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Highly Oriented Growth of Piezoelectric Thin Films on Silicon using Two-Dimensional Nanosheets as Growth Template Layer Minh D Nguyen, Huiyu Yuan, Evert P. Houwman, Matthijn Dekkers, Gertjan Koster, Johan E. ten Elshof, and Guus Rijnders ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b09470 • Publication Date (Web): 26 Oct 2016 Downloaded from http://pubs.acs.org on October 28, 2016

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ACS Applied Materials & Interfaces

Highly Oriented Growth of Piezoelectric Thin Films on Silicon using Two-Dimensional Nanosheets as Growth Template Layer Minh D. Nguyen,*,†,‡,§ Huiyu Yuan,† Evert P. Houwman,† Matthijn Dekkers,‡ Gertjan Koster,† Johan E. ten Elshof,† Guus Rijnders*,† †

MESA+ Institute for Nanotechnology, University of Twente, P.O. Box 217, 7500AE Enschede,

The Netherlands ‡

Solmates B.V., Drienerlolaan 5, 7522NB Enschede, The Netherlands

§

International Training Institute for Materials Science, Hanoi University of Science and

Technology, Dai Co Viet 1, Hanoi 10000, Vietnam KEYWORDS. PiezoMEMS, PZT thin films, Nanosheets, Microcantilevers, Piezoelectric coefficients. ABSTRACT. Ca2Nb3O10 (CNOns) and Ti0.87O2 (TiOns) metal oxide nanosheets (ns) are used as a buffer layer for epitaxial growth of piezoelectric capacitor stacks on Si and Pt/Ti/SiO2/Si (Pt/Si) substrates. Highly (001) and (110)-oriented Pb(Zr0.52Ti0.48)O3 (PZT) films are achieved by utilizing CNOns and TiOns, respectively. The piezoelectric capacitors are characterized by polarization and piezoelectric hysteresis loops and by fatigue measurements. The devices fabricated with SrRuO3 top and bottom electrodes directly on nanosheets/Si have ferroelectric and piezoelectric properties well comparable with devices using more conventional oxide buffer layers (stacks), such as YSZ, CeO2/YSZ or SrTiO3 on Si. The devices grown on nanosheets/Pt/Si with Pt top electrodes show significantly improved polarization fatigue, properties over those of

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similar devices grown directly on Pt/Si. The differences in properties are ascribed to differences in the crystalline structures and the density of the films. These results show a route towards the fabrication of single crystal piezoelectric thin films and devices with high quality, long-lifetime piezoelectric capacitor structures on non-perovskite and even non-crystalline substrates, such as glass or polished metal surfaces. INTRODUCTION The perovskite Pb(ZrxTi1-x)O3 (PZT) family has been one of the most popular materials for a wide variety of ferroelectric and piezoelectric applications,1 especially with the composition near the morphotropic phase boundary (MPB) with x = 0.52. This composition has in bulk form the largest dielectric constant and piezoelectric coefficients of the PZT family. A major step forward is the deposition of ferroelectric PZT thin films on Si substrates, opening the route towards the integration of ferroelectric devices in Si-technology,2-6 for example in microelectromechanical systems (MEMS) devices as cantilevers or membranes for sensing, actuation or energy harvesting applications. In MEMS technology, the most common substrate for micromachining is silicon due to its reliable and reproducible mechanical and electrical properties, its widespread use in electronics and relatively low cost and good availability.7,

8

PZT thin films, however, cannot be grown

directly on Si because lead diffuses into the silicon substrate to form lead silicates at the growth temperature (500-650 oC) of PZT thin films and the large lattice mismatch between them. Therefore buffer layers between the PZT thin film and the silicon substrate, such as yttriastabilized zirconia (YSZ), CeO2 and SrTiO3, are needed to prevent the interdiffusion and oxidation reactions in order to grow highly oriented (epitaxial) PZT thin films.6, 9, 10 The growth of YSZ and CeO2 buffer layers on Si substrates however generally requires a very high thermal


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budget (800 oC),11,

12

where the SrTiO3 process on silicon, generally performed by Molecular

Beam Epitaxy, is very delicate and slow. These aspects are seemingly problematic for the flawless integration of PZT films in mainstream CMOS. Therefore new buffer layers, which can be fabricated at a lower temperature, are required. Recently two-dimensional oxide nanosheets have attracted great attention for both fundamental and industrial reasons because of their novel physical properties and promising potential for a wide range of applications such as piezoelectric MEMS and high-k dielectrics.13 Nanosheets are crystalline sheets of metal oxides with a thickness of the order of a single nanometer and lateral dimensions into the micrometer range. Shibata et al.14 have demonstrated the selective growth of highly textured (100), (110) and (111) oriented SrTiO3 films on Ca2Nb3O10 (CNOns), Ti0.87O2 (TiOns) and MoO2 nanosheets buffered glass substrates, respectively. The monolayer nanosheet films were deposited on glass using the LangmuirBlodgett (LB) deposition method. CNOns and TiOns were also deposited on silicon substrates and used to control the nucleation and orientation of SrRuO3 films during growth using pulsed laser deposition (PLD).15 Moreover, in order to obtain epitaxial single-crystal PZT films, studies are mainly based on single-crystal (001)-oriented SrTiO3 (STO) substrates due to a favorable lattice match between film and substrate.16, 17 However, STO substrates are impractical for use in actual devices due to their high cost and lack of large substrate sizes. Furthermore the PZT grown on STO is difficult to fabricate into devices MEMS devices. It is therefore important to investigate the growth of PZT films on substrates more suited to industrial applications. One such instance, Pt/Si, has found widespread use in the production of microsystems with PZT films. However quite some technological challenges remain for the integration of PZT films on Pt/Si, such as obtaining


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acceptable film quality and the control of fatigue degradation of ferroelectric and piezoelectric properties. The use of nanosheets on Pt/Si is believed to tackle such challenges. In the present study, CNOns and TiOns were synthesized and transferred onto Si and Pt/Ti/SiO2/Si substrates using the LB method. Pb(Zr0.52Ti0.48)O3 films were then deposited on SrRuO3/ns/Si as well as on ns/Pt/Ti/SiO2/Si substrates using PLD. The influence of the nanosheet buffer layers on the microstructure, ferroelectric and piezoelectric properties of PZT films and devices has been investigated. EXPERIMENTAL SECTION Materials. Titanium (IV) dioxide TiO2 (≥ 99%. technical) and lithium carbonate Li2CO3 ((≥99%) were obtained from Ridel-de Haen; calcium carbonate CaCO3 (ACS reagent, standard) and niobium (IV) oxide Nb2O5 (99.99%)) were purchased from Sigma-Aldrich; anhydrous potassium carbonate K2CO3 (≥99%) was received from Fluka; nitric acid HNO3 (65% in H2O) from ACROS Organics and tetra-n-butylammonium hydroxide TBAOH (40wt% in H2O) from Alfa Aesar were used as received. Demineralized water was used throughout the experiments. Single crystal boron doped (100) silicon wafers (grown by the Czochralski process) were purchased from Okmetic. Stoichiometric targets of SrRuO3 and Pb(Zr0.52Ti0.48)O3 were obtained from Praxair electronics. Synthesis nanosheets. Ti0.87O2 and Ca2Nb3O10 nanosheet films were fabricated on Si and Pt/Ti/SiO2/Si (Pt/Si) substrates by exfoliation of layered protonated titanate, H1.07Ti1.73O4•H2O (HTO) and protonated calcium niobate, HCa2Nb3O10•1.5H2O (HCNO) followed by the LB method. The LB method can be used to deposit the nanosheet films on both hydrophilic (e.g., Si) and hydrophobic (e.g., Pt/Si) substrates.18 The details of the flux synthesized, layered precursor K0.8[Ti1.73Li0.27O4] (KLTO) and the solid state synthesized, layered precursor KCa2Nb3O10


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(KCNO) and their protonation process can be found in previous reports.14, 19-21 The exfoliation of HTO and HCNO crystals was carried out by reaction with an aqueous solution of tetra nbutylammonium hydroxide (TBAOH). In detail, 0.1 g of HTO or 0.5 g HCNO powder was mixed with water and TBAOH with a TBA+/H+ (H+ refers to the protons of HTO or HCNO powder) molar ratio of 4:1 for HTO and 2:1 for HCNO. The total volume of the HTO and HCNO solutions was 20 and 100 mL, respectively. The mixtures were stirred for 2 h and 12 h for HTO and HCNO, respectively.19 Then 2 ml resulting TiOns suspension or 5 ml resulting CNOns suspension were diluted to 500 ml by adding demineralized water. Before LB deposition the diluted suspensions were left standing for 1 h. Then, 50 ml diluted solution was separated from the middle/top part of the nanosheet suspension by using a syringe and transferred into the LB through. The nanosheet suspension was left in the LB trough for 10 min before deposition on the Si substrates started. For more information on the LB device and substrate cleaning procedures readers may find further details in ref. [20]. Film deposition. (i) Si substrates: Pb(Zr0.52Ti0.48)O3 (PZT) thin films with a thickness of 1 µm were grown on 100nm-thick SrRuO3 (SRO) bottom electrodes, that were grown on the buffered CNOns/Si and TiOns/Si substrates using pulsed laser deposition (PLD) (Lambda Physik, KrF excimer laser and 248 nm wavelength). The deposition conditions of the PZT films were: laser repetition rate 10 Hz, energy density 2.5 J cm-2, oxygen pressure 0.1 mbar and substrate temperature 600 °C. For the SRO bottom- and top-electrodes the deposition parameters were 4 Hz, 2.5 J cm-2, 0.13 mbar O2 and 600 oC. After deposition the samples were cooled down to room temperature in a 1 bar oxygen atmosphere at a ramp rate of 8 oC min-1;


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(ii) Pt/Si substrates: a 125-nm-thick Pt film and a 15-nm-thick Ti film were dc-sputtered at room temperature onto a 500-nm-thick wet thermally grown SiO2 layer to serve as the bottom electrode and as adhesion layer, respectively. The 1-µm-thick PLD-deposited PZT films were then grown on the Pt/Si, CNOns/Pt/Si and TiOns/Pt/Si substrates. In the case of the Pt/Si substrate, a 10-nm-thick PLD-deposited LaNiO3 layer was used on top of the Pt bottom electrode as a growth seed layer for the PZT. After PZT films deposition, the samples were cooled down to room temperature in a 1 bar oxygen atmosphere and at a ramp rate of 8 oC min-1. The 125-nmthick Pt top-electrodes were deposited onto the PZT films at room temperature by dc sputtering. 200×200 µm2 capacitor structures were defined using conventional lithography, Ar ion beam etching and chemical wet etching down to the bottom electrode. Piezoelectric microcantilever device fabrication. The process for fabricating piezoelectric driven Si microcantilevers has been described in a previous paper22 and is shown schematically in Figure S10 (Supporting Information). Cantilever structures, with a beam dimension of about 400 µm × 100 µm, consist of a SRO/PZT/SRO piezoelectric stack on (CNOns, TiOns)/Si, or a Pt/PZT piezoelectric stack on (CNOns, TiOns, LNO)/Pt/Si. The10-µm thick Si supporting layer was obtained by backside etching of the silicon-on-insulator (SOI) wafer. Analysis and Characterization. The crystalline structure of the thin films was analyzed by Xray diffraction θ-2θ

scans (XRD, Philips X’Pert X-ray diffractometer). Film surface and

microstructure was investigated using atomic force microscopy (AFM, Bruker Dimension ICON) and high-resolution scanning electron microscopy (HRSEM, Zeiss 1550). Electron backscatter diﬀraction (EBSD) was performed on a Merlin ﬁeld emission microscope (Zeiss 1550) equipped with an angle selective backscatter detector. The polarization hysteresis (P-E) loop measurements were performed with the ferroelectric mode of the aixACCT TF-2000


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Analyzer using a triangular ac-electric field of ±200 kV cm-1 at 1 kHz scanning frequency. The polarization-switching cycle measurements were performed with a bipolar switching pulse of 100 kV cm-1 pulse height and 5 µs rectangular pulse width at 100 kHz repetition frequency, while the P-E loops were measured at increasing cycling intervals at ±200 kV cm-1 and 1 kHz frequency. The longitudinal piezoelectric coefficient (d33,f) of the piezoelectric thin-film capacitor structures was measured by a Double Beam Laser Interferometer (aixDBLI) using a locked-in technique with a dc driving electrical field between ±200 kV cm-1 and an ac field amplitude of 5 kV/cm and 1 kHz frequency. The transverse piezoelectric coefficient d31,f was determined from the tip-displacement of the cantilevers, driven by the piezoelectric stack with an ac-amplitude of 3 V (6 Vpeak-peak) at a dc offset field of 3 V and 8 kHz frequency (see Supporting Information, Figure S11). The selected frequency was much smaller than the natural resonant frequency of the prepared cantilevers.

RESULTS AND DISCUSSION PZT thin films on nanosheets buffered Si(001) substrates Figure 1 shows typical atomic force microscopy (AFM) images of a monolayer of CNOns and TiOns on Si(001) substrates. The lateral sizes of the obtained nanosheets are about 3-5 µm for CNOns and 5-8 µm for TiOns, respectively. A relatively high surface coverage in the range of about 95-98% is achieved. All nanosheet layers exhibit a smooth surface with a root mean square roughness (RMS) of 0.78 and 0.41 nm on top of a CNOns and TiOns, respectively, within a surface area of 20×20 µm2. The measured nanosheet thicknesses are about 2.7 and 0.7 nm, respectively, for CNOns and TiOns monolayers (see Supporting Information, Figure S1). The nanosheet thickness measured by AFM depends on the fabrication conditions, such as rate and


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degree of drying, the use of exfoliation agents and the deposition method. The thicknesses of a monolayer CNOns and TiOns measured by AFM are in the range of about 1.8-4.0 nm, and 0.71.1 nm, respectively, in spite of their crystallographic thickness of 1.44 and 0.75 nm.23-28 Thus the thicknesses of our nanosheets are in the range of reported thicknesses.

Figure 1. AFM images of a monolayer of (a) CNOns and (b) TiOns deposited on Si substrates by the LB method. The dotted line box in (a) indicates an uncovered area in between the nanosheets. AFM images of SRO/PZT/SRO films on (c) CNOns/Si and (d) TiOns/Si, with a surface area of 10×10 µm2. The dotted line box in (c) indicates the morphology of SRO/PZT/SRO film on an uncovered area in between the nanosheets as shown in (a).

As can be seen from the AFM images in Figure 1c-d, the surface morphology of SRO/PZT/SRO piezoelectric stacks is affected by the type of underlying ns/Si substrate. A granular surface structure is visible on the double layer TiOns film (Figure 1d), as well as in the gaps between the single layer CNOns film (dotted line box in Figure 1c). Further on it will be shown that the granular structure reflects open columnar PZT growth in these areas. The rootmean square (Ra) surface roughness were 4.22 and 7.08 nm, respectively, for 10×10 µm2 scan areas of PZT films on CNOns and TiOns. Moreover, Figure 1a and 1c also indicate that the


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SRO/PZT/SRO film morphology follows the shape of the CNOns in lateral directions (dashed line marked as ‘B’).

Figure 2. XRD patterns and cross-sectional SEM images, respectively, of SRO/PZT/SRO piezoelectric stacks deposited on (a,c) CNOns/Si and (b,d) TiOns/Si.

The crystallinity of piezoelectric stacks on CNOns/Si and TiOns/Si was investigated by x-ray

θ-2θ scans. Figure 2a-b indicate that the PZT film on CNOns/Si is predominantly (001)-oriented with a small fraction of (110)-orientation, whereas preferentially (110)-oriented growth and a minor (001)-oriented fraction are found in the film on TiOns/Si. This demonstrates that by changing the type of nanosheet on the Si substrates, the orientation of the PZT thin films (and likewise other perovskite films) can be controlled from predominantly (110) to (001) and vice versa. We ascribe the (110)-oriented faction in the XRD scan of the SRO/PZT/SRO/CNOns/Si sample to the granular areas (columnar grain structure) along the CNO nanosheet edges and in the uncovered Si areas in between the nanosheets (dotted line box in Figure 2c and in Supporting Information, Figure S2b). The structural differences of the PZT film deposited on CNOns and on


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the uncovered area between the nanosheets are also confirmed by AFM images of defined areas of CNOns/Si and PZT film on CNOns/Si in Figure S2 (Supporting Information). In contrast, the (110)-oriented PZT film on TiOns/Si reflects the columnar structure of the underlying layers starting at the SRO electrode on the TiOns. The full-width at half maximum (FWHM) values of the rocking curves of the PZT(002) and PZT(110) peaks are 0.58o and 0.78o, respectively, for (001)- and (110)-oriented PZT films on CNOns/Si and TiOns/Si (see Supporting Information, Figure S3). The width of the rocking curves is a measure of the range over which the lattice structure in the different grains tilts with respect to the film normal[14] and is therefore an indication of the crystalline homogeneity of a film. A comparison of the FWHMs of PZT thin films grown on ns/Si with those of epitaxial PZT thin films grown on (CeO2)/YSZ/Si and STO/Si is given in Table 1. The FWHM of (001)oriented PZT film on CNOns/Si is nearly equal to that of PZT film on STO/Si, indicating a very high crystalline quality of the film. Although the FWHM of the (110)-oriented PZT film on TiOns/Si is slightly larger, but it is still significantly less than that of (001)-oriented film on CeO2/YSZ/Si and much less than of the (110)-oriented PZT film on YSZ/Si. In order to investigate the crystal orientation across the surface of SRO/PZT/SRO films deposited on CNOns/Si and TiOns/Si, part of the sample was mapped by electron backscatter diffraction (EBSD) and the results are shown in Figure 3. The uniform red colour in Figure 3b confirms that the SRO top-electrode was oriented in the [001]pseudo-cubic ([001]pc) direction out-ofplane. This observation forms a strong indication for epitaxy, where the crystallographic orientation in the films was determined by the underlying lateral CNOns. The inverse pole figure map in Figure 3c of the x- and y-direction parallel to the surface showed that some areas had the [001]pc orientation parallels to the plane of the surface while other ones had the [101]pc or [110]pc


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plane parallels to the surface. In contrast, the uniform green color in Figure 3e-f show that the SRO top-electrode of the stack on TiOns/Si has the [101]pc or [110]pc direction out-of-plane and mixed orientations parallel to the surface.

Figure 3. (a,d) SEM image, EBSD inverse pole figure maps with (b,e) the z-direction normal to the surface (out-of-plane) and (c,f) with the x- and y-direction parallel to the surface (in-plane) of the SRO/PZT/SRO piezoelectric stack grown on CNOns/Si and TiNns/Si, respectively. The legend indicates the crystal orientation corresponding to each color.

The measured polarization-electric field hysteresis (P-E) and piezoelectric hysteresis (d33,f-E) loops of SRO/PZT/SRO thin film stacks on ns/Si are shown in Figure 4. The remanent polarization (Pr) and saturation polarization (Ps), obtained from extrapolating the high field P-E loop to zero-field), maximum longitudinal piezoelectric coefficient (d33f,max) and transverse piezoelectric coefficient (d31,f) are listed in Table 1, as well as typical values from epitaxial SRO/PZT/SRO thin films on YSZ/Si, CeO2/YSZ/Si and STO/Si fabricated in our previous studies.[12] The data show that the ferroelectric properties of PZT films on nanosheets are comparable to those of the other epitaxial PZT films with the same orientation, whereas the


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piezoelectric properties show somewhat more variation. This demonstrates that nanosheets are a promising substitute for the conventional buffer layers such as YSZ, CeO2/YSZ and STO, in the fabrication of a full-oxide piezoelectric stack on silicon in MEMS applications. The (zero-field) saturation polarization of the thin films is significantly smaller than of a poled stress-free single crystal. This is ascribed to the clamping of the film to the Si substrate that causes a small tensile strain in the film (caused by difference in thermal expansions), hampering the full rotation of the polarization towards the field direction. On a SrTiO3 substrate the strain is compressive and consequently much more out-of-plane rotation of the polarization vector is observed (Figure 7), however also then the bulk value is not obtained, due to the clamping. Similarly the bulk piezoelectric coefficients are not reached in the films due to clamping to the substrate, which limits the deformation of the unit cells in the films. The substrate induced strain in the unit cells in (110) and (001) oriented films is along the in-plane ሾ11ത0ሿ and [001] directions and the [100] and [010] directions of the pseudocubic unit cell respectively. One might therefore expect differences in the out-of-plane polarization of (110) and (001) oriented films. However no significant and systematic differences are observed for films on either YSZ, STO or ns buffer layers (see Table 1). This suggests that the polarization rotation in these films is influenced approximately equally by the substrate induced strain for all these films irrespective of the growth orientation. This is a confirmation of the fact that in PZT at the MPB the energy landscape in which the polarization vector moves is fairly flat and is thus largely decoupled from the crystal structure. The polarization vector can therefore be oriented relatively freely.29 Because of the flat energy landscape one does not expect large differences in polarization for different growth directions.


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(a)

(b)

150 100

20

d33,f (pm/V)

Polarization (µC/cm2)

40

0 -20

SRO/PZT/SRO/ CNOns/Si TiOns/Si

-40

50 0 -50

SRO/PZT/SRO/ CNOns/Si TiOns/Si

-100 -150

-200

-100

0

100

-200

200

-100

0

100

200

E (kV/cm)

E (kV/cm)

Figure 4. (a) Ferroelectric hysteresis (P-E) loops and (b) Longitudinal piezoelectric (d33,f -E) loops, of SRO/PZT/SRO capacitor stacks on CNOns/Si and TiOns/Si.

(a) 150

40

(b)

0 SRO/PZT/SRO on CNOns/Si TiOns/Si

-20

d33, max (pm/V)

Pr (µC/cm )

100 20

2

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


50 0 SRO/PZT/SRO on CNOns/Si TiOns/Si

-50

-100 -40 100

102

104

106

108

1010

-150 100

102

# cycles

104

106

108

1010

# cycles

Figure 5. (a) Ferroelectric and (b) piezoelectric cycling of SRO/PZT/SRO thin film capacitor stacks on CNOns/Si and TiOns/Si substrates.

The electrical and mechanical fatigue characteristics of the SRO/PZT/SRO capacitors on CNOns/Si and TiOns/Si are shown in Figure 5. The measurements were taken at a frequency of 100 kHz. These SRO/PZT/SRO films appear fatigue-free subjected to up to at least 1010 switching cycles. We note that SRO oxide electrodes allow easy exchange of oxygen through the


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electrode films and specifically also out of or into the grain boundaries of the PZT layer, thus reducing localized charge build-up in the PZT films,30,

31

and therefore the reduction in

ferroelectric and piezoelectric properties of PZT films are prevented.

Table 1. Properties of SRO/PZT/SRO thin films on epitaxial buffer-layers/Si and nanosheets/Si, and Pt/PZT on Pt/Si and nanosheets/Pt/Si. SRO/PZT/SRO on PZT orientation FWHM(peak) (o) of PZT film Pr (µC cm-2) a) Ps (µC cm-2) a) d33f,max (pm V-1) b) d31,f (pm V-1) c) a)

Pt/PZT on

CeO2/YSZ/Si STO/Si d)

YSZ/Si

TiOns/Si

(110) 1.42(110

(110) 0.78(110)

(001) 0.96(002)

26.5 36.8 122 -86

28.8 33.5 92 -116

CNOns/Si

Pt/Si

TiOns/Pt/Si

CNOns/Pt/Si

(001) 0.56(002)

(001) 0.58(002)

Mix 4.2(002)

(110) 2.6(110)

(001) 1.2(002)

31.8 37.6 88 -130

31.2 35.9 110 -120

17.0 22.9 124 -80

17.3 23.8 132 -84

23.6 28.0 112 -102

Single crystal e)

)

24.2 36.5 100 -90

50 327 -156

Pr is obtained from the polarization hysteresis loop measured with a maximum field of ±200 kV cm-1 at

1 kHz. It is determined from the P-axis crossings of the falling and rising loop as Pr = (Pr+ −P−r ) / 2 . Similarly Ps is determined from the crossings of the linear extrapolations of the tangents at high fields with the polarization axis as Ps = (Ps+ −P−s ) / 2 . b)

d33f,max is obtained from the d33,f-E curves of the same capacitor for which the P-E hysteresis loop was

measured. A maximum dc-field of ±200 kV cm-1 and an ac-amplitude of 5 kV cm-1 at 1 kHz was applied, using a double-beam laser interferometer (DBLI). c)

d31,f is defined from the tip-displacement of microcantilevers, consisting of a piezoelectric stack on a 10

µm thick Si-supporting cantilever beam, measured at 6 Vac,peak-peak at 3 Vdc bias point and 8 kHz, using a laser Doppler vibrometer (LDV). d)

Thin SrTiO3 (STO, 8-nm-thick) buffer layer was grown on Si by molecular beam epitaxy (MBE).

e) Single crystal values were deduced for a single crystal in a single domain polarization state. The quoted values are therefore intrinsic material values, without contributions from extrinsic effects such as domain wall motion, domain wall pinning and crystal imperfections. The values are from refs. [32-35].


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PZT thin films on nanosheets buffered (111)Pt/Ti/SiO2/Si(001) substrates Figure S5 (in Supporting Information) indicates the AFM images of CNOns and TiOns on Pt/Si. Similar to the nanosheets on Si, a relatively high surface coverage of nanosheets on Pt/Si is obtained. The lateral sizes of the nanosheets are about 3-5 µm for CNOns and 5-8 µm for TiOns, respectively. Figure 6a shows the x-ray diffraction line scan for (001)-oriented PZT film grown on SRO/STO. The film grows epitaxially and is well crystallized as evidenced by the sharpness of the rocking curve of PZT(002) peak [FWHM(002) = 0.34o] and has a dense microstructure (Figure 6d). A mix of (001), (110) and (111) orientations is obtained in the PZT film grown on LNO/Pt/Si (Figure 6b). A columnar microstructure is observed in the cross-sectional SEM in Figure 6e for the PZT film. By using CNOns as a growth template layer on Pt/Si, a PZT film with predominantly (001) orientation and a dense microstructure is obtained, as shown in Figure 6c and 6f. Also the much smaller FWHM of the rocking curve (see Supporting Information, Figure S6 and Table 1) of the (002) reflection of PZT/CNOns/Pt/Si [FWHM(002) = 1.2o] as compared to the PZT/LNO/Pt/Si sample [FWHM(002) = 4.2o] indicates a much better crystallinity of the PZT film on Pt/Si utilizing the nanosheets as an intermediate layer. Moreover, the PZT thin films deposited on TiOns/Pt/Si also show a strong (110) preferential growth orientation. These results clearly indicate that the nanosheets can also be used to control the crystal orientation of PZT thin films on Pt/Si. However the FWHM values are a factor 2-3 larger for the films grown on ns/Pt/Si compared to those on ns/Si. The higher surface roughness of the Pt/Si compared to Si is the most likely explanation for this difference.


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Figure 6. (a,b,c) XRD patterns and (d,e,f) corresponding cross-sectional SEM images, respectively, of SRO/PZT/SRO/STO(001), Pt/PZT/LNO/Pt/Si and Pt/PZT/CNOns/Pt/Si. The thickness of the PZT films is 1 µm.

It is noted that the authors are aware that by process tuning the deposition of PZT on Pt/Si can result in a better crystalline quality than is presented in this work. This has been demonstrated in many other studies throughout literature. The important point we want to make here is that with similar process settings, the microstructural quality of PZT films is dramatically improved when nanosheets are used as template layers on the Pt bottom electrode.


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(a)

(b)

60 100

40

d33,f (pm/V)

2

Polarization (µC/cm )

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


20 0

-20 SRO/STO LNO/Pt/Si CNOns/Pt/Si

-40

50 0 -50 SRO/STO LNO/Pt/Si CNOns/Pt/Si

-100

-60 -200

-100

0

100

200

-200

-100

E (kV/cm)

0

100

200

E (kV/cm)

Figure 7. (a) Ferroelectric hysteresis (P-E) loops and (b) longitudinal piezoelectric (d33,f -E) loops, of SRO/PZT/SRO/STO(001), Pt/PZT/LNO/Pt/Si and Pt/PZT/CNOns/Pt/Si.

Figure 7 shows the initial P-E and d33,f-E loops of the SRO/PZT/SRO on STO and Pt/PZT thin films on LNO/Pt/Si and CNOns/Pt/Si substrates. A square hysteresis loop with a high remanent polarization of Pr = 45.8 µC/cm2 is obtained for the PZT film on a STO substrate, significantly higher than for the films on YSZ, CeO2/YSZ, STO or ns buffered Si. This is attributed to the different substrate induced strain state of films on Si versus films on STO. These films show in turn a significantly larger saturation polarization than the films on Pt/Si (with and without ns layer). We think this may be related to the less dense, more columnar structure of the latter films, arising from a different nucleation density on Pt/Si and a rougher ns/Pt/Si surface than that of ns/Si. The larger spread in tilt angle of the grains is also reflected in a much larger FWHM of the XRD reflections. We speculate that the difference in crystalline structure causes differences in polarization screening. Because of the less dense packing of the columnar grains there is freer surface of the grains (along the column length). This is likely to allow for more


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internal surface charges and thus polarization screening, leading to a decrease in the remanent polarization. The (001)-oriented PZT film fabricated on CNOns/Pt/Si shows a slightly higher polarization and a squarer loop than the mixed-oriented film on LNO/Pt/Si and also than the (110)-oriented film on TiOns/Pt/Si (see Supporting Information, Figure S7 and S8). We think that this may be attributed to less domain wall pinning in the CNOns/Pt/Si device, because of the denser microstructure. The enhanced squareness was also seen for the film on SRO/CNOns/Si. The enhanced squareness of the films with SRO electrodes is ascribed to less charge screening and consequently less domain wall pinning, as compared to the devices with Pt electrodes. The P-E loop of the film on TiOns/Pt/Si (see Supporting Information, Figure S8) is comparable to that of the film on LNO/Pt/Si directly. This is ascribed to the similar columnar microstructure and the devices having the same top and bottom Pt electrodes. Similar observations can be made for the piezoelectric loops, which are slightly more slanted for the devices on ns/Pt/Si, indicating more domain wall pinning. Overall, the d33,f values of the films on ns/Pt/Si are well comparable with those of the devices on ns/Si for the same type of nanosheets, thus with the same microstructure: for the film on CNOns/Si the piezoelectric coefficient is slightly less than that on TiOns/Si and also for the film on LNO/Pt/Si. This is attributed to the lower connectivity of grains in the columnar structure of the latter two films, causing less clamping and thus a larger d33,f.


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Figure 8. (a) Current density-voltage characteristics and (b) breakdown voltage (VBD) of the number of measured capacitor of the 1 µm-thick PZT films for the threshold current density of about 0.5 A/cm2. (c) Initial and (d) broken capacitors.

The device current density as a function of applied voltage was measured as a function of the applied voltage in 200×200 µm2 ferroelectric capacitor devices, in which the top electrode and ferroelectric film was etched down to the bottom electrode, in order to determine the maximum allowable voltage at which the capacitors are still functional. Figure 8a shows that the breakdown voltages are 68, 72 and 90 V, respectively, for the threshold maximum current density of about 0.5 A/cm2. The PZT film on CNOns/Pt/Si with a dense microstructure shows significantly higher breakdown voltage than the other devices. These devices suffer far less from structural defects and grain boundaries that facilitate conduction paths for charge carriers that induce early breakdown. A number of measurements was performed for different capacitors and


19


the results in Figure 8b present the repeatability of breakdown voltage of different capacitors for the each structure. Figure 8c-d show the top views of an exampled capacitor at before and after breakdown occurred.

(b)

(a)

150

40

100

d33, max (pm/V)

Pr (µC/cm2)

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0 -50

-100

-40 100

50

102

104

106

108

1010

-150 100

102

# cycles

104

106

108

1010

# cycles

Figure 9. (a) Ferroelectric and (b) piezoelectric cycling of SRO/PZT/SRO/STO(001), Pt/PZT/LNO/Pt/Si and Pt/PZT/CNOns/Pt/Si.

Polarization fatigue is the process by which the switchable ferroelectric polarization degrades on repetitive field cycling. It is a serious failure mechanism in ferroelectric devices and has therefore always been the subject of the characterization of ferroelectric devices.36-39 Figure 9 shows the dependence of the remanent polarization Pr and maximum piezoelectric coefficient

d33,fmax of PZT films on SRO/STO, CNOns/Pt/Si and LNO/Pt/Si with the number of switching cycles. The (001)-oriented PZT film on SRO/STO do not show any degradation of the polarization on cycling, while the (001)-oriented PZT film on CNOns/Pt/Si shows a slight reduction in the Pr values above 109 cycles, whereas the (110)-oriented PZT film on TiOns/Pt/Si (see Supporting Information, Figure S9) and the mixed-oriented PZT film on LNO/Pt/Si start degrading after 107 and 105 cycles, respectively. We speculate that the more open, columnar films of the films on Pt/Si are more susceptible to charge transport and accumulation on the


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surface of the grains, causing increased polarization screening. Further it is known that devices with Pt electrodes suffer from the creation of a passive layer near the electrodes, reducing the field seen by the ferroelectric and thus the polarization alignment to the applied field.[37] Meanwhile, the piezoelectric coefficients d33,fmax for films on STO do not show any degradation and decrease only slightly above 109 cycles for the other samples. Thus the proposed increased screening in the cases of PZT films on TiOns/Pt/Si and LNO/Pt/Si does not seem to affect the piezoelectric properties of these films. A possible explanation for the difference in aging behaviour of Pr and d33,f is in the different measurement modes. The P-E loops are measured with a field amplitude of 200 kV/cm at 1 kHz and two polarization switches in every cycle, implying fast domain wall motion over large distances. This measurement mode is expected to be sensitive to domain wall pinning and immobile screening charges even at low densities of these defects. On the other hand the d33,f-E loop is scanned slowly with on top a low field (5 kV/cm) oscillation at 1 kHz, causing short distance domain wall motion. This low amplitude oscillation should not be very sensitive to low density pinning centres, while the slow, large amplitude scan allows for sufficient time for depinning during the scan. Overall there is a clear difference in the fatigue properties of films on ns/Si and on ns/Pt/Si. The first do not show any sign of fatigue up to 1010 cycles, while the latter degrade after at the most 109 cycles, primarily at the top electrode.

CONCLUSIONS We have demonstrated the growth and control of the crystalline orientation of PZT thin films deposited on Si and Pt/Si substrates utilizing 2D oxide nanosheets as a buffer layer. Well


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oriented PZT films with (001) and (110) orientation were obtained using CNOns and TiOns, respectively, in contrast to PZT films grown directly on Pt/Si, which show multiple growth orientations. These results indicate that the highly flexible 2D nanosheets make it possible to achieve a pre-selected oriented growth of a piezoelectric film. Although the films do not have a direct epitaxial relation to the silicon substrate, the ferroelectric and piezoelectric properties are similar as obtained with the more common epitaxial buffer layers YSZ, CeO2/YSZ and STO on Si. Devices on nanosheet buffered Si show excellent fatigue properties, with no sign of the onset of fatigue up to 1010 cycles. These results show that nanosheets as growth template layer open an alternative route for integration of highly crystalline functional oxides in silicon platform.

ASSOCIATED CONTENT

Supporting Information. AFM images a monolayer of CNO and TiO nanosheets on Si(001) substrates; Cross-sectional SEM image of (001)-oriented PZT/SRO films on CNOns/Si; Rocking curves of PZT(002) and PZT(110) reflections for the PZT thin films grown on CNOns/Si and TiOns/Si; XRD patterns and cross-sectional SEM images of (110)-oriented PZT/SRO films on YSZ/Si and (001)-oriented PZT/SRO on CeO2/YSZ/Si; AFM images of CNOns and TiOns on Pt/Si substrates; Rocking curves of PZT(002) reflections of PZT films on SRO/STO, LNO/Pt/Si and CNOns/Pt/Si; XRD pattern and cross-sectional SEM images of Pt/PZT stacks on TiOns/Pt/Si; Ferroelectric hysteresis and Longitudinal piezoelectric loops of PZT film on TiOns/Pt/Si substrate; Ferroelectric and piezoelectric cycling for PZT film prepared on TiOns/Pt/Si substrate; Schematic of the surface micromachining process; Microscope image and 3-dimensional scanning actuator image of the upward displacement of a PZT/Si-beam cantilever.


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This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION

Corresponding Author *E-mail: [email protected] (MDN); [email protected] (GR).

Author Contributions MDN and GR conceived and designed the experiments. HY made the nanosheets under the supervision of JEE. MDN deposited the piezoelectric thin films, designed and made the cantilevers, performed and analysed the ferroelectric and piezoelectric characterizations under the supervision of GR. MDN, EPH, and GR wrote the paper with the help of all the other authors. All authors discussed the results, commented on the manuscript and gave their approval to the final version of the manuscript.

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT This research was supported by the project number M62.3.10404 in the framework of the Research Program of the Materials innovation institute (M2i) (www.m2i.nl) and by the NanoNextNL-a micro and nanotechnology consortium of the Government of the Netherlands and 130 partners. The authors thank M. Smithers for performing the HRSEM experiments.

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