Vanadium dioxide can be utilized as a Mott memory, where “0” and “1” states can be defined by insulator and metal states, respectively. In stoichiometric VO2, voltage or joule heating can trigger the transition and activate the volatile behavior. As a result, there is a constant need for such a stimulus to preserve the “1” state. If oxygen vacancies are introduced to the system while maintaining the crystal structure of the VO2 phase, the state “1” can be obtained/written permanently. That is, there is no need for external stimuli to read and recall the data. Here, we have shown the reversibility of the behavior and structure of the VO2 when oxygen vacancies are introduced to and removed from the system. The structure and relaxation mechanism are discussed, as well. This research paves the way for the nonvolatile application of VO2 in neuromorphic devices.
The oxygen vacancies are used widely to control the electrical properties of oxides since they act as charge carriers.1–5 In the electron-correlated systems like VO2, they play a more critical role in affecting the metal–insulator transition (MIT).2 Oxygen vacancies can affect charge degrees of freedom through meddling electron–electron interaction. This would in turn affect lattice, orbital, and spin degrees of freedom. That is, they disturb the lattice short-range ordering, change the valance state and, hence, electronic band structure, and affect the spin configuration. The effect of such a reconfiguration is changing the physics behind MIT. The combination of Peierls and Mott physics has shown to derive MIT.6,7 The MIT is triggered by dimerization of V–V ions which lifts the antibonding V–O (π*) orbital.8,9 The splitting of d-orbital into bonding and antibonding then creates a bandgap which is a result of V–V covalent bonding (Peierls physics) and electron–electron correlation (Mott physics).8,9 By introducing oxygen vacancies, we rearrange the structure to go through the transition with Mott physics domination.10–16
For VO2 as a Mott memory, the “0” and “1” states can be defined by insulator and metal states, respectively. The benefits of using Mott memories include but are not limited to the fact that phase transitions could be triggered at sub-nanosecond timescales17,18 and the state of devices can be accessed electrically. Theoretically, oxygen vacancies can be used as the “write” stimuli to derive the transition at any temperature below the actual MIT temperature.
This tunability provides us with the opportunity to enable nonvolatile operations and utilize VO2 in neuromorphic applications. For neuromorphic circuits, the synapses should be nonvolatile, whereas the neurons need to be volatile components.19 If the transition is triggered by voltage or temperature, the operation is volatile. This is due to the reason that, in both cases, the metallic state cannot be retained after the stimuli are removed. However, if the transition is induced by defects which act as charge carriers, the behavior is nonvolatile due to the persistence of defect presence.19
It is also worth mentioning that VO2 is a great memory candidate from Moor's law point of view since it is shown that the phenomenon of phase transition is achievable down to almost 4 nm.20 Recently, it was shown that oxygen vacancies stabilize the metallic phase at room temperature (RT) through the ionic liquid gating method as well as high-temperature vacuum annealing.21–23 The only remaining concern would be whether multiple cycles of adding and removing oxygen vacancies provide reliable and repeatable “0” and “1” states. To test this hypothesis, we investigate the metallicity of VO2 thin films ≈40 nm in the VO2/NiO/c-sapphire heterostructures. We introduce oxygen vacancies to the system to investigate the structural changes and behavior of electronic transition in response to that process. The oxygen annealing is then employed to investigate the reversibility of the metallicity. Moreover, the reversible behavior of VO2 in response to multiple vacuum stages and thereafter oxygen annealing cycles are presented. The annealing is conducted in two different temperatures since this is a diffusion-based process and has an exponential temperature dependency.
II. RESULTS AND DISCUSSION
We compare the results from as-deposited (pristine), vacuum annealed (VA), and oxygen annealed VO2/(111)NiO/(0001)Al2O3 heterostructures. The pristine samples have been grown above the critical thickness to study solely the effect of introducing defects without pinning the lattice degrees of freedom.23,24 The vacuum annealing process designed to introduce oxygen vacancies into the system. Thereafter, the oxygen annealing process introduces oxygen to fill out the vacancies. To study the crystallographic properties of vacuum annealed (low and high temperatures, respectively, represented as VA-low T and VA-high T) and oxygen annealed VO2 thin films, XRD patterns of the pristine and annealed samples are collected and the results are presented in Fig. 1(a). The inset figure represents the full range including NiO and sapphire peaks as well. For the pristine sample, (020) monoclinic M1 diffraction peak is observed, while for the vacuum annealed samples, a different diffraction peak shows up which is indicative of (200) M2 monoclinic planes of VO2.15 The monoclinic structure in the vacuum annealed samples has a different symmetry belonging to the C2/m space group and is similar to what is reported for the high-pressure monoclinic phase of VO2.25,26 The oxygen annealing is employed on vacuum annealed samples and interestingly the peak position switches back to (020) M1 monoclinic VO2 as shown in Fig. 1(a). The scanning transmission electron microscopy (STEM) was used to characterize these phases and sheds light on the atomic structure of these thin films. Figure 1(b) represents the low magnification high-angle annular dark-field (HAADF) image of the VO2/NiO/c-sapphire epitaxial heterostructures where VO2 is grown above the critical thickness. The atomically sharp interface and crystalline quality of the VO2/NiO layers are shown in Fig. 1(c). The HAADF image in Fig. 2(a) shows the cross section of the pristine-VO2/NiO-VO2 interface, and the pristine-VO2 higher magnification structure is shown in Fig. 2(b). These micrographs illustrate M1 monoclinic phase for pristine-VO2 as it is also shown schematically in the inset figure. Figure 2(c) is a cross section of the VO2/NiO-VO2 interface in the vacuum annealed sample at a high temperature (450 °C). The HAADF image in Fig. 2(d) shows an atomic resolution micrograph of the VA-high T. This structure is representative of the M2 monoclinic phase of VO2 as also is depicted in the inset figure and confirmed by XRD analysis. Figure 2(e) shows the interface of VO2/NiO-VO2 after vacuum annealing at a lower temperature (200 °C). Figure 2(f) captures the atomic structure of VO2 in the VA-low T sample which belongs to the M2 monoclinic phase (shown in the inset figure schematically).
The temperature-dependent electrical resistivity measurements for pristine and annealed samples are shown in Fig. 3. The pristine sample shows a typical hysteresis belonging to relaxed VO2 thin films.27 It is observed that after annealing at a low temperature, the metal–insulator transition amplitude decreases, and the resistivity of the insulating is lower compared to the pristine sample. We hypothesize two reasons for the observed behavior. Since, in the M2 phase of VO2, there are two types of V–V chains where only in one of them the dimerization occurs and takes part in changing conductivity, the electrical conductivity behavior is different from the M1 phase.28,29 Second, the presence of oxygen vacancies which act as charge carries is not neglectable. This can explain the lower resistivity in the insulating phase of this sample. After the transition to the metallic phase, also, the resistivity is lower compared to the pristine sample which is also justified with the presence of more charge carriers (oxygen vacancies) in the system.
After annealing at high temperature (450 °C), the metal–insulator transition vanishes. The vacancies introduced during the vacuum annealing derive the formation of a metallic M2 monoclinic phase. It is explained that the transition of tetragonal VO2 to M2 phase is a Mott type vs to M1 which is a combined Mott and Peierls transition.29 We believe that in M2, the role of Mott transition is dominant and the role of Peierls transition is not as strong as in the M1 phase. The smaller bandgap and the dominance of the Mott nature of transition in M2 make it easier to manipulate bandgap by introducing charge carriers. When the carrier concentration introduced to the system is above the Mott criterion at room temperature [ ≈ 3 × 1018 cm−3], energetically it is favorable for the Mott system to stay in the metallic state. The inset figure represents the full temperature range behavior for this stabilized phase which shows no signs of V2O3 or V2O5 transition upon the introduction of the vacancies. The base resistivity of the metallic M2 phase is higher than both the metallic state of the pristine and low-temperature vacuum annealed VO2 samples. This can be explained due to the presence of too many charge carriers in this sample which increases their collision leading to enhanced resistivity. The introduction of charge carriers enhances the conductivity as explained in the case of the low-temperature vacuum annealed sample. However, the introduction of too many oxygen vacancies creates high-density of scattering centers, which leads to lower conductivity in the metallic phase compared to the metallic phase derived with or without a lower amount of oxygen vacancies. In the low-temperature vacuum annealed sample, the thermodynamic equilibrium vacancy concentration is lower than the high-temperature vacuum annealed sample. Table I compares the carrier concentration in all three samples of pristine, low-temperature vacuum annealed, and high-temperature vacuum annealed. Interestingly, even at low temperatures, the carrier concentration meets the Mott criterion. However, since the transition in VO2 is not pure Mott and needs Peierls assistance, the higher concentration of charge carriers needed to derive the formation of a stable metallic phase at room temperature. This validates our hypothesis about the slight role of Peierls in MIT transition in the M2 phase of VO2. This phenomenon has been also confirmed by the report on VO2 thin films which were grown below the critical thickness.23 In this report, the VO2 structure was strained where a lower concentration of charge carriers could drive the MIT.
|Mobility ∼10 cm2 V−1 s−1 .||Pristine .||VA-low T .||VA-high T .|
|Carrier concentration (cm−3)@RT||∼1017||∼1020||∼1021|
|Mobility ∼10 cm2 V−1 s−1 .||Pristine .||VA-low T .||VA-high T .|
|Carrier concentration (cm−3)@RT||∼1017||∼1020||∼1021|
Figure 4 depicts the electron energy loss spectroscopy (EELS) of pristine, vacuum annealed samples at low and high temperatures, as well as literature extracted spectra of VO, VO2, V2O3, and V2O5 for comparision.30,31 The oxygen K-edge (O-K) in VO2 has the t2g component at 532 eV and the eg component of the hybridized orbitals which appears as a shoulder at 535 eV.32 Between these spectra, the VO2 and V2O5 spectra are the most similar to each other; however, there still exist differences. It is reported that the intensity of the first unoccupied state decreases as the number of vanadium 3d electron states increases by decreasing valance state.33 That is, the decrease in the intensity of the 532 peaks is recognized as the decrease of the transition probability of the O 1s electron to the first unoccupied hybridization state between O 2p and V 3d. That explains the higher intensity of the t2g component in V2O5 in comparison to VO2 owing lower valance state. Even though the upper hybridization states are not very sensitive to the local symmetry, the spectral feature at 544 eV is different in the case of VO2 and V2O5 with VO2 owning higher intensity. These distinct differences confirm the formation of VO2 in the pristine and vacuum annealed samples.
The oxygen pre-peak also provides information regarding oxygen vacancy concentration, which is following the trend of more oxygen vacancies leading to lower the valence states of vanadium ions.34 If the valance state of vanadium decreases, the intensity of the O-K pre-peak decreases, and it merges with the O-K main edge since the hybridization between V and O decreases. Thus, there is a shift toward higher eV. The shift of peak around 532 eV to higher eV for oxygen pre-peak in VA-high T represents the presence of a higher concentration of oxygen vacancies in this sample compared to VA-low T as shown by electrical measurements as well. The vacancy concentration is calculated using the EELS technique to be x ∼ 0.20 ± 0.02 in VO2−x.23 The chemical shift of the L-edge of 3d transition metal oxides is another indicator of changes in the oxidation state of metal atoms in the oxides. The shift of V 2p3/2 to lower eV has been reported as a decrease in the valence state which is not observed in the present work and further confirms the absence of V2O3 and V2O5 formation.30–32
The high-temperature vacuum annealed sample with metallic behavior is exposed to the oxygen annealing for the oxygen to get back into its lattice site. It is observed that after oxygen annealing at the same temperature, the sample returns to the M1 state and shows typical metal–insulator transition of the pristine sample. This behavior is illustrated in Fig. 5. The vacuum and oxygen annealing processes are repeated consecutively for four full cycles (eight annealings). No sign of degradation is observed in the hysteresis or metallic behavior as the heterostructures go through multiple cycles of annealing. The transition temperature for the hysteresis behavior, however, is observed to change to lower temperature and eventually comes back to the known transition temperature for the relaxed bulk VO2. The yellow and blue lines in Fig. 5 represent the transition temperature for the VO2 in the first and fourth cycles which is ∼342 K and ∼341 K, respectively. The transition temperatures are reported as the mean between cooling and heating transition temperatures. In the second cycle, the VO2 transition temperature has decreased to ∼336 K as the blue arrow is pointing toward it, and in the third cycle, the VO2 transition temperature has increased back toward the equilibrium temperature and equals ∼339 K. In the fourth cycle, the transition temperature is identical to the bulk strain-free structure of VO2 (341 K), which shows its structure has reached to equilibrium.
In the pristine sample, even though the transition temperature is so close to bulk behavior, there exist defects and misfit dislocations. Upon annealing at high temperatures, the system attempts to get rid of defects employing the diffusion process. We have, also, found out another relaxation process which has taken place at the interface of VO2/NiO and requires higher energy than what RT provides. The system is demanding to minimize the stress field around misfit dislocations by the formation of stacking faults at the interface as it is captured and illustrated in the HAADF image in Fig. 6. The stacking fault involves two atomic layers, and the presence of steps has provided a nucleation site for stacking fault's formation. This shows that high-temperature annealing provides enough energy to overcome the formation energy of the stacking fault and relax the system further to achieve the bulk transition temperature in the fourth cycle. After the first annealing cycle, the presence of stacking faults in addition to misfit dislocations create stress fields which could cause overall compressive strain. This would lead to a decrease in the transition temperature in VO2.35 But eventually, after multiple cycles, the dislocation movement using the diffusion process returns the system back to the relaxed state to minimize the total energy.
We successfully controlled the transition in VO2 thin films by introducing oxygen vacancies. The oxygen vacancies at low temperature annealing derive the M2 phase of VO2 which undergoes metal-to-insulator transition with dominant Mott nature. The introduction of oxygen vacancies at high temperatures and concentrations is shown to create metallic behavior in M2–VO2. This behavior is reversible through oxygen annealing. After oxygen annealing, the structure goes back to the M1 monoclinic phase of VO2. The high-temperature film relaxation mechanisms are also activated throughout these annealings such as the formation of stacking faults at the interface which affects the transition temperature. The structure eventually reached the relaxed state after three cycles.
This on and off hysteresis behavior generated by oxygen vacancies can be exploited in nonvolatile memory devices for neuromorphic applications, where there is no need for an external energy source to retrieve information (metallic and insulting behaviors). Considering the incredibly low power (10–15 W) of the human brain, energy efficiency is critical for neuromorphic computing.
A. Experimental details
Thin-film heterostructures were deposited using a KrF excimer laser with a frequency of 5 Hz and an energy of 2–2.5 J/cm2.36–38 The NiO layer was grown at 700 °C and 1 × 10−4 Torr oxygen pressure, and the VO2 film was deposited above the critical thickness at 550 °C and 1.2 × 10−2 Torr oxygen pressure.2,39 The vacuum annealing procedure was performed at 1 × 10−7 Torr in two different temperatures of 200 °C and 450 °C for 2 h. The reverse process of oxygen annealing was done at 1.2 × 10−2 Torr in two different temperatures of 200 °C and 450 °C for 2 h.
1. In situ x-ray diffraction measurement
A PANalytical Empyrean diffractometer using Cu-Kα radiation was used to collect in situ XRD measurements. The in situ structural changes during heating and cooling cycles were monitored using this technique in the 25–120 °C range. The step size of 2θ was set to 0.013° and the count time per step was set to 1 s.
2. Scanning transmission electron microscopy
The HAADF images were collected using an FEI Titan 80-300 probe aberration-corrected scanning transmission electron microscope (STEM) operated at 200 kV. Across thin-film interfaces, electron energy loss spectra (EELS) were also acquired at 200 keV with a 28 mrad collection angle and 19.6 mrad convergence angle. Spectra were collected using a 0.25 eV ch−1 dispersion with a ∼35 pA probe current and 0.07 nm pixel size to perform an elemental analysis. A routine peak fitting was applied on V-L and O-K edges using Gaussian profiles for detailed analysis.40 The noise reduction in the data is performed by fast Fourier filtering of the data, following the dark reference.
3. Electrical measurement
The physical property measurement system (PPMS) by Quantum Design was used to acquire electrical transport measurements over the temperature range of −173 to 110 °C during heating and cooling cycles. The resistivity data were collected using a step size of 0.1 at zero magnetic fields. The carrier concentration of the metallic phases are produced by Hall measurements, and for the insulating phase, it is calculated from R vs T measurements considering the mobility of the metallic state, assuming across the transition the mobility does not change considerably.41
The manuscript was written by A.M. EELS spectroscopy measurements were done by R.S. All authors have contributed to the final version of the manuscript.
This project was supported by the National Science Foundation (NSF) (Grant No. DMR-1304607). The authors acknowledge the use of the Analytical Instrumentation Facility (AIF) at the North Carolina State University, which is supported by the State of North Carolina and NSF. R.S. also acknowledges the support of faculty start-up funding at the Oklahoma State University.
The data that support the findings of this study are available within the article.