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Correlation between surface morphology and electrical properties of VO2 films grown by direct thermal oxidation method
Applied Surface Science 353 (2015) 1082–1086
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Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc
Correlation between surface morphology and electrical properties of VO2 films grown by direct thermal oxidation method Joonseok Yoon a , Changwoo Park b , Sungkyun Park c , Bongjin Simon Mun d,∗ , Honglyoul Ju a,∗∗ a
Department of Physics, Yonsei University, Seoul 120-749, Republic of Korea Division of Applied Chemistry and Biotechnology, Hanbat National University, Daejeon 305-719 and Advanced Nano Products, Sejong 339-942, Republic of Korea c Department of Physics, Pusan National University, Busan 609-735, Republic of Korea d Department of Physics and Photon Science, Ertl Center for Electrochemistry and Catalysis, Gwangju Institute of Science and Technology, Gwangju 500-712, Republic of Korea b
a r t i c l e
i n f o
Article history: Received 18 May 2015 Received in revised form 5 July 2015 Accepted 6 July 2015 Available online 14 July 2015 Keywords: VO2 thin film Metal-insulator transition Direct thermal oxidation method Surface roughness Grain size
a b s t r a c t We investigate surface morphology and electrical properties of VO2 films fabricated by direct thermal oxidation method. The VO2 film prepared with oxidation temperature at 580 ◦ C exhibits excellent qualities of VO2 characteristics, e.g. a metal-insulator transition (MIT) near 67 ◦ C, a resistivity ratio of ∼2.3 × 104 , and a bandgap of 0.7 eV. The analysis of surface morphology with electrical resistivity of VO2 films reveals that the transport properties of VO2 films are closely related to the grain size and surface roughness that vary with oxidation annealing temperatures. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Vanadium dioxide (VO2 ) exhibits a first-order metal-insulator transition (MIT) near room temperature of ∼67 ◦ C from a low-temperature insulating monoclinic M1 phase to a hightemperature metallic rutile R phase with a drastic resistivity change, as large as ∼105 in the case of VO2 single crystal [1]. This change of resistivity and structures also induces the modification of optical transition bandgap in VO2 , exhibiting significant changes in transmission and reflection of infrared [2]. In addition, the MIT in VO2 occurs at ultrafast timescales and can be triggered by various stimuli such as voltage, light, pressure, etc. [3–5]. Apparently, VO2 has been constantly studied as a promising material for diverse future applications, i.e. thermochromic coating, extreme switching devices, memory, neuromorphic computing, and ultrafast sensors [6–10]. To utilize VO2 in such practical applications, an innovative large-scale VO2 deposition technique is required with the control of surface uniformity.
∗ Corresponding author. Tel.: +82 62 7152882; fax: +82 62 7152224. ∗∗ Corresponding author. Tel.: +82 2 21232607; fax: +82 2 3921592. E-mail addresses: [email protected] (B.S. Mun), [email protected] (H. Ju). http://dx.doi.org/10.1016/j.apsusc.2015.07.036 0169-4332/© 2015 Elsevier B.V. All rights reserved.
Previously, high quality VO2 films with resistivity ratio of above 104 have been reported deposited by various techniques, such as pulsed laser deposition (PLD), chemical vapor deposition (CVD), sol–gel process, reactive magnetron sputtering [11–17]. For instance, PLD, a physical vapor deposition technique, utilizes highenergy laser ablation of target material. This technique is highly desirable for the growth of epitaxial thin film with uniform stoichiometry, yet the fabrication of the large-area VO2 thin film is difficult since PLD uses a small focused area of laser beam for evaporation. In the case of CVD, it is the chemical reactions of volatile precursors that react on substrate surface to yield desired materials. Even though it is possible to have the large-area deposition with high quality thin film with CVD technique, the process of CVD generates undesirable toxic byproducts. Recently, sol–gel method attracts much attention due to its relatively simple setup and low reaction temperature. Yet, the sol–gel method is still economically not viable. On the other hand, the simple sputtering technique has been considered as one of the promising techniques to grow high quality VO2 films over large area with high reproducibility. Of course, there exist a few problems to overcome in fabrication of ideal VO2 film with sputtering deposition technique. In the sputtering deposition, various target materials, such as vanadium metal, V2 O3 , VO2 , and V2 O5 , have been utilized under the mixture of gases
J. Yoon et al. / Applied Surface Science 353 (2015) 1082–1086
in argon and oxygen. However, the oxygen stoichiometry of the films is very sensitive to the oxygen partial pressure during deposition, i.e. the oxidation states of vanadium easily varies from +2 to +5, depending on the partial pressure of oxygen. In addition, during the sputtering process, oxygen gas often reacts non-uniformly with target surface, contributing the non-uniform stoichiometric VO2 thin film. These difficulties of identifying ideal fabrication conditions in sputtering technique make it challenging to grow high quality VO2 film with high uniformity over large area. Previously, as an alternative sputtering method to grow VO2 film, a direct thermal oxidation method has been used, i.e. depositing vanadium metal film, followed by direct thermal oxidation [18–20]. Yet, the quality of VO2 film requires further improvement in electrical properties, such as high resistivity ratio, uniform stoichiometry, and reproducible MIT temperature. In this paper, we report the improved fabrication method for growing high quality large-scale VO2 films by direct thermal oxidation method. We compare the MIT characteristics and surface morphologies of VO2 with various sample preparation conditions, i.e. oxygen annealing temperatures and oxidation pressure. With the use of X-ray diffraction (XRD), atomic force microscopy (AFM), field emission scanning tunneling microscopy (SEM), it is found that electrical transport properties are closely related to the surface grain size and surface roughness of VO2 films.
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Fig. 1. X-ray diffraction (XRD) patterns of VO2 grown at different oxygen annealing temperature condition of TA = 510, 530, 555, 580, and 600 ◦ C under oxidation pressure 2.5 Torr and that of Al2 O3 substrate before film growth. Peaks indicated by ‘x’ are coming from the Al2 O3 substrate. Small peaks from the substrates are visible due to the unfiltered X-ray.
2. Experiment VO2 thin films on Al2 O3 (1 1 2¯ 0) substrates were prepared by direct thermal oxidation method, which involved two step processes, i.e. the sputtering deposition of vanadium metal film at room temperature using vanadium metal target in pure argon environment at process pressure of 5 mTorr, followed by thermal oxidation. The thickness of vanadium metal film and the deposition rate were 100 nm and 0.1 nm/s, respectively. The oxidation pressure was adjusted from 1.5 to 4.5 Torr while substrate temperature was controlled between 510 and 600 ◦ C. Vanadium metal films were annealed at desired annealing temperatures for 10 min. The thickness of film was measured with a profilometer (AlphaStep 500 Surface Profiler, KLA-Tencor). The microstructure and morphology of VO2 thin films were characterized by SEM (JSM-6701F, JEOL) and AFM (XE-BIO, Park Systems). The surface grain size of the films was determined by SEM images while the roughness of surface grain was probed by AFM. To measure the temperature dependence of the resistivity, VO2 films were patterned via photolithography with dimensions of 2.0 mm (length) × 0.24 mm (width) and the electrical contacts were made with indium. The temperature dependence of the electrical resistivity of the VO2 films was measured from 30 to 100 ◦ C using a dc two-contact four-probe method. XRD (MiniFlex, Rigaku) was used to study the structural properties of VO2 thin films. 3. Results and Discussion Fig. 1 shows the XRD patterns of the VO2 films oxidized at various oxidation annealing temperatures. The intensities of main peaks of XRD are enhanced as oxidation annealing temperature increases. The diffraction peaks at 2 = 26.80, 27.88, 37.14, and 55.50◦ , belong to VO2 (1¯ 1 1), (0 1 1), (2 0 0), and (2¯ 2 2), respectively, according to JCPDS card No. 72-0514. On the other hand, the diffraction peaks at 2 = 26.80, 36.02, 45.44, and 55.50◦ correspond to the (0 0 l) (l = 3, 4, 5, and 6) of V6 O13 , according to JCPDS card No. 271318. Also, two diffraction peaks at 2 = 26.80 and 55.50◦ belong to both VO2 and V6 O13 . Overall, XRD results of Fig. 1 show that the films have polycrystalline structure and the mixed phase of VO2 and V6 O13 . It is estimated that the formation of V6 O13 comes from
Fig. 2. (a) Temperature dependence of the resistivity of the VO2 films oxidized at various oxidation annealing temperatures ranging from 510 to 600 ◦ C at fixed oxygen pressure (POxidation ) of 2.5 Torr. (b) Temperature dependence of the resistivity of the VO2 films oxidized at various oxidation annealing pressures ranging from 1.5 to 4.5 Torr at fixed oxidation annealing temperature of 580 ◦ C.
the oxidation of interfacial layers between vanadium metal and substrate. It is believed that the formation of V6 O13 comes from slightly over oxidizing annealing condition. One thing to note is that, the V6 O13 oxide of Fig. 1 is highly textured along the [001] direction, perpendicular to the substrate plane, and this texturing creates the high peak intensity of V6 O13 , which is larger than actual volume fraction. Of course, as will be shown later, the excess formation of this V6 O13 oxide can affect the characteristics of MIT of VO2 . The characteristics of MIT become less distinct when the formation of V6 O13 oxide becomes too excessive, such as the case of oxygen annealing at 600 ◦ C. Also, since the phase transition temperature of V6 O13 oxide is well below the MIT temperature of VO2 [21], it is the structural volume of V6 O13 that induces the degradation of MIT features. We believe that fine adjustments of interfacial
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Fig. 3. Dependence of (a) (30 ◦ C)/(100 ◦ C) (resistivity ratio), (b) the bandgap (Eg ), and (c) two transition temperature TC ↑ (transition temperature during heating) and TC ↓ (transition temperature during cooling) on TA = 510, 530, 555, 580, and 600 ◦ C at fixed oxygen pressure of 2.5 torr. Inset shows −dln (T)/dT versus oxidation annealing temperature for heating (filled red square) and cooling (filled blue circle) cycles. Dependence of (d) (30 ◦ C)/(100 ◦ C), (e) the bandgap, and (f) two transition temperature TC↑ and TC ↓ on POxidation = 1.5, 2.0, 2.5, 3.5, and 4.5 Torr at fixed oxidation annealing of 580 ◦ C. Inset shows the patterned thin film sample with indium contacts for the resistivity measurements. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)
roughness and sputtering step can improve further the uniformity of oxide films. Nonetheless, this direct thermal oxidation method for growing VO2 films provides the excellent MIT characteristics of VO2 , as will be shown below with the measurements of transport properties. Fig. 2 shows the temperature dependence of the resistivity of the VO2 films oxidized at various oxidation annealing temperatures (TA ) and oxygen pressure. All the films show the typical MIT from low temperature insulating state to high temperature metallic state with a hysteresis loop between heating and cooling cycles. First, the temperature dependence of electrical resistivity is measured as a function of TA with fixed oxygen pressure of 2.5 Torr, shown in Fig. 2(a). In Fig. 2(a), it can be clearly seen that the signature of MIT becomes more prominent as TA increases, i.e. the degrees of MIT become significant and the MIT onset temperature starts to increase, approaching to 67 ◦ C. The film grown at TA = 580 ◦ C showed the most pronounced transitions (blue line), i.e. the resistivity change at MIT is almost 4 orders of magnitude. However, the film grown at TA of 600 ◦ C shows the increase in both of high and low temperature resistivity with a slightly reduced resistivity ratio across the MIT. These modified electrical properties will be discussed later with the analysis of surface morphology. As of next step, the annealing oxidation pressure (POxidation ) is varied while TA is fixed at 580 ◦ C, shown in Fig. 2(b). Similar to the case of high annealing oxygen temperature in Fig. 2(a), the characteristics of MIT become significant only after POxidation reaches above 2.5 Torr,
i.e. sharp MIT with high changes of resistivity. However, a film grown at POxidation > 2.5 Torr shows decrease (increase) of resistivity in insulating (metallic) state. Then, based on the result of Fig. 2, the electrical properties of VO2 film at various TA and POxidation are characterized, i.e. the resistivity ratio before and after the MIT, the bandgap, and the onset temperature of MIT, shown in Fig. 3. In Fig. 3, the temperature dependence of the electrical resistivity () of the VO2 films shows strong dependence on the TA . It is believed that the resistivity ratio between the insulating and the metallic phases is most critical indicator of VO2 film quality. Fig. 3(a) shows the dependence of resistivity ratio, (30 ◦ C)/(100 ◦ C), on the oxidation annealing temperatures. The resistivity ratio stays in the range 102 –104 and increases with increasing oxidation temperature, reaching its maximum at TA = 580 ◦ C. The highest value of the resistivity ratio is 2.3 × 104 , whose value is almost comparable to that of VO2 single crystals [22]. Even for the VO2 films with mixed phases of majority VO2 and minority V6 O13 have a large resistivity ratio. This result indicates that our direct thermal oxidation method is a simple, yet very powerful technique to grow large-scale high quality VO2 films with high reproducibility. Next, assuming VO2 is an intrinsic semiconductor, the degree of bandgap is obtained using the relation; = 0 exp (Eg /2kT), where 0 , Eg , and k are the prefactor of resistivity, the bandgap, and the Boltzmann constant, respectively. Calculated bandgap for the film oxidized at TA = 580 ◦ C, based on the above relation and the resistivity curve in the insulating state of Fig. 2(a), shows the bandgap value of 0.7 eV, whose
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Fig. 4. Scanning electron microscope picture of VO2 thin film annealed at (a) 510 ◦ C, (b) 530 ◦ C, (c) 555 ◦ C, (d) 580 ◦ C, and (e) 600 ◦ C in oxygen pressure of 2.5 Torr. Atomic force microscopy image of VO2 thin film annealed at (f) 510 ◦ C, (g) 530 ◦ C, (h) 555 ◦ C, (i) 580 ◦ C, and (j) 600 ◦ C in oxygen pressure of 2.5 Torr. (k) Dependence of grain size(AG ) of the VO2 film on annealing oxidation temperature. (l) Dependence of rms surface roughness(Rq ) of the VO2 film on annealing oxidation temperature.
magnitude is the same as the value of VO2 single crystal [23] and that of taken from infrared spectroscopy [24], i.e. 0.7 eV. As shown in Fig. 3(a) and (b), it is interesting to note that the dependences of resistivity ratio and bandgap on oxidation annealing temperature are very similar, suggesting strong correlation among these quantities. Also, the onsets of MIT temperature follow the similar trend of resistivity ratio and bandgap. Fig. 3(c) shows the dependence of two transition temperature TC↑ (transition temperature during heating) and TC↓ (transition temperature during cooling) on TA . The transition temperatures during heating and cooling cycles are determined as the peak temperatures in the −dln()/dT versus temperature curve, shown in the inset of Fig. 3(c). Both of TC↑ and TC↓ increase as TA is raised up to TA = 550 ◦ C, then reach almost constant values of TC↑ = 68 ◦ C and TC↓ = 60 ◦ C for TA > 550 ◦ C. Hysteresis width (TC ), defined as the difference between TC↑ and TC↓ , was the minimum value of 7.4 ◦ C for the film oxidized at TA = 580 ◦ C. Identical analyses of Fig. 3(a)–(c) are carried out also as POxidation varies at TA of 580 ◦ C, shown in Fig. 3(d)–(f). From Fig. 3(d) and (e), the resistivity ratio and bandgap distribution indicate that oxygen pressure between 2.5 and 3.5 Torr is the critical region to grow high quality of VO2 thin film. Interesting to note is that the film with a large resistivity ratio tends to have a large bandgap, indicating a close relationship between defect states and bandgap. However, the MIT onset temperatures in Fig. 3(f) do not follow any similarity to that of Fig. 3(c), showing almost no dependence of POxidation . Overall, the comparison of Fig. 3(a)–(c) and Fig. 3(d)–(f) indicates that the crystallization of VO2 seems to strongly depend on temperature, however, POxidation will also be decisive factor for stoichiometry of the films. Finally, the surface roughness and grain size of the VO2 films are examined by AFM and SEM to locate the relation between the surface morphology and electrical properties of VO2 films. In Fig. 4, the SEM and AFM images of the VO2 films are shown only as a function
of TA . The identical measurements are taken as varying the oxygen pressure (not shown), yet it shows almost similar results as those of Fig. 4. Fig. 4(a)–(e) shows SEM of the VO2 films deposited at various TA . SEM micrograph shows granular crystallites with sizes dependent on TA . As shown in Fig. 4(k), the average grain sizes of the films are increased monotonically from ∼1 nm to 300 nm as increasing TA from 510 to 600 ◦ C. The grain sizes from SEM and AFM are in good agreements. The enhancement of grain size with increasing TA could be explained by the increased surface mobility of the growth species at the elevated temperature. Fig. 4(f)–(j) shows AFM images of VO2 films oxidized at various TA . AFM images show the strong dependence of rms roughness on the oxidation temperature. As shown in Fig. 3(j), the rms roughness of films is increased at higher oxidation annealing temperature, 5 nm at TA = 510 ◦ C and 80 nm at TA = 600 ◦ C. The rms roughness of the film oxidized at TA = 600 ◦ C is very large, whose magnitude (80 nm) is almost 30% of the VO2 film thickness (250 nm). The AFM and SEM images in Fig. 4 show that the rms roughness of the VO2 film is proportionally increased with the size of grains, suggesting strong correlation between rms roughness and grain size. Interestingly, it can be found that the main characteristics of MIT of VO2 films, i.e. resistivity ratio and sharpness of resistivity curve in Fig. 2, are directly correlated to the VO2 surface morphology of Fig. 4, i.e. the surface roughness and the grain size. In the case of insulating state of Fig. 2(a), the resistivity increases monotonically with oxygen annealing temperature. This means that defect level in the gap becomes reduced [25] and the energy gap becomes close to the value of VO2 at higher annealing temperature, which will be shown later. On the other hand, at metallic state of Fig. 2(a), the resistivity of films do not show much changes among various TA , except sample annealed at 600 ◦ C. In the case of temperature of 510 and 530 ◦ C, low resistivity of film comes mainly from the small surface roughness of Fig. 4(f) and (g). Then, as the temperature
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increase, the surface roughness starts to increase. However, at the same time, the stoichiometry of oxide become now close to that of VO2 , and the size of grain starts to become larger, which then reduces the resistivity of films at metallic state. Consequently, the resistivity becomes close to that of low temperature annealing films. This explanation is also consistent with the case of increased resistivity due to the reduced defect states at insulating state. For the films prepared at 600 ◦ C, according to the result of Fig. 4(k), the grain size is largest, which means the film should have better electrical properties, i.e. lower resistivity. That is, the film oxidized at 600 ◦ C is expected to have better MIT characteristics, such as sharp transition and high resistivity ratio, than those of the film oxidized at 580 ◦ C. However, the film oxidized at 580 ◦ C exhibited better MIT characteristics than those of 600 ◦ C. We believe this is due to both of the increased surface roughness and the growth of V6 O13 at the interface. The increase of rms roughness can easily reduce the resistivity ratio since the metallic resistivity can increase by the critical reduction in current path at very rough surface. This is in contrast to microroughness in nanoporous VO2 film, which enhances luminous transmittance and solar transmittance modulation due to decreased effective refractive index of microporous film [26]. Also, the presence of non-stoichiometric oxide, V6 O13 , can increase the resistivity of thin film, i.e. the resistivity of V6 O13 at metallic state is approximately 0.1 cm at room temperature [21]. In this letter, we present the characterization of VO2 thin films fabricated with simple direct thermal oxidation method. The properties of MIT of VO2 show strong dependence on both of annealing temperature and oxidation pressure. Furthermore, the analysis of surface morphology and electrical properties shows that large grain with smooth surface/interface roughness is required to achieve VO2 films of high performance. This report shows the evidence that the direct thermal oxidation method can be an excellent choice for growing large-scale high quality VO2 film-based devices with reproducibility. Acknowledgments H.L. Ju and B.S. Mun would like to thank the Basic Science Research Program for support through the National Research Foundation of Korea (NRF) funded by the Korean Government (MOE) (Nos. 2012R1A1A2006948 and 2012R1A1A2001745). This paper was supported by GIST 2015 TBP Research Fund and Korea Basic Science Institute Research Grant (E35800). References [1] J. Morin, Oxides which show a metal-to-insulator transition at the Neel temperature, Phys. Rev. Lett. 3 (1959) 34–36. [2] A.S. Barker, H.W. Verleur, H.J. Guggenheim, Infrared optical properties of vanadium dioxide above and below the transition temperature, Phys. Rev. Lett. 17 (1966) 1286–1289. [3] C.N. Berglund, Thermal filaments in vanadium dioxide, IEEE Trans. Electron. Device ED 16 (1969) 432–437.
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Contents lists available at ScienceDirect
Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc
Correlation between surface morphology and electrical properties of VO2 films grown by direct thermal oxidation method Joonseok Yoon a , Changwoo Park b , Sungkyun Park c , Bongjin Simon Mun d,∗ , Honglyoul Ju a,∗∗ a
Department of Physics, Yonsei University, Seoul 120-749, Republic of Korea Division of Applied Chemistry and Biotechnology, Hanbat National University, Daejeon 305-719 and Advanced Nano Products, Sejong 339-942, Republic of Korea c Department of Physics, Pusan National University, Busan 609-735, Republic of Korea d Department of Physics and Photon Science, Ertl Center for Electrochemistry and Catalysis, Gwangju Institute of Science and Technology, Gwangju 500-712, Republic of Korea b
a r t i c l e
i n f o
Article history: Received 18 May 2015 Received in revised form 5 July 2015 Accepted 6 July 2015 Available online 14 July 2015 Keywords: VO2 thin film Metal-insulator transition Direct thermal oxidation method Surface roughness Grain size
a b s t r a c t We investigate surface morphology and electrical properties of VO2 films fabricated by direct thermal oxidation method. The VO2 film prepared with oxidation temperature at 580 ◦ C exhibits excellent qualities of VO2 characteristics, e.g. a metal-insulator transition (MIT) near 67 ◦ C, a resistivity ratio of ∼2.3 × 104 , and a bandgap of 0.7 eV. The analysis of surface morphology with electrical resistivity of VO2 films reveals that the transport properties of VO2 films are closely related to the grain size and surface roughness that vary with oxidation annealing temperatures. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Vanadium dioxide (VO2 ) exhibits a first-order metal-insulator transition (MIT) near room temperature of ∼67 ◦ C from a low-temperature insulating monoclinic M1 phase to a hightemperature metallic rutile R phase with a drastic resistivity change, as large as ∼105 in the case of VO2 single crystal [1]. This change of resistivity and structures also induces the modification of optical transition bandgap in VO2 , exhibiting significant changes in transmission and reflection of infrared [2]. In addition, the MIT in VO2 occurs at ultrafast timescales and can be triggered by various stimuli such as voltage, light, pressure, etc. [3–5]. Apparently, VO2 has been constantly studied as a promising material for diverse future applications, i.e. thermochromic coating, extreme switching devices, memory, neuromorphic computing, and ultrafast sensors [6–10]. To utilize VO2 in such practical applications, an innovative large-scale VO2 deposition technique is required with the control of surface uniformity.
∗ Corresponding author. Tel.: +82 62 7152882; fax: +82 62 7152224. ∗∗ Corresponding author. Tel.: +82 2 21232607; fax: +82 2 3921592. E-mail addresses: [email protected] (B.S. Mun), [email protected] (H. Ju). http://dx.doi.org/10.1016/j.apsusc.2015.07.036 0169-4332/© 2015 Elsevier B.V. All rights reserved.
Previously, high quality VO2 films with resistivity ratio of above 104 have been reported deposited by various techniques, such as pulsed laser deposition (PLD), chemical vapor deposition (CVD), sol–gel process, reactive magnetron sputtering [11–17]. For instance, PLD, a physical vapor deposition technique, utilizes highenergy laser ablation of target material. This technique is highly desirable for the growth of epitaxial thin film with uniform stoichiometry, yet the fabrication of the large-area VO2 thin film is difficult since PLD uses a small focused area of laser beam for evaporation. In the case of CVD, it is the chemical reactions of volatile precursors that react on substrate surface to yield desired materials. Even though it is possible to have the large-area deposition with high quality thin film with CVD technique, the process of CVD generates undesirable toxic byproducts. Recently, sol–gel method attracts much attention due to its relatively simple setup and low reaction temperature. Yet, the sol–gel method is still economically not viable. On the other hand, the simple sputtering technique has been considered as one of the promising techniques to grow high quality VO2 films over large area with high reproducibility. Of course, there exist a few problems to overcome in fabrication of ideal VO2 film with sputtering deposition technique. In the sputtering deposition, various target materials, such as vanadium metal, V2 O3 , VO2 , and V2 O5 , have been utilized under the mixture of gases
J. Yoon et al. / Applied Surface Science 353 (2015) 1082–1086
in argon and oxygen. However, the oxygen stoichiometry of the films is very sensitive to the oxygen partial pressure during deposition, i.e. the oxidation states of vanadium easily varies from +2 to +5, depending on the partial pressure of oxygen. In addition, during the sputtering process, oxygen gas often reacts non-uniformly with target surface, contributing the non-uniform stoichiometric VO2 thin film. These difficulties of identifying ideal fabrication conditions in sputtering technique make it challenging to grow high quality VO2 film with high uniformity over large area. Previously, as an alternative sputtering method to grow VO2 film, a direct thermal oxidation method has been used, i.e. depositing vanadium metal film, followed by direct thermal oxidation [18–20]. Yet, the quality of VO2 film requires further improvement in electrical properties, such as high resistivity ratio, uniform stoichiometry, and reproducible MIT temperature. In this paper, we report the improved fabrication method for growing high quality large-scale VO2 films by direct thermal oxidation method. We compare the MIT characteristics and surface morphologies of VO2 with various sample preparation conditions, i.e. oxygen annealing temperatures and oxidation pressure. With the use of X-ray diffraction (XRD), atomic force microscopy (AFM), field emission scanning tunneling microscopy (SEM), it is found that electrical transport properties are closely related to the surface grain size and surface roughness of VO2 films.
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Fig. 1. X-ray diffraction (XRD) patterns of VO2 grown at different oxygen annealing temperature condition of TA = 510, 530, 555, 580, and 600 ◦ C under oxidation pressure 2.5 Torr and that of Al2 O3 substrate before film growth. Peaks indicated by ‘x’ are coming from the Al2 O3 substrate. Small peaks from the substrates are visible due to the unfiltered X-ray.
2. Experiment VO2 thin films on Al2 O3 (1 1 2¯ 0) substrates were prepared by direct thermal oxidation method, which involved two step processes, i.e. the sputtering deposition of vanadium metal film at room temperature using vanadium metal target in pure argon environment at process pressure of 5 mTorr, followed by thermal oxidation. The thickness of vanadium metal film and the deposition rate were 100 nm and 0.1 nm/s, respectively. The oxidation pressure was adjusted from 1.5 to 4.5 Torr while substrate temperature was controlled between 510 and 600 ◦ C. Vanadium metal films were annealed at desired annealing temperatures for 10 min. The thickness of film was measured with a profilometer (AlphaStep 500 Surface Profiler, KLA-Tencor). The microstructure and morphology of VO2 thin films were characterized by SEM (JSM-6701F, JEOL) and AFM (XE-BIO, Park Systems). The surface grain size of the films was determined by SEM images while the roughness of surface grain was probed by AFM. To measure the temperature dependence of the resistivity, VO2 films were patterned via photolithography with dimensions of 2.0 mm (length) × 0.24 mm (width) and the electrical contacts were made with indium. The temperature dependence of the electrical resistivity of the VO2 films was measured from 30 to 100 ◦ C using a dc two-contact four-probe method. XRD (MiniFlex, Rigaku) was used to study the structural properties of VO2 thin films. 3. Results and Discussion Fig. 1 shows the XRD patterns of the VO2 films oxidized at various oxidation annealing temperatures. The intensities of main peaks of XRD are enhanced as oxidation annealing temperature increases. The diffraction peaks at 2 = 26.80, 27.88, 37.14, and 55.50◦ , belong to VO2 (1¯ 1 1), (0 1 1), (2 0 0), and (2¯ 2 2), respectively, according to JCPDS card No. 72-0514. On the other hand, the diffraction peaks at 2 = 26.80, 36.02, 45.44, and 55.50◦ correspond to the (0 0 l) (l = 3, 4, 5, and 6) of V6 O13 , according to JCPDS card No. 271318. Also, two diffraction peaks at 2 = 26.80 and 55.50◦ belong to both VO2 and V6 O13 . Overall, XRD results of Fig. 1 show that the films have polycrystalline structure and the mixed phase of VO2 and V6 O13 . It is estimated that the formation of V6 O13 comes from
Fig. 2. (a) Temperature dependence of the resistivity of the VO2 films oxidized at various oxidation annealing temperatures ranging from 510 to 600 ◦ C at fixed oxygen pressure (POxidation ) of 2.5 Torr. (b) Temperature dependence of the resistivity of the VO2 films oxidized at various oxidation annealing pressures ranging from 1.5 to 4.5 Torr at fixed oxidation annealing temperature of 580 ◦ C.
the oxidation of interfacial layers between vanadium metal and substrate. It is believed that the formation of V6 O13 comes from slightly over oxidizing annealing condition. One thing to note is that, the V6 O13 oxide of Fig. 1 is highly textured along the [001] direction, perpendicular to the substrate plane, and this texturing creates the high peak intensity of V6 O13 , which is larger than actual volume fraction. Of course, as will be shown later, the excess formation of this V6 O13 oxide can affect the characteristics of MIT of VO2 . The characteristics of MIT become less distinct when the formation of V6 O13 oxide becomes too excessive, such as the case of oxygen annealing at 600 ◦ C. Also, since the phase transition temperature of V6 O13 oxide is well below the MIT temperature of VO2 [21], it is the structural volume of V6 O13 that induces the degradation of MIT features. We believe that fine adjustments of interfacial
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Fig. 3. Dependence of (a) (30 ◦ C)/(100 ◦ C) (resistivity ratio), (b) the bandgap (Eg ), and (c) two transition temperature TC ↑ (transition temperature during heating) and TC ↓ (transition temperature during cooling) on TA = 510, 530, 555, 580, and 600 ◦ C at fixed oxygen pressure of 2.5 torr. Inset shows −dln (T)/dT versus oxidation annealing temperature for heating (filled red square) and cooling (filled blue circle) cycles. Dependence of (d) (30 ◦ C)/(100 ◦ C), (e) the bandgap, and (f) two transition temperature TC↑ and TC ↓ on POxidation = 1.5, 2.0, 2.5, 3.5, and 4.5 Torr at fixed oxidation annealing of 580 ◦ C. Inset shows the patterned thin film sample with indium contacts for the resistivity measurements. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)
roughness and sputtering step can improve further the uniformity of oxide films. Nonetheless, this direct thermal oxidation method for growing VO2 films provides the excellent MIT characteristics of VO2 , as will be shown below with the measurements of transport properties. Fig. 2 shows the temperature dependence of the resistivity of the VO2 films oxidized at various oxidation annealing temperatures (TA ) and oxygen pressure. All the films show the typical MIT from low temperature insulating state to high temperature metallic state with a hysteresis loop between heating and cooling cycles. First, the temperature dependence of electrical resistivity is measured as a function of TA with fixed oxygen pressure of 2.5 Torr, shown in Fig. 2(a). In Fig. 2(a), it can be clearly seen that the signature of MIT becomes more prominent as TA increases, i.e. the degrees of MIT become significant and the MIT onset temperature starts to increase, approaching to 67 ◦ C. The film grown at TA = 580 ◦ C showed the most pronounced transitions (blue line), i.e. the resistivity change at MIT is almost 4 orders of magnitude. However, the film grown at TA of 600 ◦ C shows the increase in both of high and low temperature resistivity with a slightly reduced resistivity ratio across the MIT. These modified electrical properties will be discussed later with the analysis of surface morphology. As of next step, the annealing oxidation pressure (POxidation ) is varied while TA is fixed at 580 ◦ C, shown in Fig. 2(b). Similar to the case of high annealing oxygen temperature in Fig. 2(a), the characteristics of MIT become significant only after POxidation reaches above 2.5 Torr,
i.e. sharp MIT with high changes of resistivity. However, a film grown at POxidation > 2.5 Torr shows decrease (increase) of resistivity in insulating (metallic) state. Then, based on the result of Fig. 2, the electrical properties of VO2 film at various TA and POxidation are characterized, i.e. the resistivity ratio before and after the MIT, the bandgap, and the onset temperature of MIT, shown in Fig. 3. In Fig. 3, the temperature dependence of the electrical resistivity () of the VO2 films shows strong dependence on the TA . It is believed that the resistivity ratio between the insulating and the metallic phases is most critical indicator of VO2 film quality. Fig. 3(a) shows the dependence of resistivity ratio, (30 ◦ C)/(100 ◦ C), on the oxidation annealing temperatures. The resistivity ratio stays in the range 102 –104 and increases with increasing oxidation temperature, reaching its maximum at TA = 580 ◦ C. The highest value of the resistivity ratio is 2.3 × 104 , whose value is almost comparable to that of VO2 single crystals [22]. Even for the VO2 films with mixed phases of majority VO2 and minority V6 O13 have a large resistivity ratio. This result indicates that our direct thermal oxidation method is a simple, yet very powerful technique to grow large-scale high quality VO2 films with high reproducibility. Next, assuming VO2 is an intrinsic semiconductor, the degree of bandgap is obtained using the relation; = 0 exp (Eg /2kT), where 0 , Eg , and k are the prefactor of resistivity, the bandgap, and the Boltzmann constant, respectively. Calculated bandgap for the film oxidized at TA = 580 ◦ C, based on the above relation and the resistivity curve in the insulating state of Fig. 2(a), shows the bandgap value of 0.7 eV, whose
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Fig. 4. Scanning electron microscope picture of VO2 thin film annealed at (a) 510 ◦ C, (b) 530 ◦ C, (c) 555 ◦ C, (d) 580 ◦ C, and (e) 600 ◦ C in oxygen pressure of 2.5 Torr. Atomic force microscopy image of VO2 thin film annealed at (f) 510 ◦ C, (g) 530 ◦ C, (h) 555 ◦ C, (i) 580 ◦ C, and (j) 600 ◦ C in oxygen pressure of 2.5 Torr. (k) Dependence of grain size(AG ) of the VO2 film on annealing oxidation temperature. (l) Dependence of rms surface roughness(Rq ) of the VO2 film on annealing oxidation temperature.
magnitude is the same as the value of VO2 single crystal [23] and that of taken from infrared spectroscopy [24], i.e. 0.7 eV. As shown in Fig. 3(a) and (b), it is interesting to note that the dependences of resistivity ratio and bandgap on oxidation annealing temperature are very similar, suggesting strong correlation among these quantities. Also, the onsets of MIT temperature follow the similar trend of resistivity ratio and bandgap. Fig. 3(c) shows the dependence of two transition temperature TC↑ (transition temperature during heating) and TC↓ (transition temperature during cooling) on TA . The transition temperatures during heating and cooling cycles are determined as the peak temperatures in the −dln()/dT versus temperature curve, shown in the inset of Fig. 3(c). Both of TC↑ and TC↓ increase as TA is raised up to TA = 550 ◦ C, then reach almost constant values of TC↑ = 68 ◦ C and TC↓ = 60 ◦ C for TA > 550 ◦ C. Hysteresis width (TC ), defined as the difference between TC↑ and TC↓ , was the minimum value of 7.4 ◦ C for the film oxidized at TA = 580 ◦ C. Identical analyses of Fig. 3(a)–(c) are carried out also as POxidation varies at TA of 580 ◦ C, shown in Fig. 3(d)–(f). From Fig. 3(d) and (e), the resistivity ratio and bandgap distribution indicate that oxygen pressure between 2.5 and 3.5 Torr is the critical region to grow high quality of VO2 thin film. Interesting to note is that the film with a large resistivity ratio tends to have a large bandgap, indicating a close relationship between defect states and bandgap. However, the MIT onset temperatures in Fig. 3(f) do not follow any similarity to that of Fig. 3(c), showing almost no dependence of POxidation . Overall, the comparison of Fig. 3(a)–(c) and Fig. 3(d)–(f) indicates that the crystallization of VO2 seems to strongly depend on temperature, however, POxidation will also be decisive factor for stoichiometry of the films. Finally, the surface roughness and grain size of the VO2 films are examined by AFM and SEM to locate the relation between the surface morphology and electrical properties of VO2 films. In Fig. 4, the SEM and AFM images of the VO2 films are shown only as a function
of TA . The identical measurements are taken as varying the oxygen pressure (not shown), yet it shows almost similar results as those of Fig. 4. Fig. 4(a)–(e) shows SEM of the VO2 films deposited at various TA . SEM micrograph shows granular crystallites with sizes dependent on TA . As shown in Fig. 4(k), the average grain sizes of the films are increased monotonically from ∼1 nm to 300 nm as increasing TA from 510 to 600 ◦ C. The grain sizes from SEM and AFM are in good agreements. The enhancement of grain size with increasing TA could be explained by the increased surface mobility of the growth species at the elevated temperature. Fig. 4(f)–(j) shows AFM images of VO2 films oxidized at various TA . AFM images show the strong dependence of rms roughness on the oxidation temperature. As shown in Fig. 3(j), the rms roughness of films is increased at higher oxidation annealing temperature, 5 nm at TA = 510 ◦ C and 80 nm at TA = 600 ◦ C. The rms roughness of the film oxidized at TA = 600 ◦ C is very large, whose magnitude (80 nm) is almost 30% of the VO2 film thickness (250 nm). The AFM and SEM images in Fig. 4 show that the rms roughness of the VO2 film is proportionally increased with the size of grains, suggesting strong correlation between rms roughness and grain size. Interestingly, it can be found that the main characteristics of MIT of VO2 films, i.e. resistivity ratio and sharpness of resistivity curve in Fig. 2, are directly correlated to the VO2 surface morphology of Fig. 4, i.e. the surface roughness and the grain size. In the case of insulating state of Fig. 2(a), the resistivity increases monotonically with oxygen annealing temperature. This means that defect level in the gap becomes reduced [25] and the energy gap becomes close to the value of VO2 at higher annealing temperature, which will be shown later. On the other hand, at metallic state of Fig. 2(a), the resistivity of films do not show much changes among various TA , except sample annealed at 600 ◦ C. In the case of temperature of 510 and 530 ◦ C, low resistivity of film comes mainly from the small surface roughness of Fig. 4(f) and (g). Then, as the temperature
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increase, the surface roughness starts to increase. However, at the same time, the stoichiometry of oxide become now close to that of VO2 , and the size of grain starts to become larger, which then reduces the resistivity of films at metallic state. Consequently, the resistivity becomes close to that of low temperature annealing films. This explanation is also consistent with the case of increased resistivity due to the reduced defect states at insulating state. For the films prepared at 600 ◦ C, according to the result of Fig. 4(k), the grain size is largest, which means the film should have better electrical properties, i.e. lower resistivity. That is, the film oxidized at 600 ◦ C is expected to have better MIT characteristics, such as sharp transition and high resistivity ratio, than those of the film oxidized at 580 ◦ C. However, the film oxidized at 580 ◦ C exhibited better MIT characteristics than those of 600 ◦ C. We believe this is due to both of the increased surface roughness and the growth of V6 O13 at the interface. The increase of rms roughness can easily reduce the resistivity ratio since the metallic resistivity can increase by the critical reduction in current path at very rough surface. This is in contrast to microroughness in nanoporous VO2 film, which enhances luminous transmittance and solar transmittance modulation due to decreased effective refractive index of microporous film [26]. Also, the presence of non-stoichiometric oxide, V6 O13 , can increase the resistivity of thin film, i.e. the resistivity of V6 O13 at metallic state is approximately 0.1 cm at room temperature [21]. In this letter, we present the characterization of VO2 thin films fabricated with simple direct thermal oxidation method. The properties of MIT of VO2 show strong dependence on both of annealing temperature and oxidation pressure. Furthermore, the analysis of surface morphology and electrical properties shows that large grain with smooth surface/interface roughness is required to achieve VO2 films of high performance. This report shows the evidence that the direct thermal oxidation method can be an excellent choice for growing large-scale high quality VO2 film-based devices with reproducibility. Acknowledgments H.L. Ju and B.S. Mun would like to thank the Basic Science Research Program for support through the National Research Foundation of Korea (NRF) funded by the Korean Government (MOE) (Nos. 2012R1A1A2006948 and 2012R1A1A2001745). This paper was supported by GIST 2015 TBP Research Fund and Korea Basic Science Institute Research Grant (E35800). References [1] J. Morin, Oxides which show a metal-to-insulator transition at the Neel temperature, Phys. Rev. Lett. 3 (1959) 34–36. [2] A.S. Barker, H.W. Verleur, H.J. Guggenheim, Infrared optical properties of vanadium dioxide above and below the transition temperature, Phys. Rev. Lett. 17 (1966) 1286–1289. [3] C.N. Berglund, Thermal filaments in vanadium dioxide, IEEE Trans. Electron. Device ED 16 (1969) 432–437.
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