Tailoring the oxygen concentration in Ge-Sb-O alloys to enable femtojoule-level phase-change memory operations

Chalcogenide phase-change materials (PCMs), in particular, the flagship Ge2Sb2Te5 (GST), are leading candidates for advanced memory applications. Yet, GST in conventional devices suffer from high power consumption, because the RESET operation requires melting of the crystalline GST phase. Recently, we have developed a conductive-bridge scheme for low-power phase-change application utilizing a self-decomposed Ge-Sb-O (GSO) alloy. In this work, we present thorough structural and electrical characterizations of GSO thin films by tailoring the concentration of oxygen in the phase-separating GSO system. We elucidate a two-step process in the as-deposited amorphous film upon the introduction of oxygen: with increasing oxygen doping level, germanium oxides form first, followed by antimony oxides. To enable the conductive-bridge switching mode for femtojoule-level RESET energy, the oxygen content should be sufficiently low to keep the antimony-rich domains easily crystallized under external electrical stimulus. Our work serves as a useful example to exploit alloy decomposition that develops heterogeneous PCMs, minimizing the active switching volume for low-power electronics.

The crystallization (SET) or amorphization (RESET) of GST proceeds via Joule heating brought by external pulses. To amorphize the crystalline state, the rigid lattice needs to be melted down first. This step is the bottleneck in terms of energy cost, because of the high melting temperature of GST (∼900 K). Device miniaturization with proper thermal insulations is typically used to reduce power consumption [33]. But conventional GST devices still require tens of picojoules to nanojoules for each RESET operation. In 2011, Xiong et al managed to deposit GST into the nanogaps between carbon nanotubes (serving as electrodes), which drastically reduced the switching volume, leading to ∼100 femtojoules per RESET operation [34]. Further optimization was made by using GST nanowires and carbon nanotubes electrodes in 2013, reaching a RESET energy of ∼80 femtojoules [35]. In 2022, Wang et al have developed a more realistic approach to reduce the active switching volume of GST for integrated devices with an ultrascaled edge-electrode using very narrow graphene nanoribbons embedded in hexagonal boron nitride multilayer, reaching ∼54 femtojoules per RESET operation [36].
In early 2022, we have designed a conductive-bridge phase-change memory (cbPCM) scheme [73] that combines the advantages of phase-change memory with filamentaryswitching in metal oxides. The essential idea is to minimize the active switching volume in conventional devices via phase separation. Under external stimulus, e.g. heating or electrical pulses, the switching medium spontaneously decomposes into a heterogeneous network with domains having different crystallization temperature, T x . As sketched in figure 1(a), a conductive path forms in the amorphous matrix (high T x ) by crystallizing only the nanobridges between different crystalline nanodomains (low T x ). In [73], we chose to incorporate oxygen into Ge-Sb alloys [74][75][76] to increase their tendency to self-decompose into Ge-enriched and Sb-enriched regions intermingled in a network, and have achieved an average RESET energy of ∼43 femtojoules per operation using a set of Ge-Sb-O (GSO) devices. The essential ingredient that makes functional GSO devices is the oxygen content in the alloy, which however requires very careful control and calibration. In this work, we monitor the degree of oxidation during the alloy film deposition process and show that the concentration of oxygen has to be maintained at a relatively low level; otherwise, inside a heavily oxidized amorphous matrix it would no longer be feasible to create a conductive path.

Results and discussion
We deposited various GSO thin films of ∼100 nm thickness using a high-purity Ge 15 Sb 85 alloy target under O 2 flow on silicon/silica substrate through direct current magnetron sputtering at room temperature, as sketched in figure 1(b). The sputtering power was set as 30 W and deposition rate was calibrated as ∼100 nm min −1 . The O 2 flow rate was adjusted from 0 to 12.5 sccm under protective Ar flow of 50 sccm. The pressure of the chamber was 0.5 Pa. A ∼5 nm ZnS-SiO 2 capping layer was also deposited to prevent further oxidation at the surface. The five GSO thin films under O 2 flow rate = 0, 2, 5, 7.5 and 12.5 sccm were then used for electrical resistance, xray diffraction (XRD) pattern, Raman spectroscopy and x-ray photoelectron spectroscopy (XPS) and energy-dispersive spectroscopy (EDS) measurements.
The temperature-dependent electrical resistance (R-T) measurements were made using the van der Pauw method. Figure 2(a) shows the change in sheet resistance of the five thin films upon heating with a heating rate of 10 • C min -1 . With moderate oxygen doping below 7.5 sccm, the GSO thin films show a similar R-T profile as compared to the undoped Ge 15 Sb 85 thin film [77][78][79] with a clear drop at T x ∼ 235 • C, indicating the onset of crystallization. Yet, a gradual increase in resistance is observed for both the initial amorphous phase and the annealed phase at 300 • C as the oxygen concentration increases. A clear deviation in the R-T profile is observed in the heavily doped GSO thin film (12.5 sccm), which shows a higher T x ∼ 260 • C and a much larger sheet resistance value at 300 • C as compared to the other four thin films.
To assess the structural properties of the thin films after heating, we carried out XRD measurements at room temperature. As displayed in figure 2(b), the undoped Ge 15 Sb 85 thin film shows major reflections at (003), (101), (006) and (202), corresponding to the rhombohedral phase of Sb-rich domains [79]. The four GSO thin films also crystallized in the rhombohedral phase with Sb-rich grains, despite the change in crystallographic orientation, showing major reflections at (012) and (110). The (012) peak is gradually broadened as the oxygen level increases, indicating increased structural disorder upon the introduction of oxygen. No sign of pure Ge nor GeO 2 crystallites was observed in the annealed GSO thin films. Figure 2(c) shows the Raman spectroscopy measurements (using an excitation laser wavelength of 532 nm) of the annealed thin films in the range of 100-250 cm −1 . In the undoped case, the peaks in 115 cm −1 and 155 cm −1 correspond to the E g and A 1g modes of rhombohedral Sb [80]. The two peaks merged gradually with the increase of oxygen alloying. To understand this behavior, we used two Lorentzian functions to fit the spectra between 100 cm −1 and 175 cm −1 and evaluate the full-width half maximum (FWHM) values and peak positions of the two modes. The sum of fitted curve is marked by red dash-dot line, the fitted E g and A 1g modes are indicated by green and blue lines, respectively. The corresponding FWHM values and peak positions are given in figure 2(d), which shows that oxygen doping induces a pronounced broadening of the E g mode with FWHM increasing from 14 cm −1 -29 cm −1 , and also a slight broadening of the A 1g mode. In addition, opposite shift in the peak position is found in the E g mode (blue shift) and the A 1g mode (red shift), leading to a progressive merging of the two peaks.
In general, crystalline phase of high uniformity in terms of bond angles and lengths results in a limited number of vibration states, and hence sharp Raman peaks at characteristic wavenumbers. In contrast, amorphous phase with disordered atomic arrangement exhibits a much wider range of bond angles and bond lengths, and the distribution of possible states leads to broader Raman peaks [81][82][83][84]. In our measurements, the enlarged FWHM values of the E g and A 1g mode indicate an incomplete crystallization with oxygen doping, despite thermal annealing at relatively high temperatures. This behavior can be attributed to the much higher crystallization temperatures of oxides, e.g. the GeO 2 glass has a T x above 700 • C [85].
To gain a better understanding of the amorphous phase, we carried out XPS measurements on the as-deposited thin films. All binding energies were corrected with respect to the C 1 s peak. Figure 3(a) shows the Ge 2p spectra with the binding energy range between 1210-1225 eV. In undoped Ge 15 Sb 85 , the peak of Ge 2p 3/2 is detected at ∼1218 eV, corresponding to the Ge homopolar bonds (Ge-Ge) [86]. Even with a relatively low doping level of oxygen (2 sccm), the Ge-Ge peak is already weakened, and a distinct peak of Ge-O at ∼1220 eV is identified. As the oxygen doping increases to 5 sccm, the Ge-Ge peak is no longer visible, leaving only the GeO 2 peak. Figure 3(b) shows the Sb 3d spectra between 525-545 eV. The Sb 3d spectrum of undoped Ge 15 Sb 85 exhibits two contributions, 3d 5/2 and 3d 3/2 , due to the spin-orbit splitting, which are located at ∼528 eV and ∼538 eV, respectively. Small Sb-O contributions start to appear at ∼531 eV and ∼540 eV in the GSO thin film obtained under 5 sccm oxygen flux, stemming from the formation of Sb 2 O 3 [87]. The two peaks are further enhanced as oxygen doping level increases.
This two-step oxidation process can be explained by the stronger tendency to form germanium oxides, as their formation energy is much more favorable than that of antimony oxides [73]. Unlike germanium, antimony is the main component of the thin film, hence the Sb-Sb peaks still remain high even in the heavily doped GSO thin film (12.5 sccm). But if the major fraction of antimony is also oxidized, no sign of crystallization can be observed below 300 • C [76]. To design readily switchable conductive bridges, we need to keep a low T x for the Sb-rich domains so that they can be easily crystallized with external thermal or electrical stimulus. Therefore, excessive oxygen content in antimony should be prevented. From the above analyses, we regard 2-5 sccm as the suitable oxygen flow rate range for sputtering Ge 15 Sb 85 , and such flow rates should be further reduced if a Sb-richer alloy target is used. With high oxygen flow rate (12.5 sccm), the oxygen concentration in these GSO thin films was estimated to exceed 35 at% by EDS measurements, while less oxygen was incorporated into the thin films in the range of 10-16 at% at 2-5 sccm. The latter is compatible with that reported in [73], ∼16 at%, which was determined by atom-probe tomography measurements.  For validation, we fabricated two GSO devices with light and heavy oxygen doping. Figure 4(a) shows the device structure. A hole with diameter of ∼250 nm was etched in a ∼100 nm thick SiO 2 dielectric layer, and the Ge 15 Sb 85 target was sputtered under oxygen flow rate of 2 and 12.5 sccm for the two devices. The devices were covered with a ∼50 nm tungsten electrode thin layer to prevent further oxidation. Then we applied DC sweeps to the two devices to pattern the heterogeneous network. As seen in the current-voltage (I-V) profiles, the heavily O-doped GSO device showed a highly resistive behavior up to 10 µA ( figure 4(b)), while the other device was switched to a low-resistance state at ∼3.25 µA, and the subsequent DC sweep confirmed the conductive nature of the latter device ( figure 4(c)). The former device is regarded as a RESET-stuck failure, due to over-oxidation. Figure 4(d) shows a typical RESET operation of GSO device with suitable oxygen doping. To guarantee a low RESET energy E RESET = I × U × ∆t = U 2 RESET × ∆t/R SET , it is desirable to have a SET state with relatively high resistance [73], e.g. 7.5 MΩ in this case. By applying a series of voltage pulses with ∆t = 50 ns (FWHM), the resistance of the device increased to more than 40 MΩ at ∼2.5 V. The RESET energy is estimated as ∼41.6 fJ, close to the average value reported in [73]. Figure 5 summarizes how to enable the cbPCM working mode via oxidation. With low-level oxygen incorporation, amorphous GSO alloys get crystallized almost completely upon SET operation. The potential presence of amorphous domains due to oxidation can hardly alter the electrical performance of GSO devices, as the electric current can pass through the crystalline matrix easily. The switching of GSO proceeds in the conventional PCM mode. Regarding the heavily oxidized case, the vast majority of germanium and antimony atoms form strong bonds with oxygen atoms. The high crystallization temperature of GeO 2 and Sb 2 O 3 makes the glass domains insensitive to SET pulses. Therefore, GSO devices get stuck in the high-resistance state, and cannot be programmed. Only with moderate-level oxygen incorporation, GSO devices can function in the cbPCM mode with robust GeO 2 -dominant glass domains and Sb-rich crystalline domains so that the nanobridges in between can be effectively switched. To achieve this goal, the formation of antimony oxides in large fractions must be avoided, and it is important to use an alloy target with relatively high germanium concentration, such as Ge 15 Sb 85 , for sputtering to guarantee the abundance of germanium oxides glass in the heterogeneous network.

Conclusions
In summary, we have reported a thorough experimental characterization of the self-decomposed GSO alloys for low-power phase-change memory application. We have tailored the oxygen concentration in the Ge 15 Sb 85 thin films and revealed a two-step oxidation process as the doping level gradually increased. We note that the relatively rich germanium concentration and moderate oxygen doping are essential to build the heterogeneous network with an intermixed GeO 2 glass domains and Sb-rich crystalline domains, allowing the formation of conductive path under external electrical stimulus. This spontaneous decomposition in the GSO alloy reduces the active switching volume sharply despite large contact areas with electrodes, as only parts of the conductive bridge need to be melted upon RESET pulsing. We note that this cbPCM scheme inevitably introduces certain randomness for device programming, as the formation of heterogeneous network differs from device to device. In addition to binary storage, heterogeneous PCMs could also be useful for the emulation of stochastic phase-change neurons [88] and for the development of physical unclonable functions [89].