Effect of acceptor doping on phase transitions of HfO2 thin films for energy-related applications

Author’s Accepted Manuscript Effect of Acceptor Doping on Phase Transitions of HfO2 Thin Films for Energy-Related Applications Min Hyuk Park, Tony Schenk, Michael Hoffmann, Steve Knebel, Jan Gärtner, Thomas Mikolajick, Uwe Schroeder www.elsevier.com/locate/nanoenergy

PII: DOI: Reference:

S2211-2855(17)30263-X http://dx.doi.org/10.1016/j.nanoen.2017.04.052 NANOEN1932

To appear in: Nano Energy Received date: 9 February 2017 Revised date: 24 March 2017 Accepted date: 25 April 2017 Cite this article as: Min Hyuk Park, Tony Schenk, Michael Hoffmann, Steve Knebel, Jan Gärtner, Thomas Mikolajick and Uwe Schroeder, Effect of Acceptor Doping on Phase Transitions of HfO 2 Thin Films for Energy-Related Applications, Nano Energy, http://dx.doi.org/10.1016/j.nanoen.2017.04.052 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Effect of Acceptor Doping on Phase Transitions of HfO2 Thin Films for Energy-Related Applications Min Hyuk Parka*, Tony Schenka, Michael Hoffmanna, Steve Knebela, Jan Gärtnera, Thomas Mikolajicka,b, Uwe Schroedera a

NaMLab gGmbH, Noethnitzer Strasse 64, 01187 Dresden, Germany.

b

Chair of Nanoelectronic Materials, TU Dresden, 01062 Dresden, Germany

*

Corresponding author: [email protected]

Abstract Fluorite structured HfO2 or ZrO2 thin films have been intensively studied for memory- and energy-related applications since their ferroelectricity was first reported in 2011. The phase transition between the nonpolar tetragonal and the polar orthorhombic phase in these new ferroelectric materials is believed to be promising for energy harvesting, energy storage and solid state cooling. The temperature dependent phase transition and resulting strong pyroelectric and electrocaloric effect have been reported for Si-doped HfO2 and (Hf,Zr)O2 thin films. In this study, the effect of acceptor (Al and Gd) doping into HfO2 thin films on their temperature dependent phase transition was systematically examined. The phase transitions in Al- and Gd-doped HfO2 thin films were much broader compared to Si-doped HfO2 and (Hf,Zr)O2 films. The maximum adiabatic temperature change (T) values of Aland Gd-doped HfO2 film were 5.7 and 3.1 K, respectively. A giant negative electrocaloric effect with T of -7.4 K could be observed for Al-doped HfO2. The various factors which can potentially affect the phase transitions of HfO2 films, such as dopant size, grain size

distribution, spatial dopant distribution, and oxygen vacancy distribution were carefully examined to understand the different phase transition behavior. From the various factors, the distribution of oxygen vacancies is suggested as the origin of the different phase transitions of HfO2 films doped with trivalent and tetravalent dopants.

Graphic Abstract

Abbreviations T, adiabatic temperature change; S, isothermal entropy change; ESD, energy storage density; ALD, atomic layer deposition; HED, harvestable energy density; TMA, trimethyl aluminium

Keywords: Hafnium Oxide, Ferroelectricity, Phase Transition, Pyroelectricity, Electrocaloric effect, Energy Harvesting

1. Introduction The ferroelectric properties in doped HfO2 films were first reported in 2011 [1], and have since been intensively studied [2-11]. It is generally accepted that the origin of the

unexpected ferroelectric properties in undoped or doped HfO2 thin films is the formation of an orthorhombic phase (space group: Pca21) [1,2,12-15]. The field-induced-ferroelectricity which is the characteristic phenomenon in first-order phase transition materials near their transition temperature could be observed for Si- and Al-doped HfO2 as well as for (Hf,Zr)O2 films [1,7,8,16]. The phase transition in (Hf,Zr)O2 and Si-doped HfO2 films could be promising for various applications such as energy harvesting, energy storage and electrocaloric cooling [7,8]. The pyroelectric properties and the electrocaloric effect of (Hf,Zr)O2 and Si-doped HfO2 differ strongly from the ones of conventional ferroelectrics based on a perovskite structure [7,8]. Although the magnitude of the adiabatic temperature change (T = 13.4 K for (Hf,Zr)O2 and 9.5 K for Si-doped HfO2) was comparable to those of other promising candidates, the large T values could be observed within a much wider temperature range [7,8]. As a result, the cooling capacity, which was rarely examined in the field of the electrocaloric effect, of (Hf,Zr)O2 and Si-doped HfO2 films could be much larger than those of conventional perovskite ferroelectrics [7,8]. Another interesting phenomenon was the negative electrocaloric effect in Hf0.5Zr0.5O2 films [17]. This effect (T = -10.8 K) could be observed for Hf0.5Zr0.5O2 films, and a special phenomenon near the Curie temperature (Tc) was suggested [17]. It should be noted that the mechanism behind the negative electrocaloric effect is still under debate by many research groups [17-24]. The Curie-Weiss temperature of a large portion of grains in Hf0.5Zr0.5O2 thin film is lower than room temperature, meaning that the nonpolar tetragonal phase is already a metastable phase at room temperature [17]. Due to the co-existence of this nonpolar metastable phase and the ferroelectric phase, the free energy barrier between two different polarization states decreases with increasing temperature, which reduces the critical field for the phase transition. As a result, larger numbers of dipoles in the Hf0.5Zr0.5O2 film can be

poled in the same direction with increasing temperature [17]. A similar trend could also be observed in Si-doped HfO2 films, but the magnitude of the negative electrocaloric effect was not quantitatively examined in the paper by Mueller and coworkers [25]. For the case of conventional ferroelectrics such as Pb(Zr,Ti)O3 (PZT), the ferroelectric properties were known to be strongly affected by the valence of the dopants [26-30]. It was reported that acceptors such as Fe could make PZT harder, whereas donors such as Nb could soften the material [26]. For the case of PZT, when the dopant replaces the A-site (Pb) or Bsite cations (Zr or Ti), it strongly affects the ferroelectric properties of the film. Ramam et al. reported that the doping of Fe (acceptor in B-site) into Pb0.988La0.012(Zr0.53Ti0.47)O3 decreased Tc [27]. Ye et al. and Ujma et al. independently reported that the doping of Nb (donor in Bsite) into PbZrO3 and PbZr0.92Ti0.08O3, respectively, also decreased Tc [28,29]. Stenger and Burggraaf reported that the doping of La (donor in A-site) into PbZr0.30Ti0.70O3 resulted in a reduction of Tc as well [30]. Moreover, in all the aforementioned references on phase transitions of doped Pb(Zr,Ti)O3, the phase transition was broadened with increasing doping concentration [27-30]. For the case of doped HfO2 films, on the other hand, the effect of the valence number of dopants on the ferroelectric hardness has not been clearly understood yet. However, the ferroelectric hardness of doped HfO2 films is strongly affected by the dopant size [31,32]. Generally, the coercive field (Ec) [31] and/or remanent polarization (Pr) [32] of doped HfO2 films tends to increase with increasing dopant size. Therefore, examining the effect of trivalent or bivalent dopants on the temperature dependent phase transition within HfO2 thin films is an interesting topic to understand the complicated polymorphism in this material system. Al and Gd are trivalent dopants different from tetravalent ones such as Si and Zr. The phase transitions in Zr- and Si-doped HfO2 have been reported before [7,8]. Those of trivalent dopants have not been reported yet. In this study, therefore, the effect of

acceptor doping on the phase transition of HfO2 thin films is systematically examined. As dopant size seems to have a significant effect [16,33], Al is chosen as a smaller dopant and Gd as a larger dopant compared to Hf as the substituted cation in the host crystal.

2. Material and methods The Al- and Gd-doped HfO2 films were deposited using a thermal atomic layer deposition (ALD) process on TiN electrodes on p-type (100) Si substrates with a native oxide layer. The TiN bottom electrode was deposited via reactive sputtering using a Ti target in N2 atmosphere for Al-doped HfO2 films, and via plasma enhanced ALD for Gd-doped HfO2 samples. To deposit Al-doped HfO2 films, hafnium tetrachloride (HfCl4) and trimethyl aluminium (TMA) were used as precursors for Hf and Al, respectively. H2O was used as an oxygen source for both Hf and Al. The deposition temperature was fixed at 300 oC. For deposition of Gd-doped HfO2 films, HfCl4 and Gd(iPrCp)3 were used as precursors, and H2O was used as an oxidant for both Hf and Gd. The details of the ALD process were reported in the previous studies in more detail [34,35]. After deposition of the doped HfO2 films, a TiN top electrode was deposited via reactive sputtering using a Ti target in N2 atmosphere. The post-metallizationannealing (PMA) process was conducted at 800 oC for 20 s in N2 atmosphere to crystallize the films. After the PMA process, Pt electrodes were deposited via electron-beam evaporation and patterned using a shadow mask with circular holes whose size was 31400 m2. Patterned Pt top electrodes served as a hard mask during the subsequent TiN wet etch process using a SC1 solution containing NH4OH, H2O2, and H2O. The structural properties of the samples were reported in a previous study in detail [36]. The doping concentrations were calculated based on the systematic previous works on the ALD process of Al-doped HfO2 thin films [34], while those of Gd-doped HfO2 thin films were analyzed using particle induced X-ray

emission (PIXE) [37]. A more detailed analysis of the doping concentrations of Al- and Gddoped HfO2 films via PIXE and time-of-flight secondary ion mass spectrometry (TOF-SIMS) was reported elsewhere [36]. A triangular voltage was applied to the bottom electrode while the top electrode was connected to virtual ground, and a measurement frequency of 1 kHz was used. The Pr values for doped HfO2 films were taken from the polarization-electric field curves achieved with 3.5-4.0 MV/cm maximum field amplitude. For low temperature measurements, a probe station equipped with a chamber which can be evacuated down below 10-6 Torr was used. The sample stage in the chamber was cooled down using liquid nitrogen. The microstructure of the samples was analyzed using scanning electron microscopy and scanning probe microscopy, and the grain size distribution of the samples was analyzed using the watershed method implemented within the Gwyddion software [38].

3. Results and Discussion Figure 1a-c shows the polarization-electric field (P-E) curve of Al-doped HfO2 thin films with doping concentrations of 8.8, 6.9, and 5.7 % (dopant/[dopant + Hf]), respectively. As seen in figure 1a, the hysteresis was negligible for 8.8 % Al-doped HfO2 thin films, meaning that there was a negligible number of grains undergoing the field-induced phase transition from the nonpolar tetragonal to the polar orthorhombic phase within the whole temperature range. A slight opening can be seen at the regions of higher field amplitude. However, the maximum field is still below the critical field for the field-induced phase transition for the majority of grains in the polycrystalline film [7,8,39,40]. For the case of 6.9 % Al-doped HfO2 thin films, a double hysteresis with small remanent polarization (Pr ~ 2 C/cm2) could

be observed. The regions of P-E curves with steep slope (See quadrant I and III of figure 1b.) are believed to result from the field-induced phase transition. The electric field required for the phase transition from tetragonal to orthorhombic was about 2-3 MV/cm, while that required for the switching back from orthorhombic to tetragonal was about 0-1.5 MV/cm as seen in quadrant I of figure 1b. With decreasing temperature below 295 K, the critical fields for the field-induced phase transition also gradually decreases. Related to that, also the field for the reverse transition from polar to non-polar phase moves to lower fields eventually even changing sign (i.e. moving beyond the P-axis). Once this happened for a significant amount of grains, Pr starts to increase. This trend was very similar to the previous reports in Si-doped HfO2 and (Hf,Zr)O2 thin films [7,8]. For the case of the 5.7 % Al-doped HfO2 thin film it can be seen from figure 1c that the Pr (~6 C/cm2 at 300 K) was larger compared to the HfO2 doped with 6.9 % Al. This points to a higher orthorhombic phase portion in the film doped with 5.7% Al. With decreasing temperature, the Pr value also increases. However in the high field region the polarization decreases with decreasing temperature, which is generally considered as a sign of the negative electrocaloric effect [17-24]. As shown in figure 1a-c, the phase transitions in Al-doped HfO2 films are very broad compared to conventional ferroelectrics. In addition, the temperature dependent changes in Pr and Ec differ. For conventional ferroelectric perovskites, polycrystalline films with nanoscale grains or nanoscale crystallites were discussed as the root cause for broad phase transitions [41-44]. The grain size dependence of Tc was first pointed out by Hoffmann and coworkers [8]. Moreover, it is believed that the spatial distributions of dopants, oxygen vacancies [37], and local strains [45,46] can also influence the phase transition. The effect of aforementioned parameters will be discussed in more detail in figure 6 and related paragraphs. The figures 2a-c show the change of polarization with varying temperature at various

electric fields for 8.8, 6.9, and 5.7 % Al-doped HfO2 thin films, respectively. As expected from figure 1a the temperature dependent change in the 8.8 % doped sample was negligible. For the case of the 6.9 % doped sample, a strong temperature dependent polarization change could be observed, which is most pronounced in the low electric field region. For the case of the 5.7 % doped HfO2 thin film, on the other hand, the abnormal trend (positive slope) could be observed for the high electric field region, while normal pyroelectric behavior was observed only for lower electric fields. It is believed that both the normal as the negative electrocaloric effect coexist in the 5.7 % Al-doped HfO2 thin films. The increase of Pr with decreasing temperature can be attributed to the increasing relative portion of the orthorhombic phase which was suggested as a mechanism for the normal electrocaloric effect in previous studies [7,8]. Since the tetragonal phase is a higher entropy phase compared to the orthorhombic phase, the relative portion of tetragonal phase should decrease with decreasing temperature resulting in an increase of Pr [7,8,39]. In the high electric field region, on the other hand, a strong reduction in polarization with decreasing temperature could be observed. This could be attributed to the decrease in the switchable part of the orthorhombic phase, which was suggested as a mechanism for the negative electrocaloric effect in Hf0.5Zr0.5O2 films [17]. From the polarization-temperature plots in figure 2, the T and isothermal entropy change (S) values for Al-doped HfO2 films were calculated, and the results are plotted in figure 3. S shows the magnitude of the internal entropy change which can be induced by changing the electric field from E1 to E2 under isothermal conditions (at constant temperature), and T shows the magnitude of the internal temperature change due to the S change when the material is thermally isolated (in adiabatic condition). The changes in the

T and S values were calculated based on the (∂P/∂T)E values using equations 1 and 2 [4653].

T  

1 E2  P  T dE  C E1  T  E

S  

Here,

1



E2

E1

 P    dE  T  E

(1)

(2)

 and C are the density and specific heat, respectively. These equations are based on the

Maxwell relation (∂P/∂T)E = (∂S/∂E)T. A fourth order polynomial was used to fit the polarization-temperature data, and the specific heat value as a function of temperature was taken from literature [54]. As expected, the electrocaloric effect in 8.8 % Al-doped HfO2 was negligible and the maximum T was even smaller than 1 K. For the case of 6.9 % Al-doped films, the largest T and S could be observed at 295 K, and the T and S values were 5.7 K and 5.7 J/kg K, respectively. Table I compares the electrocaloric effect of Al- and Gddoped HfO2 films in this study to the results of previous reports [7,8,17-24,40,47-53]. As seen in table I, the maximum T and S value of 6.9 % Al doped HfO2 film are smaller compared to the values for Hf0.2Zr0.8O2 (13.4 K and 16.7 J/kg K) and 5.6 % Si-doped HfO2 films (9.5 K and 8.9 J/kg K). This might be due to the relatively small (∂P/∂T)E values of Al-doped HfO2 films compared to Hf0.2Zr0.8O2 and Si-doped HfO2, meaning that the spatial Tc distribution in Al-doped HfO2 film was rather broad compared to the ones for tetravalent dopants. For the case of 5.7 % Al-doped films, on the other hand, the largest negative T and S could be observed at 80 K, and the T and S values were -7.4 K and -7.8 J/kg K, respectively, which are the third largest absolute values for negative electrocaloric effect in thin films [17]. The

origin of the negative electrocaloric effect in Hf0.5Zr0.5O2 could be understood based on the Ec change near Tc. The largest values to the authors’ knowledge, namely T and S values of 52.2 K and -94.23 J/kg K were reported by Vats and coworkers [22] for nanolaminate structures based on perovskite materials. However, the largest T for 5.7 % Al-doped HfO2 thin films could be observed at 80 K making it favorable for low temperature applications. Figure 4a shows the harvestable energy density (HED) values of the Al-doped films, which is an important figure of merit for pyroelectric energy harvesting (PEH) applications using the Olsen cycle. The details of Olsen cycle can be found in previous studies [7,8]. For pyroelectric energy harvesting, a time variant temperature change is necessary, so it can be calculated between two different temperatures. For the calculation, the higher temperature was fixed at 295 K, and the lower temperature was changed from 80 to 295 K. The largest HED value of 9.2 J/cm3 per cycle could be observed for the 6.9 % Al doped HfO2 film, when the temperature range is 80 to 295 K. This value is about 45 % of the HED reported for 5.6 % Si-doped HfO2 (20.27 J/cm3 cycle) with temperature ranging from 80 to 400 K, and about 83 % of Hf0.2Zr0.8O2 thin film (11.5 J/cm3 per cycle) with temperature ranging from 298 to 423 K [7,8]. The HED of the 5.7 % Al-doped HfO2 film could be calculated based on the inverse Olsen cycle, since it showed the negative electrocaloric effect. The maximum HED value was 8.0 J/cm3 per cycle with a temperature ranging from 80 to 295 K. Figure 4b shows the electrostatic energy storage performance (energy storage density (ESD) and efficiency) calculated from the area between the upper branch of the P-E hysteresis and the polarization axis as defined in previous reports [5,55]. The 6.9 % doped sample showed the largest ESD value of 42 J/cm3 which is slightly smaller than that reported for Hf0.3Zr0.7O2 films (46 J/cm3) before [5]. However, it should be noted that the magnitude of the electric field used in this study (3.9 MV/cm) was smaller by ~13 % than that used for Hf0.3Zr0.7O2

(~4.5 MV/cm). While the ESD of the 6.9 % Al-doped film was similar to that of 5.6 % Sidoped HfO2, the efficiency (51 – 61 %) was much lower than for the Si-doped HfO2 (about 80 %) [8]. This is due to the larger loss given by the enclosed area of the P-E hysteresis for the loop on either the positive or negative side of the electric field axis. For a high efficiency, the difference between critical fields for the transition from nonpolar to polar and back has to be as small as possible. It was reported that the electric work for switching of doped HfO2 films is strongly affected by the dopant size, and Si-doped HfO2 might be the most promising for energy storage applications due to its very low energy loss [8]. Figure 5a and b show the P-E curves of 10 nm-thick Gd-doped HfO2 films with doping concentration of 3.9 and 3.0 %, respectively, measured at various temperatures. While three samples with different doping concentration (3.0, 3.9, and 6.0 %) were tested, only the P-E curves of 3.0 and 3.9 % Gd-doped HfO2 films are presented. For the case of 6.0 % Gd-doped HfO2 film, the hysteresis was negligible and only a minimal change could be observed with varying temperature. The most prominent phase transition, i.e. the highest slopes of polarization-temperature plots ((∂P/∂T)E), could be observed for the 3.9 % Gd-doped HfO2 film. Therefore, the temperature dependent Pr of this sample was taken for the dopant comparison in figure 6. A relatively large shift of the P-E curve could be observed in the 3.0 % Gd-doped HfO2 with varying temperature. However, an electric field of only up to 3.5 MV/cm could be applied due to the higher leakage current compared to the 3.9 % Gd-doped HfO2 thin film. The positive (∂P/∂T)E values, which are indicative of the negative electrocaloric effect, could be observed for the 3.0 % Gd-doped HfO2 film. However, the temperature dependent change of polarization seems much weaker compared to that of 6.9 % Al-doped HfO2 samples. Figure 6 shows the temperature dependent changes in Pr values for 5.6 % Si-doped HfO2

(9 nm), Hf0.4Zr0.6O2 (9.2 nm), 5.7 % Al-doped HfO2 (10 nm), and 3.9 % Gd-doped HfO2 (10 nm) thin films. The data for the 5.6 % Si-dopedHfO2 and Hf0.4Zr0.6O2 thin films were taken from previous studies [8,17]. For all films, the samples with the highest dPr/dT values within their measured temperature range were taken for the comparison. For all dopants, a decrease in Pr with increasing T could be observed, but the slopes were different. The slopes dPr/dT of 5.6 % Si-doped HfO2 and Hf0.4Zr0.6O2 were larger, whereas those of 5.7 % Al- and 3.9 % Gd-doped HfO2 thin films were comparatively small. This difference could not be understood based on the dopant size only, since the ionic radius of Al3+ (69 pm) is in between that of Si4+ (54 pm) and Zr (86 pm). The ionic radii were calculated for eightfold coordination based on the previous studies [56,57]. To understand the broadened phase transition in doped HfO2 films, we need to start with their theoretical phase transition. The S of field-induced-ferroelectric materials can be theoretically estimated from the difference between the entropy of the polar and nonpolar phase. From the computational calculation result by Materlik et al. [19], the theoretical S value of Hf0.5Zr0.5O2 could be estimated as ~397 mJ/cm3 K (~47.3 J/kg K) which is several times larger than the maximum S of Hf0.2Zr0.8O2 (96 mJ/cm3 K) and 5.6% Si doped HfO2 (85.5 mJ/cm3 K) films obtained from experiments [7,8]. This large S value can be achieved only when the size of all grains in the (Hf,Zr)O2 film are equivalent, so that the phase transition of every grain occurs at the same electric field. However, the phase transition of doped polycrystalline films can be hardly ideal due to the following reasons. According to the surface energy model suggested by Materlik and coworkers, the Tc of each grain should be strongly affected by its size [39]. Since the surface energy of the tetragonal phase is lower than that of the orthorhombic phase, Tc should decrease with decreasing film thickness and grain size [39]. Moreover, it was reported that the dopant and/or oxygen vacancy

concentrations can also affect the energy difference between the tetragonal and orthorhombic phase [37]. Therefore, a spatial inhomogeneity in dopant and/or oxygen vacancy concentration in doped HfO2 films can be another origin of spatial distribution in Tc. Thus, the effects of aforementioned factors such as grain size distribution and the spatial distribution of dopants and oxygen vacancies will be examined to clarify the origin of the difference in phase transition of HfO2 films doped with various dopants. Due to the surface energy effect, the Tc of a grain decreases with decreasing film thickness and grain radius. The Tc,grain of a cylindrical grain can be formulated as

T c ,grain  T c ,bulk 

 to 2 2 (  ). S to r t

(3)

In equation (3) , Tc,bulk is the Tc of the bulk t-o the surface energy difference between tetragonal and orthorhombic phase, St-o the entropy difference between tetragonal and orthorhombic phase, r the grain radius, and t the film thickness. Since t-o is negative and St-o is positive Tc,grain is generally lower than Tc,bulk. With decreasing r or t, Tc,grain decreases, meaning that the tetragonal phase is energetically more favorable. This trend is in accordance to previous reports on Hf0.5Zr0.5O2 films thinner than 8 nm, where the competition between tetragonal and orthorhombic phase could be clearly observed [58]. It was previously reported that the doped HfO2 films deposited by ALD generally have columnar grain structure [12,13]. Consequently the vertical grain size was assumed to be equivalent to the film thickness. The distribution of grain radii in the 6.9 % Al-doped HfO2 thin film was calculated from planeview SEM images of the film using the software Gwyddion [38]. The grain size distribution

of 5.6 % Si-doped HfO2 was taken from reference [8]. Since the plane-view SEM image of Hf0.4Zr0.6O2 film was not reported in the previous studies, the grain size analysis was conducted on the SEM image of a 10 nm-thick Hf0.5Zr0.5O2 film from a previous study [59]. Based on the work of Park [60] it is assumed that the difference in Zr contents of 10 % has a negligible effect on the grain size]. As seen in figure 7a, the mean grain radius of Hf0.4Zr0.6O2 film is smaller than the ones of 5.6 % Si- and 6.9 % Al-doped HfO2 thin films. From the grain size distribution in figure 7a, the distribution of (Tc,grain-Tc,bulk) was calculated using equation 3. The St-o was estimated from the temperature variant free energy in Materlik et al. [39]. The St-o and t-o values for HfO2 were taken for Si- and Al-doped HfO2 thin films, and those for Hf0.4Zr0.6O2 was calculated for Hf0.5Zr0.5O2 and ZrO2 based on Vegard’s law. As seen in figure 7b, the distribution of (Tc,grain-Tc,bulk) was much broader in Hf0.4Zr0.6O2 compared to those of 5.6 % Si- and 6.9 % Al-doped HfO2. This could be attributed to the larger t-o value (meaning a stronger size effect) and smaller grain size. However, from the experimental Pr-temperature plot in figure 6, the slope for 5.6 % Si-doped HfO2 and Hf0.4Zr0.6O2 were much larger than the one for Al-doped HfO2. Therefore, the different (dPr/dT) values cannot be understood based on the distribution of grain size only. Another possible origin for the dopant-dependence of the phase transition could be the spatial distribution of dopants. In previous studies, the non-uniform distribution of Si dopants could be detected by in-depth TOF-SIMS [36,61]. However, no such evidence was observed in Gd-doped HfO2 films by Hoffmann et al. [37]. It should be noted that similar Gd-doped HfO2 films were used in this study. So, it is believed that the distribution of Gd should be similar as well [37]. In contrast to doped HfO2 films, non-uniformity can be ruled out in (Hf,Zr)O2 samples. Considering the sequence of ALD process for the deposition of Hf0.4Zr0.6O2 film in the previous study, it should be free from inhomogeneous doping effects

[17]. The number of ALD cycles of HfO2 between each ZrO2 ALD cycle is only 1 – 2, which results in a sub-unit cell thickness [17]. To examine the depth variation of Al doping concentration, TOF-SIMS was conducted for 8.8 % Al-doped HfO2 films, and the results are shown in figure 8. For the case of Al-doped HfO2 thin films, the Al atoms seem to diffuse to the interfacial region near the TiN top and bottom electrodes, possibly resulting in the spatially inhomogeneous distribution of Tc. The diffusion of Al during the ALD process was reported for Al-doped TiO2 films in a previous study [62]. In contrast, no sign of diffusion of Al was observed in some previous studies for the ALD process [63,64]. However, in the studies reported in [63,64]. Si and Mo were used as bottom electrodes. Thus, the oxidation of TiN electrode and the resulting oxygen vacancies might affect the diffusion of Al during the ALD deposition or the annealing process. It was reported that the TiN electrode might partially reduce HfO2 forming oxygen vacancies [9,65], and oxygen vacancies are believed to affect the diffusion of Al. To summarize, the diffusivity of dopants is strongly dependent on their chemical properties. Si and Zr are expected not to diffuse even at annealing temperature (1000 oC for Si and 500 oC for Zr). On the other hand, the diffusion of Al to the interfacial region could be observed, and Gd is believed to diffuse through the whole film [37]. However, the different (dPr/dT) in figure 6 could not be solely understood by the diffusion of dopants. From the diffusion of dopants, the lower (dPr/dT) values than for (Hf,Zr)O2 and Gd-doped HfO2 are expected for Si- and Al-doped HfO2 films This is in contrast to our experimental observation in figure 6. The spatial distribution of dopants is also expected to affect the local strain in HfO2 thin films. The mechanical strain was reported to affect the polymorphism [45,46], and thus, its spatial distribution should broaden the phase transition in doped HfO2 thin films. The local strain should be related to the size difference between dopant and Hf and the spatial

distribution of dopants. Since the radii of Si4+ (54 pm) and Al3+ (68 pm) ions are smaller than an ionic radius of Hf4+ (83 pm) by about 35 and 18 %, respectively, compressive stress is expected near Si4+ and Al3+ dopants. The ionic radii were determined for eightfold coordination based on previous studies [66,67]. For the case of Gd3+, on the other hand, tensile stress is expected due to its larger size (105 pm). Zr4+ has almost equivalent ionic size (84 pm) compared to Hf4+, so the local strain induced by Zr4+ should be negligible. Moreover, the distribution of Zr4+ is spatially homogeneous, implying no spatial distribution of strain. By considering a spatial distribution of dopants and the size difference between dopant and Hf, a broader phase transition is expected for Si- and Al-doped HfO2 thin films which is not seen experimentally. Therefore, the dopant dependence of (dPr/dT) values in figure 6 cannot be understood by considering the local strain induced by dopants. The inhomogeneous distribution of oxygen vacancies can be another source of the inhomogeneous distribution of Tc. One interesting difference between the dopants with high dPr/dT (Si and Zr) and those with lower dPr/dT (Al and Gd) is their valence number. Si and Zr are tetravalent atoms, while Al and Gd are trivalent atoms which can work as acceptors. It is known that the doping of acceptors into HfO2 decreases the formation energy of oxygen vacancies, thus the concentration of oxygen vacancies increases with the doping concentration [60]. Since Al and Gd are trivalent dopants, one oxygen vacancy for every two dopant atoms might be created in the HfO2 films for reasons of charge neutrality [61]. Therefore, Al- and Gd-doped HfO2 thin films are expected to have higher oxygen vacancy concentrations compared to Si-doped HfO2 and (Hf,Zr)O2 thin films. Furthermore, these oxygen vacancies might not be uniformly distributed throughout the whole volume of the films due to the layering of dopants and the reduction of the film by the TiN electrodes. Oxygen vacancies can prefer specific positions such as the dopant sites, film surface,

interfacial regions, grain boundaries, and other defects [62]. As mentioned in the introduction, oxygen vacancies can affect the free energy of phases in HfO2 thin films [37]. Consequently, the spatial non-uniformity of the oxygen vacancy concentration is another potential cause for the spatial Tc distribution in doped HfO2 thin films. Therefore, the valence number of the dopants can be a critical factor for the phase transitions of these films. To date, various dopants were reported to induce ferroelectric properties in HfO2 thin films, and their valence number was +2 (Sr, Ca, Ba, Co, Ni, and Mg), +3 (Al, Y, Gd, La, Nd, Sm, In, Ga, and Er), or +4 (Si, and Zr), meaning that they are all acceptors or neutral [2,16,32,63]. The ferroelectricity in N-doped HfO2 was recently reported. Nitrogen is also an acceptor since it should replace oxygen [64]. Therefore, to further understand the phase transition in this leadfree fluorite ferroelectric system, it would be interesting to find possible donors as dopands that induce the ferroelectricity in HfO2 films and study their effect on the phase transition.

4. Conclusion In conclusion, the effect of the acceptor doping by Al and Gd on the phase transition of HfO2 thin films was examined and compared to previous reports on Si-doped HfO2 and (Hf,Zr)O2 films as well as perovskite based systems. Figures of merit for different electrocaloric applications were calculated for the Al- and Gd-doped HfO2 and compared to other materials. The maximum adiabatic temperature change (T) of 6.9 % Al- and 3.9 % Gd-doped HfO2 film was 5.7 and 3.1 K, respectively. A giant negative electrocaloric effect with T of -7.4 K could be observed for the HfO2 doped with 5.7 % Al. It was confirmed that the phase transition in Al- and Gd-doped HfO2 thin films occurred within broader temperature ranges compared to the phase transitions of Si-doped HfO2 and (Hf,Zr)O2 films investigated in previous studies. The temperature dependent Pr changes were much weaker

for Al- and Gd-doped HfO2 thin films compared to those in Si-doped HfO2 and (Hf,Zr)O2 thin films. The various factors which can potentially affect the broadness of the phase transition, such as dopant size, grain size distribution, spatial dopant distribution, spatial local strain distribution, spatial oxygen vacancy distribution, were comprehensively discussed. The different (dPr/dT) values could be understood based on the chemical characteristics of acceptors whose valence number is lower than Hf. The trivalent dopants might decrease the formation energy of oxygen vacancies, inducing higher oxygen vacancy concentrations in the HfO2 thin films. These higher oxygen vacancy concentrations can result in a larger spatial non-uniformity of oxygen vacancies, which then further broaden the spatial Tc distribution in HfO2 thin films. This is an important result to understand the phase transition in this new ferroelectric material system and it gives the starting point for the development of devices for energy related applications based on the temperature dependence of the phase transitions.

Acknowledgements The

authors

acknowledge

the

German

Research

Foundation

(Deutsche

Forschungsgemeinschaft) for funding part of this research in the frame of the “Inferox” project (MI 1247/11-2). MHP is supported by Humboldt postdoctoral fellowship from Alexander von Humboldt Foundation. MHP acknowledges Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education (2016R1A6A3A03012208). The authors acknowledge Dr. Christoph Adelmann and his group in IMEC for deposition of the films and PIXE analysis, and Prof. Jacob L. Jones and his group in North Carolina State University for TOF-SIMS analysis.

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30 6.9% Al:HfO2

3.9MV/cm 20 295 to 80K

3.9 MV/cm 20 295 to 80K

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(a)

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30 5.7% Al:HfO2

Polarization [C/cm ]

30 8.8% Al:HfO2

Polarization [C/cm ]

2

Polarization [C/cm ]

Figure 1. Polarization-electric field curves measured between 80 and 295 K of 10 nm-thick Al-doped HfO2 thin films with (a) 8.8, (b) 6.9, and (c) 5.7 % of Al doping, respectively.

10 0

295K

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Figure 2. Polarization-temperature plot for Al-doped HfO2 thin films doped with (a) 8.8, (b) 6.9, and (c) 5.7 % of Al. The curves were extracted from the polarization-electric field curves shown in figure 1.

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8.8% Al:HfO2

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Figure 3. (a) Adiabatic temperature change (T) and (b) isothermal entropy change (S) as a function of temperature for Al-doped HfO2 thin films doped with 5.7%, 6.9% and 8.8% Al.

6

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Figure 4. (a) Harvestable energy density (HED), and (b) energy storage density (ESD, solid symbols, left axis) as well as efficiency (open symbols, right axis) of Al-doped HfO2 thin films doped with 5.7%, 6.9% and 8.8% Al.

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Figure 5. Polarization-electric field curves of 10 nm-thick Gd-doped HfO2 thin films measured at various temperatures for (a) 3.9 and (b) 3.0 % of Gd doping.

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80K

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30 3.9% Gd:HfO2

(b)

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Figure 6. Positive remanent polarization (+Pr) of 5.6 % Si-doped HfO2 [8], Hf0.4Zr0.6O2 [17], 5.7 % Al-doped HfO2, and 3.9 % Gd-doped HfO2 thin films as a function of temperature.

5.6% Si:HfO2

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Hf0.4Zr0.6O2

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+Pr [C/cm2]

8 6 4 2 0 100

200

300

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500

Temperature [K]

Figure 7. (a) Distribution of grain size analyzed using Gwyddion and (b) distribution of the resulting difference between Curie temperature of grains (Tc,grain) and Curie temperature of bulk (Tc,bulk) for 5.6 % Si-doped HfO2, 6.9 % Al-doped HfO2, and Hf0.4Zr0.6O2 films, respectively. Tc,grain was calculated based on the surface energy model from [19].

15

(a)

(b)

Al:HfO2 Hf0.4Zr0.6O2

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Si:HfO2

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15

20

25

30

35

-400

-300

-200

-100

Tc,grain-Tc,bulk [K]

Equivalent radius [nm]

Figure 8. Time of flight secondary ion mass spectra for Si-, O2-, AlO2-, TiN-, TiO- and HfO2in 8.8 % Al-doped HfO2 thin film sandwiched between top and bottom TiN electrodes. 5

10

Al:HfO2

Intensity [a.u.]

4

10

Si-

O2

-

AlO2 3

-

10

TiN

-

TiO2

-

HfO2

2

10

1

10

0

10

20

30

Depth [nm]

40

Table I. Electrocaloric characteristics of various thin-films. (ECE: electrocaloric effect/T: temperature/T: adiabatic temperature change/E: electric field range/S: isothermal entropy change)

Positive ECE

Negative ECE

Material

T [K]

T [K]

E [kVcm-1]

T/E [K cm kV-1]

S [J K-1 kg-1]

PbZr0.95Ti0.05O2 [40]

495

12

776

0.015

8

Pb0.8Ba0.2ZrO3 [49]

290

45.3

598

0.076

46.9

PbSc0.5Ta0.5O3 [59]

341

6.2

774

0.008

6.3

0.9PbMg1/3Nb2/3O3-0.1PbTiO3[51]

348

5

895

0.006

5.6

P(VDF-TrFE)55/45 [47,48]

353

12.6

2090

0.006

60

SrBi2.1Ta2O9[52]

500

4.9

600

0.008

N/A

Hf0.2Zr0.8O2[7]

298

13.4

3260

0.003

16.7

5.6 % Si-doped HfO2[8]

298

9.5

3330

0.003

8.9

6.9 % Al-doped HfO2

295

5.7

3900

0.0015

5.7

3.9 % Gd-doped HfO2

295

3.1

4000

0.0008

3.1

(Bi0.5Na0.5)TiO3-0.1BaTiO3[23]

323

-3.3

517

-0.006

N/A

0.72Pb(Mg0.33Nb0.67)O30.28PbTiO3[19]

353

-0.15

10

-0.015

N/A

(Na0.5Bi0.5)TiO3[18]

413

-0.33

50

-0.007

N/A

(Bi0.5Na0.5)TiO3[24]

293

-1.6

70

-0.023

N/A

PbZrO3[20]

310

-1.05

100

-0.011

N/A

(Pb0.98La0.02)(Zr0.95T0.05)O2[21]

303

-5.0

308

-0.016

N/A

4 % Eu doped PbZrO3[53]

403

-6.6

709

-0.009

-5.42

Hf0.5Zr0.5O2[17]

448

-10.8

3260

-0.003

-10.9

5.7 % Al-doped HfO2

80

-7.4

3900

-0.002

-7.8

Authors biography

Min Hyuk Park received his B.S., M.S. and Ph.D degrees in Materials Science and Engineering from Seoul National University, Seoul, Korea, in 2008 and 2014, respectively. He has worked as a postdoctoral researcher in Nanoelectronic Materials Laboratory gGmbH since November in 2015. He has been supported by Alexander von Humboldt Foundation since December in 2016. His research interests include ferroelectric and antiferroelectric (Hf,Zr)O2 and doped HfO2 thin films for memory, energy storage, energy harvesting, solidstate cooling, and sensor applications. He is the (co-)author of more than 30 papers. E-mail: [email protected]

Tony Schenk received his B.Eng. degree in Microtechnology (dual) and M.Eng. degree in Nano- and Surface Technology from the Westsaxon University of Applied Sciences Zwickau in 2011 and 2012, respectively. His research interests are arranged around hafnia/zirconia based ferroelectrics. Within the frame of these activities, he (co-)authored more than 20 peerreviewed articles and received his doctoral degree from TU Dresden in 2016. E-mail: [email protected]

Michael Hoffmann received his M.Sc. degree in electrical engineering from TU Dresden in 2016. After a research visit at UC Berkeley in the same year, he is now pursuing a Ph.D. in electrical engineering at TU Dresden in collaboration with NaMLab gGmbH. His current research interests include HfO2 and ZrO2 based ferroelectrics for energy efficient electronics as well as energy conversion and storage applications. E-mail: [email protected]

Steve Knebel received the Diploma degree in microtechnology from the University of Applied Sciences, Zwickau, Germany, in 2009. From 2002 to 2009, he got his professional background as a member of the DRAM reliability department at Qimonda Dresden GmbH & Co OHG, Germany. In 2009, he moved to Nanoelectronic Materials Laboratory gGmbH, Dresden, Germany, as an engineer for electrical characterization with the focus on dielectric reliability. Since 2015, he joined X-FAB Dresden GmbH & Co. KG continuing his carrier as a reliability engineer. E-mail: [email protected]

Jan Gärtner did his apprenticeship as a microtechnologist at Infineon Dresden in 2004. He worked at the failure analysis research group “future technologies” till 2009. From 2009-2011 he was part of the failure analysis team at the CIGS Startup “Odersun”. From 2011 on he worked at Namlab gGmbH as technician for cleanroom and process. His working field includes failure analysis with physical characterizations, tool maintenance and process development. E-mail: [email protected]

Thomas Mikolajick received the Diploma (Dipl.-Ing.) in electrical engineering in 1990 and his PhD in electrical engineering in 1996 both from the University Erlangen-Nuremberg. From 1996 till 2006 he was in the semiconductor industry developing CMOS processes, Ferroelectric Memories, emerging Non-volatile Memories and Flash Memories. In late 2006 he was appointed as professor for material science of electron devices and sensors at the University of Technology Freiberg, and in October 2009 he started at Technische Universität Dresden were he holds a professorship for Nanoelectronic Materials and the position of scientific director at NaMLab gGmbH. Prof. Mikolajick is author or co-author of more than 300 publications in scientific journals or at scientific conferences and inventor or co-inventor of about 50 patents. E-mail: [email protected]

Uwe Schroeder received the Ph.D. degree from the University of Bonn, Bonn, Germany, including a research visit at UC, Berkeley. He was with the University of Chicago, Chicago, IL, as a Postdoctoral Researcher. In 1997, he joined Infineon, formerly Siemens Semiconductor, for DRAM capacitor development in the DRAM Development Alliance with IBM and Toshiba in Hopewell Junction, NY, before transferring to Infineon’s Memory Development Center, Dresden, Germany, in 2000. Here, he continued the research on high-k dielectric and its integration into DRAM capacitors. In 2009, he moved to Nanoelectronic Materials Laboratory gGmbH, Dresden, and pursued his work on high-k dielectrics and ferroelectric HfO2 layers in semiconductor devices. He is the (co-)author of more than 150 papers and conference contributions and is the holder of more than 30 patents. E-mail: [email protected]

Highlights 

-Phase transition of Al- and Gd-doped HfO2 films are reported for the first time.



-The effect of acceptor doping on the phase transition of fluorite ferroelectrics is investigated in depth.



-The energy harvesting and storage as well as electrocaloric effect of Al- and Gddoped HfO2 thin films is examined and compared to other materials.