Dopants in Atomic Layer Deposited HfO2 Thin Films

CHAPTER 3.1

Dopants in Atomic Layer Deposited HfO2 Thin Films Min Hyuk Parka, Tony Schenk, Uwe Schroeder NaMLab gGmbH, Dresden, Germany

3.1.1 Introduction In 2011, HfO2 became the first fluorite-type oxide material in which the ferroelectricity was induced by Si doping [1]. After this original report, various dopants such as Zr [2, 3], Y [4], Al [5], Gd [6, 7], Sr [8], and La [9–11] have been reported to induce ferroelectricity in atomic layer deposited HfO2 thin films. Moreover, Polakowski and Mueller reported that the ferroelectric properties can be observed even in undoped HfO2 thin film [12]. The structural origin of this unexpected ferroelectricity in fluorite-type HfO2 is the formation of the noncentrosymmetric orthorhombic phase with a space group of Pca21, which cannot be observed in the phase diagram of bulk HfO2 [13]. Although the ferroelectric doped HfO2 thin film was most frequently deposited using atomic layer deposition (ALD), the deposition technique is not limited to only ALD. There have been various reports with other deposition techniques such as chemical vapor deposition [14], physical vapor deposition [15–18], chemical solution deposition [19, 20], and pulsed laser deposition [21, 22], which will be discussed in Chapters 3.3, 3.4, and 4. Although there exist quantitative differences between the ferroelectric properties in the aforementioned previous studies with various dopants and deposition techniques, the ferroelectric properties in nanoscale undoped or doped HfO2 thin film are commonly observed, which could be ascribed to the formation of the noncentrosymmetric orthorhombic phase. In this chapter, the ferroelectricity in atomic layer deposited doped HfO2 thin films is reviewed based on the previous studies. In Section 3.1.2, the effects of doping concentration and dopant species on the structure and ferroelectric properties of HfO2 thin films are described. Although the ferroelectric phase could be formed even in undoped HfO2 a

Min Hyuk Park is currently at the School of Materials Science and Engineering, Pusan National University, Busan, Republic of Korea.

Ferroelectricity in Doped Hafnium Oxide https://doi.org/10.1016/B978-0-08-102430-0.00005-X

© 2019 Elsevier Ltd. All rights reserved.

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thin films by decreasing film thickness below 7 nm [12], in the frequently used
10 nm film thickness, a certain concentration range of doping can induce robust ferroelectricity in HfO2 thin films. Section 3.1.2 concentrates on doped HfO2 thin films while the ferroelectricity in undoped HfO2 thin films is intensively discussed in Chapter 3.5, which focuses on the effect of film thickness. The effects of doping concentration are examined for various dopants. Here, Si and La are chosen as examples for dopants of smaller and larger ionic radius than Hf, respectively, with comprehensive datasets available in previous studies [23–25]. Other dopants such as Al, Y, Gd, and Sr are included to demonstrate the generality of the derived trends as well as to emphasize additional aspects. In Section 3.1.3, the effects of annealing temperature on the structure and ferroelectricity in HfO2 thin films are reviewed. Generally, the doped HfO2 films are amorphous in an as-deposited state, and a subsequent annealing process is required to crystallize the as-deposited film into the ferroelectric orthorhombic phase. The annealing temperature is a critical factor that affects the relative fractions of crystalline phases and the resulting ferroelectricity in polycrystalline doped HfO2 thin films [7, 25–28]. Furthermore, the interface between doped HfO2 and electrodes can be strongly affected by the thermal budget during the crystallization annealing process. Especially, the (partial) oxidation of nitride electrodes such as TiN and TaN can influence the ferroelectric switching in doped HfO2 thin films.

3.1.2 Effect of Doping Concentration on Ferroelectric Doped HfO2 The effects of doping concentration on the crystalline structure and relative fractions of crystalline phases are discussed first, and La-doped HfO2 was chosen as an example system to show a general trend in doped HfO2 thin films. Fig. 3.1.1A shows the grazing incidence X-ray diffraction (GIXRD) patterns of 14 nm-thick La-doped HfO2 films with various doping concentrations from 4.4 to 34 cat% and a 30-nm-thick 8 cat% Si-doped HfO2 film deposited on TiN electrodes with reference patterns of four different crystalline phases of HfO2 (P21/c monoclinic, Pca21 ferroelectric orthorhombic, Fm3m cubic, and P42/nmc tetragonal phases). The details of the sample fabrication process can be found elsewhere [11, 23]. For the case of 4.4 cat% doped HfO2 film (black curve in Fig. 3.1.1A), the intensities of diffraction peaks from the monoclinic phase are high while weak diffraction peaks from the orthorhombic, the tetragonal, or the cubic phase

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Fig. 3.1.1 (A) GIXRD patterns for La-doped HfO2 and Si-doped HfO2 films for various doping concentrations. (B) The aspect ratio (2a/(b + c) for the orthorhombic phase) of the unit cell and the unit cell volume as a function of La content. Curves are just guides for eyes. (C) Relative phase fractions as determined by refinement of GIXRD patterns from (A). Curves are just guides for eyes (D) unit cell volume for Si, Al, Gd, and La doping in HfO2 as a function of ionic radius for the orthorhombic phase.

can also be observed. With increasing La content to 9.9 cat%, however, the intensity of the 111o/011t/111c diffraction peak at
30.5 degree becomes stronger than those of the 111m or 111m diffraction peaks at 28.5 and 31.5 degree. The hklx diffraction peak refers to the hkl diffraction peak from the x-phase, where x can be m (monoclinic), o (orthorhombic), t (tetragonal),

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or c (cubic). With increasing La content up to 12 cat%, the relative intensities of the 111m and 111m diffractions peaks decrease, and no 111m or 111m diffraction peaks could be observed above 12 cat% La doping [11]. Within this La doping concentration range, another interesting change in GIXRD patterns can be observed near 51 degree 2θ, where 220o/202o/ 022o, 112t/200t, or 220c diffraction peaks are expected. For 9.9 cat% La (red curve in Fig. 3.1.1A), the overlap of diffraction peaks can be clearly observed with a rather broad diffraction peak at 51 degree 2θ. However, when the La content increases to 34 cat% (blue curve in Fig. 3.1.1A), the 51-degree diffraction peak becomes narrower, indicating the existence of just a single diffraction peak rather than an overlap of multiple reflections. From this result, the crystalline phase of the La-doped HfO2 thin films with >15 cat% doping is believed to be the cubic phase instead of the tetragonal phase. It should be noted that this result was different from other dopants such as Si and Al. For the case of 8 cat% Si-doped HfO2 (magenta curve in Fig. 3.1.1A), the overlap of the two diffraction peaks from the tetragonal phase at 51 degree as well as a split-up of the peak at around 35 degree 2θ could be clearly confirmed, suggesting that the dominant crystalline phase of heavily Si-doped HfO2 thin films is the tetragonal phase [23]. Furthermore, the diffraction patterns in the 2θ range of 80–90 degree are also quite different for the orthorhombic, tetragonal, and cubic phases, as shown in previous studies [3–5]. Such differences can also be observed between the GIXRD patterns of 9.9 and 34 cat% La-doped HfO2 and 8 cat% Si-doped HfO2 films. The effect of the dopant species on the crystalline structure of doped HfO2 thin film will be discussed later in this section. It has been known that the phase identification in doped HfO2 thin films using XRD is highly challenging, although Sang et al. [13] experimentally proved that the ferroelectric phase is the Pca21 orthorhombic phase by using scanning transmission electron spectroscopy (STEM). In 2017, Park et al. suggested that the changes in the aspect ratio and the unit cell volume can give hints to examine the crystalline phases in ultrathin doped HfO2 films, and their changes were consistent with the changes in relative fractions of the crystalline phases analyzed using the Rietveld refinement [24]. Fig. 3.1.1B shows the changes in the aspect ratio (2a/(b+c) for the orthorhombic phase and c/a for the tetragonal or cubic phase) and unit cell volume. The method to estimate the lattice parameters can be found in a previous study [24]. The aspect ratio decreases from 1.03 to 0.98 with increasing La content, which is more evidence of the dominant crystalline phase changing from the orthorhombic to the cubic phase. The aspect ratio

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of the cubic phase is smaller than the ideal value 1, which could be attributed to a large biaxial in-plane tensile stress in doped HfO2 thin film [11, 29–31]. In the La-doped HfO2 thin film in this chapter, a large tensile stress of
2 GPa was experimentally confirmed using the sin2Ψ-method based on Bragg-Brentano XRD and the simplified but not necessarily true assumption that the out-of-plane direction is strain-free [11]. The stress in doped HfO2 thin films is discussed in Chapter 3.5 in more detail. Similar trends in the aspect ratio change with varying doping concentrations could be observed for Si, Al, and Gd-doped HfO2 thin films in a previous study [24]. Another interesting result can be found from the changes in the unit cell volume as depicted in Fig. 3.1.1B. In correspondence to that, Fig. 3.1.1C shows the phase composition of different La concentrations derived from quantitative phase analysis based on the Rietveld refinement. The dominant crystalline phase changes from the monoclinic to the orthorhombic to the cubic phase with increasing La content. The ionic radius of La is larger than that of Hf by 25% [11, 32, 33], so the increase in the unit cell volume is expected when the La content increases. However, the unit cell volume unexpectedly decreases with increasing La content up to 15 cat%, as shown in Fig. 3.1.1B. A similar trend could be observed for Si, Al, and Gd-doped HfO2 thin films [23], and such a trend originates from the phase change from the orthorhombic to the tetragonal or the cubic phase. Generally, dopants larger than Hf are expected to increase the unit cell volume of HfO2, whereas the opposite is expected for dopants smaller than Hf. However, dopants with both larger and smaller radius than Hf decreased the unit cell volume of HfO2 with increasing dopant content. For the La content higher than 15 cat%, on the other hand, the unit cell volume increases with increasing La content. In this La concentration region, the dominant crystalline phase is the cubic phase (Fig. 3.1.1C), and the lattice parameters simply increase by further dopant incorporation (Fig. 3.1.1B). The unit cell volume of the orthorhombic phase at the doping concentration for the highest orthorhombic phase fraction is plotted as a function of the ionic radius of dopants in Fig. 3.1.1D. The unit cell volume of the orthorhombic phase generally increases with increasing ionic radius. However, the unit cell volume of La-doped HfO2 was slightly smaller than that of the Gd-doped case, and the origin of that behavior is not clearly understood yet [11]. However, there might be other factors that can potentially affect the unit cell volume of atomic layer deposited HfO2. Transmission electron microscopy (TEM) has been another powerful tool to examine the crystalline structure of doped HfO2 thin films [13, 34].

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However, several special techniques, which will be discussed in Chapter 7.2, are required to detect very small differences in the atom positions in the lattice with different phases in doped HfO2 thin films [13, 34]. Fig. 3.1.2A shows the cross-sectional scanning transmission electron microscopy (STEM) image of 11 cat% La-doped HfO2 thin film sandwiched between TiN electrodes, and its inset figure depicts the Fourier transformed pattern of the red square region. The dominant crystalline phase is the ferroelectric orthorhombic phase [11]. A weak relationship between the orientation of the TiN electrode and La-doped HfO2 films could also be observed, similar with a previous report on Gd-doped HfO2 [11, 35]; this local orientation relationship is discussed in the later part of this section related to Fig. 3.1.7. Fig. 3.1.2B shows an enlarged STEM image of the interfacial region between the La-doped HfO2 and the TiN electrode. The relaxation of lattice parameters of the bulk orthorhombic phase near the TiN electrode could be observed, as shown in Fig. 3.1.2B (enclosed by dash lines) [11]. A similar result could also be confirmed for Si and Gd-doped HfO2 thin film in previous studies [23, 35]. For the cases of Si and Gd-doped HfO2, the tetragonal phase could be assigned to the interfacial layer adjacent to the electrodes [23, 35]. The existence of the nonferroelectric interfacial layer was reported to be an origin of the pinched hysteresis in the as-deposited state, and the phase transition to the ferroelectric orthorhombic phase during repetitive field cycling is accepted as a potential origin of the wake-up effect in doped HfO2 thin films [34, 36–38]. The wakeup effect during field cycling in doped HfO2 thin films is dealt with in Chapter 9.2 of this book.

Fig. 3.1.2 Orthorhombic HfO2 of the 10 cat% La (A) with [101]-zone and (B) relaxation of the orthorhombic bulk HfO2 phase at the interface to the TiN electrode. (The figure was reproduced from U. Schroeder, C. Richter, M. H. Park, T. Schenk, M. Peši
c, M. Hoffmann, F. P. G. Fengler, D. Pohl, B. Rellinghaus, C. Zhou, C.-C. Chung, J. L. Jones, T. Mikolajick, Lanthanum doped hafnium oxide: a robust ferroelectric material, Inorg. Chem. 57 (2018) 2752–2765 with permission from ACS publications).

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The oxidation of nitride electrodes such as TiN and TaN is another critical extrinsic factor that can influence the ferroelectric properties of the doped HfO2 thin films. Fig. 3.1.3A shows the time-of-flight secondary ion mass spectrometry (TOF-SIMS) result of La-doped HfO2 thin films annealed using rapid thermal annealing (RTA) for 20 s at 800°C in an N2 atmosphere. At the interfacial region between TiN and La-doped HfO2, the (partial) oxidation of the TiN electrode can be clearly confirmed. A similar result was reported for Hf0.5Zr0.5O2 thin film annealed at a temperature as low as 500°C [39]. The oxidation of the TiN electrode can be even more serious with a higher thermal budget. Fig. 3.1.3B shows the TOF-SIMS result of a TiN/La:HfO2/TiN capacitor after a hightemperature in situ x-ray diffraction study [11]. For this analysis, the capacitor was heated to 900°C within a time scale longer than 10 h, and the formation of rutile TiO2 could be confirmed from the XRD patterns [11]. As shown in Fig. 3.1.3B, the TiO 2 ion counts are much higher than  the TiN ion counts. Moreover, the diffusion of Hf into the TiN layer and that of Ti into the HfO2 layer is evident when comparing Fig. 3.1.3B to Fig. 3.1.3A. It should be noted that Fig. 3.1.3B shows an extreme example of the high thermal budget, but this result suggests that there are driving forces for intermixing of TiN and HfO2. Thus, RTA for crystallization of doped HfO2 thin films should be optimized to crystallize the doped HfO2 thin films into the ferroelectric phase, but suppress the intermixing of the HfO2 and the electrode material. The oxidation of the TiN layer can strongly degrade the properties of ferroelectric doped HfO2 thin films by shifting the polarization-electric field curves [39], creating a

Fig. 3.1.3 (A) TOF-SIMS after 800°C 20 s anneal. (B) TOF-SIMS after slow 0.2 K/s ramp during temperature dependent GIXRD measurement.

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depolarization field and increasing the required electric field for ferroelectric switching [39, 40], or making more trap sites for electric carriers [41]. The above-discussed changes in the crystalline structure and interfacial properties with varying doping concentration strongly influence the electrical properties of the fabricated metal-ferroelectric-metal capacitors. Fig. 3.1.4A and B show the P-E curves of 10 nm-thick Si-doped HfO2 films with low (2.0–4.1 cat%) and high (4.5–8.7 cat%) doping concentrations. As shown in Fig. 3.1.4A, the 2 cat% Si-doped HfO2 shows only negligible hysteresis similar to linear dielectrics, which can be attributed to the dominant monoclinic phase. With increasing Si content above 3.5 cat%, on the other hand, characteristic ferroelectric hysteresis could be observed. Within a Si content range from 3.5 to 4.1 cat%, robust ferroelectricity with Pr larger than 10 μC/cm2 could be experimentally confirmed. With further increasing Si content beyond 4.5 cat%, on the other hand, a significant change in the shape of the P-E curve could be observed, as seen in Fig. 3.1.4B. The Pr of 4.5 cat% Si-doped HfO2 was negligible, and an antiferroelectric-like double hysteresis could be observed. Such a double hysteresis can be attributed to the field-induced phase transition between the tetragonal and a polar orthorhombic phase [42–45], and details of this phenomenon will be discussed in Chapter 10.2. Fig. 3.1.4C shows the P-E curves of 14 nm-thick La-doped HfO2 films with various doping concentrations. The Pr of pure HfO2 is almost zero, and the Pr increases with increasing La content, which is attributed to the increasing orthorhombic phase fraction as presented in Fig. 3.1.1C. With 5 cat% La doping, the Pr is already larger than 15 μC/cm2, and Pr values beyond 25 μC/cm2 could be observed for 10–13 cat% La-doped HfO2 films. The largest Pr value ever reported in the doped HfO2 thin film is 45 μC/cm2 with La doping [9], but this value could not be reproduced. It should be mentioned that the Pr >25 μC/cm2 in La-doped HfO2 is higher than for many other doped HfO2 films. With increasing La content up to 20 cat%, only negligible Pr could be observed, and this is consistent with the increasing fraction of the cubic phase in Fig. 3.1.1C. Being different from Si-doped HfO2, the antiferroelectric-like properties are not clearly observed in La-doped HfO2. To date, the antiferroelectric properties have been reported for Si- [1, 45] and Al-doped HfO2 [5, 46] and Hf1 xZrxO2 [3, 44] thin films. Schroeder et al. suggested that the antiferroelectric properties can be observed in the HfO2 films doped with dopants having an ionic radius smaller than Hf [47]. However, the reason for the absence of the antiferroelectric properties in HfO2 thin films doped with dopant larger than Hf is not clearly elucidated yet.

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Fig. 3.1.4 (A)+(B) Polarization hysteresis for pristine samples with different Si dopant contents from 2 to 8.5 cat%. (C) Polarization hysteresis for samples after wake-up cycling with different La dopant contents from 6 to 20 cat%. (D) Remanent polarization Pr values after wake-up for dopants smaller Hf: Si and Al. (E) Remanent polarization Pr values after wake-up for Gd-, Y-, Sr-, and La-doped HfO2 with different dopant content (dopants larger than Hf ).

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Fig. 3.1.4D and E show the changes in Pr as functions of doping concentration for dopants smaller (Si and Al) or larger (Gd, La, Y, and Sr) than Hf. The overall trend is similar for both cases with dopants smaller and larger than Hf. The Pr is negligible for undoped or lightly doped HfO2, and the Pr increases with increasing doping concentration up to a specific value for each dopant. Such a trend can be attributed to the changes in relative fractions of various crystalline phases, as shown in Fig. 3.1.1C and other previous studies [11, 24]. The doping concentration range for a robust ferroelectricity is strongly dependent on dopant species. For Si- and Al-doped HfO2 thin films, the full width half maximum (FWHM) of the peaks of the Pr-dopant content curves are smaller than 2 cat%, suggesting that the ferroelectric orthorhombic phase can be observed in a narrow doping concentration range. For Gd-, Y-, and Sr-doped HfO2 thin films, on the other hand, the FWHM of the Pr-doping concentration curve is about 5%. Furthermore, the FWHM of the Pr-dopant content curve of La-doped HfO2 is as large as 10 cat%, which is five times larger than those of Siand Al-doped HfO2 cases. Such a broad doping concentration range for robust ferroelectricity can be attributed to the previous computational simulation result [48]. Batra et al. examined the effect of various dopant species on the relative free energy of metastable crystalline phases of HfO2, and reported that Sr, Ba, Ca, and lanthanides are expected to efficiently stabilize the ferroelectric orthorhombic phase [48]. The computational simulation result is reviewed in Chapter 6 of this book. The dielectric constant k of doped HfO2 thin film is also strongly affected by the dopant species. Fig. 3.1.5A shows the variations in dielectric constant as a function of doping concentration for Si-, Al-, La-, and Sr-doped HfO2

Fig. 3.1.5 (A) Change of the dielectric constant k for different doping concentration for Si-, Al- and La-doped HfO2 thin films. (B) Pristine remanent polarization Pr values as a function of the orthorhombic phase fraction as determined by the Rietveld refinement of GIXRD patterns.

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thin films. The maximum k value is different between the dopants with a smaller (Si and Al) and larger (La and Sr) ionic radius than Hf. The maximum k values of Si- and Al-doped HfO2 is 44 and 40, respectively, whereas those of La- and Sr-doped HfO2 are only limited to about 30. Lee et al. suggested that the dopants smaller than Hf can stabilize the tetragonal phase over the cubic phase while the dopants larger than Hf tend to stabilize the cubic phase [49]. A similar trend was also reconfirmed in other computational simulation works [50, 51]. According to the computational calculation by Materlik and coworkers [42], the k values of the orthorhombic, tetragonal, and cubic phases are in the range of 24.2–28.8, 24–56.9, and 36, respectively, and the average k values are 27.0, 45.9, and 36, respectively. Thus, the lower k values observed in HfO2 doped with larger dopants are well matched with the phase change from the orthorhombic to cubic while that observed with smaller ionic radius dopants can be attributed to the transition from the orthorhombic to the tetragonal phase. It should be noted that the dopants are expected to influence the chemical bonding state near dopants, and this will also influence the dielectric response via effective medium consideration [52–54]. The k value decrease observed in a rather high doping concentration region might be attributed to such an effect of dopants because the k values of the characterized dopant oxides are lower than that of HfO2. Thus, there might be competition between the k value increase with phase change from the orthorhombic to the tetragonal or cubic phase and the k value decrease with increasing dopant concentration. Such trends can be clearly observed for Si-, Al-, and Sr-doped HfO2 in Fig. 3.1.5A. For the case of La-doped HfO2, the k value decrease with increasing doping concentration is much weaker compared to the other three dopants, which can be attributed to a higher k value of La2O3 (30 [55]) than SiO2 (3.9 [55]), Al2O3 (9 [55]), and SrO (15 [56]). Fig. 3.1.5B shows the variations in the orthorhombic phase fraction as a function of pristine Pr which was measured before any wake-up field cycling. The orthorhombic phase fractions were analyzed using the Rietveld refinement, and the data were taken from previous studies [11, 24]. A clear linear relation can be confirmed for all the dopants in Fig. 3.1.5B, suggesting that the pristine Pr is determined by the relative fraction of the ferroelectric orthorhombic phase. The Pr expected for the 100% orthorhombic phase fraction can be estimated by extrapolating the linear fitting in Fig. 3.1.6B, and the estimated Pr values were 30 and 15 μC/cm2 for La and the other three dopants (Si, Al, and Gd). To achieve an idea of the maximum theoretical remanent polarization, two scenarios can be applied to give an indication of what is likely and what

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Fig. 3.1.6 Special texture of La:HfO2: (A) Comparison of GIXRD patterns for 10 nm thick films of the most ferroelectric hafnia compositions with different dopants and in mixture with zirconia. (B) Relative intensity ratio for the peaks in (A). GIXRD data for samples prepared as in Refs. [7, 8, 11, 23, 24, 27]. (C) PDF (#04-005-5597) reference pattern and change of peak intensity upon sample tilt. Ψ ¼ 0 means out of plane, Ψ ¼90 degree means in-plane orientation of the scattering vector and the corresponding lattice planes probed at the respective 2Θ angles. Indices of the lattice planes refer to the orthorhombic phase. 020/002 planes are more prominent in out-of-plane direction. (Panel (C) was reproduced from U. Schroeder, C. Richter, M. H. Park, T. Schenk, M. Peši
c, M. Hoffmann, F. P. G. Fengler, D. Pohl, B. Rellinghaus, C. Zhou, C.-C. Chung, J. L. Jones, T. Mikolajick, Lanthanum doped hafnium oxide: a robust ferroelectric material, Inorg. Chem. 57 (2018) 2752–2765 with permission from ACS publications).

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is less likely: A completely random orientation of the polar axis would result in 50% of the theoretical value for the spontaneous polarization [42, 57], that is, 25–28 μC/cm2 (compare to Chapter 6). Typically, the 111 reflex is very prominent in the measured GIXRD patterns (see Fig. 3.1.6A). However, a perfect 111-orientation would only result in values of 28–31 μC/cm2, depending on the lattice constants and assumed spontaneous polarization. (Here, lattice constants as extracted by Park et al. [24]) for 10 nm thick Gd:HfO2 films were chosen for the estimate). This means a strong 111 texture would not make a significant difference to a random orientation for this system. Moreover, this reflex is also the most intense in the powder diffraction files of which the intensity is
1 order of magnitude higher than the intensities of all other peaks. Thus, there is no strong indication for a strong preferential orientation in most of the hafnia-based films. Looking at the different diffractograms of La:HfO2 (Fig. 3.1.6A and B), though, there is a significantly higher peak at around 35.5 degree, which belongs to the 002/020 reflex. This peak loses intensity when the scattering vector is tilted from an out of plane (Ψ ¼0 degree) toward an orientation in the film plane (Fig. 3.1.6C) while the 111 peak around 30.5 degree increases. A 002orientation (polar axis out of plane) would indeed explain the enhanced Pr values compared to other dopants. The complete absence of the 200 peak at around 34.5 degree suggests that an in-plane orientation of the longest axis, the 100-axis, is highly preferable and there might be a causal relation to the found in-plane tensile strain (relative to the out-of-plane direction) concluded from the peak shifts upon tilting. However, for Si-doped HfO2, on the other hand, an almost random orientation with a very weak 111-texture was reported [29] and HfO2 films with many other dopants seem to show similar behavior (Fig. 3.1.6A and B). The origin of the different texture of the La-doped HfO2 film is not clearly understood yet, and it needs further study given that the in-plane tensile strain is rather similar. ALD precursors might play a role for that texture because thestrong effects of metal or nonmetal precursors on film properties are well known [58]. However, the Hf precursor TEMAHf is not different compared to the one used for the Si:HfO2 films described here. Also, other Hf precursors have been used, but no report of a special impact on texture exists. TEM work [7] gave rise to speculation about an orientation relationship between the TiN electrodes and the HfO2-based films. In addition, some relaxations of the HfO2 lattice toward the TiN electrodes were reported together with the formation of tetragonally distorted interfacial layers [34, 36]. Thus, a recent STEM study was devoted to this topic [59]. Fig. 3.1.7

62 Ferroelectricity in Doped Hafnium Oxide

Fig. 3.1.7 Examples of interface sites between TiN and hafnia samples. Detailed description of (A)–(D): see text. Top and bottom represent slightly shifted regions with different intensity scaling to account for the different intensities of the Ti and Hf columns (atomic number contrast) of the same image frame. (Used with permission from Everett Grimley. From E.D. Grimley, PhD Thesis, North Carolina State University, 2018.)

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shows examples of the results. Within the sampling capabilities of STEM, the analysis of multiple images was used to obtain a meaningful representation of the whole sample. In contrast to epitaxial single-crystalline films (see Chapter 4), misalignment of HfO2 grains toward the electron beam limits the number of regions that can be analyzed. For some regions, no clear epitaxial relation was found because the TiN grains were completely misaligned toward the beam (Fig. 3.1.7B). However, for the majority of interfaces, local epitaxial-like relations are found that can be classified as follows: The 111 TiN planes interact with 200, 020, and 002 planes of the monoclinic and orthorhombic HfO2 phases. This interaction occurs in two different ways: (1) by traditional epitaxial “stacking” (Fig. 3.1.7A) or a “pseudostacking” (Fig. 3.1.7C) of the planes, or (2) via an “angled interface” with continuous planes (Fig. 3.1.7D). The latter means that the TiN is not on zone axis for the STEM, but a direction and spacing of the planes translating into HfO2 is clearly visible. Further studies are required to clarify the role of the TiN texture on HfO2 texture as well as the impact of different dopants and why La seems to be a bit different from the others with respect to texture. The above results can only serve to identify which directions to go to further understand the complex structural interplay of the various factors in these ultrathin polycrystalline films. Another figure to compare the differently doped films to each other is the relaxed remanent polarization, which indicates the expected retention properties of the stack. The relaxed remanent polarization, Pr,relax, represents what fraction of the Pr measured during a triangular sweep can be retained after 1 s of waiting time [60]. The Pr,relax is generally strongly related to the depolarization field occurring in metal-ferroelectric-metal capacitors. Fig. 3.1.8A shows the changes in Pr,relax normalized by Pr of La-doped

Fig. 3.1.8 Normalized relaxed remanent polarization after 1 s Pr,relax (divided by remanent polarization Pr) for different (A) Si and (B) La content for the pristine sample and after wake-up cycling. The doped HfO2 films were cycled sufficiently to reach maximum Pr value for woken up states.

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HfO2 films with various doping concentrations before and after wake-up field cycling. The depolarization field in ferroelectric thin films is known to be governed by the relative fraction of nonferroelectric layers along the electrode normal direction at the electrodes or within the bulk of the ferroelectric layer. Thus, the highest normalized Pr,relax value is expected for the highest orthorhombic phase fraction. As shown in Fig. 3.1.8A, the largest normalized Pr,relax could be achieved for the 10 cat% La doping, and the La-doped HfO2 film could have the highest Pr value at the same La content, as seen in Fig. 3.1.4E, both before and after wake-up field cycling. Generally, the normalized Pr,relax is lower in the pristine state compared to the woken-up state. It was reported that the nonferroelectric layer can be transformed into the ferroelectric phase during wake-up field cycling, so the decrease in the nonferroelectric phase fraction during field cycling can increase the normalized Pr,relax. The normalized Pr,relax values after wake-up are almost 1 in a rather wide range of La content from 6 to 13 cat%, suggesting that a significant fraction of the nonferroelectric phase changed to the ferroelectric orthorhombic phase. However, this does not mean that all the nonferroelectric phase transformed into the orthorhombic phase within this La content range. There should be a remaining nonferroelectric phase even after the wake-up field cycling. However, a significant portion of the nonferroelectric layer neighboring the ferroelectric phase should have changed to the orthorhombic phase after field cycling. The details of the structural evolution during field cycling are intensively reviewed in Chapter 9.2. In Fig. 3.1.8B, the normalized Pr,relax of doped HfO2 films with various dopants was plotted as a function of the orthorhombic phase fraction at the composition of the largest Pr for each dopant. Interestingly, the normalized Pr,relax increases with increasing orthorhombic phase fraction, which is qualitatively consistent with the usual expectation of a linear trend. In this plot, the Si-doped HfO2 film shows the lowest normalized Pr,relax value with the lowest fraction of the orthorhombic phase among the four dopants. This result can be understood based on the fact that the Si is a very strong stabilizer of the tetragonal phase. The effect of dopant species on the relative free energy of the tetragonal and cubic phases was examined in several studies (Refs. [49–51], Chapter 6), and Si strongly decreases the free energy of the tetragonal phase compared to the cubic phase. It resulted in a very narrow doping concentration range for robust ferroelectricity, as shown in Fig. 3.1.4. Because the tetragonal phase can be stabilized with a rather low doping concentration, the maximum achievable orthorhombic phase

Dopants in Atomic Layer Deposited HfO2 Thin Films

65

fraction is also expected to be lower than that for the other dopants. The small ionic radius of the Si was suggested as an origin for why Si can strongly stabilize the tetragonal phase rather than cubic phase [49]. Park et al. [24] showed that a similar discussion is still valid for the competition between the tetragonal and the ferroelectric orthorhombic phases because the orthorhombic phase does not have metal-oxygen bonding as short as those in the tetragonal phase. Kuenneth et al. systematically examined the effect of four valent dopants including Si on the relative free energy of the orthorhombic and the cubic phases, and confirmed that Si tends to more strongly stabilize the tetragonal phase than the orthorhombic phase [61], which is in line with a previous result by Schenk and coworkers [8].

3.1.3 Effect of Annealing Temperature on Ferroelectric Doped HfO2 The as-deposited doped HfO2 films are mostly amorphous after the ALD process. Thus, a subsequent annealing process is essential to crystallize the film into the ferroelectric orthorhombic phase. Moreover, as discussed in the previous section, the annealing process can also affect the interface between the doped HfO2 film and the electrode. Thus, understanding the effect of the annealing process is also an important task for ferroelectric doped HfO2 thin films. Fig. 3.1.9A shows the changes in normalized polarization before and after wake-up field cycling for Si- and La-doped HfO2 thin films. For both dopants, the Pr increases with increasing annealing temperature in general, both before and after wake-up field cycling.

Fig. 3.1.9 (A) Remanent polarization Pr versus anneal temperature for a pristine in comparison to a sample after wake-up; Triangular field cycles at 4 MV/cm and 100 kHz to breakdown for different annealing temperatures. (B) Endurance field cycling as a function anneal temperature: Triangular field cycles at 4 MV/cm and 100 kHz to breakdown for different annealing temperatures.

66

Ferroelectricity in Doped Hafnium Oxide

The increase in Pr was larger for Si-doped HfO2 compared to La-doped HfO2, especially between 725°C and 800°C. The difference between Si and La can be understood based on the two potential mechanisms. First, the crystallization temperature is proportional to the difference between the ionic radius of the dopant and Hf, so Si-doped HfO2 has the highest crystallization temperature among the reported dopants, which can induce ferroelectricity in HfO2 [29]. Second, the Si-doped HfO2 film has a lower maximum fraction of the orthorhombic phase compared to the La-doped HfO2 thin film when they are annealed at the same temperature, as shown in the previous section. Because Si is a stronger stabilizer of the tetragonal phase, there should be a higher fraction of that phase compared to the La-doped HfO2 thin films. However, the tetragonal phase can be transformed to the ferroelectric orthorhombic phase by increasing annealing temperature. As a result, the Pr of Si-doped HfO2 thin film can significantly increase by increasing the annealing temperature from 725°C to 800°C. It should be noted that an annealing temperature higher than 800°C was generally used to crystallize the ferroelectric Si-doped HfO2 thin films [23, 25]. On the other hand, it was reported that La-doped HfO2 can show strong ferroelectricity after annealing at a rather low temperature [62]. Depending on the targeted integration into an existing CMOS route, each of the two can be beneficial. The endurance of doped HfO2 film is also strongly influenced by the annealing temperature. The general failure mechanism of the doped HfO2 thin films is a hard breakdown with a significant increase in the leakage current. The generation and possible percolation of defects such as oxygen vacancies are considered as a potential origin of such a failure mechanism. This point was well studied, as will be discussed in detail in Chapter 9.2. Fig. 3.1.9B shows the number of field cycles until the breakdown of the doped HfO2 thin films for the cases of Si and La doping. The endurable field cycle number decreases with increasing annealing temperature. It was suggested that TiN scavenges oxygens from doped HfO2 thin films [34, 36, 39, 63], and such a scavenging effect should become stronger with increasing annealing temperature. Thus, the total amount of the oxygen vacancies in the doped HfO2 thin film increases for higher annealing temperature, and accordingly, a permanent leakage current path formed by defects such as oxygen vacancies might be formed after a smaller amount of field cycling. The degradation of endurance with increasing annealing temperature is more severe in Si-doped HfO2 compared to La-doped HfO2 thin film. Although the exact reason for such a difference cannot be understood currently, it might be attributed to the different valence number of Si (+4) and

Dopants in Atomic Layer Deposited HfO2 Thin Films

67

La (+3). It is known that the trivalent dopant-oxygen vacancy complex can be formed to reduce the free energy near the dopant in the HfO2 matrix [64]. Therefore, the accumulation of oxygen vacancies in La-doped HfO2 might be retarded compared to the Si-doped HfO2 thin film. The crystalline structure and relative fractions of phases can also be affected by the annealing temperature. This can be identified from the GIXRD patterns of the variously doped HfO2 thin films. The details of the sample fabrication process can be found elsewhere [28]. Being similar to the discussion in the previous section, the changes in the aspect ratio, the unit cell volume, the relative phase fractions, and the pristine Pr values are examined. Fig. 3.1.10A and B show the variations in the aspect ratios (2a/(b+c) for the orthorhombic and c/a for the tetragonal phase) and unit cell volume of 10 nm-thick Gd-doped HfO2 thin films with changing

Fig. 3.1.10 The change of (A) aspect ratio, (B) unit cell volume, (C) pristine Pr, and (D) orthorhombic phase fraction with varying doping concentration for Gd-doped HfO2 thin films.

68

Ferroelectricity in Doped Hafnium Oxide

annealing temperature from 650°C to 1000°C using RTA. The aspect ratio and the unit cell volume of the Gd-doped HfO2 films increase with increasing annealing temperature near the doping concentration for the large Pr. The strongest changes in the aspect ratio and the unit cell volume could be observed at the composition at the phase boundary between the orthorhombic and tetragonal phases, which refers to 2.8 cat% for Gd-doped HfO2. A similar trend could be observed in Si- and Al-doped HfO2 thin films, which can be found elsewhere [28]. Fig. 3.1.10C shows the pristine Pr values of Gd-doped HfO2 thin films annealed at various temperatures ranging from 650°C to 1000°C. The Pr of undoped HfO2 is almost zero because the dominant crystalline phase is the monoclinic phase. With increasing dopant content to a certain value, the pristine Pr increases with the increasing orthorhombic phase fraction. After the maximum Pr value, on the other hand, the Pr decreases with increasing doping concentration, which is again attributed to the phase transition from the orthorhombic to the tetragonal phase. At the doping concentration for the maximum Pr, the pristine Pr values generally increase with increasing annealing temperature. Such an annealing temperature effect on the pristine Pr value is most evident at the boundary between the orthorhombic and tetragonal phases. Quantitative phase analysis was conducted using the Rietveld refinement. The details of the Rietveld refinement can be found in a previous study [24]. Fig. 3.1.10D shows the changes in the relative orthorhombic phase fractions of Gd-doped HfO2 thin films with varying doping concentration. The changes in the orthorhombic phase fraction with varying annealing temperature and doping concentration are well matched with the pristine Pr value changes in Fig. 3.1.10C. The maximum orthorhombic phase fraction does not show strong dopant material dependences [28], but it generally increases with the increasing ionic radius of dopant when the result for La-doped HfO2 is considered together. The maximum orthorhombic phase fraction was the smallest for Si and the largest for La, as already discussed in the previous section. However, the pristine Pr does not increase with the increasing ionic radius of the dopant. It should be noted that the pristine Pr values can also be affected by factors such as film texture and thickness of the interfacial layer [11, 39]. On the other hand, the La-doped HfO2 showed an even higher orthorhombic phase fraction in the previous study [11]. The annealing temperature as well as the dopant species should affect the doping concentration range for the maximum orthorhombic phase formation with large Pr value. Especially at the

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boundary between the orthorhombic and the tetragonal or the cubic phases, the effect of the annealing temperature is strong. Within this dopant content range, the aspect ratio, the unit cell volume, the orthorhombic phase fraction, and the pristine Pr significantly increase with increasing annealing temperature. It is believed that the small free energy difference between the orthorhombic and the tetragonal or cubic phases can be overcome by increasing the annealing temperature. From the result above, it is conceived that the high annealing temperature further stabilizes the orthorhombic phase or kinetically accelerates the phase transition from the tetragonal or the cubic phase to the orthorhombic phase. However, the exact mechanism behind this observation is not clearly elucidated yet, and it requires further study. Park et al. recently examined the structural evolution in the doped HfO2 thin film using in situ XRD in Si-, Al-, Gd-, and Sr-doped HfO2 thin films [28]. They observed a lattice parameter increase at a specific temperature for all the dopant cases, and this phenomenon is believed to be strongly related to the annealing temperature effect. Nonetheless, the origin of the unit cell parameter increase process is not clearly understood yet. The structural evolutions observed from hightemperature XRD in Si-, Al-, Gd-, and Sr-doped HfO2 thin films are discussed in Chapter 5.1 in more detail.

3.1.4 Conclusion In this chapter, the current status of ferroelectricity in atomic layer deposited doped HfO2 thin films was comprehensively reviewed. Especially, the effects of the doping concentration, the dopant species, and the annealing temperature were intensively studied. The changes in the relative fraction of the ferroelectric orthorhombic phase can be analyzed from the structural changes detected from the XRD techniques. For atomic layer deposited HfO2 thin films doped with various dopants, general characteristic changes in XRD patterns could be observed, and these could be successfully related to the variations in the electrical characteristics. With increasing doping concentration, the dominant crystalline phase changes from the monoclinic to the orthorhombic to the tetragonal (or cubie) phase. For dopants smaller than Hf (such as Si and Al), the tetragonal phase is formed with a field-induced phase transition in the high doping concentration range. For dopants larger than Hf, the cubic phase could be formed for sufficiently high dopant contents. The doping concentration range for high Pr is strongly affected by the dopant species, and the width of

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the dopant content range is fairly matched with the free energy changes induced by doping in computational simulation works. Currently, La is considered highly promising for practical applications because of its wide composition window and large Pr values both from simulations and experiments. However, further studies are required to clearly understand the effect of dopant species. The effect of the annealing temperature on ferroelectric doped HfO2 can also be understood based on the structural changes as well as the interfacial chemical reactions. The increase in annealing temperature broadens the doping concentration range for robust ferroelectricity, and this can be attributed to the structural changes with the increasing aspect ratio and volume of the unit cell. On the other hand, the increase in oxygen vacancy concentration in the doped HfO2 thin film with an increasing annealing temperature degrades its endurance, possibly due to oxygen scavenging by the TiN electrode. It is believed that there is a trade-off between the orthorhombic phase fraction and interfacial layer formation when the annealing temperature is controlled. Thus, optimization of annealing temperature is another critical task for the ferroelectric doped HfO2 thin films.

Acknowledgments M. H. Park is supported by a Humboldt postdoctoral fellowship from the Alexander von Humboldt Foundation. T. Schenk acknowledges the German Research Foundation (DFG) for funding part of this work in the frame of the project “Inferox” (project no. MI 1247/11-2).

References [1] T.S. B€ oscke, J. M€ uller, D. Br€auhaus, U. Schr€ oder, U. B€ ottger, Ferroelectricity in hafnium oxide thin films, Appl. Phys. Lett. 99 (2011) 102903. [2] J. M€ uller, T.S. B€ oscke, D. Br€auhaus, U. Schr€ oder, U. B€ ottger, J. Sundqvist, P. K€ ucher, T. Mikolajick, L. Frey, Ferroelectric Zr0.5Hf0.5O2 thin films for nonvolatile memory applications, Appl. Phys. Lett. 99 (2011) 112901. [3] J. M€ uller, T.S. B€ oscke, U. Schr€ oder, S. Mueller, D. Brauhaus, U. B€ ottger, L. Frey, T. Mikolajick, Ferroelectricity in simple binary ZrO2 and HfO2, Nano Lett. 12 (2012) 4318–4323. [4] J. M€ uller, U. Schr€ oder, T.S. B€ oscke, I. M€ uller, U. B€ ottger, L. Wilde, J. Sundqvist, M. Lemberger, P. K€ ucher, T. Mikolajick, L. Frey, Ferroelectricity in yttrium-doped hafnium oxide, J. Appl. Phys. 110 (2011)114113. [5] S. Mueller, J. Mueller, A. Singh, S. Riedel, J. Sundqvist, U. Schroeder, T. Mikolajick, Incipient ferroelectricity in Al-doped HfO2 thin films, Adv. Funct. Mater. 22 (2012) 2412–2417.

Dopants in Atomic Layer Deposited HfO2 Thin Films

71

[6] S. Mueller, C. Adelmann, A. Singh, S. Van Elshocht, U. Schroeder, T. Mikolajick, Ferroelectricity in Gd-doped HfO2 thin films, ECS J. Solid State Sci. Technol. 1 (2012) N123–N126. [7] M. Hoffmann, U. Schroeder, T. Schenk, T. Shimizu, H. Funakubo, O. Sakata, D. Pohl, M. Drescher, C. Adelmann, R. Materlik, A. Kersch, T. Mikolajick, Stabilizing the ferroelectric phase in doped hafnium oxide, J. Appl. Phys. 118 (2015) 071006. [8] T. Schenk, S. Mueller, U. Schroeder, R. Materlik, A. Kersch, M. Popovici, C. Adelmann, S. Van Elshocht, T. Mikolajick, Strontium doped hafnium oxide thin films: wide process window for ferroelectric memories, in: 2013 Proceedings of the European Solid-State Device Research Conference (ESSDERC), 2013, pp. 260–263. [9] J. M€ uller, E. Yurchuk, T. Schl€ osser, J. Paul, R. Hoffmann, S. Mueller, D. Martin, S. Slesazeck, P. Polakowski, J. Sundqvist, M. Czernohorsky, K. Seidel, P. K€ ucher, R. Boschke, M. Trentzsch, K. Gebauer, U. Schr€ oder, T. Mikolajick, Ferroelectricity in HfO2 enables nonvolatile data storage in 28 nm HKMG, in: 2012 Symposium on VLSI Technology, 2012, pp. 25–26. [10] A.G. Chernikova, D.S. Kuzmichev, D.V. Negrov, M.G. Kozodaev, S.N. Polyakov, A. M. Markeev, Ferroelectric properties of full plasma-enhanced ALD TiN/La:HfO2/ TiN stacks, Appl. Phys. Lett. 108 (2016) 242905. [11] U. Schroeder, C. Richter, M.H. Park, T. Schenk, M. Pesˇi
c, M. Hoffmann, F.P. G. Fengler, D. Pohl, B. Rellinghaus, C. Zhou, C.-C. Chung, J.L. Jones, T. Mikolajick, Lanthanum doped hafnium oxide: a robust ferroelectric material, Inorg. Chem. 57 (2018) 2752–2765. [12] P. Polakowski, J. M€ uller, Ferroelectricity in undoped hafnium oxide, Appl. Phys. Lett. 106 (2015)232905. [13] X. Sang, E.D. Grimley, T. Schenk, U. Schroeder, J.M. LeBeau, On the structural origins of ferroelectricity in HfO2 thin films, Appl. Phys. Lett. 106 (2015)162905. [14] T. Shimizu, T. Yokouchi, T. Oikawa, T. Shiraishi, T. Kiguchi, A. Akama, T.J. Konno, A. Gruverman, H. Funakubo, Contribution of oxygen vacancies to the ferroelectric behavior of Hf0.5Zr0.5O2 thin films, Appl. Phys. Lett. 106 (2015)112904. [15] T. Olsen, U. Schr€ oder, S. Mueller, A. Krause, D. Martin, A. Singh, J. M€ uller, M. Geidel, T. Mikolajick, Co-sputtering yttrium into hafnium oxide thin films to produce ferroelectric properties, Appl. Phys. Lett. 101 (2012) 082905. [16] L. Xu, T. Nishimura, S. Shibayama, T. Yajima, S. Migita, A. Toriumi, Ferroelectric phase stabilization of HfO2 by nitrogen doping, Appl. Phys. Express 9 (2016) 091501. [17] L. Xu, T. Nishimura, S. Shibayama, T. Yajima, S. Migita, A. Toriumi, Kinetic pathway of the ferroelectric phase formation in doped HfO2 films, J. Appl. Phys. 122 (2017) 124104. [18] Y.H. Lee, H.J. Kim, T. Moon, K.D. Kim, S.D. Hyun, H.W. Park, Y.B. Lee, M. H. Park, C.S. Hwang, Preparation and characterization of ferroelectric Hf0.5Zr0.5O2 thin films grown by reactive sputtering, Nanotechnology 28 (2017) 305703. [19] S. Starschich, U. Boettger, An extensive study of the influence of dopants on the ferroelectric properties of HfO2, J. Mater. Chem. C 5 (2017) 333–338. [20] S. Starschich, T. Schenk, U. Schroeder, U. Boettger, Ferroelectric and piezoelectric properties of Hf1-xZrxO2 and pure ZrO2 films, Appl. Phys. Lett. 110 (2017) 182905. [21] T. Shimizu, K. Katayama, T. Kiguchi, A. Akama, T.J. Konno, H. Funakubo, Growth of epitaxial orthorhombic YO1.5-substituted HfO2 thin film, Appl. Phys. Lett. 107 (2015) 032910. [22] T. Shimizu, K. Katayama, T. Kiguchi, A. Akama, T.J. Konno, O. Sakata, H. Funakubo, The demonstration of significant ferroelectricity in epitaxial Y-doped HfO2 film, Sci. Rep. 6 (2016) 32931. [23] C. Richter, T. Schenk, M.H. Park, F.A. Tscharntke, E.D. Grimley, J.M. LeBeau, C. Zhou, C.M. Fancher, J.L. Jones, T. Mikolajick, U. Schroeder, Si doped hafnium oxide—a “fragile” ferroelectric system, Adv. Electron. Mater. 3 (2017)1700131.

72

Ferroelectricity in Doped Hafnium Oxide

[24] M.H. Park, T. Schenk, C.M. Fancher, E.D. Grimley, C. Zhou, C. Richter, J.M. LeBeau, J.L. Jones, T. Mikolajick, U. Schroeder, A comprehensive study on the structural evolution of HfO2 thin films doped with various dopants, J. Mater. Chem. C 5 (2017) 4677–4690. [25] E. Yurchuk, J. M€ uller, S. Knebel, J. Sundqvist, A.P. Graham, T. Melde, U. Schroeder, T. Mikolajaick, Impact of layer thickness on the ferroelectric behavior of silicon doped hafnium oxide thin films, Thin Sold Films 533 (2013) 88–92. [26] M.H. Park, H.J. Kim, Y.J. Kim, W. Lee, T. Moon, C.S. Hwang, Evolution of phases and ferroelectric properties of thin Hf0.5Zr0.5O2 films according to the thickness and annealing temperature, Appl. Phys. Lett. 102 (2013)242905. [27] T. Mittmann, F.P.G. Fengler, C. Richter, M.H. Park, T. Mikolajick, U. Schroeder, Optimizing process conditions for improved Hf1xZrxO2 ferroelectric capacitor performance, Microelectron. Eng. 178 (2017) 48–51. [28] M.H. Park, C.-C. Chung, T. Schenk, C. Richter, K. Opsomer, C. Detavernier, C. Adelmann, J.L. Jones, T. Mikolajick, U. Schroeder, Effect of annealing ferroelectric HfO2 thin films: in situ, high temperature X-ray diffraction, Adv. Electron. Mater (2018) 1800091. [29] T. Schenk, Formation of Ferroelectricity in Hafnium Based Thin Films, Ph. D. Thesis, Technische Universit€at Dresden, Germany, 2016. [30] T. Shiraishi, K. Katayama, T. Yokouchi, T. Shimizu, T. Oikawa, O. Sakata, H. Uchida, Y. Imai, T. Kiguchi, T.J. Konno, H. Funakubo, Impact of mechanical stress on ferroelectricity in (Hf0.5Zr0.5)O2 thin films, Appl. Phys. Lett. 108 (2016) 262904. [31] M.H. Park, H.J. Kim, Y.J. Kim, T. Moon, C.S. Hwang, The effects of crystallographic orientation and strain of thin Hf0.5Zr0.5O2 film on its ferroelectricity, Appl. Phys. Lett. 104 (2014) 072901. [32] R.D. Shannon, Revised effective ionic radii and systematic studies of interatomie distances in halides and chaleogenides, Acta Cryst A32 (1976) 751–767. [33] G.S. Rohrer, Structure and Bonding in Crystalline Materials, The Press Syndicate of the University of Cambridge, Cambridge, United Kingdom, 2004. [34] E.D. Grimley, T. Schenk, X. Sang, M. Pesic, U. Schroeder, T. Mikolajick, J.M. LeBeau, Structural changes underlying field-cycling phenomena in ferroelectric HfO2 thin films, Adv. Electron. Mater. 2 (2016)1600173. [35] M.H. Park, T. Schenk, C. Richter, E.D. Grimley, J.M. LeBeau, M. Tallarida, C. Mariani, L. Simonelli, C.M. Fancher, J.L. Jones, R. Materlik, C. K€ unneth, A. Kersch, T. Mikolajick, U. Schroeder, Structural root causes of ferroelectricity in doped hafnium oxide, in: Presented at International Symposium of Applications of Ferroelectricity 2016 meeting, 2016. [36] M. Pesic, F.P.G. Fengler, L. Larcher, A. Padavani, T. Schenk, E.D. Grimley, X. Sang, J.M. LeBeau, S. Slesazeck, U. Schroeder, T. Mikolajick, Physical mechanisms behind the field-cycling behavior of HfO2-based ferroelectric capacitors, Adv. Funct. Mater. 26 (2016) 4601–4612. [37] H.J. Kim, M.H. Park, Y.J. Kim, Y.H. Lee, T. Moon, K.D. Kim, S.D. Hyun, C.S. Hwang, A study on the wake-up effect of ferroelectric Hf0.5Zr0.5O2 films by pulse-switching measurement, Nanoscale 8 (2016) 1383–1389. [38] M.H. Park, H.J. Kim, Y.J. Kim, Y.H. Lee, T. Moon, K.D. Kim, S.D. Hyun, F.P.G. Fengler, U. Schroeder, C.S. Hwang, Effect of Zr content on the wake-up effect in Hf1–xZrxO2 films, ACS Appl. Mater. Interfaces 8 (2016) 15466–15475. [39] M.H. Park, H.J. Kim, Y.J. Kim, W. Jeon, T. Moon, C.S. Hwang, Ferroelectric properties and switching endurance of Hf0.5Zr0.5O2 films on TiN bottom and TiN or RuO2 top electrodes, Phys. Status Solidi (RRL) 8 (2014) 532–535.

Dopants in Atomic Layer Deposited HfO2 Thin Films

73

[40] T. Schenk, E. Yurchuk, S. Mueller, U. Schroeder, S. Starschich, U. B€ ottger, T. Mikolajick, About the deformation of ferroelectric hysteresis, Appl. Phys. Rev. 1 (2014)041103. [41] M. Pesˇi
c, S. Slesazeck, T. Schenk, U. Schroeder, T. Mikolajick, Impact of charge trapping on the ferroelectric switching behavior of doped HfO2, Phys. Status Solidi A 213 (2016) 270–273. [42] R. Materlik, C. K€ unneth, A. Kersch, The origin of ferroelectricity in Hf1xZrxO2: a computational investigation and a surface energy model, J. Appl. Phys. 117 (2015) 134109. [43] S.E. Reyes-Lillo, K.F. Garrity, K.M. Rabe, Antiferroelectricity in thin-film ZrO2 from first principles, Phys. Rev. B 90 (2014) 140103. [44] M.H. Park, H.J. Kim, Y.J. Kim, T. Moon, K.D. Kim, C.S. Hwang, Toward a multifunctional monolithic device based on pyroelectricity and the electrocaloric effect of thin antiferroelectric HfxZr1xO2 films, Nano Energy 12 (2015) 131–140. [45] M. Hoffmann, U. Schroeder, C. K€ unneth, A. Kersch, S. Starschich, U. B€ ottger, T. Mikolajick, Ferroelectric phase transitions in nanoscale HfO2 films enable giant pyroelectric energy conversion and highly efficient supercapacitors, Nano Energy 18 (2015) 154–164. [46] M.H. Park, T. Schenk, M. Hoffmann, S. Knebel, J. Gartner, T. Mikolajick, U. Schroeder, Effect of acceptor doping on phase transitions of HfO2 thin films for energy-related applications, Nano Energy 36 (2017) 381–389. [47] U. Schroeder, E. Yurchuk, J. M€ uller, D. Martin, T. Schenk, P. Polakowski, C. Adelmann, M.I. Popovici, S.V. Kalinin, T. Mikolajick, Impact of different dopants on the switching properties of ferroelectric hafniumoxide, Jpn. J. Appl. Phys. 53 (2014) 08LE02. [48] R. Batra, T.D. Huan, G.A. Rossetti Jr., R. Ramprasad, Factors favoring ferroelectricity in hafnia: a first-principles computational study, Chem. Mater. 121 (2017) 4139–4145. [49] C.-K. Lee, E. Cho, H.-S. Lee, C.S. Hwang, S. Han, First-principles study on doping and phase stability of HfO2, Phys. Rev. B 78 (2008) 012102. [50] D. Fischer, A. Kersch, The effect of dopants on the dielectric constant of HfO2 and ZrO2 from first principles, Appl. Phys. Lett. 92 (2008) 012908. [51] D. Cunningham, A First-Principles Examination of Dopants in HfO2, Honors Scholar Theses, University of Connecticut, 2014. [52] A.H. Sihvola, J.A. Kong, Effective permittivity of dielectric mixtures, IEEE Trans. Geosci. Remote Sens. 26 (1988) 420–429. [53] T.C. Choy, Effective Medium Theory; Principles and Applications, Oxford University Press, Oxford, 2015. [54] A.N. Norris, P. Sheng, A.J. Callegari, Effective-medium theories for two-phase dielectric media, J. Appl. Phys. 57 (1985) 1990–1996. [55] J. Robertson, High dielectric constant oxides, Eur. Phys. J. Appl. Phys. 28 (2004) 265–291. [56] M. Galtier, A. Montaner, G. Vidal, Phonons opti
ues de cao, sro, bao au centre de la zone de Brillouin a 300 et 17 K, J. Phys. Chem. Solids 33 (1972) 2295–2302. [57] J.L. Jones, The effect of crystal symmetry on the maximum polarization of polycrystalline ferroelectric materials, Mater. Sci. Eng. B 167 (2010) 6–11. [58] S.W. Lee, B.J. Choi, T. Eom, J.H. Han, S.K. Kim, S.J. Song, W. Lee, C.S. Hwang, Influences of metal, non-metal precursors, and substrates on atomic layer deposition processes for the growth of selected functional electronic materials, Coord. Chem. Rev. 257 (2013) 3154–3176. [59] E. D. Grimley, PhD Thesis, North Carolina State University, 2018. [60] Manual of TF analyzer 3000 of Aixacct systems.

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Ferroelectricity in Doped Hafnium Oxide

[61] C. K€ unneth, R. Materlik, M. Falkowski, A. Kersch, Impact of four-valent doping on the crystallographic phase formation for ferroelectric HfO2 from first-principles: implications for ferroelectric memory and energy-related applications, ACS Appl. Nano Mater. 1 (2018) 2254–2264. [62] M.G. Kozodaev, A.G. Chernikova, E.V. Korostylev, M.H. Park, U. Schroeder, C.S. Hwang, A.M. Markeev, Ferroelectric properties of lightly doped La: HfO2 thin films grown by plasma-assisted atomic layer deposition, Appl. Phys. Lett. 111 (2017) 132903. [63] W. Weinreich, R. Reiche, M. Lemberger, G. Jegert, J. M€ uller, L. Wilde, S. Teichert, J. Heitmann, E. Erben, L. Oberbeck, U. Schr€ oder, A.J. Bauer, H. Ryssel, Impact of interface variations on J–V and C–V polarity asymmetry of MIM capacitors with amorphous and crystalline Zr(1x)AlxO2 films, Microelectron. Eng. 86 (2009) 1826. [64] L.G. Wang, H.L. Tu, Y.H. Xiong, W. Xiao, J. Du, J.W. Wang, G.J. Huang, The effect of dopants on the dielectric constant of HfO2 and ZrO2 from first principles, J. Appl. Phys. 116 (2014) 123505.