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Ferroelectric yttrium doped hafnium oxide ﬁlms from all-inorganic aqueous precursor solution ⁎
Xuexia Wang, Dayu Zhou , Shuaidong Li, Xiaohua Liu, Peng Zhao, Nana Sun, Faizan Ali, Jingjing Wang Key laboratory of Materials Modiﬁcation by Laser, Ion and Electron Beams (Ministry of Education), Dalian University of Technology, 116024, China
A R T I C LE I N FO
A B S T R A C T
Keywords: A. Films B. X-ray methods C. Ferroelectric properties E. Capacitors
We report a unique aqueous solution deposition method to prepare yttrium doped hafnium oxide (Y:HfO2) thin ﬁlms using all-inorganic reagents. The composition and chemical bonding features of the ﬁlms were investigated using X-ray photoelectron spectroscopy. The Y:HfO2 ﬁlm was integrated into metal-insulator-semiconductor (MIS) structure capacitors for electrical measurements. A transition of the polarization behavior from apparent ferroelectric-type to linear dielectric-type was observed for ﬁlms with thickness increasing from 25 nm to 80 nm, which is correlated to the dominant crystal structure change from high-symmetry phase to monoclinic phase evidenced by grazing incidence X-ray diﬀraction analysis.
1. Introduction As a leading high dielectric permittivity (high-k) material, hafnium oxide (HfO2) has been successfully used as the gate dielectric in advanced CMOS devices . In 2011, the ferroelectric (FE) properties were ﬁrstly demonstrated in silicon doped HfO2 (Si:HfO2) and Hf0.5Zr0.5O2 [2,3]. Subsequent studies showed that the intrinsic ferroelectricity can be achieved in HfO2 incorporated with various metal element dopants (Al , Y , La , Gd , and Sr  etc.) and N , HfO2–ZrO2 solid-solution  and even pure HfO2 crystallized with encapsulation of top electrode . The structural origin of FE property in thin-ﬁlm HfO2 is now primarily attributed to the formation of a non-centrosymmetric orthorhombic (o-) phase with space group Pca21, which was conﬁrmed by Sang et al. using scanning transmission electron microscopy (STEM) . In most of the published works, metal–insulator–metal (MIM) capacitors were fabricated to investigate the basic material properties of HfO2–based FE-ﬁlms [2–10,13]. The results are directly relevant to the performance of ferroelectric random access memory (FeRAM), which is constructed by transistors and FE capacitors. In the context of developing ferroelectric ﬁeld eﬀect transistors (FeFET), some studies have been particularly dedicated to understand the characteristics of FE-HfO2 layer integrated into metal–insulator–semiconductor (MIS) structure [14,15]. So far, most of the HfO2-based FE-ﬁlms reported in the literature were prepared by advanced vapor methods, such as atomic layer deposition (ALD) as the main choice [2–10], magnetron sputtering
[13,16], and pulsed laser deposition . These deposition techniques rely on sophisticated and high-cost facilities, and in many cases, the process window is quite narrow. Starschich et al. ﬁrstly reported the growth of yttrium doped HfO2 (Y:HfO2) thin ﬁlms on platinum bottom electrodes by chemical solution deposition (CSD) . The ﬁlms exhibited pronounced ferroelectricity, however, the widespread application of this method is limited by preparation of the solution precursor using expensive metal-organic reagent (e.g. hafnium ethoxide) and water-free organic solvents, and by critical requirement of all the operations in a glove box under inert gas atmosphere. It is worthy to note that, in 2011, Jiang et al. reported the use of a unique aqueous precursor solution for spin-coating of dopant-free HfO2 thin ﬁlms with a thickness ranged from < 10 to several hundred nanometers . The ﬁlms were crystallized into monoclinic (m-) structure, and therefore no ferroelectricity was observed. With regards to the ﬁlm growth technique itself, however, this aqueous solution deposition method is advantageous over the aforementioned vapor and metal-organic-reagent-based CSD methods as it is cost-eﬀective and simple due to the use of all-inorganic reagents and all the operations at ambient condition. In this work, we extend such a method to deposit yttrium doped HfO2 (Y:HfO2) thin ﬁlms sandwiched between TiN metal electrode and silicon substrate. The concentration of yttrium has been preset as ∼2 mol%, since a maximum remanent polarization of 10 μC/cm2 was demonstrated for sputtering-prepared Y:HfO2 thin ﬁlms at this composition . The MIS structure is used for the interest of future study
Corresponding author. E-mail address: [email protected]
https://doi.org/10.1016/j.ceramint.2018.04.233 Received 22 March 2018; Received in revised form 26 April 2018; Accepted 26 April 2018 0272-8842/ © 2018 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Wang, X., Ceramics International (2018), https://doi.org/10.1016/j.ceramint.2018.04.233
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was utilized as the X-ray source. The adventitious C 1 s peak at 284.6 eV was used to calibrate the measured XPS spectra. Prior to measurements, no argon sputtering was performed to prevent ﬁlm degradation. The surface morphology of the Y:HfO2 thin ﬁlms was examined by scanning electron microscopy (SEM) on a Zeiss Supra 55 microscope with an acceleration voltage of 15 kV.
on FeFET devices. It will be shown that a transition from linear dielectric to ferroelectric behavior can be achieved by lowering the ﬁlm thickness from 80 to 25 nm. 2. Material and methods 2.1. Precursor solution preparation and characterization
2.4. Electrical measurement Details of HfO2 precursor solution preparation have been reported elsewhere . HfOCl2·8H2O (Alfa Aesar, 99.9%) was ﬁrstly dissolved in deionized water (18.2 MΩ•cm) to obtain a transparent aqueous solution at room temperature. Then slight overdose of ammonia was added into the solution with vigorous stirring, until pH value of the turbid liquid was about 8.5. The precipitate resulted from centrifugation was washed with deionized water to remove chlorine and ammonia. After repeating the centrifugation and rinse three times, moderate amounts of hydrogen peroxide (H2O2) and nitric acid (HNO3) were added into the precipitate followed by continuous stirring. After 12–18 h, a clear, transparent hydrosol having pH value about 0.7 can be ﬁnally obtained. A separate yttrium stock solution was prepared by dissolving Y(NO3)3·6H2O into deionized water at room temperature and stirring continuously for several hours. Finally, two solutions were mixed to form an yttrium mixed HfO2 precursor solution with the desired yttrium concentration of 2 mol%. The thermal behavior of the precursor solution was investigated by thermogravimetric analysis (TGA, Mettler-Toledo TGA/SDTA851e) and diﬀerential scanning calorimetry (DSC, Mettler-Toledo DSC822e). The test was conducted in the temperature range of 25–630 °C under ﬂowing nitrogen. The heating rate was set as 10 °C /min.
Electrical characterization was performed on mental–insulator–semiconductor (MIS) structure capacitors. The titanium nitride (TiN) electrode dots were deposited onto the crystallized Y:HfO2 ﬁlms by radio frequency magnetron sputtering through a shadow mask. The thickness of the TiN ﬁlms is ~80 nm, and the resistivity is ~75 μΩ·cm. The polarization-electric ﬁeld (P-E) and current density-electric ﬁeld (JE) curves were measured using a Radiant ferroelectric tester (Multiferroic 100 V, Radiant Technologies, USA). 3. Results and discussion 3.1. TGA and DSC measurements TGA and DSC measurement results of the HfO2 precursor containing 2 mol% yttrium were shown in Fig. 1 to provide insights into its decomposition chemistry. Considering the resolution of analysis, the Y:HfO2 precursor solution was preheated at 80 °C for 5 min to volatilize most of the free water. There are several endothermic and exothermic peaks in the DSC curve and each corresponds to a rapid weight loss rate in the TGA curve. The ﬁrst endothermic peak at ~120 °C refers to the evaporation heat of incorporated water and the decomposition of the peroxo groups coming from the xerogel. The second endothermic peak in the vicinity of 250 °C is resulted from the decomposition of metal nitrate and hydroxyl group . The exothermic peak at about 535 °C can be ascribed to the ﬁnal densiﬁcation and crystallization of the yttrium doped hafnium oxide. An evidence for this conclusion is that almost no further weight loss is observed after 550 °C in the TGA curve. An overall weight loss of approx. 40% is observed according to the TGA curve. The TGA and DSC results were used to optimize the preheattreatment and annealing parameters of the as-deposited Y:HfO2 ﬁlms as described in material and method section, for the purpose of preventing the formation of pores and cracks in ﬁlms.
2.2. Thin ﬁlm preparation The yttrium doped hafnium oxide (Y:HfO2) thin ﬁlms were deposited onto heavily doped p-type (100) silicon substrates with a low resistivity of 10−2 − 10−3 Ω•cm. Prior to deposition, all substrates underwent a standard RCA cleaning process and then were subjected a plasma etch for 10 min to improve the wettability between precursor solution and substrates. The precursor solution was spin-coated on treated silicon substrates at a speed of 3000 rpm for 25 s, followed by a preheating on a hot plate at 150 °C for 1 min. The procedures were repeated until the ﬁlm reached a desired thickness. Following deposition, the substrates covered with incompletely decomposed ﬁlm were immediately transferred to a hot plate and ramped stepwisely to 120, 250, and 380 °C (5 °C /min) and held for 5 min at each temperature level. As shown later in the thermal analysis result of the precursor, this procedure can eﬀectively avoid fast decomposition of the xerogel, and therefore to ensure extreme smoothness, continuity and high density of the Y:HfO2 ﬁlms. After preheat treatment, all samples underwent a rapid thermal annealing at 700 °C for 30 s under a nitrogen atmosphere.
3.2. XPS analysis X-ray photoelectron spectroscopy (XPS) analyses were carried out in
2.3. Thin ﬁlm characterization The crystal structure and the thickness of the Y:HfO2 thin ﬁlms were determined by grazing incidence X-ray diﬀraction (GIXRD) and X-ray reﬂectivity (XRR) measurements, respectively, using a Bruker D8 Discover diﬀractometer with Cu Kα radiation (λ = 1.5406 Å) from a Cu tube operated at 40 kV/40 mA. For GIXRD, the incident and exit beams were conditioned with a 0.2 mm divergence slit and a Soller slit, respectively. The grazing incidence angle was ﬁxed at 0.5° and the data was collected within a 2θ range of 20–70°. For XRR, the incident and exit beams were conditioned by using a 0.2 mm divergence slit and a 0.1 mm detector slit, respectively. The detector angle was varied from 0.2° to 5° (2θ), with a step size of 0.002°. X-ray photoelectron spectroscopy (XPS, ESCALAB250, Thermo) was used to determine the yttrium content and chemical states of Hf, Y and O elements in the ﬁlms. A focused monochromatic Al Kα radiation (beam energy: 1486.6 eV)
Fig. 1. TGA/DSC curves of dried HfO2 xerogel incorporated with 2 mol% yttrium. 2
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Fig. 2. XPS analysis results of yttrium doped HfO2 ﬁlms, (a) depth proﬁle of the Hf, O, and Y element concentration in ﬁlms and (b)-(d) individual high resolution XPS spectrum of Hf 4 f, Y 3d and O 1 s core level.
spectra as shown in Fig. 2(d). Additionally, the Hf and Y contents are determined as 35 at% and 0.7 at%, respectively, indicating that the doping concentration of Y is close to 2 mol%. Fig. 2(b)-(d) show the high-resolution XPS spectrum from the Hf 4 f, Y 3d and O 1 s orbitals as a function of the binding energy. The Hf 4 f peaks can be assigned to a single HfO2 phase, where the binding energy of Hf 4 f7/2 and Hf 4 f5/2 are located at 17.3 eV and 18.9 eV with a spin-orbit splitting of 1.6 eV. The binding energies of Hf 4 f orbitals are in good agreement with XPS data of HfO2 ﬁlms prepared by ALD method . Analogously, the Y 3d5/2 and Y 3d3/2 doublet peaks observed at 157.7 eV and 159.8 eV, respectively, are consistent with values of yttrium doped HfO2 ﬁlms prepared by ALD, too . The O 1 s spectra can be deconvoluted into three peaks. The O 1 s peak at 530.5 eV can be assigned to the Hf-O bonding peak, whereas the O 1 s peak at 531.6 eV corresponds to the YO bonding peak. The observation of O 1 s peak with a lower intensity at 532.3 eV was attributed to the existence of oxygen vacancies in previous literature reports [23,24]. The inset table in Fig. 2(d) listed the area proportion of each O 1 s peak. The percentage of oxygen vacancies existed in Y:HfO2 ﬁlms is 3.9%. The generation of oxygen vacancies can be explained by the substitution of Hf4+ by trivalent Y in the HfO2 lattice and annealing of the ﬁlm under a nitrogen atmosphere.
Fig. 3. Typical surface SEM image of 2 mol% yttrium doped HfO2 ﬁlms annealed at 700 °C for 30 s.
order to investigate the chemical states and bonding environments of Y:HfO2 ﬁlms deposited on Si substrates and annealed at 700 °C. From the XPS depth proﬁle shown in Fig. 2(a), it is apparent to see that there is a higher concentration of oxygen on the surface than that in the bulk of ﬁlm. This is mainly due to the absorption of oxygen from the atmosphere. When the ﬁlm surface being cleaned using Ar ion sputtering for 60 s, the atomic ratio of Hf with respect to O (Hf/O) is slightly higher than 1:2 of stoichiometric HfO2, suggesting oxygen deﬁciency in the sample. This ﬁnding will be further corroborated by the O 1 s
3.3. SEM measurement Fig. 3 displays a typical surface SEM image of Y:HfO2 ﬁlms deposited on Si substrates and annealed at 700 °C. In general, the ﬁlm surface is rather smooth and ﬂat, no features like pores, voids, microcracks, grain and grain boundary are discernible. The result is consistent with the top-view SEM image of pure HfO2 ﬁlm (∼85 nm thick) prepared by the same aqueous solution deposition method . 3
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Fig. 4. Thickness dependent GIXRD patterns of 2 mol% yttrium doped HfO2 ﬁlms annealed at 700 °C for 30 s.
Additionally very smooth surface morphologies of the ﬁlms were also conﬁrmed by our atomic force microscopy (AFM) characterization reported elsewhere . It is worthy to note that the extreme smoothness and continuity are important features for dielectric thin ﬁlms used in microelectronic devices, whereby the critical requirements for low leakage current and high breakdown resistance can be ensured. Fig. 5. (a) polarization-electric ﬁeld (P-E) and (b) current density-electric ﬁeld (J-E) curves of 25-nm-thick Y:HfO2 ﬁlms (2 mol% Y) in pristine state and after 10000 bipolar ﬁeld cycles. The measurements were performed at 3.4 MV/cm and room temperature.
3.4. GIXRD measurements Fig. 4 illustrates the GIXRD patterns of 2 mol% yttrium doped HfO2 thin ﬁlms with thicknesses changing from 25 to 80 nm. For the 25-nmthick ﬁlm, the presence of the main diﬀraction peak near 2θ = 30.5° indicates that the majority of the ﬁlm is crystallized into higher symmetry phase. In addition, two weak reﬂections can be observed at 2θ = 28.5° and 31.5°, respectively. They are the characteristic (−111)m and (111)m peaks ascribed to lower symmetry monoclinic phase. Using the laboratory X-ray diﬀraction, it is diﬃcult to identify the dominant higher symmetry phase as an orthorhombic (o-), tetragonal (t-) or cubic (c-) structure since the positions of their main diﬀraction peaks are very close and severe broadening as well as overlap of some weak peaks. However, doping of yttrium into HfO2 ﬁlms is generally recognized to stabilize the c-phase rather than t-phase . Furthermore, the apparent ferroelectricity shown later in electrical measurements indicates doubtlessly the at least partial formation of non-centrosymmetric ophase grains in such Y:HfO2 ﬁlms. Therefore, the higher symmetry phase in the 25-nm-thick ﬁlm is termed as c/o-phase in this paper. Work is in progress to identify the crystalline structure using more detailed analysis method, like high-resolution transmission electron microscopy (HRTEM). With increasing ﬁlm thickness, the intensity of (−111)m and (111)m peaks becomes stronger. And till 80 nm, they are higher than the peak at 2θ = 30.5°, indicating the m-phase becomes dominant in the ﬁlm. The GIXRD results show clearly thickness dependence of lower-to-higher symmetry phase transition, which has been reported in previous studies for HfO2 and HfO2-ZrO2 solid solution nano-ﬁlms prepared by ALD method [11,26]. The underlying mechanism should be attributed to the surface energy eﬀect suggested ﬁrstly by Garvie . For pure HfO2 ﬁlm, the critical thickness at which the c/o-phase can be stabilized at room temperature is about 3–5 nm . For our 2 mol% yttrium doped HfO2 thin ﬁlms, the critical thickness can be extended to a value slightly lower than 25 nm. The incorporation of lower-valence Y in HfO2 results in an introduction of oxygen vacancies, which facilitate stabilization of the high symmetry c/ o-phase  and therefore lead to an increase of the critical thickness. As proposed in Ref. 5, the m-phase fraction (%) was calculated to
qualitatively describe the phase stabilization eﬀect of lower symmetry phase in polycrystalline Y:HfO2 ﬁlms by using the formula:
m − phase fraction(%) = I ( −111) m + I (111)m / I ( −111)m + I (111)m + I (111)c / o where I is the integral intensities of the main reﬂexes of the m- and c/o-phases, respectively. The calculated m-phase fraction (yellow line) is depicted as a function of ﬁlm thickness in Fig. 7. The result is correlated to the evolution of the dielectric property. 3.5. Electrical characterization Fig. 5 shows the polarization hysteresis and current density loops of 25-nm-thick Y:HfO2 ﬁlms measured at 3.4 MV/cm. For the pristine sample, the polarization hysteresis is very weak and almost no switching current peaks can be observed. It has been reported that high ﬁeld cycling helps to open initially pinched polarization loop of FEHfO2 ﬁlms prepared by ALD [30,31], and even to induce ferroelectrictype hysteresis for CSD-derived Y:HfO2 ﬁlms . Such a so-called “wake-up” treatment was also conducted for our Y:HfO2 ﬁlms. Bipolar ﬁeld cycling causes a gradual opening of the polarization hysteresis loop. After 10000 cycles, a typical ferroelectric hysteresis with clear switching induced current peaks can be observed as shown in Fig. 5. A remanent polarization (Pr) of 14.2 μC/cm2 can be achieved, which is slightly higher than the maximum Pr of Y:HfO2 ﬁlms prepared by sputtering and CSD [13,18] but lower than that recorded for ALD-derived Y:HfO2 ﬁlms . The shape of the P-E curve and coercive ﬁeld (Ec, −1.45 and +1.78 MV/cm) exhibits polarity asymmetry, which is similar to the observations reported for MIS structure capacitors [15,32]. This phenomenon may be explained by a build-in bias ﬁeld resulted from non-equivalent work function of the TiN electrode (4.5 eV) and the p-type silicon substrate (5.17 eV), and from 4
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Fig. 6. After 10000 bipolar ﬁled cycles, the polarization-electric ﬁeld (P-E) and current density-electric ﬁeld (J-E) curves of 2 mol% yttrium doped HfO2 thin ﬁlms with thickness changing from 25 to 80 nm. The measurements were performed at room temperature and 1 kHz.
accumulation of charged oxygen vacancies at the interface between HfO2 and SiO2/Si. Fig. 6 shows the polarization behavior of 2 mol% yttrium doped HfO2 thin ﬁlms with diﬀerent thicknesses. All the samples experienced the same wake-up treatment of 10000 bipolar ﬁeld cycles. An increase in ﬁlm thickness causes narrowing of the polarization hysteresis loop and quick decreases of the maximum as well as remanent polarization. The 40-nm-thick ﬁlm still exhibits weak ferroelectricity, evidenced by the appearance of switching current peaks before reaching the maximum ﬁled. The 80-nm-thick ﬁlm shows an almost linear dielectric response with negligible Pr value and indiscernible switching current peak. The transition from ferroelectric-type to linear dielectric-type polarization response originates from the changing of dominant phase in Y:HfO2. With increasing ﬁlm thickness, Fig. 7 shows the m-phase fraction increases from 40.6% to 71.3%, accompanied by decrease of the remanent polarization from 14.2 μC/cm2 to 1.2 μC/cm2 and decrease of the relative permittivity (εr) from 52 to 29. The relative permittivity was calculated from the slope of linear P-E curves measured at maximum applied ﬁeld of 1 MV/cm (data not shown). Due to the use of low frequency (1 kHz) and high ﬁeld amplitude, the polarization response includes more extrinsic contributions (e.g. space charge polarization, domain wall motion), and therefore the calculated εr is higher than the value obtained from small signal, high frequency measurements . Even so, the trend of decreasing εr with increasing m-phase fraction is quite reasonable, since the m-phase exhibits the lowest permittivity among various polymorphs of HfO2 . Similar observations of decreasing εr with increasing m-phase fraction were reported
Fig. 7. Evolution of remanent polarization, relative permittivity and monoclinic phase fraction as a function of ﬁlm thickness.
for ALD and sputtered-derived Y:HfO2 ﬁlms, too [5,13]. It is unclear the majority of the ferroelectricity observed for our Y:HfO2 ﬁlms originates from preexistence and de-pining of the noncentrosymmetric o-phase, or from the ﬁeld cycling induced c-phase to o-phase transition as reported for 5.2 mol% yttrium doped HfO2 ﬁlms prepared by CSD . Nevertheless, the experimental results show and also predict clearly that more pronounced ferroelectricity can be induced by increasing the fraction of higher symmetry phase. Referring to the methods reported for achieving the ferroelectricity in HfO2 [2–10,13,16–18], this can be realized by manipulation of the yttrium 5
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dopant concentration and ﬁlm thickness, crystallization of the ﬁlm with encapsulation of top electrode, and optimization of the annealing conditions (temperature, atmosphere, etc.).
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