Synthesis, Giant Dielectric, and Pyroelectric Response of [001]-Oriented Pr3+ Doped Pb(Mg1/3Nb2/3)O3-PbTiO3 Ferroelectric Nano-Films Grown on Si Substrates

The [001]-oriented Pr3+ doped Pb(Mg1/3Nb2/3)O3-0.30PbTiO3 (Pr-PMN-PT) thin films with a composition near the morphotropic phase boundary (MPB) were synthesized by a sol–gel method. The crystal structure was characterized using X-ray diffraction. It was found that a single perovskite phase was achieved in Pr-PMN-PT thin films annealed at 650 °C for 3 min. The dielectric constant (εr) value was 2400 in 2.5% Pr-PMN-PT thin films at room temperature, 110% higher than that of pure PMN-PT samples. Through 2.5% Pr3+ doping, remanent polarization (Pr) and coercive field (Ec) values increased from 11.5 μC/cm2 and 35 kV/cm to 17.3 μC/cm2 and 63.5 kV/cm, respectively, in PMN-PT thin films. The leakage current densities of pure and 2.5% Pr-PMN-PT thin films were on the order of 1.24 × 10−4 A/cm2 and 5.8 × 10−5 A/cm2, respectively, at 100 kV/cm. A high pyroelectric coefficient (py) with a value of 167 μC/m2K was obtained in 2.5% Pr-PMN-PT thin films on Si substrate, which makes this material suitable for application in infrared detectors.


Materials and Methods
Before sol-gel processing, a TiO 2 diffusion barrier layer was first fabricated by oxidizing the Ti metal layer as follows: Ti films of 60 nm thickness were deposited on SiO 2 /Si (100) substrate by DC magnetron sputtering in an Ar gas environment at room temperature. The sputtering power was 175 W. The base vacuum of Ti sputtering was higher than 8 × 10 −4 Pa. The ratio of Ar:O 2 was 90:10. After deposition, the Ti layer was annealed in a rapid thermal processing (RTP) furnace at 700 • C for 1 h. Then, a Pt bottom electrode of 80 nm thickness was deposited on the TiO 2 diffusion barrier layer. The sputtering parameters were the same as the Ti metal layer deposition. The sheet resistance was about 2.85 Ω/sq, which was measured by the four probe method.
After that, Pr 3+ doped Pb(M 1/3 Nb 2/3 )O 3 -0.30PbTiO 3 (xPr-(1 − x)(PMN-0.30PT), x = 1%, 2%, 2.5%, and 3%) precursor solutions were prepared by the sol-gel method. The lead acetate trihydrate (Pb(CH 3 COO) 2 ·3H 2 O), magnesium acetate tetrahydrate (Mg(CH 3 COO) 2 ·4H 2 O), niobium ethoxide (Nb(OCH 3 CH 2 ) 5 ), tetrabutyl titanate (Ti(OC 4 H 9 ) 4 ), and praseodymium acetyl acetone (C 15 H 21 PrO 6 ·xH 2 O) were used as starting materials to synthesize the Pr-PMN-PT sol-gel solution. An additional 20 mol% of lead acetate trihydrate was added to equalize the Pb volatilization and 5 mol% more magnesium acetate tetrahydrate was mixed to boost the production of perovskite phase. Next, 2-methoxyethanol (C 5 H 8 O 2 ) and glacial acetic acid (CH 3 COOH) were added as solvent and catalyst respectively. Acetyl acetone (C 5 H 8 O 2 ) was added as a stabilizer in the Pb-Pr-Mg-Nb-Ti solution. The concentration of the Pr-PMN-PT solution was attenuated to 0.4 mol/L. Then, the Pr-PMN-PT thin films were deposited by spin coating on Pt/TiO 2 /SiO 2 /Si substrates at 2000 rpm for 9 s before the spin coating was subsequently sped up to 4000 rpm for 20 s. To evaporate the solvent in solution, the sample was dried at 220 • C for 3 min after each deposition and then pyrolyzed at 450 • C for 5 min. Finally, Pr-PMN-PT films were annealed at 600, 650, and 700 • C for 1-5 min in air by RTP at ramp rates +20 • C/s and −5 • C/s respectively. In order to get 500 nm thick film, this entire process was repeated 10 times. For further electric measurement, Au top electrodes were deposited on Pr-PMN-PT samples using magnetron sputtering through a shadow mask with a 1 mm diameter.
The crystalline structure of Pr-PMN-PT thin films was analyzed by X-ray diffraction (XRD) (Shimadzu X-6000, Kyoto, Japan) with Cu-Kα radiation in the 2theta range of 20 • -60 • and a scan step length of 0.02 at room temperature. The surface morphologies of the samples were characterized by a field emission scanning electron microscope (FE-SEM) (ZEISS Gemini SEM 500, Jena, Germany), atomic force microscope (AFM) (Bruker Multimode 8, Billerica, MA, USA), and piezoresponse force microscopy (PFM) (Bruker dimension icon, Billerica, MA, USA). For electric measurement, the samples were all of a pure perovskite phase structure due to the high electric performance. The dielectric properties were measured using an impedance analyzer (HP 4284, Hewlett Packard, Palo Alto, CA, USA) with a precision LCR meter connected to a heating stage (Linkam THMSE 600, Waterfield, UK). The P-E hysteresis loops and leakage current were examined by a standard ferroelectric system (aixACCT TF-2000, Aachen, Germany). The pyroelectric coefficient (p y ) was calculated by the equation as below: where I is the pyroelectric current, A is the surface area of the samples, and dT/dt is the rate of temperature cooling.

Phase Structure
Avoiding the formation of a pyrochlore phase is a key factor for the synthesis of PMN-PT thin films as the addition of this phase will reduce the piezoelectric and ferroelectric properties of the resulting thin films. In order to restrict the formation of a pyrochlore phase, the annealing conditions for thin film preparation were first investigated. Figure 1a Figure 1c,d present the phase structure of 2.5% Pr-PMN-PT thin films deposited on Pt/TiO 2 /SiO 2 /Si substrates annealed at 650 • C for 1 min, 3 min, and 5 min respectively. It was found that 2.5% Pr-PMN-PT thin films were mostly in the perovskite phase with a little pyrochlore phase when annealed for 1 min, which was further crystallized to a single perovskite phase structure after one more minute annealing. However, a peak of the pyrochlore phase was found again in the sample annealed for 5 min, which can be ascribed to the greater lead loss caused by the longer time annealing [19].
The degree of (001) texture in 2.5% Pr-PMN-PT thin film was evaluated by preferred orientation parameter α hkl , which can be calculated by the following formula [32]: where I hkl is the relative intensity of the corresponding diffraction peaks (hkl). As shown in Figure 1a,c, α 001 values are 86.6% and 73.6% in Pr-PMN-PT thin films annealed at 650 • C for 2 and 3 min respectively, presenting the highly (001)-oriented textured structure. The preferred orientation was related to many factors, such as substrate materials [33], growth method, doping elements [31], annealing temperatures [34], and so forth. For Pr-PMN-PT thin films, the nucleation was in random direction at the start of thin film growth on the surface of Pt. With further temperature increases up to 650 • C, the (001) direction became the preferential growth direction.  Figure 2a,b shows the FE-SEM images of surface and cross-sectional micrographs of Pr-PMN-PT thin films. From these images, it can be seen that the thin films have a dense structure without any cracks. It can further be seen that the quite uniform Pr-PMN-PT, Pt, and TiO2 films are around 450 nm, 80 nm, and 80 nm in thickness respectively, which are in agreement with the values measured by a surface profilometer. Figure 2c presents surface grain size distributions calculated from the image in Figure 2a, showing an average grain size of 100 nm. A typical out-of-plane image with a 3 × 3 μm 2 scan area at the surface of 2.5% Pr-PMN-PT thin films is shown in Figure 2d. The grain size is about 100 nm, in accordance with the FE-SEM observation. However, nano-domains could not be found in the picture, which can be analyzed by macroscale level ferroelectric properties, similar to the results in Ref. [35,36]. The 3D AFM image of 2.5% Pr-PMN-PT thin films is shown in Figure 2e. The root mean square roughness (Rrms) is about 2.59 nm with a scan area of 2 × 2 μm 2 , indicating low interference of electrode contact in electronic measurement. A schematic of the device for further electric investigation is shown in Figure 2f. Two silver wires for electric measurement were bonded to the surface of the Au electrodes.

Surface Morphology and Electric Properties at Room Temperature
The dielectric, pyroelectric, and ferroelectric properties in 2.5% Pr-PMN-PT thin films at room temperature are summarized in Table 1. The values of εr, py, and pr in 2.5% Pr-PMN-PT thin films were on the order of 2400, 167 μC/m 2 K, and 17.3 μC/cm 2 , showing 110%, 123%, and 50% higher factors respectively than those in pure PMN-PT thin films. Compared with PMN-0.30PT grown by a chemical solution deposition (CSD) method [19], the εr in PMN-0.30PT thin films on Si substrate was lower than that of the sample synthesized on perovskite structure single crystal substrates. A similar tendency of pyroelectric property was found in PMN-0.30PT thin film grown on various substrates. The pr value in 2.5% Pr-PMN-PT thin films was still higher than that of PMN-0.32PT thin films deposited on Ba0.5Sr0.5RuO3/NdScO3 and La0.5Sr0.5CoO3/LaAlO3 crystals using pulsed laser deposition. In contrast, the py value showed an opposite trend, which is attributed to the lattice misfit. Compared with PVDF and normal ferroelectric material, the relaxor-PT thin films showed higher dielectric and pyroelectric properties. In contrast, the response time of the PMN-0.30PT pyroelectric sensor was longer than that of PVDF [41] and LiTaO3 [43].   Figure 2d. The grain size is about 100 nm, in accordance with the FE-SEM observation. However, nano-domains could not be found in the picture, which can be analyzed by macroscale level ferroelectric properties, similar to the results in Ref. [35,36]. The 3D AFM image of 2.5% Pr-PMN-PT thin films is shown in Figure 2e. The root mean square roughness (R rms ) is about 2.59 nm with a scan area of 2 × 2 µm 2 , indicating low interference of electrode contact in electronic measurement. A schematic of the device for further electric investigation is shown in Figure 2f. Two silver wires for electric measurement were bonded to the surface of the Au electrodes.

Surface Morphology and Electric Properties at Room Temperature
The dielectric, pyroelectric, and ferroelectric properties in 2.5% Pr-PMN-PT thin films at room temperature are summarized in Table 1. The values of ε r , p y , and p r in 2.5% Pr-PMN-PT thin films were on the order of 2400, 167 µC/m 2 K, and 17.3 µC/cm 2 , showing 110%, 123%, and 50% higher factors respectively than those in pure PMN-PT thin films. Compared with PMN-0.30PT grown by a chemical solution deposition (CSD) method [19], the ε r in PMN-0.30PT thin films on Si substrate was lower than that of the sample synthesized on perovskite structure single crystal substrates. A similar tendency of pyroelectric property was found in PMN-0.30PT thin film grown on various substrates. The p r value in 2.5% Pr-PMN-PT thin films was still higher than that of PMN-0.32PT thin films deposited on Ba 0.5 Sr 0.5 RuO 3 /NdScO 3 and La 0.5 Sr 0.5 CoO 3 /LaAlO 3 crystals using pulsed laser deposition. In contrast, the p y value showed an opposite trend, which is attributed to the lattice misfit. Compared with PVDF and normal ferroelectric material, the relaxor-PT thin films showed higher dielectric and pyroelectric properties. In contrast, the response time of the PMN-0.30PT pyroelectric sensor was longer than that of PVDF [37] and LiTaO 3 [38].  εr dielectric constant, tanδ dielectric loss, py pyroelectric coefficient, pr remnant polarization, pulsed-laser deposition (PLD), chemical solution deposition (CSD).
In order to evaluate the Pr 3+ doping effect in PMN-PT thin films, the elements distribution was measured, as shown in Figure 3, corresponding to the surface micrograph presented in Figure 2a. The Pb, Mg, O, Pr, Nb, Ti, Si, and Pt maps were achieved by performing standardless quantification of the energy dispersive X-ray spectrometry (EDS) signals in the Lα1 and Kα1 peaks. From these pictures, it can be seen that the Pb, Mg, Nb, Ti, Si, and Pt elements were homogeneously distributed in the local area. It is interesting to note that the Pr element is mostly distributed along the grain boundaries of thin films, compared to Figure 2a.  In order to evaluate the Pr 3+ doping effect in PMN-PT thin films, the elements distribution was measured, as shown in Figure 3, corresponding to the surface micrograph presented in Figure 2a. The Pb, Mg, O, Pr, Nb, Ti, Si, and Pt maps were achieved by performing standardless quantification of the energy dispersive X-ray spectrometry (EDS) signals in the Lα1 and Kα1 peaks. From these pictures, it can be seen that the Pb, Mg, Nb, Ti, Si, and Pt elements were homogeneously distributed in the local area. It is interesting to note that the Pr element is mostly distributed along the grain boundaries of thin films, compared to Figure 2a.  Materials 2018, 11, x FOR PEER REVIEW 6 of 11  Figure 4a presents the ferroelectric properties of Pr-PMN-PT thin films with different Pr 3+ ratios of 0, 1.5%, and 2.5%. The addition of Pr 3+ in PMN-PT thin films increased the value of remanent polarization (Pr) and coercive field (Ec) from 11.5 μC/cm 2 and 35 kV/cm to 17.3 μC/cm 2 and 63.5 kV/cm respectively. The leakage current densities of pure and 2.5% Pr-PMN-PT thin films were approximately 1.24 × 10 −4 A/cm 2 and 5.8 × 10 −5 A/cm 2 respectively at 100 kV/cm, as shown in Figure  4b. A donor dopant in PMN-PT samples was poled at the higher electric field due to the higher Ec in Pr-PMN-PT thin films, indicating that the domain structure was easily reoriented owing to the presence of lead vacancies near the grain boundaries [44,45], corresponding to the Pr 3+ ion distribution shown in Figure 3. Those effects also led to a higher dielectric constant of Pr 3+ doped compositions compared to the undoped counterpart, as shown in Table 1.  Figure 4a presents the ferroelectric properties of Pr-PMN-PT thin films with different Pr 3+ ratios of 0, 1.5%, and 2.5%. The addition of Pr 3+ in PMN-PT thin films increased the value of remanent polarization (P r ) and coercive field (E c ) from 11.5 µC/cm 2 and 35 kV/cm to 17.3 µC/cm 2 and 63.5 kV/cm respectively. The leakage current densities of pure and 2.5% Pr-PMN-PT thin films were approximately 1.24 × 10 −4 A/cm 2 and 5.8 × 10 −5 A/cm 2 respectively at 100 kV/cm, as shown in Figure 4b. A donor dopant in PMN-PT samples was poled at the higher electric field due to the higher E c in Pr-PMN-PT thin films, indicating that the domain structure was easily reoriented owing to the presence of lead vacancies near the grain boundaries [44,45], corresponding to the Pr 3+ ion distribution shown in Figure 3. Those effects also led to a higher dielectric constant of Pr 3+ doped compositions compared to the undoped counterpart, as shown in Table 1.  Figure 5a shows dielectric constant (εr) and dielectric loss (tanδ) of Pr-PMN-PT thin films in the frequency range of 20 Hz to 200 kHz at room temperature. We see that the εr value decreased slightly with an increasing frequency, which is induced by the polarization decrease due to charge accumulation at low frequency. Meanwhile, the dielectric loss showed a very quick increase in the range 10 4 -2 × 10 5 Hz, indicating conduction through the electrode surface and grain boundary capacitance interference in poled samples at high frequency. Figure 5b presents the dielectric properties dependent on temperature for Pr-PMN-PT thin film at 100 Hz and 1 kHz. The εr curve presents a broad peak centered at around 100 °C, which corresponds to the diffuse ferroelectric to paraelectric phase transition [10,17]. The value of dielectric loss values is about 1.8% in the temperature range from 25 to 100 °C, which indicates good electrode bonding for further device integrating.  Figure 6a shows the pyroelectric coefficient (py) measured at room temperature for the obtained PMN-PT thin films as a function of Pr 3+ doping content. The py in the thin films with 2.5% dopant presented the maximum value on the order of 167 μC/m 2 K, which is almost a two-fold increase over that of pure PMN-PT thin films, offering a new way for enhancing the pyroelectric performance in relaxor-PT thin films for use in infrared detectors. The temperature dependence of py values is  Figure 5a shows dielectric constant (ε r ) and dielectric loss (tanδ) of Pr-PMN-PT thin films in the frequency range of 20 Hz to 200 kHz at room temperature. We see that the ε r value decreased slightly with an increasing frequency, which is induced by the polarization decrease due to charge accumulation at low frequency. Meanwhile, the dielectric loss showed a very quick increase in the range 10 4 -2 × 10 5 Hz, indicating conduction through the electrode surface and grain boundary capacitance interference in poled samples at high frequency. Figure 5b presents the dielectric properties dependent on temperature for Pr-PMN-PT thin film at 100 Hz and 1 kHz. The ε r curve presents a broad peak centered at around 100 • C, which corresponds to the diffuse ferroelectric to paraelectric phase transition [10,17]. The value of dielectric loss values is about 1.8% in the temperature range from 25 to 100 • C, which indicates good electrode bonding for further device integrating.

Temperature Dependent of Dielectric and Pyroelectric Properties
Figure 5a shows dielectric constant (εr) and dielectric loss (tanδ) of Pr-PMN-PT thin films in the frequency range of 20 Hz to 200 kHz at room temperature. We see that the εr value decreased slightly with an increasing frequency, which is induced by the polarization decrease due to charge accumulation at low frequency. Meanwhile, the dielectric loss showed a very quick increase in the range 10 4 -2 × 10 5 Hz, indicating conduction through the electrode surface and grain boundary capacitance interference in poled samples at high frequency. Figure 5b presents the dielectric properties dependent on temperature for Pr-PMN-PT thin film at 100 Hz and 1 kHz. The εr curve presents a broad peak centered at around 100 °C, which corresponds to the diffuse ferroelectric to paraelectric phase transition [10,17]. The value of dielectric loss values is about 1.8% in the temperature range from 25 to 100 °C, which indicates good electrode bonding for further device integrating.  Figure 6a shows the pyroelectric coefficient (py) measured at room temperature for the obtained PMN-PT thin films as a function of Pr 3+ doping content. The py in the thin films with 2.5% dopant presented the maximum value on the order of 167 μC/m 2 K, which is almost a two-fold increase over that of pure PMN-PT thin films, offering a new way for enhancing the pyroelectric performance in relaxor-PT thin films for use in infrared detectors. The temperature dependence of py values is  Figure 6a shows the pyroelectric coefficient (p y ) measured at room temperature for the obtained PMN-PT thin films as a function of Pr 3+ doping content. The p y in the thin films with 2.5% dopant presented the maximum value on the order of 167 µC/m 2 K, which is almost a two-fold increase over that of pure PMN-PT thin films, offering a new way for enhancing the pyroelectric performance in relaxor-PT thin films for use in infrared detectors. The temperature dependence of p y values is illustrated in Figure 6b. In the temperature range of −50-100 • C, the pyroelectric coefficient showed a slight fluctuation (~10%), indicating good thermal stability for application in infrared detectors integrated into Si semiconductor devices. illustrated in Figure 6b. In the temperature range of −50-100 °C, the pyroelectric coefficient showed a slight fluctuation (~10%), indicating good thermal stability for application in infrared detectors integrated into Si semiconductor devices.

Conclusions
The pure and Pr 3+ doped PMN-PT ferroelectric thin films with a single perovskite structure were prepared by a sol-gel processing and investigated by X-ray diffraction. The microstructure of Pr-PMN-PT samples was characterized by scanning electron microscopy, atomic force microscopy, and piezoresponse force microscopy. The average grain size was about 100 nm. It was found that Pr 3+ ions were located at the edge of grain boundaries measured by energy dispersive X-ray spectroscopy. For 2.5% Pr 3+ doping, the dielectric constant, pyroelectric coefficient, and remenant polarization all reached a maximum. Importantly, the dielectric constant in PMN-PT thin films increased from 1140 to 2400 through 2.5% Pr 3+ doping. Such enhancement was ascribed to the domain wall motion induced by the Pr 3+ ion distribution around the grain boundary. Finally, this rare earth dopant presents a useful method for developing high performance relaxor-PT thin films for use in energy transducers and photodetectors.