ar pressures of (a) 0.6 Pa, (b) 1.6 Pa and (c) 1.8 Pa, respectively. The peaks marked with asterisks are associated with the substrate.
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Fig. 3. Top-view and cross-sectional SEM images of the FeSb2 films grown at 350 °C for 3 h by sputtering a Type A target but with different Ar pressures of (a) 0.6 Pa, (b) 1.6 Pa and (c) 1.8 Pa, respectively. The scale bar applies to all the SEM images.
presumably due to that at higher Ar pressures a higher flux of energetic Ar+ ions (being accelerated from the plasma by the floating substrate potential) reaches the substrate, thereby creating more nucleation centers and hence smaller crystallites. The EDX results reveal that the Fe:Sb ratios of the three FeSb2 films shown in Fig. 2a, b and c are about 33:67, 38:62 and 39:61, respectively. We suggest that the lower Sb content at high Ar pressures originates from resputtering of weakly bound Sb atoms in the films due to a high flux of reflected Ar neutrals (from the target) reaching the substrate. In summary, with increase of the Ar pressure from 0.6 Pa to 1.8 Pa, the FeSb2 films grown at 350 °C evolve from stoichiometric FeSb2 films with b101N-orientation to Fe-rich FeSb2 films with b002N-orientation. Further study revealed that the b002N-orientation of the FeSb2 films does not depend on the composition of the films, since some Sb-rich FeSb2 films also have b002N-orientation (shown in Fig. 4a). Detailed studies on orientation control of FeSb2 films will be discussed below.
3.2. Effect of the substrate temperature The substrate temperature is one of the most important parameters for the growth of high quality FeSb2 films. Since the vapor pressure of
Fig. 4. Normalized XRD patterns of film samples produced with Ar pressure of 0.6 Pa for 15 min at (a) 250 °C, (b) 300 °C, (c) 350 °C, and (d) 400 °C. A Type A target was employed to produce these films. The peaks marked with asterisks are associated with the substrate. The insets show top view SEM images of the FeSb2 films. The scale bar applies to all the SEM images.
Fe is much lower than that of Sb, the composition of the films depends strongly on the substrate temperatures, yielding lower Sb concentrations at higher temperatures. In addition, our work reveals that the substrate temperature also influences the orientations of FeSb2 films. Fig. 4a, b, c and d shows normalized XRD patterns of FeSb2 films produced with Ar pressure of 0.6 Pa for 15 min at 250 °C, 300 °C, 350 °C and 400 °C, respectively. The FeSb2 (002) peak dominates the XRD pattern of the film grown at 250 °C. The existence of a small Sb peak in Fig. 4a indicates that the sample grown at 250 °C is antimony rich, which may be due to the lower Sb vapor pressure at 250 °C. In contrast, only FeSb peaks are observed in Fig. 4d, showing the higher evaporation rate of Sb at 400 °C. The FeSb2 films grown at 300 °C and 350 °C have clear b101N-orientation. Weak Sb and FeSb2 (002) peaks are observed in Fig. 4b, but both peaks can barely be identified in the b101N-textured film shown in Fig. 4c. The evolution of the XRD patterns of the film samples with substrate temperatures clearly reflects the effect of substrate temperature on the compositions and orientations of FeSb2 films. It is notable that the sample presented in Fig. 4a is an Sb-rich b002N-textured FeSb2 film. An Fe-rich b002N-textured FeSb2 film was also prepared at 350 °C with a high Ar pressure (shown in Fig. 2c) which suggests that the growth of the b002N-oriented FeSb2 films does not directly depend on the compositions of the films. b002Ntextured FeSb2 films were formed at 350 °C with Ar pressure of 1.8 Pa or at 250 °C with Ar pressure of 0.6 Pa; while a substrate temperature of 350 °C and an Ar pressure of 0.6 Pa lead to growth of b101Ntextured FeSb2 films. This indicates that either a reduced average energy of the reflected Ar atoms reaching the substrate (due to a high Ar pressure) or a low substrate temperature, or both, can be responsible for the formation of the b002N-textured FeSb2 films. In contrast, the (101) surface appears at low Ar pressure (corresponding to an increase in the energy of reflected Ar neutrals from the target while assuming that the contribution from a lower flux is less significant) and high temperature (350 °C). In both cases, contributions to increase the “effective” surface temperature are made. This suggests that the (101) surface has the lowest energy, but that the (002) surface appears because of kinetic restrictions at lower “effective” temperatures. Top view SEM images of the film samples grown at different substrate temperatures are shown in the insets of Fig. 4. With increasing substrate temperature from 250 °C to 350 °C, the morphology of the FeSb2 film samples evolves from high density islands to continuous films. The crystallite sizes in these FeSb2 thin films are in the range of 20– 100 nm, lower than the crystallite sizes of the FeSb film grown at 400 °C (~200 nm). The very thin FeSb2 or FeSb films contain a large number of uniform small crystallites, which could offer nucleation sites, and therefore could be used as templates to not only guide the orientations but also the crystallite sizes of the following growth of the FeSb2 films, as will be discussed below.
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Y. Sun et al. / Thin Solid Films 519 (2011) 5397–5402 on silica at 375°C for 3h (a) Grown (120)(101)
Normalized Intensity
(111)
(b)
(002)
(202)
Grown on a template layer (200°C) at 375°C for 3h (002) (101) (120)
on a template layer (350°C) at 375°C for 3h (c) Grown(101)
(120) on a template layer (400°C) at 375°C for 3h (d) Grown(120)
(101)
20
30
40
50
60
70
80
2θ (°) Fig. 5. Normalized XRD patterns of FeSb2 films grown at 375 °C for 3 h with an Ar pressure of 0.6 Pa by sputtering a Type B target (produced with Fe chips and Sb powder in stoichiometric ratio of 1:4) on (a) quartz substrate, (b) thin template layer formed at 200 °C, (c) thin template layer formed at 350 °C and (d) thin template layer formed at 400 °C.
3.3. Effect of template layers In our previous work, a pre-deposited FeSb2 thin layer grown at 200 °C for 10 min has been used as template to control the orientation of FeSb2 films grown at 350 °C. As a result, b002N-textured FeSb2 films on the template layer and b101N-textured FeSb2 films on quartz substrates were formed under identical growth conditions [21]. These two FeSb2 films have nearly the same composition but different orientations and crystallite sizes. The orientation control of FeSb2 films could be critical for the future work to realize epitaxial growth of FeSb2 films as well as the property optimization of FeSb2 films.
In the present work, sputtering of the Type A targets leads to b101N-textured stoichiometric FeSb2 films on quartz wafers at 350 °C with Ar pressure of 0.6 Pa. In order to produce stoichiometric FeSb2 films at a substrate temperature higher than 350 °C, Type B targets produced with Fe and Sb in atomic ratio of 1:4 were employed. As a result, nearly phase pure FeSb2 films were produced at 375 °C with an Ar pressure of 0.6 Pa. The normalized XRD pattern of a stoichiometric FeSb2 film grown on quartz at 375 °C is shown in Fig. 5a. The similar intensities of the FeSb2 (120) and FeSb2 (101) peaks in Fig. 5a indicate that the films directly grown on quartz wafers at 375 °C are not highly textured films. To produce highly textured FeSb2 films at 375 °C, three different pre-deposited layers grown at 200 °C, 350 °C and 400 °C with pressure of 0.6 Pa for 10 min were used as template layers. After deposition of a thin template layer on the quartz substrate, the substrate temperature was adjusted to 375 °C and kept at this temperature for 1 h, and then the film deposition was started again at 375 °C for 3 h. Normalized XRD patterns of the FeSb2 films grown on 200 °C, 350 °C and 400 °C template layers are shown in Fig. 5b, c and d, respectively. As was reported in our previous work, an FeSb2 thin layer deposited at 200 °C has b002N-orientation and a thin layer deposited at 350 °C has b101N-orientation [21]. Thus it is reasonable that the FeSb2 (002) peak dominates the XRD pattern in Fig. 5b and the FeSb2 (101) peak dominates the XRD pattern in Fig. 5c. It is worth noting that the XRD pattern presented in Fig. 5d indicates formation of a highly b120N-textured FeSb2 film on a template layer produced at 400 °C, which is practically an FeSb film. Because the lattice plane spacing of FeSb2 (120) is 0.2916 nm, which is very close to that of FeSb (101) (0.2900 nm), it is not surprising that the FeSb crystallites with b101N-orientation can guide the subsequent growth of b120N-textured FeSb2 films. Top view SEM images of the four FeSb2 films presented in Fig. 5 are shown in Fig. 6. Although these four films were produced under quite similar growth conditions and their film thicknesses are all about 400 nm, it can be seen that the average crystalline sizes of the highly textured films presented in Fig. 6b, c and d are larger than the FeSb2 film
Fig. 6. Top view SEM images of FeSb2 films grown at 375 °C for 3 h with Ar pressure of 0.6 Pa by sputtering a Type B target on (a) quartz substrate, (b) thin template layer formed at 200 °C, (c) thin template layer formed at 350 °C and (d) thin template layer formed at 400 °C. The scale bar applies to all the SEM images.
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grown directly on quartz substrates. In addition, some rectangular-like crystal faces can be observed in Fig. 6a, c and d, while absent in the SEM image of b002N-textured FeSb2 films (Fig. 6b). These SEM results confirm the formation of FeSb2 films with different orientations and the important influence of template layers on it. 3.4. Thermoelectric properties In our previous work, the thermoelectric properties of FeSb2 films as well as their anisotropy were reported [20,21]. It was shown that the thermoelectric properties of the FeSb2 films are dominated by the intrinsic properties of FeSb2 with, however, high charge carrier concentrations [20]. Anisotropy of FeSb2 films were demonstrated by comparing in-plane thermoelectric properties of b101N- and b002Ntextured FeSb2 films grown at 350 °C [21]. In this work, b002N-, b101Nand b120N-textured FeSb2 films were produced at 375 °C. Here, we report on the in-plane thermoelectric transport properties of three film samples formed with Ar pressure of 0.6 Pa for 3 h but at different substrate temperatures: an Sb-rich FeSb2 film grown at 250 °C, a stoichiometric FeSb2 film grown at 350 °C and an FeSb film grown at 400 °C. The data potentially reveal the effect of Sb and FeSb impurities on the thermoelectric properties of FeSb2 films, and they improve our understanding on thermoelectric properties of FeSb2. Further studies on anisotropy of FeSb2 films, especially the measurement on cross-plane thermoelectric properties of FeSb2 films are ongoing. Normalized XRD patterns and top view SEM images of these three samples are shown in Fig. 7. An Sb peak together with the FeSb2 (101) and (111) peaks dominate the XRD pattern of the Sb-rich FeSb2 film (shown in Fig. 7a). EDX reveals that the atomic ratio of Fe:Sb is about 27:73. Both the EDX and XRD results confirm the formation of an Sb-rich FeSb2 film at 250 °C. The SEM image in the inset of Fig. 7a shows that this Sb-rich FeSb2 film consists of a large number of small crystallites with sizes in the range of 20–200 nm. It is notable that the FeSb2 films grown at 250 °C for 15 min has b002N-orientation (shown in Fig. 4a), while the 3 h sample presented in Fig. 7a has lost its b002N-orientation. The excess amount of Sb may segregate to the grain-boundaries thereby limiting coarsening and local epitaxial growth of individual crystallites
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during coalescence and film growth. Thus, renucleation may lead to a film structure consisting of small crystallites with random orientations, as observed in Fig. 7a. The film sample grown at 350 °C is a highly b101N-textured, stoichiometric FeSb2 film. The FeSb2 (101) peak dominates the XRD pattern and an Fe:Sb atomic ratio of 33:67 measured by EDX is consistent with all other samples produced using this optimized growth condition. The crystallite sizes of this continuous film are in the range of 400–600 nm (shown in the inset of Fig. 7b). Fig. 7c shows the XRD pattern of the sample grown at 400 °C. The FeSb (110) peak has the strongest intensity in contrast to the very weak FeSb2 (101) peak. The Fe:Sb atomic ratio of this FeSb2 film is measured as 55:45; and its crystallite sizes are in the range of 400–1000 nm. In-plane S(T) and ρ(T) of these three film samples are shown in Fig. 8a and b, respectively. At T b 165 K, S of the stoichiometric FeSb2 film is negative and the maximum absolute value of S ~ 160 μV K− 1 occurs at ~ 50 K. At T N 165 K, S changes to positive value and increases with T to ~ 34 μV K− 1 at 300 K. The measured S(T) of the Sb-rich FeSb2 film grown at 250 °C present similar temperature dependence to that of the stoichiometric FeSb2 film. The maximum absolute value of S ~ 62 μV K− 1 occurs at ~ 60 K, and S changes to positive value at T N 200 K and increases to ~ 30 μV K− 1 at 300 K. However, |S| of the FeSb film is very small and almost within the measurement uncertainties in the whole temperature range of the measurements. This indicates that the FeSb film produced in this work have very poor thermoelectric performance. ρ(T) of the film samples was measured simultaneously with S(T). Fig. 8b shows the in-plane ρ(T) of the three film samples. ρ(T) of the stoichiometric FeSb2 film decrease with increasing temperature. A band gap energy of ~ 22.6 meV at 100–250 K is estimated based on the Arrhenius plot (presented in the inset of Fig. 8b), which is lower than previously published values of 26–36 meV on polycrystalline samples and single crystals of FeSb2 [11,22]. ρ(T) of the Sb-rich FeSb2 film is smaller than that of the stoichiometric FeSb2 film in the whole temperature range of the measurements. It increases with temperature up to ~ 50 K and then subsequently decreases. The estimated band gap energy of the Sb-rich FeSb2 film is ~ 18.2 meV at 100–250 K. ρ of the FeSb film presents similar value to that of the stoichiometric FeSb2 film at 300 K, but it only increases slightly with decreasing temperature, and
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(a)
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0 -60 Sb-rich FeSb 2 film Stoichiometric FeSb2 film
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T (K) Fig. 7. Normalized XRD patterns and top view SEM images of film samples grown for 3 h with an Ar pressure of 0.6 Pa by sputtering a Type A target at (a) 250 °C, (b) 350 °C, and (c) 400 °C. These three films are characterized as (a) an Sb-rich FeSb2 film, (b) a stoichiometric FeSb2 film, and (c) an FeSb film, respectively. The peak marked with an asterisk is associated with the substrate. The scale bar applies to all the SEM images.
Fig. 8. In-plane (a) S(T) and (b) ρ(T) of the Sb-rich FeSb2 film, the stoichiometric FeSb2 film, and the FeSb film presented in Fig. 7. The inset of b shows the Arrhenius plots of the ρ(T) curves of the Sb-rich and stoichiometric FeSb2 films. Their band gap energies were estimated as ~ 18.2 meV and ~ 22.6 meV at 100–250 K, respectively.
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ρ(20 K) of the FeSb film is about one-fourth of that of the stoichiometric FeSb2 film. Obviously, excess Sb in the Sb-rich films should account for the different thermoelectric properties of stoichiometric and Sb-rich FeSb2 films. S(T) and ρ(T) of thin polycrystalline antimony films have been reported by several groups [24,25]. It was shown that S(T) of the Sb films increases with temperature above 80 K. S(80 K) and S (300 K) of Sb films are reported to be ~ 5 and ~ 30 μV K− 1, respectively [24]. Since stoichiometric FeSb2 films present n-type thermoelectric properties at T b 165 K, the existence of p-type Sb impurities in the FeSb2 films should reduce the thermopower of the FeSb2 films at low temperature. While, Sb and FeSb2 films have similar positive values of S at room temperature, and the room temperature S of Sb-rich FeSb2 films should be similar to that of stoichiometric FeSb2 films. ρ(T) of antimony films was reported to increase with temperature at T N 80 K [24]. The small ρ of Sb (~ 1 × 10− 6 Ωm [24]) will result in the Sb-rich FeSb2 film having a smaller ρ(T) than that of stoichiometric FeSb2 film (~ 1–7 × 10− 5 Ωm). As a correlated semiconductor with a narrow band gap, the physical properties of FeSb2 itself, including the band structure, could be strongly influenced by slight impurities and defects, which may also lead to metallic-like conductivity of FeSb2 film at T b ~ 50 K. The thermopower of FeSb has not been reported previously. In order to confirm the thermoelectric properties of the FeSb films in this work, FeSb bulk samples were produced by heating 99.99% Fe chips and 99.9999% Sb powder in a stoichiometric ratio of 1:1 in an induction furnace followed by pressing by Spark Plasma Sintering. According to the XRD patterns FeSb dominate the bulk samples with trace amounts of FeSb2 impurity. |S| of the FeSb bulk samples was measured to be smaller than 5 μV K− 1 throughout the temperature range of measurement (T b 300 K), which is consistent with the results of the FeSb films. Based on the thermoelectric properties of FeSb films measured in this work, the effect of FeSb impurities on the thermoelectric properties of FeSb2 films has been revealed. Due to its very low thermopower of ~ 0, FeSb impurities will reduce the thermopower of FeSb2 films in the whole temperature range of the measurements. On the other hand, FeSb impurities are not supposed to strongly affect ρ(T) of FeSb2 films due to their similar ρ at T b 300 K. Since both Sb and FeSb impurities will reduce the thermoelectric performance of FeSb2 films at low temperature, it is essential to grow stoichiometric FeSb2 films with high quality and purity. Slight impurities and crystal defects of FeSb2 films will result in largely reduced thermopower at low temperature compared with FeSb2 single crystals. The realization of epitaxial growth of FeSb2 films and ultimately production of single crystalline FeSb2 films clearly merit further study. The current studies on the growth of FeSb2 films, especially orientation control of FeSb2 films should be important for growth of single crystalline FeSb2 films and their property optimization. 4. Conclusion FeSb2 films were formed on quartz (0001) substrate by means of magnetron sputtering. The influence of substrate temperature and Ar working gas pressure on the structures, compositions and orienta-
tions of the film samples were explored. The orientations of the FeSb2 films could be controlled by adjusting the substrate temperature, the Ar working gas pressure or employing thin template layers. The template layers grown at different temperature facilitate synthesis of b101N-, b120N- and b002N-textured FeSb2 films under identical growth conditions. These results are important for studies of anisotropy and property optimization of FeSb2 films. The thermoelectric properties of FeSb films, stoichiometric and Sb-rich FeSb2 films produced at different growth temperature were studied and compared. The effects of Sb and FeSb impurities on the thermoelectric properties of FeSb2 films were discussed. These findings should serve to strengthen interest in application of FeSb2 films in thermoelectrics and electronics. Acknowledgments The authors are grateful to the Danish Strategic Research Council (Center for Energy Materials) and The Danish National Research Foundation (Center for Materials Crystallography) for financial support. References [1] H. Holseth, A. Kjekshus, Acta Chem. Scand. 23 (1969) 3043. [2] H. Holseth, A. Kjekshus, A.F. Andresen, Acta Chem. Scand. 24 (1970) 3309. [3] A.K.L. Fan, G.H. Rosenthal, H.L. McKinzie, A. Wold, J. Solid State Chem. 5 (1972) 136. [4] J. Steger, E. Kostiner, J. Solid State Chem. 5 (1972) 131. [5] C. Petrovic, Y. Lee, T. Vogt, N.D. Lazarov, S.L. Bud'ko, P.C. Canfield, Phys. Rev. B 72 (2005) 045103. [6] A. Bentien, G.K.H. Madsen, S. Johnsen, B.B. Iversen, Phys. Rev. B 74 (2006) 205105. [7] A. Perucchi, L. Degiorgi, R.W. Hu, C. Petrovic, V.F. Mitrovic, Eur. Phys. J. B 54 (2006) 175. [8] A. Bentien, S. Johnsen, G.K.H. Madsen, B.B. Iversen, F. Steglich, Eur. Phys. Lett. 80 (2007) 17008. [9] P. Sun, N. Oeschler, S. Johnsen, B.B. Iversen, F. Steglich, Phys. Rev. B 79 (2009) 153308. [10] P. Sun, N. Oeschler, S. Johnsen, B.B. Iversen, F. Steglich, Appl. Phys. Express 2 (2009) 091102. [11] P. Sun, N. Oeschler, S. Johnsen, B.B. Iversen, F. Steglich, Dalton Trans. 39 (2010) 1012. [12] R. Hu, V.F. Mitrovic, C. Petrovic, Appl. Phys. Lett. 92 (2008) 182108. [13] R. Hu, K.J. Thomas, Y. Lee, T. Vogt, E.S. Choi, V.F. Mitrović, R.P. Hermann, F. Grandjean, P.C. Canfield, J.W. Kim, A.I. Goldman, C. Petrovic, Phys. Rev. B 77 (2008) 085212. [14] G.D. Mahan, Solid State Phys. 51 (1998) 81. [15] S. Paschen, Thermoelectric Handbook, in: D.M. Rowe (Ed.), CRC Press, Taylor & Francis Group, Boca Raton, 20068, Chap. 15. [16] X.F. Qiu, L.N. Austin, P.A. Muscarella, J.S. Dyck, C. Burda, Angew. Chem. Int. Ed. 45 (2006) 5656. [17] R. Venkatasubramanian, E. Siilvola, T. Colpitts, B. O'Quinn, Nature 413 (2001) 597. [18] M. Takashiri, K. Miyazaki, S. Tanaka, J. Kurosaki, D. Nagai, H. Tsukamoto, J. Appl. Phys. 104 (2008) 084302. [19] I. Chowdhury, R. Prasher, K. Lofgreen, G. Chrysler, S. Narasimhan, R. Mahajan, D. Koester, R. Alley, R. Venkatasubramanian, Nat. Nanotechnol. 4 (2008) 235. [20] Y. Sun, S. Johnsen, P. Eklund, M. Sillassen, J. Bøttiger, N. Oeschler, P. Sun, F. Steglich, B.B. Iversen, J. Appl. Phys. 106 (2009) 033710. [21] Y. Sun, E. Zhang, S. Johnsen, M. Sillassen, P. Sun, F. Steglich, J. Bøttiger, B.B. Iversen, J. Phys. D: Appl. Phys. 43 (2010) 205402. [22] C. Petrovic, J.W. Kim, S.L. Bud'ko, A.I. Goldman, P.C. Canfield, Phys. Rev. B 67 (2003) 155205. [23] M. Mayer, Nucl. Instrum. Meth. Phys. Res. B 194 (2002) 177. [24] F. Volklein, Thin Solid Films 191 (1990) 1. [25] A. Boyer, D. Deschacht, E. Groubert, Thin Solid Films 76 (1981) 119.