Block Copolymer-Assisted Nanopatterning of Porous Lead Titanate

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Block Copolymer-Assisted Nanopatterning of Porous Lead Titanate Thin Films for Advanced Electronics Alichandra Castro, Paula Ferreira, and Paula M. Vilarinho J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.6b02581 • Publication Date (Web): 04 May 2016 Downloaded from http://pubs.acs.org on May 10, 2016

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The Journal of Physical Chemistry

Block Copolymer-Assisted Nanopatterning of Porous Lead Titanate Thin Films for Advanced Electronics

Alichandra Castro1, Paula Ferreira1*, Paula M. Vilarinho1*

1

CICECO – Aveiro Materials Institute, Department of Materials and Ceramic Engineering, University of Aveiro, 3810-193 Aveiro, Portugal

Corresponding Authors: PF: phone: +351 234401419; fax: +351 234401470; e-mail: [email protected] PMV: phone: +351 234370259; fax: +351 234425300; e-mail: [email protected]

1 ACS Paragon Plus Environment


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Abstract Chemical self-assembly of functional ferroelectric PbTiO3 using amphiphilic block copolymers is used to prepare low-cost nanopatterned porous thin films with nanostructure control down to ~ 22 nm of lateral size. The approach developed in this work is highly effective in patterning PbTiO3 and avoids costly high-resolution lithography techniques and harmful etching processes. Nanopatterned films presenting a honeycomb-like arrangement of pores with diameter between 80 and 140 nm are obtained. The ordering and thickness of the films are discussed as a function of the precursors concentration and withdrawal rate used for films dip-coating. The nanopatterned porous films with thickness of 22 nm exhibit tetragonal perovskite crystallographic phase and ferroelectric response at the nanoscale. These films display vertical open porosity from the film surface to the substrate, leading to nearly 50% of the substrate surface accessible to further functionalization towards low-cost sustainable multifunctional composites for microelectronics industry.


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Introduction Despite of the presence of lead in its composition, PbTiO3 is one of the best room temperature ferroelectric perovskite with high spontaneous polarization (>50 µC/cm2 at 23 °C), large piezoelectric displacement and dielectric constant (ɛ = 190), good pyroelectric properties, high Curie Temperature (Tc = 490 ºC) and large tetragonality (c/a = 1.06).1–3 PbTiO3 thin films are of interest in microelectronics applications such as multilayer capacitors, infrared pyroelectric sensors, nonvolatile and dynamic random access memories, and ultrasonic transducers, for which the ferroelectric behavior is one of the most important requisites.4–9 For many of these applications it would be a challenge to prepare ferroelectric thin films with lateral sizes well below 100 nm.10 Furthermore, the design of nanofeatures, uniformed in size and shape at a reasonable large-range order, i.e. “nanopatterning”, would extend their utility for electronic devices and integrated circuits, which require that each pixel feature can be individually addressable. Additionally, nanopatterned porous ferroelectric thin films may be of interest to develop vertical composite structures with perfect strain coupling at the interface. The transition from lateral (substrate) strain control to vertical strain control occurs when the vertical interface area between both phases of the composite exceeds that of the substrate.11 In this context important questions arise such as: i) how to reduce the lateral dimension? and ii) what is a suitable technique to nanopattern ferroelectric multimetallic oxide materials, preserving the ferroelectric properties?; and iii) how to enhance the surface and interface areas? Based on Devonshire-Ginzburg-Landau (DGL) theory, as the physical dimensions of a ferroelectric material are reduced the stability of the ferroelectric state is altered. This leads to the existence of a critical size below which ferroelectricity is suppressed.12 For



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PbTiO3 thin films, it was predicted that the c lattice parameter and, consequently, the tetragonality (c/a), substantially decrease below a thickness size of 20 nm.13 Lichtensteiger et al.13 established through an effective Hamiltonian approach that a decrease of c/a is related with a progressive reduction of the polarization due to an imperfect screening of the depolarizing field. Although their polarization is significantly reduced, epitaxial grown thin films with 2.4 nm (about 6 unit cells) of thickness remain ferroelectric at room temperature.13 Indeed the most recent results show that the ferroelectric critical size values vary considerable and are directly dependent on the methodology of preparation.14,15 Nanopatterning in microelectronics is currently conducted by lithographic techniques. Nevertheless, most of these techniques are time consuming, difficult to apply to large areas and have limitations for feature sizes below 100 nm, thereby becoming extremely expensive.16–19 Honeycomb ordered arrays can be prepared using amphiphilic block copolymers depending on the chemical compositions of the block copolymers.20 The interfacial tension between polymer solutions and water droplets is a key point to prevent water droplets coalescence, thus contributing to preparing ordered honeycombpatterned films.20 Nguyen et al.21 reported the replication of the porosity of block copolymer films to achieve a sílica-based porous network. The use of block copolymers as masks together with lithography has been highlighted as promising for future integrated circuit technology patterning applicable to large areas.18,22 However, fabrication methodologies involving chemical etching in ferroic materials should be avoid as they can not be etched easily, and etching creates important structural defects degrading the electric and magnetic properties. Therefore, it is interesting to use the liquid-crystal template mechanism developed by the Scientists of Mobil Research and Development Corporation, in 1992,23 to directly self-assembly the ferroic materials into


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ordered array of porous. Sol-gel coupled with amphiphilic block copolymers has been widely studied for the preparation of porous binary oxide films.24–28 In 2004, a high thermally stable block copolymer was used to prepare nanocrystalline multimetallic inorganic mesoporous thin films of ternary oxide films of SrTiO3, MgTa2O6 and CoxTi(1-x)O(2-x).29 A similar method was applied to prepare ~ 100 nm thick piezo and ferrolectric porous thin films of BaTiO3 and PbTiO3, proving the functional properties at the nanoscale.30 Recently, we have reported that the crystallization of tetragonal PbTiO3 takes place at a lower temperature in porous films when compared with dense counterparts.31 Moreover, for the same thermal conditions, porous films exhibit enhanced piezoelectric coefficients, switchable polarization and low local coercivity.31 The drawback of the procedure we have used before is the partial loss of the organization of the porous structure during the thermal treatment required to achieve the tetragonal crystalline phase. Here we report an optimized method using a different block copolymer, solution concentration, withdrawal rate and temperature of processing, in which the nano ordered porosity is preserved. The prepared ferroelectric nanopatterned PbTiO3 thin films exhibit an ordered hexagonal-like array of pores perpendicular to the substrate. The importance of the inorganic precursor concentration, withdrawal rate and thermal treatment on the crystallographic phase formation, porosity size and ordering, film thickness and accessibility to the substrate underneath is discussed. Local piezoelectric and ferroelectric properties are reported. To the best of our knowledge, highly hexagonal ordered PbTiO3 porous thin films prepared by direct block copolymer self-assembly have never been achieved before for a film thickness below 100 nm.

Experimental



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Nanopatterned porous PbTiO3 thin films were prepared using a diluted sol-gel solution and evaporation induced self-assembly methodology (EISA). In the present work, we employed a different amphiphilic block copolymer from our previous works.30,31 Poly(butadiene(1,4 addition)-b-ethylene oxide) (PB51-b-PEO62) with MWPB = 51000 gmol-1 and MWPEO = 62000 gmol-1, was used as a micellar templating agent. Three solutions were prepared. In solution A, PB51-b-PEO62 block copolymer (45 mg, Polymer Source) was dissolved in absolute ethanol (98.15 mmol, Riedel-de Haën) at 70 ºC. Solution B was prepared by the dissolution of lead (II) acetate trihydrate (0.60 mmol, Fluka, purity ≥ 99.5% w/w) in glacial acetic acid (8.93 mmol, Merck) at room temperature. Solution C resulted from mixing 2,4-pentanedione (0.43 mmol, Fluka, purity 99.3% w/w) with titanium (IV) n-butoxide (0.60 mmol, Merck, 98.0% w/w) under stirring at room temperature. Afterwards, solutions B and C were added to solution A, forming the final solution. The films prepared with the inorganic quantities described above were named High C. In order to optimize the inorganic precursors concentration, we have reduced the concentration of inorganic sources (lead (II) acetate trihydrate and titanium (IV) n-butoxide) in solutions B and C to half of the quantity. The films derived from these solutions were denoted Low C. Nanopatterned PbTiO3 thin films were deposited at room temperature and 30% relative humidity by dipcoating onto platinized silicon (Pt/TiO2/SiO2/Si) (Radiant Inc.) at 0.76 and 1.6 mm/s. All films were thermally treated in air at 350 °C during 5 min in order to complete the inorganic condensation (mesostructuration) of the matrix and to partially decompose the organic content. The films were also annealed at 550 and 600 ºC during 5 min in order to achieve the crystallization. The film microstructure was investigated by highresolution SEM using a SU-70 Hitachi microscope. To follow the phase formation process, thermal gravimetric and differential thermal analyses (TGA-DTA) were carried


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out on dried nanoporous and dense powders obtained by drying the solutions described above in open vessels at 60 ºC for a few days. A Setaram Labsys™ TGA-DTA16 system was used with a heating rate of 10 °C/min under flowing air up to 700 °C. The crystalline phases in the films were identified by XRD using a Philips X’Pert MPD Xray diffractometer with Cu Kα radiation and 2° grazing incidence angle. Raman spectroscopy was performed in JY Horiba LabRam model HR800 equipment, with a high-resolution 800 mm focal length spectrometer. An argon ion laser beam at a wavelength of 325 nm was utilized. Atomic Force Microscopy (AFM) and Vertical Piezo Force Microscopy (VPFM) were carried out on Nanoscope III, Digital Instruments system with a lock-in amplifier, SRS Stanford Research Systems, using Tap300 cantilevers with Al-coated tips (Budget Sensors, resonant frequency of 200−400 kHz, force constant of 20−75 N/m) and DPE-18 cantilevers with Pt-coated tips (Mikromasch, resonant frequency of 60−100 kHz, force constant of 1.1−5.6 N/m), respectively. For the VPFM measurements, the topography signal of the film surface was taken simultaneously with the mixed signal. Both were collected in contact mode. The AFM measurements were also collected in contact mode. Since the results were obtained with the same type of cantilevers and under identical scanning and acquisition conditions, a comparison between films can be made. Several hysteresis loops with bias from -30 to +30 V were obtained to ensure the reproducibility of the results, and representative loops are presented. The piezoelectric and ferroelectric properties are mean values taken from several (at least ten) hysteresis loops. The critical voltage is defined as the necessary voltage for the nucleation of a new domain. The imprint is defined as Im = (V+ + V-)/2, where V+ correspond to the positive coercive bias and V- to the negative coercive bias. Switchable polarization corresponds to the difference between the positive saturated piezoresponse and negative ones (Rm = (RS)+ - (RS)-). As



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the effective piezoelectric coefficient ((d33)eff) is proportional to the amplitude signal, this can be defined as (d33)eff ∝ (amplitude signal·cos(phase signal)) / Vac, where Vac is the ac voltage applied, and its value can be taken from the mixed signal at zero voltage.

Results and discussion Dense and nanopatterned porous PbTiO3 thin films were prepared by a chemical-based methodology using evaporation-induced self-assembly. Figure 1 shows a top view scanning electron microscopy (SEM) micrographs of nanopatterned PbTiO3 thin films deposited starting from solutions with high and low inorganic precursor concentration (High C and Low C), using two distinct withdrawal rates (0.76 and 1.6 mm/s) and heated at 350, 550 and 600 ºC. These micrographs are representative of the entire covered surface. All films present well-ordered arrays of pores forming honeycomb-like patterns. After heat treatment at 350 ºC for 5 min, films appear amorphous and some pores seem to remain closed due to the incomplete removal of the block copolymer (Figure 1a, 1d and 1g). The complete PB-b-PEO decomposition occurs between 200 and 450 °C as observed through the TGA-DTA analyses (Figure S1 of supporting information – SI). The void motifs result from the thermal decomposition of the block copolymer. Independently from the solution concentration, the films treated at 550 and 600 ºC present arrays of hexagonal-like ordered porous structures (Figure 1b, 1e, 1f and 1c, 1f, 1g, respectively). The High C film deposited at 0.76 mm/s and treated at 600 ºC (Figure 1c) has pores with larger diameters than the High C film deposited at 1.6 mm/s and treated at the same temperature (98.1 ± 13.3 nm versus 86.9 ± 13.7 nm, calculated from the SEM micrographs) (Figure 1c and 1f; and Table 1). The latter film presents thicker walls than the former one (91.9 ± 13.1 versus 77.2 ± 14.4 nm). According to literature, the high withdrawal rate provides more matter at the substrate, thereby 8 ACS Paragon Plus Environment

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generating thicker films with thicker walls and reducing the pore diameter.24,32 At 600 °C, the film deposited with the highest concentrated solution and withdrawal rate (Figure 1f) shows different contrast areas in the SEM top view (bright areas). Despite this fact, the order and periodicity of the framework remain after this high temperature treatment. The variation on the brightness contrast may be related with the presence of different crystalline phases as observed in our previous work.31 Bright areas may be related with the presence of a tetragonal perovskite phase. The imaging of the Low C film treated at 600 ºC is difficult, this is probably due to the reduction of the film thickness and high exposure of the platinum of the substrate (very bright contrast). The average pore diameter and wall thickness, calculated from the SEM image, are 127.0 ± 11.6 and 66.1 ± 14.8, respectively (Table 1). Under these conditions, some of the pores become interconnected (Figure 1i). At Low C, the inorganic precursors concentration seems to be insufficient to allow the formation a homogeneous continuous layer after heat treatment. According to the literature, the inorganic/organic ratio of the initial precursors solution has a determinant influence on the final mesostructure of the mesoporous and nanopatterned thin films.25,28,33,34 Ethanolic solutions containing PB-b-PEO were reported to form high quality monolayers of oxides with homogeneous pore sizes and ordered arrangements.33–35 The formation of the well-ordered monolayers is mainly dependent on the nature of the amphiphilic block copolymer, solvent, solution dilution and withdrawal rate used.33–35 High C films deposited with the same withdrawal rate and treated at 600 ºC exhibit a higher thickness than the films prepared from Low C. The high content of inorganic precursors enhances interactions with PEO chains, leading to thicker walls and turning pore sizes more heterogeneous (wider distribution of pore sizes) as observed when comparing Figure 1e and 1h.



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When the concentration of precursors increases as a consequence of the viscous drag regime35 a slight increase on the film thickness is expected. In order to acquire more data on the effect of the inorganic precursor concentration and withdrawal rate on the nanopatterned films, AFM was carried out in the High C films deposited at both withdrawal rates (0.76 and 1.6 mm/s) and in Low C films deposited at 1.6 mm/s, all treated at 600 ºC (Figure 2). The AFM profiles (not presented here) taken from the AFM images allowed the calculation of: pore diameter (dpore), wall thickness, film thickness, fraction of accessible substrate surface (SPt) and root mean square surface roughness (RMS). Average values are presented in Table 1. For thin films deposited with High C, the AFM images present a good periodicity of pores with wide distribution of pore sizes, confirming the SEM results. The AFM images of Low C films (Figure 2e and 2f) show narrow distribution of pore sizes and a homogeneous wall thickness. Indeed, from the values presented in Table 1, dpore is higher in Low C (140.9±6.2 nm) than in High C for 1.6 mm/s (82.3±11.4 nm) or 0.76 mm/s (93.8±13.6 nm). A decrease of dpore when the withdrawal rate increases was observed, which is consistent with the SEM results. Simultaneously, the wall thickness between the pores increases and the SPt slightly decreases. These results are only consequence of the withdrawal rate. At 600 ºC, as a result of the phase formation, the nanopatterned films deposited at 1.6 mm/s present a significant number of well-defined nanograins and the hexaganol array is not well-defined. As the AFM measurements were performed in contact mode with high-resolution tips (tip radius < 10 nm and tip height of 17 µm), the pore depth may be related with film thickness. According to the values of Table 1, the film thickness increases as the inorganic precursors concentration and withdrawal rate increase. The increase of the inorganic precursor concentration increases the solution viscosity improving film thickness. In fact, as the highest value of film thickness is


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below to the tip height, the pore depth can be assumed to be close to the value of the film thickness. Furthermore, the thickness value obtained for the thin film deposited with High C and 1.6 mm/s is in good agreement with the value measured by the SEM cross-section image (not presented here). Thus, the porosity seems to be completely open and practically directed to the substrate. It is worth noting the remarkable porous order of all studied films, despite the different film thicknesses (varying from ~ 5 to 22 nm). Figure 3 shows the XRD patterns of nanopatterned PbTiO3 thin films deposited with different inorganic precursors concentrations (High C and Low C), different withdrawal rates (0.76 and 1.6 mm/s) and thermally treated at 600 °C. Solid black and dashed vertical lines correspond to the tetragonal (JCPDS no. 00-003-0721) and cubic (JCPDS no. 00-040-0099) crystalline PbTiO3 phases, respectively. The gray lines at 2θ 40.4 and 47.0° correspond to reflections of the Pt layer of the substrate (JCPDS no. 04-0105118). Only the nanopatterned thin films prepared with High C of inorganics and deposited with the faster withdrawal rate presents the main diffraction peaks associated to crystalline PbTiO3 phase. In all the other cases, the XRD patterns do not show the typical reflections of the perovskite phase. A potential explanation may be the size of the crystallites that can be too small to be detected by XRD diffraction. For the films prepared with High C and a withdrawal rate of 1.6 mm/s, it is not possible to differentiate between the cubic and tetragonal structure of PbTiO3 (JCPDS no. 00-0400099 and 00-003-0721, respectively) by XRD. As the thickness of the nanopatterned thin films is low (~22 nm, see Table 1), the intensity of the diffraction pattern is weak. Another issue to consider is the relation between thickness and tetragonality. One must consider, that it is possible that the thickness of the film affects the tetragonality. As reported by Lichtensteiger et al.13, the tetragonality of PbTiO3 thin films progressively



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decreases as thickness of the film decreases below 20 nm due to the in-plane strain relaxation and the constraining effect of the substrate. To verify the presence of the tetragonal phase, Raman spectroscopy was performed in PbTiO3 thin films deposited with High C and 1.6 mm/s of withdrawal rate, treated at 600 ºC (Figure 4). These films present the typical Raman modes of the tetragonal phase (E+B1, A1 (2TO), E(3TO) and A1(3TO) bands) at 287, 324, 503 and 599 cm-1, as previously reported for other PbTiO3 and lead zirconate titanate thin films.36,37 The Raman peaks are slightly shifted to low wavenumber in relation to single crystal values,38 probably due to strains induced by the substrate. The Raman results clarify the XRD results. As the presence of tetragonal phase leads to the appearance of ferroelectricity in ABO3 perovskite type materials, the local piezoelectric and ferroelectric behaviour of the High C film deposited at 1.6 mm/s and treated at 600 ºC was investigated through vertical piezoresponse force microscopy (VPFM) and piezoresponse force spectroscopy (PFS). Figure 5 shows the topographic image, VPFM mixed signal, and the three-dimensional image of the topography with mixed signal. The dark domains in the VPFM mixed signal image correspond to domains in which the polarization is oriented towards the substrate (phase = -180 °), while bright regions correspond to domains with polarization oriented towards the free surface of the films (phase = 180°). Grains with in-plane polarization exhibit an intermediate contrast. The VPFM mixed signal image presents two distinct piezoelectric behaviours (Figure 1f). In analogy to our previous work,31 the areas with strong piezoelectric response correspond to the bright areas observed in SEM micrograph and can be assigned to crystalline PbTiO3 tetragonal phase. The extension of those areas may be enlarged by treating the film at a temperature above 600 ºC or increasing the heating time.31 However, experimentally, it was difficult to maximize the


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areas with strong piezoelectric behaviour without losing the porosity order (see Figure S2 of SI). To investigate the ferroelectric behaviour of these nanopatterned films, we measured the local piezoelectric response of films as a function of an applied bias between the platinum substrate and the conducting AFM tip. Figure 6 shows a representative example of the local remanent out-of-plane piezoelectric hysteresis loop obtained directly from the PFM equipment/software and measured from individual domains with bias from −30 V to +30 V. The hysteresis loop reveals that the switched polarization remains after bias removal and clearly confirms the ferroelectric behaviour at room temperature of these films. For a quantitative analysis, the ferroelectric properties such as critical voltage, coercivity, imprint, switchable polarization and (d33)eff, were calculated from several mixed hysteresis loops and are presented in Table 2. The values obtained in our preview work27 for dense and nanoporous 100 nm thick films are also presented in Table 2 for comparison proposes. The critical voltage (2.0 ± 0.5 V) and coercivity value (5.3 ± 0.5 V) of the nanopatterned nanoporous PT thin films are higher than the values observed in ~ 100 nm thick nanoporous PbTiO3 films prepared in our previous work (1.0 ± 0.1 V and 2.4 ± 0.1 V, respectively).31 These results clearly demonstrate how the ferroelectric behaviour is affected by the low thickness of the nanopatterned thin films. The hysteresis loops obtained for these films show an imprint effect in terms of coercive voltage (1.3 ± 0.5 V). This effect is usually caused by the preference for a certain polarization state over the other and provides a measure of the internal field.39 Imprint behaviour can also be related with the self-polarization that depends, in a large extent, on the film deposition technology. In the case of sol–gel deposited films, it was reported that the self-polarization effect is thickness dependent, suggesting that the alignment of domains occurs locally near the film–bottom electrode



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interface.40 In the nanopatterned thin films prepared in the current work the imprint value is slightly higher than the value obtained in the 100 nm thick PbTiO3 nanoporous films of our preview work (1.0 ± 0.1 V).31 This variation can be associated with the incipient crystallization degree of the tetragonal phase of the nanopatterned thin film and/or surface/interface defects (lattice distortion due to different thermal expansion coefficient of the film and substrate). The switchable polarization and (d33)eff piezoelectric coefficient of the nanopatterned porous thin films (137.6 ± 6.7 pm/V and 49.6 ± 0.7 pm/V, respectively) are lower than the ones of the 100 nm thick PbTiO3 nanoporous films (261.3 ± 6.6 and 148.8 ± 5.5 pm/V, respectively). The degradation of the electrical properties may result from the decrease of the film thickness with consequent reduction of the tetragonality. The decrease of the tetragonality leads to a drop on the spontaneous polarization of the films, due to a residual unscreened depolarization field. The depolarization field increases as the thickness of the ferroelectric film decreases and consequently it has a detrimental effect on the homogeneous polarization state, leading to either a reduction of the spontaneous polarization or formation of ferroelectric domains.41,42 Moreover comparing the values obtained for nanopatterned films with the values reported for dense films in our previous work,31 it is possible to observe that the nanopatterned films present higher switchable polarization and (d33)eff piezoelectric coefficient than dense thin films (137.6 ± 6.7 pm/V and 49.6 ± 0.7 pm/V versus 85.1±2.0 pm/V and 35.7±0.8 pm/V, respectively). These results clearly demonstrate that the porosity even in films with low thickness (20 nm) and consequent low tetragonality degree leads to enhanced ferroelectric properties. Although, nanopatterned films present higher coercivity than nanoporous 100 nm thick films,31 the nanopatterned ones can be activated at lower power thus being interesting for applications in memory devices. The piezoelectric


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coefficient (d33) value obtained for the nanopatterned films is approximately 50 pm/V, which is slightly better than the (d33)eff obtained for the single crystal nanoislands (resulting from complex preparation approaches) where the crosstalk effect provoked by the domain movement or thermal diffusion is inherently excluded (39.4 pm/V).40 Thus, comparing the nanopatterned with porous films of PbTiO3 synthesized by different methods7,43 or with PbTiO3 nanoislands44, nanopatterned films here prepared by an easy and low-cost method present enhanced ferroelectric properties. Furthermore, the porosity is completely open and directed to the substrate, being the platinum underneath almost 50% accessible.

Conclusions Well-ordered nanopatterned porous PbTiO3 thin films were prepared using an easy and low-cost block copolymer sol-gel solution. Despite the 20 nm thickness, the nanopatterned PbTiO3 films present a tetragonal crystallographic phase, clearly exhibiting PFM domains contrast and a well-defined hysteresis loops as proof of their piezoelectric and ferroelectric behaviour. Film and wall thicknesses and pore diameter can be tuned by tailoring the withdrawal speed and inorganic solution dilution. Moreover, as the substrate is highly accessibly and readily available, these porous thin films open an avenue of opportunities for the development of advanced multifunctional electronic devices by functionalization of the pores.

Supporting Information Available: TGA/DTA analyses of powders obtained from drying the deposition solution used on the preparation of the nanopatterned porous thin films. SEM micrograph of the nanopatterned porous film treated at 650 ºC with white



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arrows indicating areas of degradation of the porous structure. This material is available free of charge via the Internet at http://pubs.acs.org/.

Acknowledgements This work was developed in the scope of the project CICECO-Aveiro Institute of Materials POCI-01-0145-FEDER-007679 (Ref. FCT UID/CTM/50011/2013), financed by national funds through the FCT/MEC and when applicable co-financed by FEDER under the PT2020 Partnership Agreement. Authors are grateful to FCT and POPH/FSE, for

doctoral

fellowship

SFRH/BD/67121/2009

and

investigator

fellowship

IF/00327/2013.

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Figure and Tables Captions

Figure 1. SEM micrographs illustrating the typical morphology of nanopatterned PbTiO3 films deposited with high concentration (High C) (a-f) and low concentration (Low C) (g-i) of inorganic precursors; using withdrawal rates of 0.76 mm/s (a-c) and 1.6 mm/s (d-i) and treated at 350 (a, d and g), 550 (b, e and h) and 600 ºC (c, f and i). As the inorganic precursors concentration increases, the distribution of pores sizes becomes wider. The thin film deposited with High C solution and fast withdrawal rate presents different contracts in the SEM image after treatment at 600 °C (bright areas indicated with arrows in Fig. 1i). All films present a well-ordered hexagonal arrangement of pores as shown in the inset of Fig 1g.

Figure 2. AFM topographical images of nanopatterned PbTiO3 films (5 µm x 5 µm and 1.5 µm x 1.5 µm) obtained with High C (a-d) and Low C (e-f) solutions, deposited at 0.76 mm/s (a and b) and 1.6 mm/s (c-f) and thermal treated at 600 ºC. All nanopatterned PbTiO3 films present ordered hexagonal array of pores. Figure 3. X-ray diffraction patterns of nanopatterned PbTiO3 films prepared with High C and Low C solutions, deposited at 0.76 or 1.6 mm/s and treated at 600 ºC. Solid black and dashed vertical lines correspond to the tetragonal (JCPDS no. 00-003-0721) and cubic (JCPDS no. 00-040-0099) crystalline phases of PbTiO3, respectively. The gray lines correspond to the Pt layer of the substrate. Crystalline PbTiO3 phase only appears in thin films deposited with High C and 1.6 mm/s withdrawal rate.

Figure 4. Raman spectra of nanopatterned PbTiO3 films deposited with High C solution at 1.6 mm/s and treated at 600 ºC. Raman spectroscopy proves the presence of the


Page 23 of 33

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tetragonal PbTiO3 phase in nanopatterned films deposited with High C and fast withdrawal rate.

Table 1. Nanopatterned PbTiO3 films features: pore diameter (dpore) and wall thickness between the pores (both calculated using AFM and SEM images); and film thickness, fraction of accessible platinum substrate surface (SPt) and RMS (root mean square) surface roughness (calculated from AFM images). As the inorganic precursor concentration and withdrawal rate increase, film thickness and wall thickness also increase, with concomitant decrease of the pore diameter.

Figure 5. Scanning probe microscopy images of nanopatterned PbTiO3 film deposited with High C at 1.6 mm/s and treated at 600 ºC: a) topography image; b) VPFM mixed signal and c) three-dimensional image of the topography with mixed signal. The scale of the VPFM mixed signal image ranges between -10 to 10 V. This 20 V scale corresponds to 360º, thus opposite domains oscillate 180º out of phase.

Figure 6. Representative remanent local hysteresis loop obtained in nanopatterned PbTiO3 film deposited with High C solution at fast withdrawal rate and treated at 600 ºC. The dashed vertical line indicates the bias = 0 V.

Table 2. Average values of critical voltage, coercivity, imprint, switchable polarization and (d33)eff calculated from at least ten phase, amplitude and mixed hysteresis loops (Piezoresponse Force Microscopy loops) for dense and nanoporous with 100 nm of thickness of our preview work [27] and nanopatterned porous thin films of this work. The higher coercivity values present in nanopatterned thin films show how the ferroelectric behaviour is affected by the low thickness of the nanopatterned thin films. 23 ACS Paragon Plus Environment


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However in terms of the vertical shift, nanopatterned thin films present reduced imprint effect probably due to the high porosity area of this kind of structure. Consequently the switchable polarization and (d33)eff coefficient are slightly higher than those obtained for dense ones where the crosstalk effect provoked by the domain movement or thermal diffusion is not excluded.


Page 25 of 33

Figures and tables

350 ºC

550 ºC

600 ºC

b)

c)

d)

e)

f)

g)

h)

i)

0.76 mm/s

a)

1.6 mm/s

1.6 mm/s

High C Low C

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


Figure 1. SEM micrographs illustrating the typical morphology of nanopatterned PbTiO3 films deposited with high concentration (High C) (a-f) and low concentration (Low C) (g-i) of inorganic precursors; using withdrawal rates of 0.76 mm/s (a-c) and 1.6 mm/s (d-i) and treated at 350 (a, d and g), 550 (b, e and h) and 600 ºC (c, f and i). As the inorganic precursors concentration increases, the distribution of pores sizes becomes wider. The thin film deposited with High C solution and fast withdrawal rate presents different contracts in the SEM image after treatment at 600 °C (bright areas indicated with arrows in Fig. 1i). All films present a well-ordered hexagonal arrangement of pores as shown in the inset of Fig 1g.



600 ºC 38.0046.00 nm

46.0046.00 nm

a)

High C

0.76 mm/s

b)

0.00 0.00 nm

0.00 0.00 nm

50.0050.00 nm

46.0050.00 nm

d)

1.6 mm/s

c)

0.000.00 nm

0.000.00 nm

36.0040.00 nm

27.0040.00 nm

f)

1.6 mm/s

e) Low C

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Page 26 of 33

0.000.00 nm

0.00 0.00 nm

Figure 2. AFM topographical images of nanopatterned PbTiO3 films (5 µm x 5 µm and 1.5 µm x 1.5 µm) obtained with High C (a-d) and Low C (e-f) solutions, deposited at 0.76 mm/s (a and b) and 1.6 mm/s (c-f) and thermal treated at 600 ºC. All nanopatterned PbTiO3 films present ordered hexagonal array of pores.


Pt (111) (002)

(111)

(110)

(101)

(100)

Low C 1.6 mm/s

(001)

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Intensity (a.u.)

Page 27 of 33

High C 1.6 mm/s High C 0.76 mm/s

20 25 30 35 40 45 50 2θ (°)

Figure 3. X-ray diffraction patterns of nanopatterned PbTiO3 films prepared with High C and Low C solutions, deposited at 0.76 or 1.6 mm/s and treated at 600 ºC. Solid black and dashed vertical lines correspond to the tetragonal (JCPDS no. 00-003-0721) and cubic (JCPDS no. 00-040-0099) crystalline phases of PbTiO3, respectively. The gray lines correspond to the Pt layer of the substrate. Crystalline PbTiO3 phase only appears in thin films deposited with High C and 1.6 mm/s withdrawal rate.


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Raman Intensity (a.u)


200

400 600 800 Wavenumber (cm-1)

1000

Figure 4. Raman spectrum of nanopatterned PbTiO3 films deposited with High C solution at 1.6 mm/s and treated at 600 ºC. Raman spectroscopy proves the presence of the tetragonal PbTiO3 phase in nanopatterned films deposited with High C and fast withdrawal rate.


Page 29 of 33

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Table 1. Nanopatterned PbTiO3 films features: pore diameter (dpore) and wall thickness between the pores (both calculated using AFM and SEM images); and film thickness, fraction of accessible platinum substrate surface (SPt) and RMS (root mean square) surface roughness (calculated from AFM images). As the inorganic precursor concentration and withdrawal rate increase, film thickness and wall thickness also increase, with concomitant decrease of the pore diameter. High C

Low C

0.76 mm/s

1.6 mm/s

1.6 mm/s

dpore – AFM (nm)

93.8±13.6

82.3±11.4

140.9±6.2

dpore – SEM (nm)

98.1 ± 13.3

86.9 ± 13.7

127.0 ± 11.6

Wall thickness – AFM (nm)

76.2±13.6

89.4±11.5

40.8±10.6

Wall thickness – SEM (nm)

77.2 ± 14.4

91.9 ± 13.1

66.1 ± 14.8

Film thickness – AFM (nm)

18.9±2.1

22.4±1.7

5.7±1.4

SPt – AFM (%)

52.1±2.4

47.7±2.1

59.9±2.6

RMS surface roughness – AFM (nm)

5.6

6.7

2.9



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

10.00 V

43.00 nm

43.00

a)

1.50

c)

b)

0.000.00 nm

Page 30 of 33

-10.00 V -1.50

Figure 5. Scanning probe microscopy images of nanopatterned PbTiO3 film deposited with High C at 1.6 mm/s and treated at 600 ºC: a) topography image; b) VPFM mixed signal and c) three dimension image of the topography with mixed signal.. The scale of the VPFM mixed signal image ranges between -10 to 10 V. This 20 V scale corresponds to 360º, thus opposite domains oscillate 180º out of phase.


Page 31 of 33

80

d33 signal (pm/V)

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40 0 -40 -80

-30 -20 -10 0 10 20 30 Bias (V) Figure 6. Representative remanent local hysteresis loop obtained in nanopatterned

PbTiO3 film deposited with High C solution at fast withdrawal rate and treated at 600 ºC. The dashed vertical line indicates the bias = 0 V.



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Page 32 of 33

Table 2. Average values of critical voltage, coercivity, imprint, switchable polarization and (d33)eff calculated from at least ten phase, amplitude and mixed hysteresis loops (Piezoresponse Force Microscopy loops) for dense and nanoporous with 100 nm of thickness of our preview work [27] and nanopatterned porous thin films of this work. The higher coercivity values present in nanopatterned thin films show how the ferroelectric behaviour is affected by the low thickness of the nanopatterned thin films. However in terms of the vertical shift, nanopatterned thin films present reduced imprint effect probably due to the high porosity area of this kind of structure. Consequently the switchable polarization and (d33)eff coefficient are slightly higher than those obtained for dense ones where the crosstalk effect provoked by the domain movement or thermal diffusion is not excluded.

Films

Critical voltage (V)

Coercivity (V)

Imprint (V)

Switchable polarization (pm/V)

(d33)eff (pm/V)

Dense–10027

1.8±0.1

2.5±0.1

1.2±0.1

85.1±2.0

35.7±0.8

27

Porous-100

1.0±0.1

2.4±0.1

1.0±0.1

261.3±6.6

148.8±5.5

Nanopatterned

2.0±0.5

5.3±0.5

1.3±0.5

137.6±6.7

49.6±0.7


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Table of Contents Graphic


Block Copolymer-Assisted Nanopatterning of Porous Lead Titanate

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