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[T. Onaya and T. Nabatame, Jpn. J. Appl. Phys. 65, 080804 (2026)..pdf](https://mdr.nims.go.jp/filesets/bf54ea8d-cb9c-427d-acd5-e0831bc3a998/download)

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[Takashi Onaya](https://orcid.org/0000-0002-2710-8623), [Toshihide Nabatame](https://orcid.org/0000-0002-5973-0230)

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[Interface engineering of HfO                    <sub>2</sub>                    -based ferroelectric devices for crystal phase control and enhanced ferroelectricity using atomic-layer-deposited ZrO                    <sub>2</sub>                    nucleation layers](https://mdr.nims.go.jp/datasets/94441818-697a-48e4-9a27-47035f323fa3)

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Interface engineering of HfO2-based ferroelectric devices for crystal phase control and enhanced ferroelectricity using atomic-layer-deposited ZrO2 nucleation layersJapanese Journal ofApplied Physics      PROGRESS REVIEW • OPEN ACCESSInterface engineering of HfO2-based ferroelectricdevices for crystal phase control and enhancedferroelectricity using atomic-layer-deposited ZrO2nucleation layersTo cite this article: Takashi Onaya and Toshihide Nabatame 2026 Jpn. J. Appl. Phys. 65 080804 View the article online for updates and enhancements.You may also likeEnhanced polarization and enduranceproperties of ZrO2-based ferroelectriccapacitor using HfO2 interfacial layerWei Zhang, Yuxuan Shi, Bowen Zhang etal.-Preparation and characterization offerroelectric Hf0.5Zr0.5O2 thin films grownby reactive sputteringYoung Hwan Lee, Han Joon Kim,Taehwan Moon et al.-(Invited) Advances in Atomic LayerProcessing of Hafnia-ZirconiaFerroelectricsPatrick D. Lomenzo, Ruben Alcala, MonicaMaterano et al.-This content was downloaded from IP address 144.213.253.16 on 07/05/2026 at 01:48https://doi.org/10.35848/1347-4065/ae5662/article/10.1088/1361-6528/ad6871/article/10.1088/1361-6528/ad6871/article/10.1088/1361-6528/ad6871/article/10.1088/1361-6528/ad6871/article/10.1088/1361-6528/ad6871/article/10.1088/1361-6528/ad6871/article/10.1088/1361-6528/aa7624/article/10.1088/1361-6528/aa7624/article/10.1088/1361-6528/aa7624/article/10.1088/1361-6528/aa7624/article/10.1088/1361-6528/aa7624/article/10.1088/1361-6528/aa7624/article/10.1088/1361-6528/aa7624/article/10.1149/MA2022-02311135mtgabs/article/10.1149/MA2022-02311135mtgabs/article/10.1149/MA2022-02311135mtgabsaaaInterface engineering of HfO2-based ferroelectric devices for crystal phasecontrol and enhanced ferroelectricity using atomic-layer-deposited ZrO2nucleation layersTakashi Onaya* and Toshihide NabatameResearch Center for Materials Nanoarchitectonics (MANA), National Institute for Materials Science (NIMS), 1-1 Namiki, Tsukuba, Ibaraki 305-0044,Japan*E-mail: ONAYA.Takashi@nims.go.jpReceived January 31, 2026; revised March 18, 2026; accepted March 23, 2026; published online May 6, 2026Interface engineering is a key for high-performance HfO2-based ferroelectric devices. In this study, the fabrication techniques for HfO2-based thinfilms using atomic-layer-deposited ZrO2 nucleation layers [atomic layer deposition (ALD)-ZrO2-NLs], which are inserted at ferroelectric/electrodeinterfaces, are systematically summarized. ALD-ZrO2-NLs were found to promote crystallization and formation of the ferroelectric orthorhombic(O) phase in HfO2-based films, resulting in superior ferroelectricity. This is attributed to two critical roles of ALD-ZrO2-NLs: providing a seed layerthat induces epitaxial-like grain growth because of the high degree of lattice matching between the pre-crystallized ALD-ZrO2-NL and the O phaseof HfO2, and acting as a stressor layer that induces tensile stress due to the difference in thermal expansion coefficients during annealing process.In addition, improved reliability and a low-temperature fabrication process for HfO2-based thin films are demonstrated using ALD-ZrO2-NLs.These findings suggest that interface engineering using ALD-ZrO2-NLs is a promising approach for fabricating next-generation HfO2-basedferroelectric devices. © 2026 The Author(s). Published on behalf of The Japan Society of Applied Physics by IOP Publishing Ltd1. IntroductionDoped HfO2 thin films were reported to show ferroelectricityin 2011, and ferroelectric HfO2-based thin films have sinceattracted significant attention for use in future non-volatilememory device applications.1,2) HfO2-based materials arecompatible with complementary metal-oxide-semiconductor(CMOS) technology because they have already been well-established as gate dielectrics in conventional CMOSdevices.3) In addition, HfO2-based thin films can be depositedby atomic layer deposition (ALD), which is a requirement forfilms to be conformally deposited on complex three-dimen-sional structures of cutting-edge semiconductor devices.4)HfO2-based materials also exhibit a unique advantage inthickness scalability, maintaining robust ferroelectricity evenwhen their film thickness is less than 10 nm.5) Such scalabilityremains a significant challenge for conventional perovskiteferroelectric materials such as Pb(ZrxTi1−x)O3 because of theirsevere size effect.6,7) In addition, HfO2-based thin films withrobust ferroelectricity can be achieved even at a low processtemperature of less than 400 °C, which is compatible with thethermal budget requirements of back-end-of-line (BEOL)processes.8–11) Therefore, ferroelectric HfO2-based materialsare considered promising candidates for use in future ferro-electric devices with higher performance and lower powerconsumption, such as ferroelectric random access memory,ferroelectric field-effect transistors (FeFETs), and ferroelectrictunnel junctions.6,12,13)The origin of ferroelectricity in HfO2-based materials isthought to be displacement of oxygen atoms in a unit cell ofthe non-centrosymmetric orthorhombic (O) phase with spacegroup Pca21.14) Thus, to achieve superior ferroelectricity,crystallization of a HfO2-based thin film and an increase inthe ratio of the ferroelectric O phase to the non-ferroelectricphases in the film are required. However, numerous re-searchers have reported that the fraction of the O phase inHfO2-based thin films is affected by numerous factors,including dopants, electrodes, thickness, and annealingconditions, because this phase is metastable, whereas thestable phase is a paraelectric monoclinic (M) phase.13,15–17)Therefore, the selection of those conditions is critical fordetermining the crystal structure of HfO2-based thin films.TiN is typically used as an electrode material to fabricateHfO2-based metal–ferroelectric–metal (MFM) capacitors be-cause TiN electrodes provide mechanical tensile stress due tothe difference in thermal expansion coefficients between TiN(9.4× 10−6 K−1) and HfO2 (6.5× 10−6 K−1); TiN electrodesalso induce oxygen vacancies (VO) in HfO2-based thin filmsbecause of the scavenging effect of TiN during the annealingprocess, promoting the formation of the O phase.18–22)HfO2-based thin films have been reported to exhibit a highremanent polarization (2Pr = Pr+ − Pr−) of 48 μC cm−2 evenafter an annealing process at a low temperature of 400 °Cwhen a HfO2-based thin film was sandwiched with TiN topand bottom electrodes (TE- and BE-TiN) for the fabrication ofMFM capacitors.18,23) Many other electrode materials, in-cluding Si,9,20,24) TaN,25) Ir,26) Pt,27) W,19) and RuO2,28) havebeen used for HfO2-based MFM capacitors; however, thecrystallinity and ferroelectricity of the HfO2-based thin filmswere found to be strongly dependent on the choice of electrodematerial. Considering practical applications of ferroelectricdevices composed of various electrode materials, engineeringof the interface between the ferroelectric HfO2-based thin filmand the electrode is required to achieve superior ferroelec-tricity irrespective of the electrode materials.HfO2-based thin films generally show fatigue phenomena (i.e.degradation of their 2Pr values with increasing number ofswitching cycles).29,30) This unstable switching behavior de-pending on the number of switching cycles is a major reliabilityissue that needs to be overcome for HfO2-based thin films to beused in practical applications. The origin of fatigue has beenattributed to pinning of domains caused by VO formed duringContent from this work may be used under the terms of the Creative Commons Attribution 4.0 license. Any further distribution ofthis work must maintain attribution to the author(s) and the title of the work, journal citation and DOI.080804-1© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJapanese Journal of Applied Physics 65, 080804 (2026) PROGRESS REVIEWhttps://doi.org/10.35848/1347-4065/ae5662https://crossmark.crossref.org/dialog/?doi=10.35848/1347-4065/ae5662&domain=pdf&date_stamp=2026-05-06mailto:ONAYA.Takashi@nims.go.jphttps://creativecommons.org/licenses/by/4.0/https://doi.org/10.35848/1347-4065/ae5662field cycling.29) In addition, Onaya et al. reported that additionalVO should be formed in HfO2-based films because of theinterfacial reaction induced by electric field cycling between aHfO2-based film and a TiN electrode.31–33) To overcome thesereliability issues, the precise design of the interface between aHfO2-based film and electrodes is necessary.Various interfacial layers (ILs) inserted between ferro-electric HfO2-based films and electrodes have been reportedto improve the ferroelectricity (Table I).9,34–62) For theAl2O3-IL case, the ferroelectricity of HfO2-based thin filmshas been reported to be enhanced through optimization of thethickness and the insertion position.36,46,49–51) In addition, theleakage current for the overall films was reduced when anAl2O3-IL was used because of its amorphous structure andlarger bandgap (∼8.8 eV) compared with that for HfO2(∼5.8 eV).63–65) However, increasing the total capacitanceequivalent thickness (CET) has raised concerns because thedielectric constant (k) for Al2O3 (k ≈ 9) is lower than that forferroelectric HfO2 (20 ⩽ k ⩽ 40).63,64,66,67) In addition,because most of the applied voltage drop occurs across theAl2O3-IL, which has a lower k value, the Al2O3-IL can preventHfO2-based thin films from fully exhibiting theirferroelectricity.34) Consequently, a thicker Al2O3-IL cannotbe used without sacrificing device performance. Previousstudies have also shown that a TiO2-IL promotes the forma-tion of the ferroelectric O phase, driven by the difference inthermal expansion coefficient between TiO2 (5× 10−5 K−1)and HfO2 (6.5× 10−6 K−1).41,68) Although TiO2 has theadvantage of a higher k value (40 ⩽ k ⩽ 110) than that forferroelectric HfO2 (20 ⩽ k ⩽ 40), its smaller bandgap(∼3.5 eV) may cause increased leakage current.63,64,66,69–71)HfO2 itself has also been investigated as an IL to enhance theferroelectricity of HfO2-based thin films.41,43,56–58) Onayaet al. reported that HfO2-based thin films stacked with aHfO2-IL predominantly formed the stable M phase after anannealing process.39) This behavior is attributed to the crystalstructure of HfO2-based thin films being stabilized by theHfO2-IL with a large M-phase fraction because HfO2 formsthe stable M phase preferentially over the metastable O phase.Therefore, precise optimization of film thickness and an-nealing conditions is required when using HfO2-ILs. On theother hand, Onaya et al. were the first to report that using anALD-ZrO2-IL as a nucleation layer (NL) to promotecrystallization and O-phase formation resulted in enhancedferroelectricity;9,34,35) the details of these studies will bediscussed later. Numerous authors have also reported thatsuperior ferroelectricity can be attained using aZrO2-IL.36,37,40–47) The physical properties of ZrO2 (e.g. kvalue, bandgap, and lattice parameters) are similar to those ofHfO2, which is expected to lead to high material compatibilitywithin devices employing an ALD-ZrO2-IL.In this paper, we comprehensively review interfaceengineering using ALD-ZrO2-NLs to achieve crystal phasecontrol and superior ferroelectricity in HfO2-based thinfilms. This paper discusses the effect of ALD-ZrO2-NLsand their insertion position by comparison with other typesof ILs. The improved reliability of HfO2-based thin filmsfabricated using ALD-ZrO2-NLs is also discussed.Furthermore, the low-temperature fabrication of aHfO2-based metal–ferroelectric–semiconductor (MFS) struc-ture, which is a basic structure of FeFETs, is demonstrated at300 °C using ALD-ZrO2-NLs. This review focuses onstudies in which HfxZr1−xO2 (HZO) thin films were mainlyused as ferroelectric films because, compared with the otherHfO2-based thin films with various dopants, they providestable ferroelectricity even with a low thermal budget (<400 °C)8–11) and even over a wide composition rangecentered on a Hf/Zr ratio of 1:1.25,66,72) In addition, weemployed a unique ALD approach for HZO films using a(Hf/Zr)[N(C2H5)CH3]4 (Hf:Zr = 1:1) cocktail precursor, inwhich Hf[N(C2H5)CH3]4 and Zr[N(C2H5)CH3]4 precursorsare pre-mixed in a ratio of 1:1.10,11,34,35)2. Crystal phase control using ALD-ZrO2-NLs2.1. Role of ALD-ZrO2-NLs in crystal phase controlFirst, we investigated NLs to promote the formation of theferroelectric O phase and enhance the ferroelectricity in HZOthin films. The crystallization onset temperature for ZrO2 isknown to be lower than that for HfO2. One of our previousstudies also revealed that a ZrO2 film deposited by ALD at300 °C using (C5H5)Zr[N(CH3)2]3 as a precursor formed apolycrystalline structure consisting predominantly of O,tetragonal (T), and cubic (C) phases even after the ALDprocess, whereas an ALD-HZO film formed an amorphousstructure, as determined by X-ray diffraction (XRD) analysis(Fig. 1(a)).34) In addition, the as-grown ALD-ZrO2 film withFig. 1. (a) XRD patterns for as-grown ZrO2 and HZO films deposited by ALD at 300 °C. (b) Cross-sectional TEM and EDS mapping images of as-grown ALD-ZrO2 film with thickness of 2 nm deposited on TiN substrate. Reproduced from Ref. 34 with permission of IOP Publishing. Copyright 2017The Japan Society of Applied Physics.080804-2© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJpn. J. Appl. Phys. 65, 080804 (2026) PROGRESS REVIEWTable I. Various ILs for interface engineering between HfO2-based thin film and electrodes.9,34–58)IL HfO2-based film AnnealingMaterial Thickness Material Thickness Temperature Annealing type TE/BE materialDevicetype Year ReferencesZrO2 0.5–5 nm HZO 10 nm 600 °C PDA TiN/TiN MFM 2017 Onaya et al.34)2 nm HZO 10–25 nm 400 °C–700 °C PDA TiN/TiN MFM 2019 Onaya et al.35)2 nm HZO 9.5 nm 900 °C PMA TiN/Si MFS 2019 Lee et al.36)1.5 nm HZO 12 nm 550 °C PMA TaN/Si MFS 2019 Xiao et al.37)10 nm HZO 10–30 nm 600 °C PDA TiN/TiN MFM 2020 Onaya et al.38)10 nm HZO 10 nm 600 °C PDA TiN/TiN MFM 2020 Onaya et al.39)2 nm HZO 10 nm 500 °C–800 °C PMA TiN/TiN MFM 2020 Lee et al.40)1 nm HZO 10 nm 300 °C–400 °C PMA TiN/TiN MFM 2021 Gaddam et al.41)1 or 2 nm La-doped HZO 10 nm 550 °C PMA TiN/TiN MFM 2021 Popovici et al.42)2 or 10 nm HZO 10 nm 300 °C PMA TiN/Si MFS 2022 Onaya et al.9)1 nm HZO 4.5 or 9.5 nm 300 °C–500 °C PMA TiN/TiN MFM 2023 Yu et al.43)1 nm HZO 13 nm 450 °C–650 °C PDA TiN/TiN MFM 2023 Song et al.44)1–2 nm HZO 12 nm 350 °C PDA Au/Ti/TiN MFM 2023 Liu et al.45)1 nm HZO 10 nm 450 °C PMA TiN/TiN MFM 2024 Li et al.46)2 nm HZO 8 nm 450 °C–600 °C PMA TiN/TiN MFM 2024 Yang et al.47)Al2O3 2 nm HZO 10 nm 600 °C PDA TiN/TiN MFM 2017 Onaya et al.34)2 nm HZO 9.5 nm 900 °C PMA TiN/Si MFS 2019 Lee et al.36)4–20 nm HZO 20 nm 500 °C PMA TiN/TiN MFM 2019 Si et al.48)0.8–3.2 nm HZO 21.6 nm 450 °C PMA TiN/TiN MFM 2021 Chen et al.49)1 nm HZO 10 nm 550 °C PDA TiN/W or TiN/TiN MFM 2021 Liu et al.50)1 nm HZO 4.5 or 9.5 nm 300 °C–500 °C PMA TiN/TiN MFM 2023 Yu et al.43)1 nm HZO 10 nm 800 °C Unknown TiN/TiN MFM 2023 Hsain et al.51)1 nm HZO 10 nm 450 °C PMA TiN/TiN MFM 2024 Li et al.46)TiO2 1 nm HZO 10 nm 300 °C–400 °C PMA TiN/TiN MFM 2021 Gaddam et al.41)2–3 nm HZO 10 nm As-grown or 500 °C PMA TiN/TiN MFM 2021 Qi et al.52)1–5 nm HZO 10 nm 400 °C Deposition temperature of TE-TiN TiN/TiN MFM 2022 Koroleva et al.53)1 nm HZO 10 nm 500 °C PMA W/W MFM 2024 Wang et al.54)1 nm HZO 10 nm 450 °C PMA TiN/TiN MFM 2024 Li et al.46)1 nm HZO 10 nm 300 °C–600 °C PMA TiN/TiN MFM 2025 Liu et al.55)HfO2 1–20 nm HZO 10 nm 600 °C PMA TiN/TiN MFM 2020 Gaddam et al.56)10 nm HZO 10 nm 600 °C PDA TiN/TiN MFM 2020 Onaya et al.39)1 nm HZO 10 nm 300 °C–400 °C PMA TiN/TiN MFM 2021 Gaddam et al.41)2 nm HZO 12 nm 550 °C PMA TaN/Si MFS 2021 Liu et al.57)3 nm HZO 15 nm 450 °C PMA TiN/Si MFS 2022 Liu et al.58)1 nm HZO 4.5 or 9.5 nm 300 °C–500 °C PMA TiN/TiN MFM 2023 Yu et al.43)Continued on next page.080804-3©2026TheAuthor(s).PublishedonbehalfofTheJapanSocietyofAppliedPhysicsbyIOPPublishingLtdJpn.J.Appl.Phys.65,080804(2026)PROGRESSREVIEWTable I. Continued.IL HfO2-based film AnnealingMaterial Thickness Material Thickness Temperature Annealing type TE/BE materialDevicetype Year References1 nm HZO 10 nm 450 °C PMA TiN/TiN MFM 2024 Li et al.46)HZO 2 nm HZO 10 nm 600 °C PDA TiN/TiN MFM 2017 Onaya et al.34)2.2 nm HZO 6.6 nm 550 °C–800 °C Deposition temperature of HZO film Pt/La0.67Sr0.33MnO3 MFM 2023 Song et al.59)2.2 nm HZO 6.8 nm 550 °C PMA W/W or W/Si MFM orMFS2025 Park et al.60)SiO2 1 nm HZO 10 nm 300 °C–400 °C PMA TiN/TiN MFM 2021 Gaddam et al.41)Ta2O5 1–5 nm HZO 10 nm 600 °C PMA TiN/TiN MFM 2020 Gaddam et al.61)HfOxNy 2–200 nm HZO 7–25 nm 500 °C PMA TiN/TiN MFM 2022 Kim et al.62)080804-4©2026TheAuthor(s).PublishedonbehalfofTheJapanSocietyofAppliedPhysicsbyIOPPublishingLtdJpn.J.Appl.Phys.65,080804(2026)PROGRESSREVIEWa thickness of 2 nm was found to be uniformly formed on aTiN substrate and to have crystallized without an annealingprocess, as determined by cross-sectional transmission elec-tron microscopy (TEM) and energy dispersive X-ray spectro-scopy (Fig. 1(b)). The ALD-ZrO2 film was determined to beadvantageous as an NL because the lattice constants for theO, T, and C phases in ZrO2 well match those for theferroelectric O phase in HfO2. Thus, the ALD-ZrO2 film isexpected to promote the formation of the ferroelectric Ophase as an NL in HZO films when an ALD-ZrO2 film isinserted between the HZO film and the electrode.Here, MFM capacitors were fabricated using three dif-ferent ALD-ILs with different crystal structures: crystallizedZrO2 (w/ZrO2-NL), amorphous Al2O3 (w/Al2O3), and amor-phous HZO (w/HZO), which were inserted between a 10-nm-thick amorphous HZO film and the BE-TiN (Fig. 2(a)).A TiN/HZO/TiN MFM capacitor without an ALD-IL (w/o)was also prepared as a reference. Post-deposition annealing(PDA) was performed at 600 °C for 1 min under a N2atmosphere before the formation of the TE-TiN. The 2Prvalues extracted from polarization–electric field (P–E) hys-teresis loops increased in the order w/Al2O3 (2.1 μC cm−2)< w/HZO (12.1 μC cm−2) < w/o (16.0 μC cm−2) <w/ZrO2-NL (22.1 μC cm−2) (Fig. 2(b)). In addition, in thew/ZrO2-NL case, the maximum 2Pr value was obtainedwhen the thickness of the ZrO2-NL was 2 nm (Fig. 2(c)).Notably, P–E measurements were carried out in the rangefrom −3.0 to 3.0 MV cm−1 because the breakdown electricfields for the w/o, w/HZO, w/Al2O3, and w/ZrO2-NLcapacitors were 3.3, 3.1, 3.2, and 3.1 MV cm−1, respectively,as evaluated by leakage current density–electric field (J–E)measurements (data not shown). The w/Al2O3 and w/HZOcapacitors fabricated using amorphous Al2O3 and HZO-ILs,respectively, did not show an improvement in 2Pr valuecompared with the w/o capacitor. In the w/Al2O3 case, theAl2O3-IL maintained an amorphous structure even after thePDA at 600 °C because of the higher crystallization onsettemperature of Al2O3 (>700 °C).73) This result suggests thatamorphous Al2O3-IL inhibited the crystallization and for-mation of the O phase in HZO films during the PDA process.In addition, most of the electric field (E) was applied not tothe ferroelectric HZO film but to the Al2O3-IL, which led toa further reduction in the 2Pr value because the k value foramorphous Al2O3 (k ≈ 9) is substantially lower than that forpolycrystalline HZO (20 ⩽ k ⩽ 40).63,64,66,67) In the w/HZOcase, the amorphous HZO film was crystallized after thePDA process but the 2Pr value did not improve. This lack ofimprovement in the 2Pr value might be attributable toenhanced formation of the stable M phase as a result ofthe increase in total thickness of the HZO films withadditional HZO-IL.26,74–76) By contrast, the w/ZrO2-NLcapacitor fabricated using a crystallized ZrO2-IL consistedpredominantly of O/T/C phases and exhibited a 2Pr valueapproximately 1.4 times higher than that for the w/ocapacitor.Two possible roles of the crystallized ZrO2-NL in enhan-cing the ferroelectricity of HZO films have been proposed.The first is that of a seed layer for the crystallization andformation of the ferroelectric O phase in HZO films.34,35)Fig. 2(d) shows cross-sectional TEM images of PDA-treatedFig. 2. (a) Schematics of four types of MFM capacitors: w/o, w/ HZO, w/ Al2O3, and w/ ZrO2-NL capacitors. (b) P–E hysteresis loops for four typesof MFM capacitors. (c) 2Pr and Ec values as functions of ZrO2-NL thickness. (d) Cross-sectional TEM images of PDA-treated w/o and w/ ZrO2-NLsamples. Reproduced from Ref. 34 with permission of IOP Publishing. Copyright 2017 The Japan Society of Applied Physics.080804-5© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJpn. J. Appl. Phys. 65, 080804 (2026) PROGRESS REVIEWHZO/BE-TiN (w/o) and HZO/ZrO2/BE-TiN (w/ZrO2-NL)samples. For the w/o sample, the HZO film consisted ofnumerous separate grains with a size of 5–10 nm. For thew/ZrO2-NL sample, on the other hand, no grain boundarieswere observed between the HZO film and the ZrO2-NL andthe lattice fringes of the HZO film were aligned with those ofthe ZrO2-NL. The ZrO2-NL crystallized even after the ALDprocess and formed the O/T/C phases (Fig. 1), which haveapproximately the same lattice constant as the O phase ofHZO. These results indicate that the amorphous HZO filmwas crystallized on the surface of the pre-crystallizedZrO2-NL, leading to epitaxial-like grain growth of theHZO film with the O/T/C phases during the PDA process.Consequently, the HZO/ZrO2-NL grain size (10–18 nm) islarger than that in the w/o case. A larger grain size isconsidered to promote phase transformation from the me-tastable O phase to the stable M phase as a result of thereduced influence of surface energy.26,77–80) On the otherhand, on the basis of first-principles calculations, Huanget al. reported that the ZrO2-NL effectively stabilizes the Ophase in the HfO2 film.81) Thus, the w/ZrO2-NL samplemaintained a higher fraction of O/T/C phases and exhibited ahigher 2Pr value because of the crucial roles the ZrO2-NLplayed in both the formation and stabilization of the O phasein HZO films. ZrO2 thin films with an amorphous structurecan be used as an IL. In this case, because the crystallizationonset temperature for ZrO2 is lower than that for HfO2, theamorphous ZrO2-IL crystallizes prior to HfO2 during theannealing process, thereby potentially acting as a NL.The second role of a ZrO2-NL is that of a stressor layer forthe formation of the O phase in HZO films during thePDA process, which is a result of a difference in thermalexpansion coefficients between the ZrO2-NL and theHZO films.9,41) As previously mentioned, Kim et al. reportedthat mechanical tensile stress during the annealing processis a key factor in promoting the formation of the metastableO phase.18) The thermal expansion coefficient for ZrO2(10.3× 10−6 K−1) is greater than that for HfO2(6.5× 10−6 K−1).68,82) Therefore, the ZrO2-NL can inducetensile stress in the HZO film during the PDA process,originating from the difference in thermal expansion coeffi-cients. To provide evidence supporting the effect of stress,Gaddam et al. investigated the relationship between thethermal expansion coefficient for ILs and the ferroelectricityof HZO films and found that the ZrO2-NL successfullypromoted the formation of the O phase in a HZO film.41) Onthe basis of the aforementioned experiments, we concludedthat ALD-ZrO2-NL enables crystal phase control inHfO2-based films because of its two critical roles as a seedlayer and a stressor layer.As previously explained, a 2 nm thick ALD-ZrO2-NLeffectively improved the 2Pr value for a 10 nm thick HZOfilm (Fig. 2(c)). For ZrO2-NL with a thickness up to 2 nm,the 2Pr value increased with increasing ZrO2-NL thickness,possibly because the ALD-ZrO2-NL, which has a polycrys-talline structure on the BE-TiN, transitioned from island-likegrowth to a continuous layer as the number of ALD cyclesincreased. In addition, the mechanical tensile stress origi-nating from the difference in thermal expansion coefficientscould be enhanced during the PDA process becauseof the increasing ZrO2-NL thickness. By contrast, when theZrO2-NL thickness exceeded 2 nm, both the 2Pr value andthe coercive electric field (Ec) decreased with increasingZrO2-NL thickness. Given that a ZrO2 single layer generallyexhibits antiferroelectric behavior, this decrease in the 2Prand Ec values is attributed to an increased proportion of theantiferroelectric component of the ZrO2-NL within theHZO/ZrO2-NL stack. Therefore, to achieve superior ferroe-lectricity, the thicknesses of HfO2-based films and ZrO2-NLsare required to be carefully selected.2.2. Effect of the insertion position of theALD-ZrO2-NLThe insertion position of the ALD-ZrO2-NL also affects thecrystallinity and ferroelectricity of HfO2-based films. Fourdifferent HZO-based MFM capacitors were fabricated without(w/o) and with a bottom ZrO2-NL (B-ZrO2), top ZrO2-NL(T-ZrO2), and double ZrO2-NLs (D-ZrO2), as shown inFig. 3(a). The thicknesses of the HZO film and ZrO2-NLwere 10 and 2 nm, respectively. PDA was carried out at 600 °Cfor 1 min under a N2 atmosphere before the TE-TiN wasformed. The 2Pr value for the MFM capacitors increased inthe order w/o (12 μC cm−2) < B-ZrO2 (15 μC cm−2) < T-ZrO2(23 μC cm−2) < D-ZrO2 (29 μC cm−2). These values wereextracted from the P–E hysteresis curves (Fig. 3(b)). The 2Prvalues for the MFM capacitors were improved using theZrO2-NL. A comparison of the 2Pr values for the MFMcapacitors fabricated using only a single ZrO2-NL revealsthat the T-ZrO2 capacitor exhibited a higher 2Pr value thanthe B-ZrO2 capacitor. Notably, the D-ZrO2 capacitor withtop and bottom ZrO2-NLs showed the highest 2Pr value,which was approximately 2.4 times greater than that for thew/o capacitor.To clarify the differences in 2Pr values for HZO films, weevaluated their crystallinity by TEM and XRD analysis(Figs. 3(c) and 3(d)). From cross-sectional TEM images,epitaxial-like grain growth between the HZO film and theZrO2-NLs was found to occur for the B-, T-, and D-ZrO2samples with ZrO2-NLs during the PDA process; bycontrast, random grain growth of the HZO film was observedfor the w/o sample. In addition, for the D-ZrO2 sample, grainboundaries were partially observed in the lateral direction inthe HZO film; these grain boundaries are considered to be aninterface where HZO grains growing from the top andbottom ZrO2-NLs collided. A grazing-incidence XRD (GI-XRD) pattern showed overlapped diffraction peaks at 2θ ≈30.5°, which are related to a mixture of the O, T, and Cphases (Fig. 3(d)); deconvolution of these peaks is difficultbecause of the extremely similar lattice constants for the O,T, and C phases and the limited resolution of conventionalXRD instruments that use a laboratory-based X-raysource.10) To compare the crystal phase fraction for theHZO films for four types of MFM capacitors, the relativeratio of O/T/C phases (rO/T/C) was estimated using therelation rO/T/C = AO(111)/T(101)/C(111)/{AO(111)/T(101)/C(111) +AM(−111) + AM(111)}, where AO(111)/T(101)/C(111), AM(−111),and AM(111) are the integrated peak areas of O(111)/T(101)/C(111), M(−111), and M(111) from the GI-XRD patterns,respectively.35) When ZrO2-NLs were used, the rO/T/Cincreased in the order w/o (52%) < B-ZrO2 (54%) <T-ZrO2 (62%) < D-ZrO2 (64%).To investigate the relationship between the crystallinityand ferroelectricity of HZO films, we plotted the 2Pr values080804-6© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJpn. J. Appl. Phys. 65, 080804 (2026) PROGRESS REVIEWas a function of rO/T/C (Fig. 3(e)). In this figure, the data forthe four types of MFM capacitors are plotted as a function ofthe HZO thickness from 10 to 25 nm. A clear positivecorrelation is observed, indicating that the enhancement ofthe O/T/C phase fraction directly contributed to the im-proved 2Pr value, consistent with previously reportedtrends.72) In addition, although the rO/T/C includes contribu-tions from the non-ferroelectric T and C phases, the 2Prvalue of ∼41 μC cm−2 estimated by extrapolating the hor-izontal axis of rO/T/C to 100% is in reasonable agreementwith the theoretical value (40–53 μC cm−2) for randomlyoriented polycrystalline HfO2-based films on the basis offirst-principles calculations.83) This result suggests that theseferroelectric films predominantly consisted of the ferro-electric O phase rather than the non-ferroelectric T and Cphases. As previously explained, the 2Pr and rO/T/C valuesfor the w/o sample are higher than those for the B-ZrO2sample because of the role the ZrO2-NL plays in O/T/Cphase formation. A comparison of the samples preparedusing a single ZrO2-NL reveals that the T-ZrO2 sampleexhibited a higher 2Pr and rO/T/C than the B-ZrO2 sample.Thin films are widely recognized as tending to crystallizepreferentially from the surface rather than from the bulkregion because of the lower activation energy for nucleationat the free surface. This behavior indicates that the topZrO2-NL plays a more dominant role than thebottom-ZrO2-NL in determining the crystallinity of HZOfilms, resulting in a higher 2Pr and rO/T/C for the T-ZrO2sample. The D-ZrO2 sample showed the highest 2Pr andrO/T/C because of the effect of both top and bottomZrO2-NLs. On the basis of these results, we found that thecrystallinity and ferroelectricity of HfO2-based films wereaffected by the insertion position of ZrO2-NLs, where bothtop and bottom ZrO2-NLs promote O/T/C phase formation.The dependence of the PDA process temperature on theferroelectricity was also investigated (Fig. 3(f)). All ofthe MFM capacitors exhibited their highest 2Pr valueswhen the PDA temperature was 600 °C. The D-ZrO2capacitors maintained a higher 2Pr value than the othercapacitors for PDA temperatures of 400 °C–700 °C. Inaddition, for the D-ZrO2 capacitor, a twofold higher 2Prvalue (18 μC cm−2) was obtained compared with that for thew/o capacitor (9.0 μC cm−2) even at the low annealingtemperature of 400 °C, which is compatible with the thermalbudget of BEOL processes.On the basis of these results, ZrO2-NLs play a crucial rolein promoting the formation of the O/T/C phases inHfO2-based films; in particular, the D-ZrO2 capacitor withboth top and bottom ZrO2-NLs exhibited the highest 2Pr andrO/T/C.2.3. CET scaling using ALD-ZrO2-NLVarious dopants for ferroelectric HfO2-based materials havebeen reported, including Si,1) Al,84,85) and Zr.2) Among suchdoped materials, the HZO film consisting of HfO2 and ZrO2Fig. 3. (a) Schematics of four types of MFM capacitors: w/o, B-ZrO2, T-ZrO2, and D-ZrO2 capacitors. (b) P–E hysteresis loops, (c) cross-sectionalTEM images, and (d) GI-XRD patterns for four types of MFM capacitors. (e) Relationship between rO/T/C and 2Pr value for four types of MFMcapacitors as function of HZO thickness. (f) 2Pr value for four types of MFM capacitors as function of PDA temperature. Reproduced from Ref. 35with permission of AIP Publishing, licensed under a Creative Commons Attribution (CC BY) license.080804-7© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJpn. J. Appl. Phys. 65, 080804 (2026) PROGRESS REVIEWwith a higher k value of approximately 20–40 is expected toexhibit a low CET value, whereas the lower k values for SiO2and Al2O3 of 3.9 and ∼9, respectively, may lead to anincrease of the CET value for Si- and Al-doped HfO2films.63,64,66,67) HfO2 and ZrO2 are well known to typicallyform four crystal phases of M, O, T, and C phases. Inaddition, the O, T, and C phases of HfO2 and ZrO2 exhibit kvalues greater than that for the stable M phase because the kvalue increases with increasing density according to theClausius–Mossotti relation.86) Thus, the formation of theHZO film with the O/T/C phases improves not only itsferroelectricity but also its k value, thereby leading to a lowCET value. In addition, the reduction of the CET value forferroelectric HZO films allows for a relatively greaterphysical thickness, which consequently leads to a reductionin the leakage current. In this subsection, we thereforediscuss the electrical characteristics (e.g. k, CET, leakagecurrent, and 2Pr) of the ferroelectric HZO films fabricatedusing the top and bottom ZrO2-NLs.The w/o, B-ZrO2, T-ZrO2, and D-ZrO2 MFM capacitorswere fabricated with different HZO thicknesses from 7.5 to25 nm while the ZrO2-NL thickness was maintained at 2 nm(Fig. 3(a)). The k–E properties of the four types of MFMcapacitors with a 10 nm thick HZO film are shown inFig. 4(a). The k value was extracted using the capacitance at0 MV cm−1. All capacitors exhibited butterfly-like hysteresisloops, which are a typical characteristic of ferroelectricmaterials. The overall k value for the ferroelectric film thatincluded both the HZO film and the ZrO2-NLs increased inthe order w/o (24) < B-ZrO2 (26) < T-ZrO2 (29) < D-ZrO2(32). This order is the same as that for the 2Pr values(Fig. 3(b)), indicating that the rO/T/C extracted from GI-XRDpatterns is also correlated with the k value. The D-ZrO2capacitor exhibited the highest k value, which was approxi-mately 1.3 times higher than that for the w/o capacitor.Figure 4(b) shows the overall k and CET values for thefour types of MFM capacitors as a function of the HZOthickness. In addition, the relationship between the rO/T/C andthe overall k value as a function of the HZO thickness isshown in Fig. 4(c). Various k values were obtained bychanging the insertion position of the ZrO2-NLs and thethickness of the HZO film; however, the overall k value forall of the capacitors increased linearly with increasing rO/T/C.These results indicate that the k value is closely correlatedwith the fraction of the O/T/C phases in the HZO film as wellas with the 2Pr value. The B-, T-, and D-ZrO2 capacitorswith the ZrO2-NLs exhibited higher k values than the w/ocapacitor because of a higher rO/T/C. The k value for all of thecapacitors linearly decreased with increasing HZO thicknessbecause a larger grain size caused by the thicker HZO filmspromotes the formation of the stable M phase rather than theO/T/C phases.26,74–76) Consequently, the CET value for all ofthe MFM capacitors increased linearly with increasing HZOthickness. Notably, the B-, T-, and D-ZrO2 capacitorsexhibited approximately the same CET values as the w/ocapacitor despite the additional thickness of the ZrO2-NLs,Fig. 4. (a) k–E properties of four types of MFM capacitors: w/o, B-ZrO2, T-ZrO2, and D-ZrO2 capacitors, with 10-nm-thick HZO film. (b) Overall kand CET values for four types of MFM capacitors as functions of HZO thickness. (c) Relationship between rO/T/C and overall k value and (d)relationship between overall CET and 2Pr value for four types of MFM capacitors as function of HZO thickness. (e) J–Eeff characteristics of four typesof MFM capacitors with 10 nm thick HZO film. (f) Relationship between overall CET value and J value at 0.6 V for four types of MFM capacitors asfunction of HZO thickness from 7.5 to 15 nm. The reference samples of ZrO2- and ZAZ-based MIM capacitors were compared with these four types ofMFM capacitors.67) © The Electrochemical Society. Reproduced from Ref. 95 with permission of IOP Publishing. All rights reserved.080804-8© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJpn. J. Appl. Phys. 65, 080804 (2026) PROGRESS REVIEWespecially in the HZO thickness region <15 nm, because oftheir higher k values. For the capacitors with a 10 nm thickHZO film, the overall CET value for the D-ZrO2 capacitor(1.7 nm) was approximately the same as that for the w/ocapacitor (1.6 nm), whereas the thickness of the overallferroelectric film (14 nm) was 40% greater than that for thew/o capacitor (10 nm).Figure 4(d) shows the relationship between the CET and2Pr value for the overall ferroelectric film for four types ofMFM capacitors. The four types of MFM capacitorsexhibited higher 2Pr values when the HZO thickness was<15 nm; however, the 2Pr value decreased dramatically withincreasing HZO thickness because of a lower rO/T/C.Notably, the D-ZrO2 capacitor with a 10 nm thick HZOfilm achieved the highest 2Pr value (29 μC cm−2) among thefour types of MFM capacitors. The 2Pr values for the fourtypes of MFM capacitors decreased with decreasing HZOthickness from 10 to 7.5 nm. This behavior might beattributable to the HZO film forming more T or C phasethan O phase. On the other hand, the D-ZrO2 capacitor with a7.5-nm-thick HZO film (CET = 1.4 nm) exhibited a 2Prvalue (25 μC cm−2) approximately 3.3 times greater than that(7.5 μC cm−2) for the w/o capacitor (CET = 1.2 nm).To further demonstrate the effects of the ZrO2-NL, theleakage current density–effective electric field (J–Eeff) char-acteristics of four types of MFM capacitors with a 10 nmthick HZO film are shown in Fig. 4(e). The Eeff wascalculated using the applied voltage and the overall CETvalue. The J values for the MFM capacitor fabricated usingthe ZrO2-NL were lower than that for the w/o capacitor. TheEeff at 10−7 A cm−2 increased in the order w/o(4.0 MV cm−1) < B-ZrO2 (5.1 MV cm−1) < T-ZrO2(6.1 MV cm−1) < D-ZrO2 (6.4 MV cm−1). The Eeff for theT-ZrO2 capacitor was greater than that for the B-ZrO2capacitor even though the capacitors had the same totalphysical thickness of 12 nm. This greater Eeff for the T-ZrO2capacitor is attributed to its lower CET value (1.8 nm)compared with that for the B-ZrO2 capacitor (1.6 nm), whichis a result of the different roles of the top and bottomZrO2-NLs, as discussed in the previous subsection. The Jproperties of the T- and D-ZrO2 capacitors were approxi-mately the same even though the physical thickness of theD-ZrO2 capacitor (14 nm) was greater than that of theT-ZrO2 capacitor (12 nm). This result can be explained bythe minimal difference in CET values between the T- andD-ZrO2 capacitors (1.6 and 1.7 nm, respectively) as a resultof the higher k value for the D-ZrO2 capacitor.Finally, the relationship between the overall CET valueand the J value at 0.6 V for four types of MFM capacitors asa function of HZO thickness from 7.5 to 15 nm is shown inFig. 4(f). For comparison, data for metal–insulator–metalcapacitors with a ZrO2 single layer (ZrO2) (k = 28) andZrO2/Al2O3/ZrO2 multilayers (ZAZ) (21 ⩽ k ⩽ 26), whichare typically used as insulators for dynamic random accessmemory capacitors, are also plotted.67) The J values for theB-, T-, and D-ZrO2 capacitors were lower than those for thew/o capacitor with almost the same profiles as the ZrO2capacitor. In addition, the D-ZrO2 capacitors showed lower Jvalues than the B- and T-ZrO2 capacitors. These results areattributed to a greater thickness of the overall ferroelectricfilms with both the top and bottom ZrO2-NLs. The D-ZrO2capacitor with a 7.5 nm thick HZO film exhibited a low Jvalue of 3.8× 10−8 A cm−2 while maintaining a lower CETvalue of 1.4 nm, which is comparable to that for the ZAZcapacitor. This lower CET value was achieved by main-taining a higher k value as a result of the ZrO2-NL effect,even as the physical thickness increased. On the basis ofthese experimental data, we concluded that electrical proper-ties such as the 2Pr, k, and J of the HfO2-based films weresubstantially improved by the incorporation of the top andbottom ZrO2-NLs.3. Improvement of reliability using ALD-ZrO2-NLsSince the first report of ferroelectric HfO2-based materials in2011, these materials have moved closer to being used inpractical applications. Currently, to ensure reliability ofHfO2-based ferroelectric devices, numerous researchers arefocusing on the critical challenges regarding reliability. Toachieve a higher breakdown voltage (Vbd) and superiorendurance properties while maintaining high 2Pr values,MFM capacitors with thick HZO films were fabricated usingALD-ZrO2-NLs.For the fabrication of HZO-based MFM capacitors, a topZrO2-NL was inserted between the HZO film and TE-TiNand the thickness of the HZO films was varied from 10 to30 nm. To achieve a higher Vbd, a thick top ZrO2-NL with athickness of 10 nm was used because HZO-based MFMcapacitors with a 10 nm thick top ZrO2-NL maintained ahigh 2Pr value while suppressing the influence of theantiferroelectric behavior of ZrO2 according to the relation-ship between the 2Pr value and ZrO2-NL thickness (data notshown).38) A PDA process was performed at 600 °C for1 min under a N2 atmosphere before the formation of the TE-TiN. The MFM capacitors with a HZO single film (w/o) andHZO/ZrO2-NL bilayer (w/ZrO2-NL) are represented as“HZ#” and “HZ#Z#,” respectively, where HZ or Z and theattached number represent the film type and thickness,respectively. For example, HZ20Z10 represents an MFMcapacitor with bilayer consisting of a 10 nm thick ZrO2-NLon top of a 20 nm thick HZO film.Figure 5(a) shows P–E hysteresis loops for the w/o andw/ZrO2-NL MFM capacitors. The relationship between thetotal thickness and 2Pr values extracted from these P–Ehysteresis loops is summarized in Fig. 5(b). For the w/ocapacitor cases, the 2Pr value was 10 μC cm−2 for the HZ10capacitor; however, the 2Pr value decreased substantially asthe HZO thickness increased, consistent with previouslyreported trends.26,74–76) Notably, the 2Pr value for the HZ25capacitor is approximately half that for the HZ10 capacitor.This low 2Pr value is attributed to a decrease in theformation of the O phase and a concomitant increase in thefraction of the stable M phase with increasing thickness ofthe HZO films, as evaluated by GI-XRD analysis (data notshown). As the grain size increases with increasing thicknessof the polycrystalline HZO films, a phase transformationfrom the metastable O phase to the stable M phase typicallyoccurs because of the diminishing contribution of surfaceenergy.26,77–80) This mechanism makes simultaneouslyachieving both increased HZO thickness and robust ferroe-lectricity difficult. For the w/ZrO2-NL capacitor cases, theHZ10Z10 capacitor exhibited a hybrid hysteresis loop with amixture of ferroelectricity and antiferroelectricity because of080804-9© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJpn. J. Appl. Phys. 65, 080804 (2026) PROGRESS REVIEWa higher proportion of antiferroelectric ZrO2, resulting in asmall 2Pr value. On the other hand, clear hysteresis loopswere observed when the HZO thickness was increased to15 nm or more. Notably, the HZ15Z10 capacitor achieved amaximal 2Pr value of 14 μC cm−2, which is higher than thatfor the HZ10 capacitor (10 μC cm−2), even though theHZ15Z10 capacitor had a relatively large total thickness of25 nm. This result suggests that the ZrO2-NL could alsoexhibit ferroelectricity when the thickness ratio between theHZO film and the ZrO2-NL was properly adjusted. Atthe same total thickness of 25 nm, the 2Pr value for theHZ15Z10 capacitor (14 μC cm−2) was approximately 2.5times higher than that for the HZ25 capacitor (5.5 μC cm−2).Even for the w/ZrO2-NL capacitors, a decrease in 2Pr valuewith increasing HZO thickness was unavoidable; however,the w/ZrO2-NL capacitors maintained a higher 2Pr valuethan the w/o capacitors. Thus, the ALD-ZrO2-NL enabled anincrease in HZO thickness without compromising its high2Pr value.Figure 5(c) shows the relationship between the totalthickness and the Vbd, which was extracted from the J–Vcharacteristics shown in the inset, for the w/o and w/ZrO2-NLcapacitors. The Vbd for the HZ10 capacitor was 3.6 V(E= 3.6 MV cm−1), which is consistent with the typicallyreported breakdown E for HZO films (3–4 MV cm−1).25) Forboth types of MFM capacitors, Vbd increased monotonicallywith increasing HZO thickness. The relationship between theVbd and the 2Pr value for both capacitors, as obtained fromFigs. 5(a) and 5(c), is shown in Fig. 5(d). The profile for thew/ZrO2-NL capacitors is shifted toward the higher Vbd regioncompared with that for the w/o capacitors. In particular, acomparison of capacitors with the same HZO thickness of15 nm reveals that the 2Pr value and Vbd for the HZ15Z10capacitor (14 μC cm−2 and 6.7 V, respectively) were approxi-mately 1.4 and 1.6 times higher than those for the HZ15capacitor (10 μC cm−2 and 4.3 V, respectively). Thesefindings demonstrate that both a higher 2Pr and a higher Vbdcan be achieved by thick HZO films fabricated using theALD-ZrO2-NL.We next investigated the endurance properties of thesecapacitors, which remain a major challenge for their use inpractical devices. Figure 5(e) shows the endurance propertiesof the HZ15 and HZ20 capacitors for the w/o capacitor casesand the HZ15Z10 and HZ20Z10 capacitors for thew/ZrO2-NL capacitor cases, as evaluated by positive-upnegative-down measurements. The switching polarization(Psw) for the HZ15 and HZ20 capacitors clearly increasedbecause of the wake-up effect up to ∼104 cycles, reachingvalues approximately 15% and 11% higher than initial Pswvalues of these capacitors, respectively. By contrast, theHZ15Z10 and HZ20Z10 capacitors with the ZrO2-NLshowed minimal wake-up, with an increase rate of lessthan 5%. The wake-up effect has been reported to originatefrom a field-induced phase transformation from the non-ferroelectric phase to the ferroelectric O phase.29,87–90) Thus,the suppressed increase rate during wake-up field cycling forFig. 5. (a) P–E hysteresis loops for w/o and w/ZrO2-NL capacitors. (b) 2Pr values and (c) Vbd, extracted from J–V characteristics shown in inset, ofw/o and w/ZrO2-NL capacitors as function of HZO thickness. (d) Relationship between Vbd and 2Pr values for w/o and w/ZrO2-NL capacitors. (e)Endurance properties of HZ15 and HZ20 capacitors for w/o capacitor cases and of HZ15Z10 and HZ20Z10 capacitors for w/ZrO2-NL capacitor cases,as evaluated by PUND measurements. Reproduced from Ref. 38 with permission of AIP Publishing.080804-10© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJpn. J. Appl. Phys. 65, 080804 (2026) PROGRESS REVIEWthe HZ15Z10 and HZ20Z10 capacitors is attributed to ahigher initial fraction of the O phase compared with that inthe HZ15 and HZ20 capacitors, as evaluated by GI-XRDanalysis (data not shown).38) In addition, the wake-up effectis speculatively attributed to domain depinning inducedmainly by the redistribution of VO, which are generallyformed around grain boundaries and interfaces between theHZO film and TiN electrodes.29,91,92) Therefore, we con-sidered that larger grains, formed through the epitaxial-likegrain growth of the HZO film on the ZrO2-NL for thew/ZrO2-NL capacitors, reduced the density of grain bound-aries, thereby suppressing the increase in the Psw caused bythe wake-up effect. Subsequent to the wake-up state, afatigue phenomenon (i.e. the degradation of Psw withincreasing number of switching cycles) were observed. ThePsw values for the HZ15 and HZ20 capacitors at 108 cyclesdecreased by 42% and 34%, respectively, compared withtheir peak values at 104 wake-up cycles. By contrast, thereduction rate for the Psw value from 104 to 108 cycles for theHZ15Z10 and HZ20Z10 capacitors was approximately 30%,indicating that the fatigue for the w/ZrO2-NL capacitors withthe ALD-ZrO2-NL was suppressed compared with that forthe w/o capacitors. New defects, such as VO, have beenreported to be preferentially generated along grain bound-aries during field cycling, leading to domain-pinning-in-duced degradation of ferroelectricity.29,93) For thew/ZrO2-NL capacitors, the insertion of the ZrO2-NL pro-moted continuous grain growth across the HZO film and theZrO2-NL, leading to an increase in grain size. This isconsidered to have reduced the density of grain boundaries,thereby suppressing the degradation of the Psw values. Thisinterpretation is also consistent with the observation that thePsw degradation was suppressed in the HZ20 capacitor with athicker HZO film (which causes an increase in the grain size)compared with that in the HZ15 capacitor. In addition,Onaya et al. reported that one of the origins of VO formationin HZO films is the field-induced interface reaction, whereoxygen atoms are scavenged from the HZO film into TiNelectrodes.31–33,94) Furthermore, the MFM capacitor with thetop and bottom ZrO2-NLs (D-ZrO2) showed higher fatigueresistance than that without the ZrO2-NL (w/o) becausethe ALD-ZrO2-NL could function as a blocking layerfor oxygen diffusion, supporting the suppression of suchinterface reactions during field cycling (Figs. 6(a) and 6(b)).32,95) Accordingly, the ZrO2-NL was considered tofurther effectively suppress Psw degradation for thew/ZrO2-NL capacitors. These findings indicate that theALD-ZrO2-NL is highly effective in improving reliabilityas well as in promoting formation of the O phase inHfO2-based thin films.4. Low-temperature fabrication process usingALD-ZrO2-NLsRecently, fabrication processes with a low thermal budget forHfO2-based thin films have attracted widespread attention inattempts to enhance the versatility of these films for variousapplications such as BEOL integration and flexible deviceswith thermally sensitive substrates.8–11,96) In addition, low-temperature fabrication can effectively suppress the formationof unexpected ILs between HfO2-based thin films andsubstrates or electrodes. In particular, the formation of a thickand low quality SiO2-IL between the HfO2-based film and theSi substrate is a critical issue for the fabrication of ferroelectricHfO2-based FeFETs because such SiO2-ILs severely degradedevice performance and reliability.24,97–99) In previous re-search, Onaya et al. reported the relationship between processtemperature and SiO2-IL formation for ALD-HZO/SiO2-IL/Sisamples evaluated using X-ray photoelectron spectroscopyanalysis (Fig. 7(a)).9,100) The peak attributed to Si–O bonds at∼103.5 eV was negligibly small even after the ALD growth ofthe HZO film at 300 °C, showing a peak area comparable tothat for a Si substrate after treatment with buffered hydro-fluoric acid. Similarly, the peak area remained equivalent tothat for the as-grown sample after the post-metallizationannealing (PMA) process at 300 °C. By contrast, when thePMA temperature was 400 °C or greater, the Si–O peak areaincreased linearly with increasing PMA temperature. Theseresults indicate that no substantial growth of the SiO2-ILoccurs during the fabrication process when both the ALD andthe annealing temperature are 300 °C or less. In addition,cross-sectional TEM images reveal that a SiO2-IL with athickness of one or two monolayers (0.3–0.6 nm) was formedbetween the HZO film and the Si substrate after the PMAprocess at 300 °C (Fig. 7(b)). This SiO2-IL is remarkably thincompared with those typically reported in the literature(⩾1 nm).101–105)Fig. 6. (a) Endurance properties and (b) normalized Psw values for w/o and D-ZrO2 MFM capacitors with HZO and ZrO2-NL thicknesses of 10 and2 nm, respectively.080804-11© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJpn. J. Appl. Phys. 65, 080804 (2026) PROGRESS REVIEWOn the basis of this background, a 300 °C PMA processwas carried out following the fabrication of TE-TiN tofabricate an MFS structure, which is a fundamental structureof FeFETs. In addition, to promote the crystallization andformation of the O phase and enhance the fatigue resistance,the ALD-ZrO2-NL was integrated into the HZO-based MFSstructures. Figure 8(a) shows cross-sectional TEM images ofthe TiN/HZO (10 nm)/SiO2-IL/p+-Si (w/o) and TiN/ZrO2-NL(10 nm)/HZO (10 nm)/SiO2-IL/p+-Si (ZrO2–10 nm) MFS ca-pacitors after the PMA process at 300 °C. The GI-XRDpatterns for both capacitors are shown in Fig. 8(b). Notably,a suppressed SiO2-IL with a thickness of one or twomonolayers (0.3–0.6 nm) for both capacitors was formedbetween the HZO film and the Si substrate, consistent withthe results in Fig. 7. For the w/o capacitor, a few nanocrystalswith a grain size of 5–10 nm partially formed in the HZO film,whereas most of the film remained amorphous after the PMAprocess at 300 °C. This result is attributed to the thermalexpansion coefficient for Si (2.5× 10−6 K−1) being smallerthan that for TiN (9.4× 10−6 K−1), whereas the HZO filmsandwiched with TE- and BE-TiN was crystallized even afterthe PMA process at 300 °C.10,11,19,82) On the other hand, theHZO film in the ZrO2-10 nm capacitor was fully crystallizedwith the same orientation as the ZrO2 grains and the grain sizewas 10–20 nm. In addition, the O/T/C phases were dominantlyformed for the ZrO2–10 nm capacitor. Consequently, a higher2Pr value was obtained for the ZrO2-10 nm capacitor(15 μC cm−2) than for the w/o capacitor (2.2 μC cm−2), asshown in Fig. 8(c). Furthermore, the ZrO2-10 nm capacitorexhibited a higher 2Pr value than the 2 nm thick ZrO2-NL(ZrO2-2 nm). This result might be attributable to a thickerZrO2-NL enhancing the tensile stress resulting from thedifferences in the thermal expansion coefficient during thePMA process, which promotes the crystallization and forma-tion of the O phase in the HZO film.18–22,106) These resultsindicate that the ZrO2-NL plays a critical role in the crystal-lization and formation of the ferroelectric O phase of HZOfilms during the low-temperature PMA process at 300 °C,even on a Si substrate. Notably, the fabrication temperature of300 °C achieved in this work is considerably lower than thosetypically reported in previous studies (⩾400 °C).8,11,24) Theendurance properties of the capacitors were also evaluated(Fig. 8(d)). No wake-up effect was observed for theZrO2-2 nm and ZrO2-10 nm capacitors, whereas the Pswvalue for the w/o capacitor clearly increased with in-creasing number of switching cycles up to ∼103 cycles.These wake-up-free characteristics are attributable to thepreferential formation of the O phase and to the reducedgrain boundary density resulting from the increase in grainsize due to the ZrO2-NL effect, as noted previously.9) Inaddition, the degradation of the Psw value was suppressedcompared with that observed for the w/o capacitor becauseof the role the ZrO2-NL played in the growth of largergrains and because the ZrO2-NL as a blocking layerprevented oxygen movement at the TE-TiN/HZO interface,as explained previously.9,31,32) On the basis of these results,we found that wake-up-free properties and higher fatigueresistance, as well as superior ferroelectricity for HZO-based MFS capacitors, were achieved using the ZrO2-NL,even at a low thermal budget of 300 °C.Fig. 7. (a) Relationship between process temperature and area of Si–O peak in XPS Si 2p spectra for as-grown and PMA-treatedALD-HZO/SiO2-IL/Si samples. The inset shows XPS spectra for the Si 2p core level. (b) Cross-sectional TEM images of as-grown and 300 °C-PMA-treated ALD-HZO/SiO2-IL/Si samples. Reproduced from Ref. 9 with permission of AIP Publishing, licensed under a Creative Commons Attribution(CC BY) license.080804-12© 2026 The Author(s). Published on behalf ofThe Japan Society of Applied Physics by IOP Publishing LtdJpn. J. Appl. Phys. 65, 080804 (2026) PROGRESS REVIEW5. ConclusionsEngineering of the interface between ferroelectric films andelectrodes has attracted considerable attention for the prac-tical application of HfO2-based ferroelectric devices. In thepresent study, we focused on the fabrication of HfO2-basedthin films using ALD-ZrO2-NLs inserted between a ferro-electric film and electrodes. The HZO films fabricated usingALD-ZrO2-NLs promoted the formation of and stabilized theferroelectric O phase, resulting in high 2Pr values. The rolesof the ALD-ZrO2-NLs in controlling the crystal phase ofHZO films can be explained by two primary mechanisms.First, the ALD-ZrO2-NLs were crystallized even after theALD process, forming crystal phases whose lattices wellmatch that of the ferroelectric O phase of HfO2, whereasHfO2-based thin films typically formed an amorphousstructure. Thus, the ALD-ZrO2-NL serves as a seed layerthat promotes the crystallization and formation of the Ophase in HfO2-based thin films during annealing processes.Second, the ALD-ZrO2-NL is considered to play a role as astressor layer: the difference in thermal expansion coeffi-cients between HfO2 and ZrO2 leads to mechanical tensilestress on HfO2-based thin films during an annealing process,thereby promoting the formation of the ferroelectric O phase.The effects of the ALD-ZrO2-NL were found to varydepending on their insertion position. The top ZrO2-NLinserted at the TE-TiN/HZO interface is more effective forthe O phase formation in HfO2-based thin films than thebottom ZrO2-NL inserted at the BE-TiN/HZO interface.Notably, the HZO films fabricated using a combination ofboth the top and bottom ZrO2-NLs exhibited a morepronounced O phase formation and higher 2Pr valuescompared with films fabricated using either the top or bottomZrO2-NL. A higher k value was also obtained because of theformation of a larger portion of the O/T/C phases, resultingin a lower CET value. As a result, when ALD-ZrO2-NLswere used, a lower J value was achieved while maintaining alow CET value, even in cases of a thicker ferroelectric film.In addition, by incorporating the ALD-ZrO2-NL, the typicaldegradation of 2Pr values with increasing HfO2-based filmthickness was successfully suppressed, allowing for robust2Pr values even in thicker films with a high Vbd. With respectto the endurance properties of HfO2-based ferroelectricdevices, the insertion of ALD-ZrO2-NLs led to superiorfatigue resistance because of the reduced grain boundarydensity resulting from the growth of larger grains; it also ledto the prevention of interface reactions during field cycling.Low-temperature fabrication of the HZO films at 300 °C wasalso demonstrated using the ALD-ZrO2-NL, resulting in thesuppression of growth of the SiO2-IL with a thickness of oneor two monolayers (0.3–0.6 nm) at the HZO film and the Sisubstrate.On the basis of these results, interface engineering usingALD-ZrO2-NLs is a promising approach for the fabricationof high-performance HfO2-based ferroelectric devices.AcknowledgmentsThis study was supported in part by JSPS KAKENHI (Nos.JP24K17304, JP21J01667, JP20H02189, and JP18J22998),MEXT Leading Initiative for Excellent Young Researchers(No. JPMXS0320220213), and The Samco Foundation. 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Introduction 2. Crystal phase control using ALD-ZrO2-NLs 2.1. Role of ALD-ZrO2-NLs in crystal phase control 2.2. Effect of the insertion position of the ALD-ZrO2-NL 2.3. CET scaling using ALD-ZrO2-NL 3. Improvement of reliability using ALD-ZrO2-NLs 4. Low-temperature fabrication process using ALD-ZrO2-NLs 5. Conclusions Acknowledgments