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We first present the scaling trends of HZO capacitors and the main switching characteristics of the thinnest (5 nm) ferroelectric HZO obtained. Thickness scaling and proper choice of electrode materials are both important for operation of single layer FTJ. Semiconductor bottom electrodes such as Ge or Nb-doped SrTiO₃ (NSTO) are used because of their lower carrier densities compared to metals which allow for uncompensated polarization charges, the latter being very important for the FTJ functionality. Ge produces sharp, crystalline interfaces with HZO¹⁷ therefore, it is preferred over Si which typically forms low-k interfacial oxides or silicates. On the other hand, NSTO has been used successfully for the fabrication of BaTiO₃ (BTO)/NSTO FTJs⁴, where the giant tunneling electroresistance effect was obtained by modulating the Schottky barrier thickness at the interface, for various Nb concentration¹⁸, therefore it could be a good candidate also for HZO/NSTO FTJs. HZO is deposited by plasma assisted molecular beam deposition^17,19,20 and crystallization annealing is performed via Rapid Thermal Annealing (RTA)²⁰ or Flash Lamp Annealing (FLA)¹⁷ as described elsewhere^17,19,20,21 and in Methods.

The main structural characteristics of the thinnest TiN/HZO/Ge capacitor structures are shown in Fig. 1 and Supplementary Note 1. Grazing Incidence X-ray Diffraction (GIXRD) scans before (gray line, amorphous) and after annealing (blue line) are shown in Fig. 1a. The latter shows no evidence of the non-polar monoclinic phase (green lines at 28.4 and 31.7 deg²²) while the main peak at 30.7 deg (Fig. 1a) is attributed to a combination of the ferroelectric orthorhombic phase Pca2₁²³) and the tetragonal phase P42/nmc²².

a Grazing Incidence X-Ray Diffraction (GIXRD) spectra before (gray line, amorphous) and after annealing (blue line), gaussian fit (cyan line) and reference powder diffraction patterns of tetragonal t-ZrO2, monoclinic m-HZO and orthorhombic o-HZO. b X-Ray Reflectometry (XRR) and Fast Fourier Transform (FFT) scan. c High-resolution Transmission Electron Microscopy (HRTEM) image of the Ge/ HZO /TiN heterostructure. d FFT pattern of the crystallite marked by the square in (c). e simulation of the FFT in (d) assuming Pca2₁ orthorhombic ferroelectric structure along the [-23-1] zone axis. f, g Show the crystallite in detail and an HRTEM simulation with the orthorhombic ferroelectric phase Pca2₁ embedded in the real image in (g). A defocus of −20 was assumed, which is the same as the one used to obtain the experimental real image in (f).

From the Scherrer fit of the peak at 30.7 deg, with k-factor ranging from 0.7 to 1, the average crystalline grain size along the direction ω−2θ/2 = 15 deg (approximately in-plane) is estimated to be between 4 and 5.6 nm. A HZO thickness of 5 nm is accurately measured by X ray reflectivity (Fig. 1b) and cross-sectional Transmission Electron Microscopy (TEM) (Fig. 1c and Supplementary Fig. 1). No interfacial layers at the top and bottom surface of HZO are detected by TEM (Fig. 1c and Supplementary Fig. 1), although a non-discernible amorphous TiO_xN_y layer at the top interface cannot be excluded. Based on filtering of the Fast Fourier Transform (FFT) pattern (Fig. 1d)), the grain size of the HZO polycrystalline film is estimated between 5 nm and 20 nm (Supplementary Fig. 1), noting that the lower limit is compatible with the GIXRD results discussed above. Based on FFT (Fig. 1d, e) and real image (Fig. 1g) simulations, the observed nanocrystallites (Fig. 1f) are compatible with the ferroelectric orthorhombic phase Pac2₁. However, the distinction between orthorhombic and tetragonal at the macroscale based on GIXRD (Fig. 1a) is difficult. It is anticipated that possible tetragonal phase will be present in rather small quantities since, after wake-up, the remanent polarization P_r is high (>20 μC cm⁻²) (see Supplementary Note 2 and discussion below) indicating the dominance of the ferroelectric orthorhombic phase.

Figure 2 shows the evolution of the switching characteristics as a function of decreasing HZO thickness for Ge and NSTO Metal Ferroelectric Semiconductor (MFS) capacitors. The general trend is that thick HZO capacitors show full ferroelectric hysteresis loops in the pristine state, while as the thickness is reduced, the hysteresis loops become distorted and at 5 nm, they are pinched (AFE-like) yielding low remanent polarization. Full hysteresis loops are recovered (“waked-up”) after field cycling while the two switching peaks in the displacement current nearly merge.

a–h Polarization and Displacement current as a function of applied voltage, for several thicknesses of Hf_0.5Zr_0.5O₂ (HZO) films on Ge substrate. i–p Polarization and Displacement current as a function of applied voltage, for several thicknesses of HZO films on Nb-doped STO (NSTO) substrate.

For thick HZO films (Fig. 2a and b), only few cycles are needed but for thinner films, larger number of cycles are required to wake-up the devices. In some cases (as in Fig. 2c, g), the devices are not affected by cycling at all and remain pinched. It should be noted that for 4 nm HZO and thinner (Fig. 2d, h, l, p), ferroelectricity becomes obsolete and HZO behaves like a normal dielectric.

The waked up devices down to the thinnest 5 nm HZO show high remanent polarization P_r > 20 μC cm⁻² after correction, with only small parasitic (non-switchable) polarization (see PUND measurements in Supplementary Fig. 2) and a low coercive voltage V_c ~ 1 V. These performance characteristics are very promising for low voltage/low power FTJ memory devices. Moreover, the high P_r indicates the predominance of the ferroelectric orthorhombic phase as already mentioned above.

The capacitance-voltage measurements on p-Ge/HZO/TiN capacitors at 1 kHz are shown in Fig. 3 featuring the characteristic “butterfly” structure due to ferroelectricity in HZO. Due to the “butterfly” structure which masks the capacitance variations with bias, and the low energy gap of Ge (consequently the high minority carrier concentration) the C-V at room temperature looks more like an MFM capacitor rather than MFS. Note however, that at lower temperatures where the minority carrier concentration is dramatically reduced, a more typical MOS-like behavior is recovered with a clear distinction between accumulation and depletion¹⁷. The accumulation capacitance C_acc (measured at > −2V) of the 5 and 15 nm capacitors are measured to be 3.1 and 1.7 μF cm⁻², respectively. If a single layer HZO is assumed, the measured C_acc yields a value of HZO relative dielectric constant ε_F = 18 which is not compatible with the orthorhombic ferroelectric phase, since the latter has an expected value of ε_F in the range between 30 and 40 [refs. ^{20,24,25,26,27,28}]. It is therefore necessary to consider an interfacial capacitance C_int in series with the FE capacitance C_F in order to account for the measured C_acc. Taking into account both 5 nm and 15 nm capacitors and by assuming the same C_int and ε_F, values of C_int ~ 6 μF cm⁻² and ε_F ~ 39 are extracted. This value of C_int is considerably smaller than the values of 30–110 μF cm⁻² reported for Metal-Ferroelectric-Metal (MFM) capacitors^27,29 with very thin or no interfacial layers.

ac Capacitance (C) -Voltage (V) curves of 15 nm and 5 nm thick films at a frequency of 1 KHz.

In our case, C_int could be dominated from the semiconductor region near the interface due to the finite screening length l. Using an expression³ for \(l=\frac{{{{{{\rm{h}}}}}}}{q}* \sqrt{\frac{{\varepsilon }_{0}{\varepsilon }_{{Ge}}}{3{m}^{* }}}* {\left(\frac{3}{8\pi }\right)}^{\frac{1}{3}}* {\left(\frac{1}{{n}_{h}}\right)}^{\frac{1}{6}}\) (h is the Planck constant and q is the electron charge), derived for metallic systems and assuming a hole concentration n_h = 1.6*10¹⁷ cm⁻³, effective mass m* = 0.29m_e³⁰ and ε_s (Ge) = 16, a value of l ~ 3.5 nm and a corresponding capacitance C_s~ε₀ε_s/l~3.6 μF cm⁻² are estimated. Although C_s is only a rough estimate of the semiconductor capacitance, it does not differ much from the C_int derived from capacitance measurements. In fact, a capacitance C_it due to interface defect charging could be added in parallel to C_s so the total capacitance C_int = C_s + C_it could better match the measured value of 6 μF cm⁻² (see above). In conclusion, the semiconductor capacitance, including possible contribution from interface defects, could account for the rather small value of measured C_int even in the absence of interfacial (dead) dielectric layers.

Metastable ferroelectricity and pinched pristine hysteresis curves

To analyze the switching characteristics of the thin HZO MFS capacitors,

we assume a 1st order ferroelectric with Gibbs free energy [Eq. 1].

where P is the polarization order parameter and α < 0, β < 0 and γ > 0 are the Landau coefficients. From simple electrostatics (see Supplementary Note 3, for details and refs. ^31,32), the electric field E_F in the ferroelectric is given by

Where V is the applied bias on the MFS, and C_F, C_int are the capacitances of the ferroelectric and interfacial layers, respectively and t_F is the FE thickness. The first term in the right-hand side of the equation accounts for the reduction of the volage in the ferroelectric by a factor \({\left[1+\frac{{C}_{F}}{{C}_{{int}}}\right]}^{-1}\), since part of the voltage drops on interfacial (dead layers), or, on the semiconductor near the interface. The second term in the right-hand side of Eq. (2) is the electric field built as a result of uncompensated charge. The surface charge σ is the charge trapped at pre-existing defects at the interface; it has a sign opposite to that of P and assists in the compensation of polarization charges thus stabilizing P. Charge σ becomes important only during field cycling (wake-up) when charge injection and trapping occur and it will be discussed in subsequent sections. Here in this section, we study the pristine case where σ ≈ 0. It should be noted that when V = 0, then P = P_s (spontaneous polarization) and E_F = E_dep= \(-\frac{{P}_{s}}{{\varepsilon }_{0}{\varepsilon }_{F}}{\left[1+\frac{{C}_{{int}}}{{C}_{F}}\right]}^{-1}\), that is the DF.

Using (1) and (2) with σ ≈ 0, the FE thickness dependent expression for G becomes

In the last equation, C_F has been expressed as C_F = ε₀ε_F/t_F, where t_F is the FE thickness. A corresponding expression for C_int cannot be given due to the lack of knowledge about the interfacial (dead) layer characteristics. Instead, experimental C_int values are deduced from C–V analysis (see Fig. 3 and related discussion above) which yield C_int.

The polarization switching dynamics under the influence of a time varying bias V are described by the Landau-Khalatnikov (L-K) equation \(\delta \frac{{dP}}{{dt}}=-\frac{{dG}}{{dP}}\) ^33,34 where δ is the kinetic (or viscosity) coefficient. Using Gibbs free energy from Eq. (3), the L-K equation becomes:

The L-K Eq. (4) is solved by applying time-varying, triangular waveforms for V (see Supplementary Note 4, Supplementary Note 5) to simulate the experimental P-V curve while equation \(J=\frac{{dD}}{{dt}}={\varepsilon }_{0}\frac{d{\rm E}}{{dt}}+\frac{{dP}}{{dt}}\) is complementarily used to calculate the displacement current J–V curves. The simulated curves (Fig. 4a and b) using a kinetic coefficient δ = 2.5*10⁶ s cm F⁻¹ and Landau parameters α = −23*10¹⁰ cm F⁻¹, β = −21.4*10¹⁹ cm⁵ F⁻¹ Cb⁻² and γ = 15.5*10²⁸ cm⁹ F⁻¹ Cb⁻⁴, are in semiquantitative agreement with the experiment.

a,b Experimental data (dotted) and simulation (red line) of Polarization (P) and Displacement current (I) respectively, as a function of Voltage (V) at 1KHz for 5 nm Hf_0.5Zr_0.5O₂ (HZO) films. c Gibbs free energy landscape drawn using the Landau coefficients which are determined from the simulated P–V and I–V.

The deviations are partly due to non-idealities of the experimental curves. A small offset in the voltage (~0.15 V) was set in the calculations in order to account for rigid shifts to higher bias due to small workfuncion differences between top TiN and bottom Ge electrodes. The simplified phenomenological LGD model which ignores extrinsic factors affecting switching dynamics (see discussion section) also accounts for the discrepancies between for the experimental and simulated curves.

At a fixed FE thickness of t_F = 5 nm, C_int and ε_F are the two parameters which control the depolarization ~P² term in Eq. (3), hence the pinching of the hysteresis curve. It should be emphasized that the pinching characteristics (reduced P_r) are best simulated with C_int = 5.5 μF cm⁻² and ε_F = 39 which compare well with the values of C_int and ε_F obtained from the analysis of the C-V measurements (Fig. 3). Our modeling indicates that the positive depolarization term (~P²) in the Gibbs free energy counteracts the negative ¼ βP⁴ (and the negative ½ αP²) term and results in the triple potential well (Fig. 4c) which describes a metastable HZO ferroelectric. The stable paraelectric phase (P = 0) is then responsible for the pinching of the P-V hysteresis reducing the remanent polarization to near zero values. More insights are given in the discussion and conclusions section below.

To better illustrate the influence of C_int on the ferroelectric stability and the switching characteristics, the Gibbs free energy landscapes and the P-V curves for two different C_int values of 10 and 20 μF/cm² are given in Fig. 5. When C_int is has an intermediate value ~10 μF cm⁻², the 6 nm film shows metastable ferroelectricity (pink curve) and a weakly pinched loop with reduced P_r. The 4 and 2 nm HZO show very small remanent polarization and an ordinary dielectric behavior. When C_int becomes larger (20 μF cm⁻²), the situation is improved since the 6 nm HZO has stable ferroelectricity and a fully open hysteresis loop with high P_r.

Gibbs free energy landscapes and calculated polarization P – voltage V hysteresis for different ferroelectric thickness t_F for C_int = 10 μF cm⁻² (a, b) and C_int = 20 μF cm⁻² (c, d).

The 4 nm HZO is marginally stable ferroelectric (blue curve), with a nearly open ferroelectric loop and only the very thin 2 nm film shows metastable ferroelectricity resembling an ordinary dielectric behavior. Therefore, it can be concluded that by increasing the C_int above 20 μF cm⁻² it is possible to drive the system to a stable ferroelectric state even for very thin films <4 nm.

In MFM capacitors, an infinitely high C_int corresponds to perfect metal electrodes forming sharp interfaces^27,29 with the HZO (without dead layers). However as already briefly mentioned in the introduction, at least one of the electrodes should be less than perfect to allow for incomplete compensation of the polarization charge, which is essential for the operation of ultrathin FTJs. Therefore, in most cases of interest related to FTJs, C_int is expected to be small resulting in loop pinching. Nevertheless, severely pinched loops at small thickness <6 nm can be recovered by injecting charge during field cycling, as described in the following section.

Field cycling (“wake-up”) in the framework of Landau-Ginzburg-Devonshire theory

Charge injection³⁵ can occur during field cycling since pre-existing defect states are filled with trapped electrons which could tunnel through interfacial layer barriers, if such barriers exist. This mechanism (Fig. 6) provides efficient screening of polarization charges mitigating E_dep thus stabilizing polarization. Both, polarization P and applied bias V, control the states available for trapping and consequently the total trapped charge.

a Unpolarized MIFIM with charge neutral interfaces showing flat bands in equilibrium. b Polarized MIFIM in equilibrium (V = 0) showing depolarization and interfacial fields E_dep and E_int respectively. P is the polarization and Φ is the electron potential lowering at the interface due to the positive polarization charges at the interface on the right. c The MIFIM at a finite bias V establishing a field parallel to the polarization. V_int is the voltage drop at the interfacial layers resulting in a total Φ+V_int electron potential lowering at the interface on the right. The potential lowering creates a trapped charge −|σ | and +|σ | at the opposite interfaces due to the set of trap levels, indicated by red (green) squares, moving below (above) the Fermi energy level E_F.

Although the analysis presented below is made for the case of identical I and M-layers on either side (Fig. 6a), with appropriate adaptations it could be valid more generally for a structure with different insulator (I) and metal (M) layers on the opposite sides including the cases where one I-layer is missing or one M-layer is a semiconductor. As a first approximation, a constant energy density of states is considered here.

Once established (Fig. 6b), the polarization lowers the potential for electrons at the interface by an amount Φ proportional to P according to Φ = (t_F/2)|E_dep | = t_FP/2ε₀ε_F(1 + C_int/C_F). As a consequence, a portion of interface states within an energy ΔE_Φ = eΦ moves below E_F and becomes accessible to electrons for trapping.

An applied external bias V further lowers the potential (Fig. 6c) at the interface by an amount V_int = V/2(1 + C_int/C_F) making more states accessible to electrons in an energy interval ΔE_V = eV_int. The total states accessible for charge trapping are within the energy interval ΔΕ = e(Φ + V_int) therefore, the total trapped charge, in absolute value, is

In Eq. (5), n_tr is the trapped electrons density per unit energy which is determined by solving the Shockley Read Hall^36,37 rate equation ∂n_tr/∂t = c_n(D_it −n_tr) −e_nn_tr. D_it is the density of interface states per unit energy and c_n, e_n are the capture and emission rates³⁸.

In Fig. 7a the trapped carrier density per unit energy is plotted as a function of field cycles and for two different D_it, while the details near the origin are shown in Fig. 7b. The carriers are captured in the first half of the cycle, while a smaller portion is emitted back to the electrodes during the second half of the cycle. A positive balance of carriers remains captured after each complete cycle so gradually the trapped density increases and finally saturates at a value that depends on the D_it: a larger D_it allows for more trapped charge density as can be seen from the comparison between the green and red curves in Fig. 7a.

a The carrier density per unit energy n_tr as determined from the Shockley Read Hall (SRH) rate equation. Density of interface traps D_it is the interface energy density of states. b Detail near the origin (marked by the square in (a)), showing the trapping and de-trapping in each cycle. c The total charge σ trapped in all interface energy states which are made available by polarization P and bias V for D_it = 1.4 × 10¹³ eV⁻¹ cm⁻². d Detail at the origin (marked by square in (c)) showing separately the contribution of polarization (red) and voltage (blue) to the total trapped charge.

The corresponding total charge density σ captured in all energy levels made available by P and V is calculated using Eq. (5) and is plotted in Fig. 7c. The detail at the origin in Fig. 7d shows the different contributions from the polarization P (red) and the bias V (blue) according to Eq. (5) as well as the total σ (black). It should be noted that σ has always opposite sign to both, the polarization charges and the bias therefore it changes sign in each half period when P and V also change sign. The two interfaces have charges of opposite sign which alternates in each half period. Setting a relatively high value of D_it  ~ 1.4 × 10¹³ eV⁻¹ cm⁻², the calculated charge is accumulated to both interfaces increasing in amplitude and saturating after about 160 cycles at a value of about 10 μC cm⁻² which is a substantial fraction of the polarization charge (~20 μC cm⁻²) (Fig. 7c).

This indicates that after cycling, the polarization charge is sufficiently screened, thus minimizing the DF and stabilizing ferroelectricity.

A large D_it determines the screening efficiency of the polarization charges and ensures the opening of the ferroelectric loop. On the other hand, e_n (consequently c_n) determine how fast the traps are filled up to the saturation value. Higher e_n (and c_n) ensures faster filling so wake-up and loop opening is easier occurring only after a small number of cycles.

Introducing the contribution of trapped charge σ (Eq. (5) and Fig. 7c)) in the expression for the electric field E_F (Eq. (2)), additional P² and P terms appear in the Gibbs free energy (Eq. (3)) and L-K equations (Eq. (4), and Supplementary Fig. 5)), respectively, that counteract the depolarization terms and make ferroelectricity stable again alleviating the pinched characteristics after cycling is completed. The results are summarized in Fig. 8 above.

a Experimental polarization P – voltage V curves, (b) polarization P – voltage V simulation and (c) Gibbs free energy landscape for different number of cycles. The simulation is performed with D_it = 1.4 × 10¹³ eV⁻¹ cm⁻².

Although a direct fitting of the experimental curves is not possible due to non-idealities in the experimental data, a qualitative agreement between experimental data (Fig. 8a) and simulation (Fig. 8b) is observed. Using D_it = 1.4 × 10¹³ eV⁻¹ cm⁻² and e_n = 20 sec⁻¹ the simulation reproduces the trend of the P-V curves which shows a progressive opening of the hysteresis loop as the number of cycles increases. The loop opening is accompanied by a corresponding transition from a metastable ferroelectric state (pristine, black curve) to a stable ferroelectric state (160 cycles, pink curve) as seen from the Gibbs free energy plots in Fig. 8c.

It should be noted that both D_it and e_n are compatible with experimental observations in the literature and in the present work. Briefly, we refer to our electrical admittance spectroscopy (Supplementary Note 6): the D_it value at Ge/HZO interface is estimated from ac conductance losses to be 0.93–3.5 × 10¹³ eV⁻¹ cm⁻² (Table S1), which is very similar to the D_it value deduced from the simulation results in Fig. 8. It should be noted that a D_it which is an order of magnitude smaller (see Supplementary Note 7), is insufficient to open the loop inducing only marginal changes to the pristine curves. The low D_it could explain the behavior of some of the thinnest samples where the pristine, pinched characteristics do not recover after extensive cycling (See Fig. 2, and Supplementary Fig. 7).