Christoph
Baeumer
a,
Thomas
Heisig
a,
Benedikt
Arndt
a,
Katharina
Skaja
a,
Francesco
Borgatti
b,
Francesco
Offi
c,
Federico
Motti
d,
Giancarlo
Panaccione
e,
Rainer
Waser
af,
Stephan
Menzel
a and
Regina
Dittmann
*a
aPeter Gruenberg Institute 7, Forschungszentrum Juelich GmbH, JARA-FIT, 52425 Juelich, Germany. E-mail: r.dittmann@fz-juelich.de
bCNR, Istituto per lo Studio dei Materiali Nanostrutturati (ISMN), Via P. Gobetti 101, 40129 Bologna, Italy
cCNISM, Dipartimento di Scienze, Università di Roma Tre, Via della Vasca Navale 84, 00146, Rome, Italy
dDipartimento di Fisica, Università di Milano, Via Celoria 16, I-20133 Milano, Italy
eIstituto Officina dei Materiali (IOM)-CNR, Laboratorio TASC, in Area Science Park, 34149 Trieste, Italy
fInstitute of Materials in Electrical Engineering and Information Technology, RWTH Aachen University, 52056 Aachen, Germany
First published on 3rd September 2018
Resistive switching oxides are highly attractive candidates to emulate synaptic behaviour in artificial neural networks. Whilst the most widely employed systems exhibit filamentary resistive switching, interface-type switching systems based on a tunable tunnel barrier are of increasing interest, since their gradual SET and RESET processes provide an analogue-type of switching required to take over synaptic functionality. Interface-type switching devices often consist of bilayers of one highly mixed-conductive oxide layer and one highly insulating tunnel oxide layer. However, most tunnel oxides used for interface-type switching are also prone to form conducting filaments above a certain voltage bias threshold. We investigated two different tunnel oxide devices, namely, Pr1−xCaxMnO3 (PCMO) with yttria-stabilized ZrO2 (YSZ) tunnel barrier and substoichiometric TaOx with HfO2 tunnel barrier by interface-sensitive, hard X-ray photoelectron spectroscopy (HAXPES) in order to gain insights into the chemical changes during filamentary and interface-type switching. The measurements suggest an exchange of oxygen ions between the mixed conducting oxide layer and the tunnel barrier, that causes an electrostatic modulation of the effective height of the tunnel barrier, as the underlying switching mechanism for the interface-type switching. Moreover, we observe by in operando HAXPES analysis that this field-driven ionic motion across the whole area is sustained even if a filament is formed in the tunnel barrier and the device is transformed into a filamentary-type switching mode.
Interface-type switching devices usually consist of bilayers of one highly mixed-conductive oxide layer and one highly insulating tunnel oxide layer. It is commonly supposed that resistive switching in these devices is induced by a laterally homogeneous change in the tunnel barrier height induced by the field-driven motion of oxygen ions. However, the details of the underlying processes are not well understood so far. The scenario is further complicated by the fact that most tunnel oxides used for interface-type switching are also prone to form conducting filaments above a certain voltage bias threshold.5 Here, we investigated two different tunnel oxide devices, namely, Pr1−xCaxMnO3 (PCMO) with yttria-stabilized ZrO2 (YSZ) tunnel barrier and substoichiometric TaOx with HfO2 tunnel barrier, by hard X-ray photoelectron spectroscopy (HAXPES), in order to gain insights about the chemical changes in different heterostructures during filamentary and interface-type switching.
Thanks to the larger probing depth (>10 nm) achievable with respect to conventional, soft X-ray photoelectron spectroscopy, HAXPES provides an excellent, nondestructive approach to probe chemical changes and charging effects of interfaces buried inside layered heterostructures, in addition to the bulk electronic structure of the materials.6–14 In particular, in operando HAXPES provides an opportunity to correlate the resistive switching effects to the electronic structure and the electrical potential of the probed volume.15–17
For the PCMO/YSZ heterostructures, we observed a reversible shift of the Zr 3d core-level between the two resistive states during interface-type switching, which we attribute to a shift of negative charge into the YSZ layer during the RESET process.16 This is consistent with an exchange of oxygen ions between PCMO and YSZ that causes an electrostatic modulation of the effective height of the YSZ tunnel barrier. For this material stack, the devices show a permanent breakdown when a certain bias value is exceeded and no filamentary switching mode can be stabilized.
For interface-type switching TaOx/HfO2 heterostructures, we were able to also turn the devices into a filamentary switching mode with the same switching polarity. We recorded the Ta and Hf core-level peaks after the application of positive and negative bias, respectively, by in operando HAXPES. The observed shifts of the Hf and Ta core-level spectra and the varying amounts of substoichiometric TaOx are consistent with a reversible shift of oxygen vacancies between the TaOx layer and the HfO2 layer during filamentary switching. However, due to the comparably small estimated filament size, the relative change of the electronic structure would be insufficient in order to be detected on the large device area of 60 × 200 μm2. We therefore attribute our observations to changes of the electronic structure across the whole device area taking place as a side effect during electrical biasing. We could thereby evidence a co-existence of strongly temperature-accelerated ionic motion within the filament and near room temperature ionic motion taking place over the whole device area. We complemented our measurements with finite-element calculations of the field and temperature distribution in our devices and discuss its implications on the co-existence of the two processes.
In order to take advantage of the high bulk sensitivity of HAXPES in grazing incidence geometry, which results in a beam footprint exceeding by far the single device area, we prepared special array-type samples (see also ref. 9). The arrays consisted of single devices 50 × 50 μm2 in size and arranged in arrays about 7 mm long and 0.5 mm wide, as shown in Fig. 1(a). The spacing between the devices was 10 μm. The large surface area of one array allows the adoption of a grazing incidence geometry with a grazing angle of 5° for the X-ray beam. This way, about 50 devices of the over one thousand devices present in one stripe of devices were probed simultaneously. The device resistance states were programmed ex situ, using a semi-automatic probe station, in advance. The devices were set to the LRS and HRS by voltage sweeps. One array was not programmed and was left in the initial or as-deposited state.
The core-level spectra for Rh 3d and Zr 3d are presented in Fig. 1(c) and (d). Both panels show the spectra for arrays in the initial state (as deposited) and after programming to the HRS and LRS. No changes of the line shape are evident in the core-level spectra of Rh and Zr, indicating the absence of changes in the chemical environment within the detection level of our measurements. The reproducibility of the Rh 3d line shapes indicate that the Rh remains metallic throughout switching and that no oxidation takes place.18 In contrast, clear shifts can be observed in the position of the Zr 3d core-level. In particular, both the LRS and the HRS are shifted to lower binding energies compared to the initial state, with the HRS shifted to about 0.1 eV lower binding energies than the LRS.
We interpret this result as the build-up of an electrostatic charge within the YSZ rather than a change in the Zr oxidation state, because the line shape does not change. We propose that the electrostatic charge originates from O2− ions (ionic space charge), which migrate into and out of the YSZ but are not charge-compensated by a valence change of the Zr ions.
Concerning the switching between the LRS and HRS, we conclude that oxygen ions move into the YSZ layer upon application of a positive bias to the top electrode, leading to an increase of the tunnel barrier height and hence to a decreased tunnel current in the HRS state, as sketched in Fig. 1(e). Conversely, applying a negative bias leads to the LRS state by removing oxygen ions from the YSZ. This modulation of the effective tunnel barrier height is caused by the excess negative charge that accumulates in the YSZ by the migration of oxygen ions. This description can be reproduced via I–V curve simulation of both resistance states based on asymmetric tunnel barriers with variable effective barrier heights.16 At the same time, it is consistent with our HAXPES observations, as negative charges increase the kinetic energy of the photoelectrons, shifting the spectrum towards lower binding energies. In the same sense, by removing negative charge from the interface, the spectrum shifts towards higher binding energies. This model is in general similar to the mechanism suggested for PCMO/tunnel oxide devices in ref. 19, however it slightly differs since we observed only a less negatively charged YSZ barrier in the LRS and not a positively charged region as suggested by the authors.
The samples fabricated for in operando measurements possess an edge length of 60 × 200 μm2 and the same sample geometry as the YSZ/PCMO samples, as sketched in Fig. 4(a). The surroundings of the devices are covered by SiO2 to guarantee electrical insulation and prevent any photoemission contribution from the bottom electrode. The I–V curves of the devices looked similar to the small test devices shown in Fig. 2(a), besides that the current increased as expected from the area scaling behaviour.
We started HAXPES measurement on our device while applying a sequence of negative voltages 0 V → −1 V → −2 V → −0.2 V → 0 V to our devices and recorded the resistance of the devices as depicted in Fig. 3(a). We observed a sudden drop in the resistance at 120 s while applying constant voltage. After increasing the voltage to −2 V, the resistance decreased further in several sudden jumps until reaching a resistance of ∼2000 Ω after 400 s. Due to the current limit of 100 μA, the cell voltage dropped to −0.2 V simultaneously. Fig. 3(b) shows the I–V curve that was recorded subsequently. The I–V curve has a completely different shape than the I–V curves recorded from the device before (similar curves are shown in Fig. 2(a)) and looks very similar to filamentary switching HfO2 devices.20,21 In particular, the I–V curve in Fig. 3(b) exhibits a sharp SET process in the negative branch and a resistance ratio between the HRS and LRS of about three orders of magnitude, which is about two orders of magnitude higher than in the interface-type switching mode of the cell observed before. We can therefore conclude that a filament was formed within the HfO2 layer during DC electrical biasing in the vacuum chamber. The exposure of the sample to the synchrotron beam might have induced oxygen vacancies in the HfO2 layer similarly as reported for SrTiO3 single crystals,22 and might thereby have promoted this forming process.23
Fig. 3 Details of the electrical biasing procedure during in operando measurements. (a) First biasing scheme which resulted in a transformation to filamentary switching. Timing of the voltage steps (red) and the corresponding device resistance (blue). (b) I–V curve recorded after applying the voltage sequence shown in (a). (c) Biasing scheme employed for the in operando HAXPES measurements shown in Fig. 4. Timing of the voltage steps (red) and resulting resistance (blue). The different devices states, namely (A), (B), and (C) at which the core-level spectra at 0 V were recorded are marked with black points. |
Fig. 4(b) and (c) show the HAXPES spectra collected with an X-ray photon energy of 3.2 keV with an incidence angle of 45°, illuminating the device area of the selected device. The photoelectrons were detected at 45° emission geometry. The overall spectral resolution was determined to be 0.40 eV. With these settings, we were able to detect a significant signal intensity for the core level peaks Rh 3d (top electrode), Hf 4f (tunnel barrier), Ta 4f (conductive oxide layer), and O 1s (HfO2, Ta2O5 and surrounding SiO2). In order to determine the peak position and the spectral intensity of each term, all the core level spectra were analysed with CasaXPS Version 2.3.17. A Shirley background profile was subtracted from all of them. The peaks of the oxides (Ta 4f, Hf 4f) were fitted using Voigt functions while the core level spectra of the metal (Rh 3d) were fitted using a Doniach–Sunjic lineshape. The spectra clearly indicate that we were able to probe both the Rh/HfO2 as well as the HfO2/TaOx/Ta interfaces with a good signal to noise ratio. The first Ta 4f and Hf 4f spectra recorded at 0 V (corresponding to state (A) in Fig. 3(c)) are displayed in Fig. 4(b). The spectra can be nicely reproduced by a fit with Ta5+, Ta4+, Ta0 and Hf4+ components. The main contribution to the Ta 4f spectrum can be assigned to Ta5+, however, a significant amount of Ta4+ and metallic Ta0 is present as well. It is important to note that there is no evidence of a suboxide component in the HfO2 layer.
Fig. 4 (a) Sketch of the layout for the in operando measurement of the TaOx/HfO2 layer stacks. (b) Ta 4f and Hf 4f core-level spectra recorded at 0 V in state (A) of Fig. 3(b) together with the fit of the data. (c) Ta 4f and Hf 4f core-level spectra recorded at different sample voltages. (d) Binding energy of the Rh 3d and Hf 4f peak for the different states (A)–(C) of the device. (e) Binding energy of the different Ta 4f spectral components in the different states (A)–(C) of the device. (f) Ta/Rh peak area ratio of Ta5+ and (Ta4+ + Ta0) species in the different states (A)–(C) of the device. |
The core-level position of the Rh spectra (not shown) shifts linearly with the voltage as observed for the PCMO/YSZ devices. Fig. 4(c) shows the voltage dependence of the Ta 4f and Hf 4f core level spectra recorded during in situ biasing with voltages in a range of 4 V to −2 V according to the biasing scheme shown in Fig. 3(c). The spectra shifted linearly to higher (lower) kinetic energy for negative (positive) applied voltage, proving that the devices are electrically biased in a proper way. The shift of the core-level spectra is different from the magnitude of the applied voltage. This is due to a significant portion of voltage drop at the series resistance of the measurement equipment and the electrical connections. However, the actual voltage drop across the device is provided directly by the relative shift of the core-level spectra for the Rh and Ta layers.
For a closer inspection, we determined the peak positions of the spectra collected at 0 V, i.e., after initial negative bias (state (A)), after positive biasing (state (B)) and after negative biasing (state (C)). The determined shifts of the Rh 3d and Hf 4f core-level peaks and, for the Ta 4f, core-level peaks of the different Ta valence state contributions in the three different states ((A)–(C)) are shown in Fig. 4(d) and (e), respectively. The Rh 3d peak position is constant in the three states, indicating that no changes are taking place in the Rh during electrical biasing. In contrast, the Hf 4f binding energy decreases slightly from state (A) to state (B) and increases again in state (C). All three Ta binding energy positions show the opposite trend, namely they increase from state (A) to state (B) and decrease again in state (C). Assuming a shift of the relative position of the Fermi-level within the band gap of the TaOx layer as the origin of this shift, this indicates a reduction from state (A) to state (B) and an oxidation of the layer from state (B) to state (C). This is fully consistent with the changes in the spectral weight of the different Ta valences shown in Fig. 4(f). In particular, the Ta5+ contribution is higher at state (A) and (C) and the lower valence state contributions increase in state (B).
Since we did not see a change in the spectral weight of HfO2 in the different states, we attribute the observed shift of the kinetic energies to a negative charging of the HfO2 layer from state (A) to state (B) and a positive charging from state (B) to state (A). Keeping in mind that positive voltage steps were applied from state (A) to state (B) and negative voltage steps between state (B) and state (C) (see Fig. 3(c)), the observed effect would be consistent with a shift of negatively charged oxygen ions between the TaOx and the HfO2 layer. The suggested mechanism is sketched in Fig. 5. By applying a negative voltage to the top electrode, oxygen moves from the HfO2 layer into the TaOx layer, leading to a less negatively charged HfO2−δ tunnel barrier and a more oxidized TaOx layer (Fig. 5(a)). In turn, by applying a positive voltage to the top electrode, the oxygen moves from the TaOx layer into the HfO2−δ layer, leading to a more negatively charged HfO2 tunnel barrier and a more reduced TaOx−δ layer (Fig. 5(b)). Thus, the HfO2−δ tunnel barrier decreases (increases) due to positive (negative) charging induced by the motion of the oxygen ions (cf.Fig. 5).
This process is the same as that proposed for the interface switching PCMO/YSZ devices. However, one should keep in mind that the electrical behaviour of the TaOx/HfO2 devices is completely different. The devices have been transformed to a filamentary switching mode during the first DC biasing step as shown in Fig. 2(b). In order to assign the systematic change of the observed spectral fingerprint to the resistance changes of the device, we will now consider in detail the resistance changes observed in the three different states. In state (A), the device is in an LRS and is expected to switch into an HRS in state (B) after the application of a positive bias sequence. However, as can be seen in Fig. 3(c), the device resistance further decreased in state (B). Afterwards, we applied a negative voltage, which would nominally switch the device back to the LRS. However, it can be seen in Fig. 3(c) that the resistance remains unchanged during the last biasing step. Therefore, we can conclude that we have induced a permanent breakdown of the device in the filament region (sketched as the shaded region in Fig. 5), which determines the device resistance. Therefore, the observed systematic changes of the core-level shifts and the spectral distribution between the different Ta valence states are not correlated with the resistance state of the device. Assuming a typical filament size in the nm regime, the relative change of the electronic structure would be insufficient in order to be detected within the device area of 60 × 200 μm2 (even a very large filament of 1 μm diameter would contribute only ∼0.01% of the total HAXPES intensity). Therefore, we are unable to draw any conclusions about the chemical composition of the filament and we have to assume that the observed chemical changes take place at the interface region of the entire device area. These observations imply that the reversible ionic motion at the whole interface region (which presumably gives rise to the interface-type resistive switching in both PCMO/YSZ and TaOx/HfO2 devices) is sustained despite the filament formation and the final breakdown process. This has important implications for the HAXPES analysis of filamentary resistive switching devices, since the parallel processes might lead to misinterpretations of the observed chemical changes if they are attributed to the nanoscale filament region within 100 μm-sized device areas.
To evaluate the temperature distribution in the TaOx/HfO2 layer stack, an electro-thermal simulation model was used. To reduce the numerical complexity, we considered the two-dimensional axisymmetric structure shown in Fig. 6(a). It consists of a 430 nm-thick SiO2 bottom layer, a 50 nm-thick Ta electrode, a 5 nm-thick TaOx layer, a 2 nm-thick HfO2 layer, and a 2 nm-thick Rh electrode. The HfO2 layer is divided into a filamentary HfO2−γ region with radius rfil at the symmetry axis and a less conducting HfO2 matrix region. An insulating 100 nm-thick SiO2 layer is located on top of the Rh at the outer rim of the stack covering 11 μm of the Rh electrode. A 130 nm-thick Au/Pt electrode covers 9 μm of the Rh electrode and the SiO2 spacer.
To calculate the temperature distribution in this structure, the static heat transfer equation
(1) |
−∇J = ∇σ∇φ = 0 | (2) |
σmatrix = J(Vapp)dHfO2/Vapp. | (3) |
The conductivity of the filament σfil is calculated by assuming that the total current during filamentary switching flows through the filament. The total cell resistance Rcell can then be approximated by the resistance of the filament Rfil and the spreading resistance Rspread in the TaOx film according to
(4) |
Solving the equation for σfil yields
(5) |
This system of equations is completed by an appropriate set of boundary conditions for solving the partial differential eqn (1) and (2), which are shown in Fig. 6(a). The voltage Vapp is applied to the top electrode, while the bottom electrode is set to ground. For all other boundaries a Neumann boundary condition is used setting the normal charge flow to zero. The temperature at the bottom of the SiO2 layer is set to T = 300 K, as it is in contact with the measurement stage. For all other boundaries the heat flux is set to zero.
As the exact filament radius is not known, the electro-thermal simulations are performed for different filament radii. To simulate the temperature distribution during the filamentary SET transition, the applied voltage is set to Vapp = −1.25 V and the current Icell is varied from −10 μA to −1.5 mA. The latter leads to a change of the filament conductivity according to eqn (5).
Fig. 6(a) shows the simulated temperature distribution at Icell = −1.5 mA and rfil = 50 nm. The highest temperature evolves in the filament itself, while the matrix remains at a low temperature. This observation is valid for all simulations performed as shown in Fig. 6(b). The temperature within the filament is a strong function of the current as well as the filament radius. Based on the assumption that the abrupt current increase during the filamentary SET transition is due to a thermal runaway process, a minimum temperature increase of 5 K should occur at the lower current level of the current jump.25 This is only true for filament radii smaller than 50 nm. Another interesting result is the average voltage drops over the filament and the surrounding HfO2 matrix. The voltage drop over the filament increases with increasing filament radius. The reason is the increased spreading resistance of the TaOx layer below the filament for small filament radii. In contrast, the voltage drop over the HfO2 matrix is constant and equals the applied voltage despite the presence of the highly conducting filament.
In a second simulation study, the temperature during interface-type switching was estimated. For this purpose, we considered only a homogenous HfO2 layer without a filament, which corresponds to the situation in our device before the forming step in Fig. 3(a). According to Fig. 2 the highest electrical power is dissipated at the turning point of the SET transition, i.e. Vapp = −4 V and J = −2 × 106 A m−2. Using this data pair, the conductivity of the HfO2 layer is recalculated using eqn (3). The simulation result shows a slight temperature increase of the whole sample of about 10 K. The voltage drop over the HfO2 layer again equals the applied voltage.
Our simulations show that this type of switching goes along with local Joule-heating, leading to very high filament temperatures, especially for small filament radii (note that for the thermal runaway leading to filamentary resistive switching, filament radii below 50 nm are required). At the same time, our simulations reveal that the temperature in the tunnel oxide matrix remains close to RT for the interface-type switching, and even if filamentary switching with high local temperatures occurs in parallel. In both cases, the voltage drop over the tunnel oxide layer is constant and equals the applied voltage. As the HAXPES signal used to derive the switching mechanism originates from the entire device area, this means that the oxygen ion motion for the whole interface area occurs close to RT.
To examine whether such purely field-driven ionic motion may lead to the observed changes in the electronic structure and the corresponding resistance change, we consider the Mott–Gurney equation, which can be used as a first approximation to calculate the field-induced velocity of oxygen ions:29
(6) |
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