Coexistent VO₂ (M) and VO₂ (B) Polymorphous Thin Films with Multiphase-Driven Insulator–Metal Transition

Abstract

Reversible insulator–metal transition (IMT) and structure phase change in vanadium dioxide (VO₂) remain vital and challenging with complex polymorphs. It is always essential to understand the polymorphs that coexist in desired VO₂ materials and their IMT behaviors. Different electrical properties and lattice alignments in VO₂ (M) and VO₂ (B) phases have enabled the creation of versatile functional devices. Here, we present polymorphous VO₂ thin films with coexistent VO₂ (M) and VO₂ (B) phases and phase-dependent IMT behaviors. The presence of VO₂ (B) phases may induce lattice distortions in VO₂ (M). The plane spacing of (011)_M in the VO₂ (M) phase becomes widened, and the V-V and V-O vibrations shift when more VO₂ (B) phase exists in the VO₂ (M) matrix. Significantly, the coexisting VO₂ (B) phases promote the IMT temperature of the polymorphous VO₂ thin films. We expect that such coexistent polymorphs and IMT variations would help us to understand the microstructures and IMT in the desired VO₂ materials and contribute to advanced electronic transistors and optoelectronic devices.

1. Introduction

Polarizable metal–oxygen bonds [1] and strong correlations among localized valance electrons in transition metal oxides [2,3] have sparked intensive interest, and they have shown promise in optoelectronics [4,5,6,7], sensors [8,9,10,11] and catalysis [12,13,14,15]. As one of the correlated electronic oxides, VO₂ has several polymorphs, including VO₂ (R) [16], VO₂ (B) [17] and VO₂ (M) [18]. Various crystalline symmetries and electronic structures in these polymorphs have shown versatile potential applications [19,20,21]. Wherein, the stable monoclinic VO₂ (M) (P2₁/c) phase has attracted widespread attention due to a reversible IMT near 340 K. Lattice variations are often connected with the transformation between delocalized and localized states during IMT in VO₂. Meanwhile, pure electronic behaviors have been observed due to strain-engineering and lattice-freezing during IMT processes [22,23]. Companied with IMT, structural distortion generally occurs within a first-order phase change from a low-temperature monoclinic VO₂ (M) phase to a high-temperature rutile VO₂ (R) phase [24]. In the insulating VO₂ (M) phase, adjacent vanadium atoms form zigzag chains with two types of vanadium–vanadium bond lengths along the c_R axis, and the shorter one may form a vanadium–vanadium dimer, which makes the electrons localized [25].

The polymorphous lattices and crystalline phases often affect electrical and optical behaviors in VO₂, which are closely related to the vanadium–vanadium bonds. Different from the VO₂ (M) phase, the VO₂ (B) (C2/m) phase is a Wadsley phase with a quasi-layered monoclinic structure [26] and undergoes IMT over a wide temperature range from 180 K to 300 K [27], while the vanadium–vanadium pairs induce layered alignments. It is noted that the edge-sharing VO₆ octahedra in the VO₂(B) phase aid ion diffusion and gas reduction in ion batteries and energy storage devices [28]. When VO₂ (B)-like vanadium–vanadium bonds exist in VO₂ (M) lattices, the microstructures, the electronic states and IMT behaviors in VO₂ (M) may be modified due to the variations in crystal field energy and orbital occupancies near the Fermi level [29]. Actually, it is still of significance and challenging to understand microstructural changes and IMT in the coexistent polymorphs in VO₂ towards versatile applications [30].

Herein, we described the lattice variations and IMT behaviors in VO₂ thin films with coexistent VO₂ (M) and VO₂ (B) phases. The widened plane spacing of (011)_M and varied V-V and V-O phonon vibrations were induced by the coexisting VO₂ (B) phase in the VO₂ (M) thin films, as confirmed via Raman spectra and X-ray diffraction (XRD). The IMT temperature for the VO₂ (M) thin films increased when more VO₂ (B) phases were presented in the polymorphous thin films. We expected such phase-dependent IMT behaviors to aid our understanding of polymorphous lattice variations and electronic phase changes in correlated vanadium oxides and enable the fabrication of advanced optical, electronic and even energy devices.

2. Results and Discussion

According to previous work [31], we exploited atomic layer deposition and post-annealing processes to obtain coexistent VO₂ (M) and VO₂ (B) phases in the VO₂ thin films. Figure 1a displays a cross-sectional TEM image for a ~32 nm thick VO₂ thin film. When the XRD measurement was performed on the VO₂ thin film, diffraction peaks were obvious in the XRD patterns, as shown in Figure 1b. The diffraction peaks near 2θ = 28.00° were ascribed to monoclinic VO₂(PDF #09-0142) [29], indicating the presence of a VO₂ (M) phase in the resulting thin film. In the TEM image (Figure 1c), obvious lattice fringes were also found with a d-spacing of 10d ≈ 3.17 nm, which further confirmed the existence of a VO₂ (M) phase [32]. As expected for the VO₂ (M) phase, the VO₂ (B) phase was also found in the resulting VO₂ thin film, as shown in Figure 1. XRD peaks near 2θ = 14.40° and 2θ = 29.00° in Figure 1b corresponded to VO₂ (B) (PDF #81-2392) [33], which was also verified by a d-spacing of 10d ≈ 6.14 nm for the lattice fringes in VO₂ (B) [34] in Figure 1d. All these results suggested the presence of coexistent VO₂ (M) and VO₂ (B) phases in the as-prepared VO₂ thin film.

As mentioned above, the microstructures of VO₂ (M) thin films may be changed due to the existence of the VO₂ (B) phase. Figure 2 displays the microstructure variations in VO₂ (M) thin films when the content of VO₂ (B) phases increased in the as-prepared polymorphous VO₂ thin films. Figure 2a,b displays the XRD patterns of the polymorphous VO₂ thin films with various VO₂ (B) contents, referred to as D_M (VO₂ thin film with few VO₂ (B) phases), D_M-B (VO₂ thin film with some VO₂ (B) phases) and D_B-M (VO₂ thin film with more VO₂ (B) phases). The XRD peaks around 2θ = 28.00° appeared in all D_M, D_M-B and D_B-M thin films and were assigned to the (011) plane of VO₂ (M). Notably, XRD peaks near 2θ = 14.40° and 2θ = 29.00° were obvious and found in the D_M-B and D_B-M thin films, and they were assigned to the (001) and (002) planes in VO₂ (B) (PDF #81-2392), respectively. For the D_M thin film, XRD peaks near 2θ = 14.40 ° were hardly observed, which indicated the main lattice was the VO₂ (M) phase in the D_M thin film, and few VO₂ (B) phases might have existed. We also found the intensity ratios at the XRD peaks near 2θ = 28.00° and 2θ = 29.00° were different in the case of the D_M-B and D_B-M thin films. For the D_B-M thin film, a larger ratio implied that there was more VO₂ (B) phases in the VO₂ thin films with coexistent VO₂ (M) and VO₂ (B) phases. Figure 2a suggests that the XRD patterns were different and related to the content of VO₂ (B) phases. More VO₂ (B) phases would have enhanced the intensity at about 2θ = 29.00°, while the intensity declined at about 2θ = 28.00° for VO₂ (M). The color changed in the temperature-dependent optical images for the D_M-B thin film upon heating (Figure S1 in SI), which indicated that the VO₂ (M) phase was dominant in the D_M-B thin film. In contrast, the invariant color in the optical images (Figure S1 in SI) above room temperature for the D_B-M thin film implied that the VO₂ (B) phase dominated in the D_B-M thin film, because the IMT process associated with the VO₂ (B) phase occurred between 180 K and 300 K. All measurements and analyses above suggested that the contents of the VO₂ (B) phase in the D_M, D_M-B and D_B-M thin films increased gradually.

Subsequently, we fitted the XRD curves and checked the peak variations for the (011)_M plane in the VO₂ (M) phase, as shown in Figure 2b. For the D_M, D_M-B and D_B-M thin films, the corresponding 2θ were about 28.05°, 28.00° and 27.98°, respectively. According to Bragg’s law [35], the d-spacings for the (011)_M plane in VO₂ (M) were calculated to be about 3.179 Å, 3.184 Å and 3.186 Å by using the above 2θ values for the D_M, D_M-B and D_B-M thin films, respectively. The d-spacing of (011)_M plane directly impacted the distance of the V-V lattice along the c_R axis, which was related to electron–electron and electron–phonon interactions and the IMT behaviors of VO₂ (M) [25]. The above results suggested that the VO₂ (B) phases may affect the lattice, especially the d-spacing of the (011)_M plane in VO₂ (M) in the VO₂ thin films with coexistent VO₂ (M) and VO₂ (B) phases. VO₂ (B) usually formed in an oxygen-excess environment, while VO₂ (M) formed in an oxygen-deficient environment. An appropriate amount of V could support an environment where VO₂ (B) was formed [36]. The thin films with more VO₂ (B) phases annealed at a slower heating rate, resulting in the formation of VO₂ (B) at a lower temperature and insufficient activation energy to fully convert to stable VO₂ (R), which would transform to VO₂ (M) below about 340 K [37]. X-ray photoelectron spectroscopy (XPS) measurements were used to analyze the binding energy and oxidation state related to vanadium [38] (Figure 2c and Figure S2). Figure S2 indicates the fitted peaks for varied oxidation states involving V⁺⁴ (with binding energies at about 516.2 eV and 523.4 eV) and V⁺⁵ (with binding energies at about 517.4 eV and 525.2 eV) for the V 2p core-level peaks [38]. We made calculations according to the XPS peaks and found that the valence of V⁺⁴ was dominant in the D_B-M (~57.5%), D_M-B (~59.2%) and D_M (~67.6%) thin films, while the V⁺⁵ contents in the D_B-M and D_M-B thin films were higher than that in the D_M thin film. The presence of V⁺⁵ might have been induced by surface oxidation and adsorption or the precursor of vanadium pentoxide [38]. Noticeably, the binding energy between vanadium atoms had a tendency to decrease from the D_M to D_M-B, and then to D_B-M thin films (Figure 2c). As verified in the XRD patterns, the XPS spectra suggested that the increased binding energy possibly corresponded to more VO₂ (B) existing in the polymorphous thin films. The layered lattice alignments and surface-induced V⁺⁵ components might have increased the binding energy and stabilized the VO₂ (B) phases in the D_M-B and D_B-M thin films [36,39].

Raman spectra were adopted to identify the phases and lattice distortion in the D_M, D_M-B and D_B-M thin films, as shown in Figure 3 and Figure S3. The distinguishable Raman shifts at about 137, 195, 224 and 615 cm⁻¹ corresponded to the VO₂ (M) in the D_M and D_M-B thin films [40], while the one at about 520 cm⁻¹ was from silicon substrate. For the D_B-M thin film, the Raman vibration modes at around 96, 191, 403 and 668 cm⁻¹ were obvious and assigned to VO₂ (B). The strong Raman peaks (520 cm⁻¹) from silicon substrate indicated a deeper detective depth, which implied that the vibration modes of the whole thin film could be obtained along the vertical direction. The vibration mode at about 96 cm⁻¹ corresponded to the translations of adjacent layers, and the vibration modes at around 191 and 403 cm⁻¹ were related to the bending and stretching of the V-O-V structure in VO₂ (B) [40]. Generally, the Raman shift at about 195 cm⁻¹ (referred as ω_V1) corresponded to the tilting motion of the V-V dimer, while the one at around 615 cm⁻¹ (referred as ω_V-O) was used to identify the V-O bond vibrations in VO₂ [41]. Line-scanned Raman mapping in Figure 3 further displays the variations in the phonon vibration modes in the D_M, D_M-B and D_B-M thin films. Here, about 100 statistical points were obtained with a scan length of 100 µm and a step of 1 µm during micro-Raman mapping. The thermal effect of the laser could be ignored because no obvious shift was found at 520 cm⁻¹ for the silicon substrates [42]. Figure 3a,c show the Raman details at each scan position around 194 cm⁻¹ (dashed blue box in Figure 3b) and 640 cm⁻¹ (dashed yellow box in Figure 3b) for the D_M, D_M-B and D_B-M thin films. In Figure 3a, Raman peaks shift towards low frequencies from around 198.5 cm⁻¹ in the D_M, to around 195 cm⁻¹ in the D_M-B, and then to around 191 cm⁻¹ in the D_B-M thin film. The Raman peak around 224 cm⁻¹ for the VO₂ (M) phase was observed in both the D_M and D_M-B thin films, while this peak was hardly observed in the case of the D_B-M thin film. As shown in Figure 3c, the ω_V-O peak around 668 cm⁻¹ was obvious in the D_B-M thin film, while the ones around 629 cm⁻¹ and 619 cm⁻¹ were for the D_M and D_M-B thin films, respectively. The above peaks indicated the variations in the ω_V1 and ω_V-O modes in the D_M, D_M-B and D_B-M thin films. When comparing the varied ω_V1 and ω_V-O modes in the D_M to D_M-B thin films, we found an obvious shift because of the presence of VO₂ (B), as shown in Figure 3 and Figure S3. In the D_M thin film, the ω_V-O peaks were dispersed and changed with the scanning position, while the ω_V-O peaks in the D_M-B thin film shifted and were concentrated. Meanwhile, the VO₂ (B) domains were distributed in the D_M-B film, which may have induced an abrupt shift [43]. For the D_M and D_M-B thin films, the variations in the ω_V-O modes were different, which may have been attributed to the content or grain size of the VO₂ (B) phase [43,44] in the corresponding thin films. In addition, the characteristic peaks around 668 cm⁻¹ for VO₂ (B) kept in the D_B-M thin film. In general, the phonon behavior in the D_B-M thin film was mainly attributed to the VO₂ (B) phase. All these results further verified the role of VO₂ (B) in the D_M, D_M-B and D_B-M thin films, which were consistent with those indicated during the above XRD measurements.

As shown in Figure 4 and Figure S4, variable-temperature Raman measurements were taken to detect the thermal-induced structural evolutions in the D_B-M, D_M-B and D_M thin films. Figure 4a displays the peak variations in the temperature-dependent Raman spectra for the D_B-M. The two modes at around ~191 and 400 cm⁻¹ were related to the bending and stretching vibration of the V-O-V structure in the VO₂ (B) phase. When the temperature rose from 220 K to 280 K, the mode at ~394 cm⁻¹ shifted to ~403 cm⁻¹ (Figure 4a’). The vibration mode around ~195 cm⁻¹ at 200 K shifted to ~193 cm⁻¹ at 280 K, and then to ~191 cm⁻¹ at 320 K. Meanwhile, the intensity of the Raman peak around ~195 cm⁻¹ increased obviously from 200 K to 320 K (Figure 4a). There was no significant change in Raman spectra for the D_B-M above 320 K. Due to the semimetal nature of the VO₂ (B) phase at room temperature [44], we could still observe a relatively resolvable Raman shift. In the D_M-B thin film, the ω_V1 vibration mode of ~195 cm⁻¹ at 343 K shifted to ~201 cm⁻¹ at 358 K, while the ω_V-O vibration mode changed from ~620 cm⁻¹ to ~647 cm⁻¹ due to the thermal-induced lattice distortion. Above 358 K, the ω_V1 and ω_V-O modes vanished due to the formation of metallic rutile VO₂ during IMT (Figure 4b, b’ and Figure S4b). The metallic rutile phase did not contribute to a discriminable Raman signal, and thus, a flat Raman spectrum indicated its presence. Meanwhile, in the D_M thin film, the ω_V1 vibrations suddenly changed from ~195 cm⁻¹ at 343 K to ~201 cm⁻¹ at 353 K, and the ω_V-O mode at 620 cm⁻¹ shifted to 638 cm⁻¹. In Figure 4c,c’ and Figure S4c, both two modes disappeared above 353 K, which also resulted from the formation of rutile VO₂ during IMT. Although the ω_V1 and ω_V-O modes in the D_M-B and D_M thin films were similar at room temperature, the structural evolution in the D_M-B thin film finished at a higher temperature (at 358 K) than that in the D_M thin film (at 353 K). We believed that the higher evolution temperature was attributed to the increased content of VO₂ (B) in the coexistent VO₂ (M) and VO₂ (B) thin films.

As indicated by the XRD and Raman measurements, the coexistent VO₂ (M) and VO₂ (B) phases were likely to modify the microstructure and the temperature-dependent structure evolutions. Subsequently, the temperature dependent resistance of the D_M, D_M-B and D_B-M thin films is shown in Figure 5 and Figure S5. The resistance of the D_B-M and D_M-B thin films changed and underwent a change of over 2 orders of magnitude when the temperature went from 200 K to 300 K, which was related to the IMT process of the VO₂ (B) phase [27]. However, no such decrease occurred in the D_M thin film, as displayed in Figure 5b. The following drop in the resistance of the D_M, D_B-M and D_M-B thin films at around 350 K was related to VO₂ (M) (Figure 5c) [45]. For the D_M-B thin film, the small change in resistance from 200 K to 300 K shown in Figure 5b may have resulted from the lower VO₂ (B) content and the poorer crystallinity [32]. Figure 5c further suggests the variations in resistance changed from 300 K to 400 K in the D_M, D_M-B and D_B-M thin films [46]. Notably, the increased content of VO₂ (B) would have greatly reduced the resistance variations in the as-prepared coexistent VO₂ (M) and VO₂ (B) thin films from 300 K to 400 K [29]. By taking the derivation curves and using a Gaussian fit, we obtained the transition temperature (T_IMT) above room temperature of about 345, 357 and 360 K for the D_M, D_M-B and D_B-M thin films, respectively (Figure 5d). As we obtained in the temperature-dependent Raman spectra, the increased contents of VO₂ (B) led to an increase in T_IMT [44]. In addition, the T_IMT increased with the widened plane spacing from the D_M to D_M-B and then to D_B-M thin films. These results implied that coexistent polymorphs and lattice distortion accompanied and caused such changes in the VO₂ thin films with the coexistent VO₂ (M) and VO₂ (B) phases [45].

3. Conclusions

We have demonstrated the lattice variations and phase-dependent IMT behaviors in the polymorphous thin films with coexistent VO₂ (M) and VO₂ (B). XRD and Raman measurements revealed that the increased (011)_M plane spacing and the red-shifted V-V and V-O vibrations in VO₂ (M) were triggered by the presence of more VO₂ (B) phases. Temperature-dependent electrical and Raman analyses indicated that increased VO₂ (B) phases would promote the transition temperature and reduce the strength of IMT in the polymorphous VO₂ thin films. We believe that such a study will contribute to uncovering the coexistent polymorphous phases and the underlying physics behind IMT in VO₂ towards advanced electronic and optical devices.

4. Experimental Section

The polymorphous VO₂ thin films were obtained with a combination of ALD and post-annealing procedures according to the processes previously reported [31]. The temperature-dependent optical images were obtained via a optical microscope with a home-made heating stage. Raman spectra (Nanofinder 30, Tokyo Instruments, Inc., Tokyo, Japan, the spot size of laser was about 500 nm, power was 1 mW and the grating was 1800 g/mm), micro-XRD (Bruker D8 Discover, Bruker, Karlsruhe, Germany), XRD (Bruker D8, Bruker, Karlsruhe, Germany), X-ray photoelectron spectroscopy (XPS, ESCALAB Xi+, Thermo Fischer Scientific Inc., Waltham, MA, USA) and transmission electron microscopy (TEM, JOEL JEM-2100F, JOEL Ltd., Tokyo, Japan) were used to characterize the VO₂ thin films. Electrical curves were recorded at variable temperatures from 77 K to 450 K on a cryogenic probe station (Lakeshore TTPX, Lake Shore Cryotronics, Inc., Woburn, MA, USA) with a semiconductor characterization system (Keithley 2636B, Tektronix, Inc., Beaverton, OR, USA).