# Improvement of the Parameter Distribution of TiN/HfO2/CeOx/TiN Memristors by Rapid Thermal Annealing

Authors: Guihua Yao; Zhaozhu Qu; Changfang Li; Yuntao Peng; Qixin Li; Zhaohui Zeng; Jianghui Dong; Baolin Zhang
Journal: Journal of Nanomaterials
 (2022)

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## Abstract

TiN/HfO2/CeOx/TiN memristors were prepared by magnetron sputtering. To further improve their performance, the devices were rapidly thermally annealed at different temperatures for different times. Compared with those of unannealed devices, the coefficients of variation (CVs) of the set voltage (VSET) and the reset voltage (VRESET) were reduced by 35.1% and 59.4%, respectively, and the CVs of the resistances in low and high resistance states (RLRS and RHRS) were reduced by 70.2% and 52.7%, respectively, after annealing at 400°C for 2 min in air. Through X-ray diffraction, X-ray photoelectron spectroscopy, and I‐V curves of the devices before and after annealing, we propose that the combined effect of grain growth (i.e., grain boundary reduction) and decreased oxygen vacancy content in the switching film resulting from annealing is responsible for the improvement in the switching parameter distribution of TiN/HfO2/CeOx/TiN devices. This work presents a simple way to enhance the performance of memristors.


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## Body

## 1. Introduction

Memristors have great potential as next-generation memory devices because of their high scalability, simple structure, rapid resistance switching and compatibility with CMOS technology [1–4]. In general, the resistance switching mechanism of memristors relies on the formation/disruption of conductive filaments (CFs) resulting from oxygen vacancy or oxygen ion movements in oxide layers, and on the interface effects between electrodes and oxides [5–7]. Although the formation of CFs after activation by electroforming and resistance switching has been observed [8] and quantitative information (diameter, composition, and gap width) of CFs has been investigated [9], the formation/disruption of CFs in the oxide layer still needs further elucidation and adjustment. Oxygen vacancies can reduce the bandgap and enhance the conductivity of oxide materials, leading to the formation of CFs [10, 11]. The intrinsic conductivity of oxide materials leads to different electrical activation modes, such as field-induced electroforming and current bifurcation-induced electroforming [12], which affect the switching characteristics of memristors. At the same time, because of the variation in or nonuniform structure of resistive switching materials [13, 14], the resistances of memristor devices in high/low resistance (RHRS/RLRS) states vary greatly, which restricts their further development and application. The fundamental causes of variations in memristors parameters and the methods to improve these properties need to be further studied [15].Ceria has been studied as a switching material for memristors [1, 2, 16–18]. Its excellent oxygen storage/release capacity is because the two most common cation valence states, Ce4+ and Ce3+, are prone to redox reactions [19–21]. However, memristors based on CeOx still have severe problems, such as high switching parameter dispersion and high operating current. Annealing can decrease the number of defects and improve the crystal quality of films. It has been reported that proper annealing can improve the switching endurance, on/off ratio, data retention and uniformity in the operational parameters of memristors [16, 22]. However, there is still no insight into the reasons why annealing improves the resistance switching properties. In this study, TiN/HfO2/CeOx/TiN memristors were prepared. A HfO2 layer was used to increase the initial resistances of CeOx-based memristor devices and to modulate the separation and recombination of oxygen ions/vacancies during resistance switching [7]. The effects of annealing the devices at different temperatures and times on the characteristics of resistive switching were studied. The stability of the switching parameters of TiN/HfO2/CeOx/TiN was obviously improved after annealing at 400°C for 2 min. We propose that the combined effect of grain growth and decrease in oxygen vacancy content accounts for this improvement. This work shows that annealing can be a general and effective way to reduce the dispersion of memristor switching parameters.


## 2. Experiment

### 2.1. Materials

A TiN target was used to deposit a 100 nm thick TiN bottom electrode (BE) on SiO2 (300 nm)/Si (450 μm) substrates by direct current (DC) sputtering with a sputtering time of 2000 s, a sputtering atmosphere of 10% N2, and a sputtering power of 150 W. Next, a CeO2 target was used to deposit a 30 nm thickness CeOx layer by radio frequency (RF) sputtering for 4000 s in an atmosphere with an Ar : O2flowrateratio=11.2sccm:5.6sccm with a sputtering power of 80 W. Then, a HfO2 target was used to deposit a 10 nm thick HfO2 layer by RF sputtering (100 W) for 300 s in 16.8 sccm Ar. Finally, TiN with a thickness of 100 nm and diameter of 50 μm was deposited as the top electrode (TE) by DC sputtering and was patterned by ultraviolet lithography to expose the AZ4620 photoresist-coated wafer. TiN/HfO2/CeOx/TiN devices were obtained. Some of the TiN/HfO2/CeOx/TiN were rapidly thermally annealed in air at different temperatures (200 to 400°C) for different times (2 to 20 min).


### 2.2. Characterization

Scanning electron microscopy (SEM, Hitachi S-4800), X-ray diffraction (XRD, SmartLab Mini flex600) and X-ray photoelectron spectroscopy (XPS, ESCALAB 250Xi) were utilized to investigate the cross-section morphology, chemical composition, and crystal structure ofCeOx and HfO2. For XRD, a 0.5° grazing angle, 1 divergence slit, and 20°–80° scanning angles were adopted. For XPS, Avantage software was used to analyze the data, and the binding energy scale was calibrated with the C 1s peak at 284.8 eV.A Keithley 4200-SCS analyzer was used to measure the electrical properties of all devices. Bias voltages were applied to the TiN TE, the TiN BE was grounded during electrical property measurements, and the sweep rate was 0.01 V/step. All tests for electrical property determination were performed at room temperature.


## 2.1. Materials

A TiN target was used to deposit a 100 nm thick TiN bottom electrode (BE) on SiO2 (300 nm)/Si (450 μm) substrates by direct current (DC) sputtering with a sputtering time of 2000 s, a sputtering atmosphere of 10% N2, and a sputtering power of 150 W. Next, a CeO2 target was used to deposit a 30 nm thickness CeOx layer by radio frequency (RF) sputtering for 4000 s in an atmosphere with an Ar : O2flowrateratio=11.2sccm:5.6sccm with a sputtering power of 80 W. Then, a HfO2 target was used to deposit a 10 nm thick HfO2 layer by RF sputtering (100 W) for 300 s in 16.8 sccm Ar. Finally, TiN with a thickness of 100 nm and diameter of 50 μm was deposited as the top electrode (TE) by DC sputtering and was patterned by ultraviolet lithography to expose the AZ4620 photoresist-coated wafer. TiN/HfO2/CeOx/TiN devices were obtained. Some of the TiN/HfO2/CeOx/TiN were rapidly thermally annealed in air at different temperatures (200 to 400°C) for different times (2 to 20 min).


## 2.2. Characterization

Scanning electron microscopy (SEM, Hitachi S-4800), X-ray diffraction (XRD, SmartLab Mini flex600) and X-ray photoelectron spectroscopy (XPS, ESCALAB 250Xi) were utilized to investigate the cross-section morphology, chemical composition, and crystal structure ofCeOx and HfO2. For XRD, a 0.5° grazing angle, 1 divergence slit, and 20°–80° scanning angles were adopted. For XPS, Avantage software was used to analyze the data, and the binding energy scale was calibrated with the C 1s peak at 284.8 eV.A Keithley 4200-SCS analyzer was used to measure the electrical properties of all devices. Bias voltages were applied to the TiN TE, the TiN BE was grounded during electrical property measurements, and the sweep rate was 0.01 V/step. All tests for electrical property determination were performed at room temperature.


## 3. Results

### 3.1. SEM Analyses of TiN/HfO2/CeOx/TiN

Figure1(a) shows the structural diagrams of TiN/HfO2/CeOx/TiN. Figure 1(b) shows the morphology of the cross-section of the devices. The layers of the different compounds with uniform thicknesses are clearly shown. The thickness of each layer realized by the magnetron sputtering parameters is in accordance with the design.Figure 1
(a) Schematic diagram of TiN/HfO2/CeOx/TiN. (b) SEM image of the cross section of TiN/HfO2/CeOx/TiN.
(a)(b)


### 3.2.I‐V Curves of Annealed and Unannealed TiN/HfO2/CeOx/TiN Devices

The forming voltages of TiN/HfO2/CeOx/TiN before annealing were 8.83±0.70V (the average results of 16 devices) with a compliance current of 1 mA. After forming the devices were switched from the LRS (low resistance state) to the HRS (high resistance state) and vice versa with voltage sweeping from 0 to ±2 V (Figure 2(a)). The forming voltages of TiN/HfO2/CeOx/TiN after annealing at 400°C for 2 min were 8.63±0.70V (the average results of 16 devices) with a compliance current of 1 mA. After forming the devices were switched from the LRS to the HRS and vice versa with voltage sweeping from 0 to ±2 V (Figure 2(b)). The cumulative distributions of the VSET (set voltage), VRESET (reset voltage), RLRS, and RHRS of TiN/HfO2/CeOx/TiN devices before and after annealing at 400°C for 2 min are shown in Figures 2(c) and 2(d). Table 1 shows that the dispersions of the VSET, VRESET, RLRS, and RHRS of the annealed devices were significantly less than those of the devices before annealing. The coefficients of variation (CV=standarddeviation/meanvalue) of VSET decreased from 15.68% to 10.18% (a decrease of 35.1%), and the CVs of VRESET decreased from 4.56% to 1.85% (a decrease of 59.4%). For TiN/HfO2/CeOx/TiN, the CVs of RLRS decreased from 36.95% before annealing to 10.99% after annealing (a decrease of 70.2%), and the CVs of RHRS decreased from 47.56% before annealing to 22.48% after annealing (a decrease of 52.7%).Figure 2
I‐V curves of unannealed TiN/HfO2/CeOx/TiN devices (a) and annealed TiN/HfO2/CeOx/TiN devices (b) with voltage sweeping from 1⟶2 and 3⟶4 and cumulative probabilities of the VSET and VRESET (c) and RHRS and RLRS (d) of TiN/HfO2/CeOx/TiN, after the forming process (the insets show the forming process).
(a)(b)(c)(d)Table 1
Mean values, standard deviation, and CV∗ (in the brackets) of the VSET, VRESET, RLRS, and RHRS of TiN/HfO2/CeOx/TiN devices before and after annealing at 400°C for 2 min.
VSET (V) [CV]VRESET (V) [CV]RLRS (Ω) [CV]RHRS (Ω) [CV]TiN/HfO2/CeOx/TiN (unannealed)1.19±0.18 [15.68%]−1.82±0.08 [4.56%]129.56±47.87 [36.95%]1100.04±523.19 [47.56%]TiN/HfO2/CeOx/TiN(400°C-2 min)1.05±0.11 [10.18%]−1.82±0.03 [1.85%]93.57±10.28 [10.99%]521.70±117.3 [22.48%]∗CVcoefficientofvariation=standarddeviation/meanvalue.The current-voltage (I‐V) curves of TiN/HfO2/CeOx/TiN annealed at different annealing temperatures and times with voltage sweeping from 0 V to ±5 V to 0 V (forming was not applied) are shown in Figure 3. The conductance of the devices increased with increasing annealing temperature and time. However, with increasing annealing time, the relative conductance of the annealed devices decreased (but was still higher than that of the unannealed devices). The conductance changes may be caused by the combined effect of the oxygen vacancy concentration and microstructure of the films modified by annealing, and we will clarify this in the discussion section.Figure 3
I‐V curves of TiN/HfO2/CeOx/TiN devices with voltage sweeping from 0 V to ±5 V to 0 V (forming was not applied) at different annealing temperatures and times.


### 3.3. XRD and Grain Size Analyses of Unannealed and Annealed TiN/HfO2/CeOx/TiN

Figure4 shows the XRD patterns of TiN/HfO2/CeOx/TiN and the change in grain sizes of CeOx thin films at different annealing temperatures and times. From Figure 4(a), the (103) and (121) diffraction peaks show the HfO2 orthorhombic crystal structure (JCPDS#40-1173), and (111) the diffraction peaks show the CeOx cubic crystal structure (JCPDS#34-0394). With increasing annealing temperature and time (Figures 4(a) and 4(b)), the CeO2 (111) diffraction peak intensity gradually increases, and the diffraction peaks of CeO2 (220) and CeO2 (311) appear and gradually grow. All the diffraction peak intensities of CeOx gradually increased, and the full width at half maximum (FWHM) of the peaks decreased. However, the diffraction peaks of CeOx annealed for 5 min, 10 min, and 20 min did not change obviously.Figure 4
Small-angle grazing incidence XRD patterns of TiN/HfO2/CeOx/TiN at different annealing temperatures (annealing time: 2 min) (a) and with different annealing times (annealing temperature: 400°C) (b). CeOx grain size changes under different annealing conditions (c).
(a)(b)(c)According to the Scherrer formula [23], D=kγ/Bcosθ (B is the FWHM, D is the grain size, K is the Scherrer constant, θ is the Bragg diffraction angle, and γ is the X-ray wavelength), we calculated the grain sizes from the FWHM of CeO2 (111) peaks for the CeOx films, as shown in Figure 4(c). The CeOx grain sizes increased with increasing annealing temperature. However, the CeOx grain sizes almost did not change with increasing annealing time. No obvious grain growth of HfO2 was observed from the calculation according to the Scherrer formula. From these results, we speculated that the decrease in the RLRS and RHRS dispersions of TiN/HfO2/CeOx/TiN devices after annealing (400°C, 2 min) may be related to the increase in CeOx grain sizes.


### 3.4. Oxygen Vacancy Concentrations in TiN/HfO2/CeOx/TiN Analyzed by XPS

The analysis depth of XPS analyses was 2–5 nm, and the region that was not covered with the top TiN electrode was focused on by an electron beam to analyze the O 1s signals of the HfO2 films under different annealing conditions. Figures 5(a) and 5(b) show the XPS spectra of the Hf 4f and Ce 3d peaks of TiN/HfO2/CeOx/TiN (unannealed). The binding energies of 16.58 eV (Hf 4f7) and 18.18 eV (Hf 4f5) correspond to the Hf4+ oxidation state. The binding energies of 882.58 eV, 889.08 eV, 898.48 eV, 901.08 eV, 907.08 eV, and 916.48 eV correspond to the Ce4+ oxidation state. The binding energies of 885.68 eV and 903.68 eV correspond to the Ce3+ oxidation state. The O 1s XPS peak at 529.98 eV (Figure 5(c)) of the HfO2 films under different annealing conditions is ascribed to lattice oxygen in HfO2, and the O 1s peak at 531.88 eV is ascribed to nonlattice oxygen, as reported in the literature [24, 25]. The oxygen vacancy concentration is proportional to the ratio of the nonlattice oxygen concentration [26], so nonlattice oxygen concentration can be used to qualitatively analyze the oxygen vacancy concentration. The oxygen vacancy concentration percentages are calculated by the areas below the peaks, which are adjusted by sensitivity factors [27]. Figures 5(d) and 5(e) show the oxygen vacancy concentrations in HfO2 films annealed in air at different temperatures and times. With increasing annealing temperature and time, the oxygen vacancy concentration in the HfO2 films decreased.Figure 5
XPS spectra showing the Hf 4f peaks (a) (the inset shows the C 1s peak at 284.8 eV for calibration) and Ce 3d peaks (b) of TiN/HfO2/CeOx/TiN (unannealed). XPS spectra showing the O 1s peaks for the HfO2 films of TiN/HfO2/CeOx/TiN under different annealing conditions (c). Change in the oxygen vacancy concentration of the HfO2 films after annealing at 400°C for different times (d) and after annealing at different temperatures for 2 min (e).
(a)(b)(c)(d)(e)Annealing in air can also remarkably reduce the oxygen vacancy concentration inCeOx, which has been proven in our previous work [28]. After annealing in an air atmosphere, the Ce3+in CeOx lost an electron and was oxidized to Ce4+, which was accompanied by the elimination of oxygen vacancies. With increasing annealing temperature and time, the oxygen vacancy concentration in CeOx films decreased.


## 3.1. SEM Analyses of TiN/HfO2/CeOx/TiN

Figure1(a) shows the structural diagrams of TiN/HfO2/CeOx/TiN. Figure 1(b) shows the morphology of the cross-section of the devices. The layers of the different compounds with uniform thicknesses are clearly shown. The thickness of each layer realized by the magnetron sputtering parameters is in accordance with the design.Figure 1
(a) Schematic diagram of TiN/HfO2/CeOx/TiN. (b) SEM image of the cross section of TiN/HfO2/CeOx/TiN.
(a)(b)


## 3.2.I‐V Curves of Annealed and Unannealed TiN/HfO2/CeOx/TiN Devices

The forming voltages of TiN/HfO2/CeOx/TiN before annealing were 8.83±0.70V (the average results of 16 devices) with a compliance current of 1 mA. After forming the devices were switched from the LRS (low resistance state) to the HRS (high resistance state) and vice versa with voltage sweeping from 0 to ±2 V (Figure 2(a)). The forming voltages of TiN/HfO2/CeOx/TiN after annealing at 400°C for 2 min were 8.63±0.70V (the average results of 16 devices) with a compliance current of 1 mA. After forming the devices were switched from the LRS to the HRS and vice versa with voltage sweeping from 0 to ±2 V (Figure 2(b)). The cumulative distributions of the VSET (set voltage), VRESET (reset voltage), RLRS, and RHRS of TiN/HfO2/CeOx/TiN devices before and after annealing at 400°C for 2 min are shown in Figures 2(c) and 2(d). Table 1 shows that the dispersions of the VSET, VRESET, RLRS, and RHRS of the annealed devices were significantly less than those of the devices before annealing. The coefficients of variation (CV=standarddeviation/meanvalue) of VSET decreased from 15.68% to 10.18% (a decrease of 35.1%), and the CVs of VRESET decreased from 4.56% to 1.85% (a decrease of 59.4%). For TiN/HfO2/CeOx/TiN, the CVs of RLRS decreased from 36.95% before annealing to 10.99% after annealing (a decrease of 70.2%), and the CVs of RHRS decreased from 47.56% before annealing to 22.48% after annealing (a decrease of 52.7%).Figure 2
I‐V curves of unannealed TiN/HfO2/CeOx/TiN devices (a) and annealed TiN/HfO2/CeOx/TiN devices (b) with voltage sweeping from 1⟶2 and 3⟶4 and cumulative probabilities of the VSET and VRESET (c) and RHRS and RLRS (d) of TiN/HfO2/CeOx/TiN, after the forming process (the insets show the forming process).
(a)(b)(c)(d)Table 1
Mean values, standard deviation, and CV∗ (in the brackets) of the VSET, VRESET, RLRS, and RHRS of TiN/HfO2/CeOx/TiN devices before and after annealing at 400°C for 2 min.
VSET (V) [CV]VRESET (V) [CV]RLRS (Ω) [CV]RHRS (Ω) [CV]TiN/HfO2/CeOx/TiN (unannealed)1.19±0.18 [15.68%]−1.82±0.08 [4.56%]129.56±47.87 [36.95%]1100.04±523.19 [47.56%]TiN/HfO2/CeOx/TiN(400°C-2 min)1.05±0.11 [10.18%]−1.82±0.03 [1.85%]93.57±10.28 [10.99%]521.70±117.3 [22.48%]∗CVcoefficientofvariation=standarddeviation/meanvalue.The current-voltage (I‐V) curves of TiN/HfO2/CeOx/TiN annealed at different annealing temperatures and times with voltage sweeping from 0 V to ±5 V to 0 V (forming was not applied) are shown in Figure 3. The conductance of the devices increased with increasing annealing temperature and time. However, with increasing annealing time, the relative conductance of the annealed devices decreased (but was still higher than that of the unannealed devices). The conductance changes may be caused by the combined effect of the oxygen vacancy concentration and microstructure of the films modified by annealing, and we will clarify this in the discussion section.Figure 3
I‐V curves of TiN/HfO2/CeOx/TiN devices with voltage sweeping from 0 V to ±5 V to 0 V (forming was not applied) at different annealing temperatures and times.


## 3.3. XRD and Grain Size Analyses of Unannealed and Annealed TiN/HfO2/CeOx/TiN

Figure4 shows the XRD patterns of TiN/HfO2/CeOx/TiN and the change in grain sizes of CeOx thin films at different annealing temperatures and times. From Figure 4(a), the (103) and (121) diffraction peaks show the HfO2 orthorhombic crystal structure (JCPDS#40-1173), and (111) the diffraction peaks show the CeOx cubic crystal structure (JCPDS#34-0394). With increasing annealing temperature and time (Figures 4(a) and 4(b)), the CeO2 (111) diffraction peak intensity gradually increases, and the diffraction peaks of CeO2 (220) and CeO2 (311) appear and gradually grow. All the diffraction peak intensities of CeOx gradually increased, and the full width at half maximum (FWHM) of the peaks decreased. However, the diffraction peaks of CeOx annealed for 5 min, 10 min, and 20 min did not change obviously.Figure 4
Small-angle grazing incidence XRD patterns of TiN/HfO2/CeOx/TiN at different annealing temperatures (annealing time: 2 min) (a) and with different annealing times (annealing temperature: 400°C) (b). CeOx grain size changes under different annealing conditions (c).
(a)(b)(c)According to the Scherrer formula [23], D=kγ/Bcosθ (B is the FWHM, D is the grain size, K is the Scherrer constant, θ is the Bragg diffraction angle, and γ is the X-ray wavelength), we calculated the grain sizes from the FWHM of CeO2 (111) peaks for the CeOx films, as shown in Figure 4(c). The CeOx grain sizes increased with increasing annealing temperature. However, the CeOx grain sizes almost did not change with increasing annealing time. No obvious grain growth of HfO2 was observed from the calculation according to the Scherrer formula. From these results, we speculated that the decrease in the RLRS and RHRS dispersions of TiN/HfO2/CeOx/TiN devices after annealing (400°C, 2 min) may be related to the increase in CeOx grain sizes.


## 3.4. Oxygen Vacancy Concentrations in TiN/HfO2/CeOx/TiN Analyzed by XPS

The analysis depth of XPS analyses was 2–5 nm, and the region that was not covered with the top TiN electrode was focused on by an electron beam to analyze the O 1s signals of the HfO2 films under different annealing conditions. Figures 5(a) and 5(b) show the XPS spectra of the Hf 4f and Ce 3d peaks of TiN/HfO2/CeOx/TiN (unannealed). The binding energies of 16.58 eV (Hf 4f7) and 18.18 eV (Hf 4f5) correspond to the Hf4+ oxidation state. The binding energies of 882.58 eV, 889.08 eV, 898.48 eV, 901.08 eV, 907.08 eV, and 916.48 eV correspond to the Ce4+ oxidation state. The binding energies of 885.68 eV and 903.68 eV correspond to the Ce3+ oxidation state. The O 1s XPS peak at 529.98 eV (Figure 5(c)) of the HfO2 films under different annealing conditions is ascribed to lattice oxygen in HfO2, and the O 1s peak at 531.88 eV is ascribed to nonlattice oxygen, as reported in the literature [24, 25]. The oxygen vacancy concentration is proportional to the ratio of the nonlattice oxygen concentration [26], so nonlattice oxygen concentration can be used to qualitatively analyze the oxygen vacancy concentration. The oxygen vacancy concentration percentages are calculated by the areas below the peaks, which are adjusted by sensitivity factors [27]. Figures 5(d) and 5(e) show the oxygen vacancy concentrations in HfO2 films annealed in air at different temperatures and times. With increasing annealing temperature and time, the oxygen vacancy concentration in the HfO2 films decreased.Figure 5
XPS spectra showing the Hf 4f peaks (a) (the inset shows the C 1s peak at 284.8 eV for calibration) and Ce 3d peaks (b) of TiN/HfO2/CeOx/TiN (unannealed). XPS spectra showing the O 1s peaks for the HfO2 films of TiN/HfO2/CeOx/TiN under different annealing conditions (c). Change in the oxygen vacancy concentration of the HfO2 films after annealing at 400°C for different times (d) and after annealing at different temperatures for 2 min (e).
(a)(b)(c)(d)(e)Annealing in air can also remarkably reduce the oxygen vacancy concentration inCeOx, which has been proven in our previous work [28]. After annealing in an air atmosphere, the Ce3+in CeOx lost an electron and was oxidized to Ce4+, which was accompanied by the elimination of oxygen vacancies. With increasing annealing temperature and time, the oxygen vacancy concentration in CeOx films decreased.


## 4. Discussion

### 4.1. The Conduction Mechanism Fitting of TiN/HfO2/CeOx/TiN

To elucidate the resistive switching behavior of TiN/HfO2/CeOx/TiN, the I‐V curves of the devices were fitted with conductive mechanisms (Figures 6(a) and 6(b)). When the unannealed and annealed devices were in the HRS (the lower curves in Figures 6(a) and 6(b)), the I‐V curves possess linear parts with a slope of ~1, which indicates Ohmic conduction (I∝V) [29] due to the partial filling of traps during the weak injection of charge carriers from the bottom/top TiN electrode; the I‐V curves also possess rapid increasing current parts with a slope of ~2, which indicates Child’s square law (I∝V2) [30] due to the injected electrons with applied bias which are trapped/detrapped in the CeOx layer. These results suggest that the devices mainly followed the SCLC conduction mechanism in the HRS. When the unannealed and annealed devices are in the LRS (the upper parts of the curve in Figures 6(a) and 6(b)), the I‐V curves are linear with slopes of ~1, implying that the conduction filament is the main conduction mechanism for the devices at the LRS. These results imply that the conduction mechanisms of TiN/HfO2/CeOx/TiN in the HRS and LRS did not changed after annealing.Figure 6
Log-logI‐V curves and their fit for determining the conduction mechanism of the unannealed TiN/HfO2/CeOx/TiN devices (a) and TiN/HfO2/CeOx/TiN devices annealed at 400°C for 2 min (b).
(a)(b)


### 4.2. Effect of Annealing on the Resistance Switching Characteristics of TiN/HfO2/CeOx/TiN

The conduction ofCeOx switching materials is composed of ionic conduction and electronic conduction [31]; both ionic conduction and electronic conduction are influenced by grain size (and, therefore, grain boundaries), and oxygen vacancies and their distribution. The preferential accumulation of oxygen vacancies along the grain boundary sites has been reported [32]. It was suggested that although there were oxygen vacancies in the grains, the oxygen vacancy concentration near the grain boundaries was generally higher than that in the inner grains, and furthermore, the oxygen vacancies in the grains were generally in linear but discontinuous arrangements, making it difficult to form a conduction channel in the grains. On the other hand, it was more likely that oxygen vacancies formed chains at the grain boundaries to construct the ionic conduction channel, and electrons also tended to accumulate in the vicinity of the grain boundaries and hop along the oxygen vacancies to contribute to the conductivity [33].In this work, annealing significantly reduced the dispersion of theVSET and VRESET and RLRS and RHRS of TiN/HfO2/CeOx/TiN devices (Table 1, Figures 2(c) and 2(d)), which could be ascribed to the combined effect of the increased grain size in the CeOx films (no obvious grain size change was observed in HfO2) (Figure 4(c)) and the reduced oxygen vacancy concentration in the HfO2 and CeOx films caused by annealing (Figure 5). Figure 7 clarifies the possible conduction paths in the CeOx of TiN/HfO2/CeOx/TiN. It can be speculated that the oxygen vacancies accumulate along CeOx grain boundaries, and that the oxygen vacancy channel forms and ruptures along the grain boundaries. For the unannealed devices, the channel forms and ruptures at more abundant boundaries (Figure 7(a)), while for the annealed devices the channel forms and ruptures at fewer positions because of the reduction in the number of grain boundaries (Figure 7(b)), and smaller variations in the conduction paths result in decreased dispersions of the VSET and VRESET and RLRS and RHRS of the annealed devices.Figure 7
Possible conduction paths in theCeOx of TiN/HfO2/CeOx/TiN before annealing (a) and after annealing at 400°C for 2 min (b).
(a)(b)To support our proposition that the grain size (and, therefore, grain boundaries) and oxygen vacancy content jointly influence the conduction of the TiN/HfO2/CeOx/TiN devices, we measured the conductance of the devices before forming. The conductance in the HRS increased with increasing annealing temperature and time, as shown in Figure 3. Rapid thermal annealing at first results in grain growth (and, therefore, grain boundary reduction), which makes the oxygen vacancies that act as ion carriers [31] redistribute and concentrate in the new grain boundaries, tending to form a broad oxygen vacancy channel that facilitates ionic/electronic conductivity, resulting in an increase in conductivity in this regime. However, the relative conductance of the devices decreased with a further increase in annealing time. It can be speculated that the major factor influencing conductance is the oxygen vacancy concentration in this regime, because with the increase in annealing time, the grain size of the CeOx films did not change much (Figure 4(c)), while the oxygen vacancy concentration decreased continually (Figures 5(d) and 5(e)), so the conductance of the devices with voltage sweeping from 0 V to ±5 V to 0 V decreased in this regime.In addition, the interface effect might also be considered [34–36]. It was reported that the resistance of the HRS would decrease after annealing when a metal with high oxygen affinity was used as the TE [37], which is consistent with our result (Table 1). Moreover, the roughness between CeOx/TiN interfaces increased with increasing grain size, which would enhance the uneven distribution of the local electric field to a certain extent, enhance the electric field intensity [38], and promote the formation of CFs. These combined effects might contribute to the reduction of the VSET and VRESET of the device and the dispersion of VSET and VRESET.


## 4.1. The Conduction Mechanism Fitting of TiN/HfO2/CeOx/TiN

To elucidate the resistive switching behavior of TiN/HfO2/CeOx/TiN, the I‐V curves of the devices were fitted with conductive mechanisms (Figures 6(a) and 6(b)). When the unannealed and annealed devices were in the HRS (the lower curves in Figures 6(a) and 6(b)), the I‐V curves possess linear parts with a slope of ~1, which indicates Ohmic conduction (I∝V) [29] due to the partial filling of traps during the weak injection of charge carriers from the bottom/top TiN electrode; the I‐V curves also possess rapid increasing current parts with a slope of ~2, which indicates Child’s square law (I∝V2) [30] due to the injected electrons with applied bias which are trapped/detrapped in the CeOx layer. These results suggest that the devices mainly followed the SCLC conduction mechanism in the HRS. When the unannealed and annealed devices are in the LRS (the upper parts of the curve in Figures 6(a) and 6(b)), the I‐V curves are linear with slopes of ~1, implying that the conduction filament is the main conduction mechanism for the devices at the LRS. These results imply that the conduction mechanisms of TiN/HfO2/CeOx/TiN in the HRS and LRS did not changed after annealing.Figure 6
Log-logI‐V curves and their fit for determining the conduction mechanism of the unannealed TiN/HfO2/CeOx/TiN devices (a) and TiN/HfO2/CeOx/TiN devices annealed at 400°C for 2 min (b).
(a)(b)


## 4.2. Effect of Annealing on the Resistance Switching Characteristics of TiN/HfO2/CeOx/TiN

The conduction ofCeOx switching materials is composed of ionic conduction and electronic conduction [31]; both ionic conduction and electronic conduction are influenced by grain size (and, therefore, grain boundaries), and oxygen vacancies and their distribution. The preferential accumulation of oxygen vacancies along the grain boundary sites has been reported [32]. It was suggested that although there were oxygen vacancies in the grains, the oxygen vacancy concentration near the grain boundaries was generally higher than that in the inner grains, and furthermore, the oxygen vacancies in the grains were generally in linear but discontinuous arrangements, making it difficult to form a conduction channel in the grains. On the other hand, it was more likely that oxygen vacancies formed chains at the grain boundaries to construct the ionic conduction channel, and electrons also tended to accumulate in the vicinity of the grain boundaries and hop along the oxygen vacancies to contribute to the conductivity [33].In this work, annealing significantly reduced the dispersion of theVSET and VRESET and RLRS and RHRS of TiN/HfO2/CeOx/TiN devices (Table 1, Figures 2(c) and 2(d)), which could be ascribed to the combined effect of the increased grain size in the CeOx films (no obvious grain size change was observed in HfO2) (Figure 4(c)) and the reduced oxygen vacancy concentration in the HfO2 and CeOx films caused by annealing (Figure 5). Figure 7 clarifies the possible conduction paths in the CeOx of TiN/HfO2/CeOx/TiN. It can be speculated that the oxygen vacancies accumulate along CeOx grain boundaries, and that the oxygen vacancy channel forms and ruptures along the grain boundaries. For the unannealed devices, the channel forms and ruptures at more abundant boundaries (Figure 7(a)), while for the annealed devices the channel forms and ruptures at fewer positions because of the reduction in the number of grain boundaries (Figure 7(b)), and smaller variations in the conduction paths result in decreased dispersions of the VSET and VRESET and RLRS and RHRS of the annealed devices.Figure 7
Possible conduction paths in theCeOx of TiN/HfO2/CeOx/TiN before annealing (a) and after annealing at 400°C for 2 min (b).
(a)(b)To support our proposition that the grain size (and, therefore, grain boundaries) and oxygen vacancy content jointly influence the conduction of the TiN/HfO2/CeOx/TiN devices, we measured the conductance of the devices before forming. The conductance in the HRS increased with increasing annealing temperature and time, as shown in Figure 3. Rapid thermal annealing at first results in grain growth (and, therefore, grain boundary reduction), which makes the oxygen vacancies that act as ion carriers [31] redistribute and concentrate in the new grain boundaries, tending to form a broad oxygen vacancy channel that facilitates ionic/electronic conductivity, resulting in an increase in conductivity in this regime. However, the relative conductance of the devices decreased with a further increase in annealing time. It can be speculated that the major factor influencing conductance is the oxygen vacancy concentration in this regime, because with the increase in annealing time, the grain size of the CeOx films did not change much (Figure 4(c)), while the oxygen vacancy concentration decreased continually (Figures 5(d) and 5(e)), so the conductance of the devices with voltage sweeping from 0 V to ±5 V to 0 V decreased in this regime.In addition, the interface effect might also be considered [34–36]. It was reported that the resistance of the HRS would decrease after annealing when a metal with high oxygen affinity was used as the TE [37], which is consistent with our result (Table 1). Moreover, the roughness between CeOx/TiN interfaces increased with increasing grain size, which would enhance the uneven distribution of the local electric field to a certain extent, enhance the electric field intensity [38], and promote the formation of CFs. These combined effects might contribute to the reduction of the VSET and VRESET of the device and the dispersion of VSET and VRESET.


## 5. Conclusions

This work demonstrates that rapid thermal annealing can significantly reduce the coefficients of variation (CVs) of theVSET, VRESET, RLRS, and RHRS of TiN/HfO2/CeOx/TiN devices. Annealing increased the grain size and reduced the oxygen vacancy concentration of the switching layers of the devices. We suggest that the combined effect of grain size growth and oxygen vacancy reduction and redistribution is responsible for the improvement in the distributions of the VSET, VRESET, RLRS, and RHRS of annealed devices. This work shows the potential of rapid thermal annealing in optimizing the performance of memristive devices.



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*Source: 6522007-2022-04-11.xml*