H-treatment impact on conductivefilament. formation and stability in

Transcription

1 H-treatment impact on conductivefilament formation and stability in Ta 2 O 5 -based resistive-switching memory cells L. Goux 1,*, J. Y. Kim 2, B. Magyari-Kope 2, Y. Nishi 2, A. Redolfi 1, M. Jurczak 1 1: Imec, Kapeldreef 75, B-3001 Leuven, Belgium 2: Department of Electrical Engineering, Stanford University, Stanford, California 94305, USA * Electronic mail: gouxl@imec.be (Submitted to J. Appl. Phys., 2014)

2 Abstract In this article we evidence the lower formation energy and improved stability of the conductive filament (CF) formed in TiN\Ta 2 O 5 \Ta resistive-switching memory cells treated in NH 3 atmosphere at 400 C. This annealing treatment results in (i) lower forming voltage, (ii) lower CF resistance, and (iii) longer retention lifetime of the oxygen-vacancy (V o ) chain constituting the CF. Atomistic insights into these processes are provided by ab initio calculations performed for Hydrogen (H) species incorporated in non-stoichiometric Ta 2 O 5 supercells: (i) V o formation energy is reduced by the presence of H, (ii) V o chain CF conductivity is increased by V o +OH complex formation, and (iii) V o chain retention is strengthened by the stable V o +OH complex. As a result, efficient CF formation and excellent state stability is obtained after 15 days at 250 C. 2

3 I. Introduction The resistive-switching memory (RRAM) technology has recently gained a lot of interest due to promising scalability, low-voltage and fast programming, as well as impressive progress in terms of reliability [1-6]. Typical cells consist of an oxide layer sandwiched between metal electrodes. After a socalled forming process generally required to create a conductive filament (CF) through the oxide, the cell may be reversibly reset-switched to a highresistance state (HRS) and set-switched to a low-resistance state (LRS). It is generally accepted that the CF consists of a chain of oxygen-vacancy defects (V o ) bridging both electrodes. For an optimum bipolar-switching operation, asymmetric cells are typically used whereby one of the electrodes shows high oxygen affinity in order to increase the oxygen-vacancy profile from the low-affinity electrode towards the so-called oxygen-scavenging electrode [6,7]. Based on this guideline, excellent switching control has been reported in HfO 2 -based cells using either Ti or Hf scavenging metals [1,2,7]. Recently, promising switching and reliability properties have also been demonstrated in Ta 2 O 5 -based cells [3-5]. In this article we use a TiN\Ta 2 O 5 \Ta stack where the Ta layer plays the role of scavenging electrode. In spite of a remarkable technological progress, the microscopic origin of the switching is still debated. In particular, the role of Hydrogen in the V o generation/annihilation is not clear yet. Previous experimental studies suggested a beneficial impact of catalytic anode material on the reset 3

4 operation in NiO [8] and HfO 2 cells [9]. On the other hand, it was experimentally evidenced that the incorporation of H in the oxide layer during deposition allows decreasing the forming voltage [10-12]. Other reports demonstrated an active role of H in the reset operation of TiO 2 [13] or HfO 2 cells [14]. In Ref. 14, we showed that H-incorporation in HfO 2 layers is achieved by post-deposition annealing in H 2 /Ar plasma or NH 3 atmosphere, and that the HfO 2 layer is more aggressively reduced by the former treatment. As on the other hand we evidenced elsewhere [15] that Ta 2 O 5 layers are more sensitive to reducing environment than HfO 2 layers, we exploited the softer anneal treatment in NH 3 atmosphere to incorporate H species in Ta 2 O 5 layers. Using this method, we evidence in this article a clear impact of the NH 3 treatment on forming voltage, LRS resistance level and retention lifetime. These results are finally interpreted based on ab initio calculations for H and V o species introduced in Ta 2 O 5 supercells. II. Experimental The investigated TiN\Ta 2 O 5 \Ta cells are crossbar-patterned RRAM cells stacked on the drain of a select transistor, in a so-called 1-Transistor\1- Resistor (1T1R) integration scheme [figure 1(a)]. After sputter-deposition and patterning of the TiN bottom-electrode (BE), a 6 nm-thick amorphous Ta 2 O 5 layer is deposited by atomic layer deposition (ALD) at 250 C using TaCl 5 as precursor and H 2 O as oxidant [figure 1(b)]. Except for reference cells, a postdeposition rapid-thermal anneal (RTA) treatment in NH 3 is applied at 400 C 4

5 for 5 min [figure 1(c)]. Finally, a top-electrode (TE) stack consisting of 10 nmthick Ta and 30 nm-thick TiN is sputter-deposited and patterned, followed by passivation and bond-pad processing. The size of the RRAM cell is 30x30 nm 2. Current-Voltage (I-V) characteristics are obtained using a conventional parameter analyzer. Forming characteristics are investigated on 1R cells (no select transistor), by applying voltage ramps to the TE while the BE is grounded. Resistive switching is studied in a 1T1R configuration, using voltage ramps applied to the bit-line (BL), contacting the cell TE, while the source line (SL) is kept grounded and the word-line (WL) potential is appropriately chosen to limit the current to I c = 50 µa during forming/set ramps. III. Results Figures 1(d) and (e) show respectively typical forming and switching properties obtained on the reference cells. The cells submitted to the NH 3 treatment have very similar electrical behavior, except for the following aspects: both the forming voltage (V f ) extracted from I-V traces and the LRS resistance (R LRS ) obtained after forming/set operation are lower than for reference cells [figure 2(a,b)]. From X-Ray Reflectometry characterization, neither a densification of the Ta 2 O 5 layer nor significant generation of V o species is induced by the NH 3 5

6 treatment. Although it is not excluded that this treatment generates a low amount of discrete V o species, possibly accounting for lower V f characteristics, this alone cannot explain the lower R LRS data obtained for NH 3 -treated cells, because the V o -density in the neighborhood of the CF after forming and reset is several decades larger than before forming. Thus, here we discuss the possible scenario that the lower R LRS data is induced by the incorporation of H into the Ta 2 O 5 layer during the NH 3 treatment. In order to get some insight into the origin of these electrical effects, we performed ab initio simulations using density functional theory (DFT) with H and V o species introduced in large Ta 2 O 5 supercells. The DFT calculations were carried out using the Vienna ab initio simulation package [16-19], employing the projector augmented-wave method and the GGA+U approximation with corrections applied on oxygen p orbitals U p in addition to metal d electrons, U d. This method had been shown previously to accurately predict the electronic properties of non-stoichiometric TiO 2 and Ta 2 O 5 [20,21]. The O 2s 2 2p 4 and Ta 5p 6 5d 3 6s 2 were considered as valence electrons. All atoms were allowed to relax with the energy convergence tolerance of 10-6 ev/atom and the ground state was obtained by minimizing the force on each atom to be less than ev/å. The V o and H were introduced within a 2x2x4 supercell of a pseudo-hexagonal Ta 2 O 5 consisting of 64 Ta and 159 O atoms. This structure has been predicted previously to be one of the possible structures stabilized in thin film Ta 2 O 5 samples [21]. To assess the role of H in the forming process, we first calculated the formation energy of V o species both in an isolated position and in the presence of H, considering various distances and atomic arrangements 6

7 between V o and H [figure 2(c)]. V o formation in the presence of H was found to be favorable over isolated V o, which explains the V f characteristics trends shown in figure 2(a). With the H incorporation the overall concentration of V o increases in the sample. This holds all the more under forming field, which may results in larger V o density in the CF after forming and possibly accounting for the lower R LRS data obtained, as observed in figure 2(b). Secondly, the H species are also observed to preferably locate in the immediate or next nearest neighborhood of V o sites. Moreover, the formation of the energetically favorable O-H complexes causes local atomic relaxations and electron transfer between the neighboring Ta and O atoms. The electron localization function is shown in figure 2(d) for H placed near a V o. H prefers to form the O-H complex in the immediate vicinity of V o, inducing charge delocalization around the V o, and affecting the local electronic structure. Figure 3 shows the impact of the O-H complex formation on the total density of states. In the absence of H impurity, the defect state generated by the V o is formed mostly by 5d orbitals of Ta atoms with a small contribution of O 2p states, and this neutral V o defect is 0.25 ev wide in the middle of the band gap [figure 3(a)]. With the O-H complex formation, new defect states are introduced also in the mid-gap region, in the close proximity of the V o induced defect, resulting in a widening of the midgap defect state region to ~1 ev [figure 3(b)]. These defect states are occupied by electrons delocalized around the V o and the O-H complex as shown by the electron localization function in figure 2(d). Therefore, the incorporation of H in the CF would increase the conductivity of this very CF. Based on these insights, the 7

8 lower R LRS data obtained for NH 3 -treated cells [figure 2(b)] may be due to increased V o mobility during switching, or due to increased CF conductivity, or both. Further studies are needed to elucidate on this aspect, along the incorporation of V o dynamics and electron transport of metallic character in CF. Finally, we carried out isothermal retention tests on the reference and NH 3 -treated cells. After programming to HRS and LRS states, the cells were baked for 15 days at 250 C. For the reference cell, figure 4(a) shows that both LRS and HRS resistances have drifted towards higher resistances levels after baking. This effect has previously been observed in Al 2 O 3 \Hf and HfO 2 \Hf cells and has been associated to O-V o recombination due to thermally-activated back-diffusion of O-ion species into the V o -rich CF from the intermetallic HfO 2- layer created during forming at the Hf interface [6]. In Al 2 O 3 \Hf stack for example, this O back-diffusion from this intermetallic layer was claimed to be driven by the stronger Al-O chemical affinity [6]. Consistently with this scenario, we have recently shown for Ta 2 O 5 cells that the scavenging electrode material and thickness allows to modulate the oxygen chemical potential and to directly control the retention properties [15]. These results confirmed that the retention properties in TiN\Ta 2 O 5 \Ta cells are less controlled by the lateral diffusion than by the oxygen motion along the CF, as resulting from the chemical potential profile set by the scavenging layer [15]. Based on these findings, figure 4(b) shows for Ta 2 O 5 \Ta cells the O-ions species generated during forming/set operations and at the origin of the created intermetallic Ta 2 O 5- layer at the CF\Ta interface, together with the forming-generated V o species making the CF. Note these two types of 8

9 oxygen species are shown for sake of clarity, however they are not independent species and the actual defect-motion mechanism along the CF may be understood by V o motion along the CF. Interestingly, for cells submitted to the NH 3 treatment, the drift of LRS and HRS resistances towards higher resistances is clearly reduced [figure 4(a)]. These retention properties are excellent as compared to state-of-the-art reported retention characteristics for Ta 2 O 5 cells operated at even larger current [4,5,15]. Figure 4(b) shows a schematic accounting for the improved LRS retention of H-treated Ta 2 O 5 \Ta cells. According to our calculations, the spontaneous formation of O-H complex in the presence of H species stabilizes V o defects in Ta 2 O 5 [figure 2(c)]. We found that these results also hold for H species at non-interacting distances of V o s, as the energy rises by only slightly less than 0.2eV [figure 5]. Hence, the improved LRS retention may be explained, not only by the O-H complex formation close to V o s in the CF, but also because O-H complex formation will also be favored in the intermetallic layer with O-ions generated at forming/set. In this situation, the back-diffusion of these O-ions should be reduced, resulting thus into retarded retention loss. IV. Discussion The aim of this section is to discuss these findings from the broader perspective of past works on H-incorporation into oxide materials in general, which have emerged in particular since early 1980s [22-29]. These works have 9

10 been mostly driven by the rising interest in proton conductivity in oxides for emerging applications such as electrochemical sensors and fuel cells. Therefore, they have been mainly conducted on materials showing large proton conductivity, i.e. on complex perovskites [22-24] or rare-earth oxides [24]. However, H incorporation has also been studied in binary oxides in general [25-29], and in particular, effective H incorporation has been physically evidenced in Ta 2 O 5 layers [28]. From our results, H species introduced in Ta 2 O 5 supercells readily form O-H complexes. In Kroger-Vink notation, this may be written as follows: 1 2 H ( g) O HO x 2 o o e (1) This relation (1) shows that O-H formation has a role of donor, as it leads to the generation of electron defect (e - ). This source of n conductivity would add to the donor role of V o defects generated in the CF, as follows from the relation (2): O x o V 2 o 1 O ( g) 2 2e 2 (2) This understanding is in agreement with the increased CF conductivity discussed from results in figures 2 and 3. The donor behavior of H impurity has already been reported in SiO 2 [25] and ZnO [26]. In ZnO in particular, it was also similarly found that H prefers to locate near V o defects [26]. In general the bond energy of hydroxide defect HO + in oxides is moderate, in the range below 0.8eV [22,24,26]. As said earlier, we found small energy change between V o +H and V o +OH complexes, even when O-H complex is formed far from V o. In fact, an energy increase of less than 0.2eV is obtained 10

11 when O-H complex is moved to a location non-interacting with V o [figure 5], suggesting that these complexes are metastable and that O-H bonds are likely easy to break under high local temperature and field and reform at location under lower stress. Based on these calculations and possible scenario, we propose that V o +OH bonds are stable in temperature ranges up to those of retention baking tests (250 C corresponding to <50meV), and in turn stabilize V o s in these conditions. However these complexes would break during set/reset switching due to the rise of the current-induced temperature of the CF up to ranges >400 C. Indeed, such ranges were previously shown to be favorable to O-H bond breaking and to lead to proton conduction [22,23,29]. In the high local temperature and field stresses during set/reset, high H and V o mobility would be expected, explaining the resistance switching and in particular the results in figure 2(b). V. Conclusion To summarize, in this article we demonstrate that the annealing of Ta 2 O 5 layers in NH 3 atmosphere at 400 C allows lowering the forming voltage and the resistance of the conductive filament (CF) formed in TiN\Ta 2 O 5 \Ta resistive-switching memory cells. This annealing also improves significantly the retention lifetime of the CF. In agreement with these electrical characteristics, ab initio calculations performed for Hydrogen (H) species incorporated in non-stoichiometric Ta 2 O 5 supercells resulted in lower V o 11

12 formation energy in the presence of H and increased CF conductivity due V o +OH complex formation. In addition, the improved CF stability was related to the stabilization of the V o chain by the stable V o +OH complex. Acknowledgments We acknowledge the partial funding by IMEC s Industrial Affiliation program on RRAM. The NNIN Computing Facility of Stanford University, the XSEDE Computing Environment, and the Center for Nanoscale Materials (CNM) Computational Cluster Carbon at Argonne National Laboratory are acknowledged for computational time. Use of the CNM was supported by the U. S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH

13 1 B. Govoreanu, G. S. Kar, Y.-Y. Chen, V. Paraschiv, S. Kubicek, A. Fantini, I. P. Radu, L. Goux, S. Clima, R. Degraeve, et al., Tech. Dig. - Int. Electron Devices Meet. 2011, C.-H. Wang, Y.-H. Tsai, K.-C. Lin, M.-F. Chang, Y.-C. King, C.-J. Lin, S.-S. Sheu, Y.-S. Chen, H.-Y. Lee, F. T. Chen, and M.-J. Tsai, Tech. Dig. - Int. Electron Devices Meet. 2010, M.-J. Lee, C. B. Lee, D. Lee, S. R. Lee, M. Chang, J. H. Hur, Y.-B. Kim, C.-J. Kim, D. H. Seo, S. Seo, U-In Chung, I.-K. Yoo, and K. Kim, Nature Materials 10, 625 (2011). 4 T. Ninomiya, T. Takagi, Z. Wei, S. Muraoka, R. Yasuhara, K. Katayama, Y. Ikeda, K. Kawai, Y. Kato, Y. Kawashima, S. Ito, T. Mikawa, K. Shimakawa and K. Aono, Symposium on VLSI Technology Digest of Technical Papers 2012, 73 5 S. Muraoka, T. Ninomiya, Z. Wei, K. Katayama, R. Yasuhara, and T. Takagi, Symposium on VLSI Technology Digest of Technical Papers 2013, 62 6 L. Goux, A. Fantini, R. Degraeve, N. Raghavan, R. Nigon, S. Strangio, G. Kar, D. J. Wouters, Y. Y. Chen, M. Komura, et al., Symposium on VLSI Technology Digest of Technical Papers 2013,

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17 Figure captions FIG. 1. (a) 1T1R RRAM integration scheme; (b) Cross-section TEM picture of the TiN\Ta 2 O 5 \Ta RRAM cell; (c) schematic of the RTA treatment in NH 3 applied after Ta 2 O 5 ALD deposition; Typical I-V forming (d) and switching (e) traces obtained on the reference cell (no RTA treatment). FIG. 2. Impact of the NH 3 treatment on the forming voltage V f (a) and on the Cumulative Distribution Function (CDF) of the LRS resistance R LRS obtained using I c =50µA (b), for a statistical sample of >30 cells; (c) ab-initio calculations of the formation energy of isolated V o defects with or without H incorporation; (d) electron localization function (ELF) calculated in the pseudo-hexagonal Ta 2 O 5 with one neutral V o and one H impurity; the ELF shows the missing electrons in the blue regions and the highly localized ones in the red regions. FIG. 3. Total density of states calculated in the pseudo-hexagonal Ta 2 O 5 with one neutral V o (a) or a neutral V o combined with O-H complex (b), showing a widening of the defect state in the gap after H incorporation. FIG. 4. (a) Impact of the NH 3 treatment on the retention properties of LRS and HRS states after 15 days at 250 C, showing improved state stability; (b) schematics of the thermally-activated back-diffusion of O-ion species during 17

18 baking, leading to O-V o recombination. For H-treated cells, not only the O-H complex formation in the CF stabilizes V o s in the CF, but also the favored O- H complex formation with O-ions generated at forming/set in the intermetallic layer formed at the interface with the scavenging Ta layer lowers the driving force for back-diffusion of O-ions from the Ta layer, resulting thus into retarded retention loss. FIG. 5. Thermodynamic stability dependence on the H-V o distance, showing the better stability obtained, by slightly less than 0.2eV, when V o and H are close to each other (conditions A and B) as compared to non-interacting V o and H (condition C) 18

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