Manipulating Ferroelectrics through Changes in ... - ACS Publications

Subscriber access provided by UNIVERSITY OF LEEDS

Article

Manipulating Ferroelectrics through Changes in Surface and Interface Properties Nina Balke, Ramamoorthy Ramesh, and Pu Yu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b10747 • Publication Date (Web): 23 Oct 2017 Downloaded from http://pubs.acs.org on October 25, 2017

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

ACS Applied Materials & Interfaces is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 35

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

ACS Applied Materials & Interfaces

Manipulating Ferroelectrics through Changes in Surface and Interface Properties Nina Balke1*, Ramamoorthy Ramesh2, Pu Yu3,4* 1

2

Oak Ridge National Laboratory, 1 Bethel Valley Road, Oak Ridge TN, 37831, USA University of California, Berkeley, 210 Hearst Memorial Mining Building, Berkeley, CA

94720, USA 3

Tsinghua University, and Collaborative Innovation Center of Quantum Matter, Beijing 100084,

China 4

RIKEN Center for Emergent Matter Science (CEMS), Wako, Saitama 351-0198, Japan

Keywords: ferroelectrics, nanodomains, retention, scanning probe microscopy, electrode interface, oxygen vacancies

Abstract Ferroelectric materials are used in many applications of modern technologies including information storage, transducer, sensor, tunable capacitor and other novel device concepts. In many of these applications, the ferroelectric properties, such as switching voltages, piezoelectric constants, or stability of nanodomains are crucial. For any application, even for material characterization, the material itself needs to be interfaced with electrodes. Based on the structural, chemical, and electronic properties of the interfaces, the measured material properties can be determined by the interface. This is also true for surfaces. However, the importance of interfaces and surfaces and their effect on the experiments are often neglected which results in 1 ACS Paragon Plus Environment


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

many dramatically different experimental results for nominally identical samples. Therefore, it is crucial to understand the role of the interface and surface properties on internal bias fields and the domain switching process. Here, the nanoscale ferroelectric switching process and the stability of nanodomains for Pb(Zr,Ti)O3 thin films are investigated by using scanning probe microscopy. Interface and surface properties are modulated through the selection/redesign of electrode materials as well as tuning the surface-near oxygen vacancies, which both can result in changes of the electric fields acting across the sample, and consequently controls the measured ferroelectric and domain retention properties. By understanding the role of surfaces and interfaces, ferroelectric properties are tuned to eliminate the problem of asymmetric domain stability by combining the effects of different electrode materials. This study forms an important step towards integrating ferroelectric materials in electronic devices.

1. Introduction Information storage is one of the most important aspects of today’s information technology. While many concepts for information storage exist, including magnetic, resistive and phase change, the active demand for new concepts is ubiquitous due to the growing requirements of higher information density, lower power consumption as well as better device performance. Newer concepts including ferroelectric random access memories (FeRAM),1 and more recently probe-based data storage systems,2-11 possess promising application potentials. While FeRAM is promising with the advantages of lower power usage and faster writing performance, its application is only limited to some niche markets, due to the lower storage density and higher cost. Alternatively, recently developed concepts of probe-based ferroelectric memory reignited the research interests of ferroelectric memories due to the great improvement of the storage

2 ACS Paragon Plus Environment

Page 2 of 35

Page 3 of 35

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


densities owing to the reduced dimension size of atomic force probe comparing with the sandwiched capacitor structures of the FeRAM. However, to further develop the application of ferroelectric in information technology, various key reliability issues, such as fatigue, imprint and retention loss, remain to be fully understood and solved.

Imprint, which is due to the built-in bias voltage of the thin films, results in an asymmetric hysteresis loop and thus makes one of the states (the as-grown preferred polarization direction) more stable over the other. So far, research work on the imprint effect has mostly focused on sandwiched structures with both top and bottom electrodes, which are of relevance to only FeRAM applications,12-16 while less attention has been paid to the probe-based ferroelectric memory when scanning probe microscopy (SPM) probes are used to measure the internal bias fields.17, 18 In the past, different concepts have been explored to change bias fields including postannealing treatment at different temperatures and oxygen partial pressures to change the oxygen vacancy concentrations,16,

19-21

cooling rates,22 biaxial strain,23 or application of DC electrical

fields.15, 24, 25 In the first case, it has been demonstrated indirectly that the polarization direction can be switched without the application of electrical fields by changing the oxygen vacancy concentration through post-annealing as a result of a modulated bias field acting across the ferroelectric film.20,

21

However, this experiment is only limited to indirect structure

measurements; There has been no report for the direct probe of the correlation between annealing process with the ferroelectric states as well as its influence on the ferroelectric dynamic properties, such as bias field and retention, of the ferroelectric heterostructures.



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

Page 4 of 35

Retention loss is described as the loss of information over time. In the specific case of atomic force probe based ferroelectric memory, this means the instability of artificial constructed nanometer-sized domains (or called nanodomains are used to store information with their polarization orientation. A lot of work has been done experimentally23, 26-30 and theoretically31, 32 to determine the size of switched domains as a function of the amplitude and time of the applied bias pulse. From these measurements, domain wall velocities are often extracted. Ultimately, the domain size depends not only on the applied voltage pulses, but also on the bias field described above, the sample/electrode contact properties

33, 34

and defect chemistry. Also, factors such as

surface charges and humidity in the environment can also strongly affect the results.35,

36

Although the retention properties of nanodomains have been investigated by experiment37, 38 and theory,39 an overarching picture is still on how retention is correlated with the bias field and how to manipulate the retention properties to overcome the technological reliability issues for specific applications.

Although ferroelectric materials break the inversion symmetry, the two opposite ferroelectric polarization directions are energetically identical. However，this is no longer the case when the thin film/bottom electrode interface and thin film/air surface are considered, which can result in very asymmetric domain stabilities. Therefore, interfaces and surfaces are the key locations to manipulate ferroelectric retention properties. Here, a comprehensive study of bias fields and retention properties of Pb(Zr0.2Ti0.8)O3 (PZT) thin films after a series of annealing experiments and design of different bottom electrodes to tune the surface and the interface properties is presented. The PZT films used in this study were grown on (001) –oriented SrTiO3 (STO) substrates with SrRuO3 (SRO) and (La0.7Sr0.3)MnO3 (LSMO) as bottom electrodes by pulsed


Page 5 of 35

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


laser deposition (PLD) method, as described previously.40-42 Piezoresponse force microscopy (PFM) technique is employed, which has been used widely to image physical properties in ferroelectrics—such as domain structure, local switching voltage, etc. in ferroelectric thin films as well as relaxor ceramics and single crystals.43-51 SRO as well as LSMO are good electric conductors and we can assume that the applied electric voltage results in comparable electric fields.33 Here, the ferroelectric properties are measured by 1) determination of the as-grown polarization direction by so called PFM double square poling, 2) extracting the coercive voltages and bias fields from locally measured and averaged PFM hysteresis loops, and 3) switching small domains with short voltage pulses applied to the tip and then subsequently imaging them as a function of time to extract the retention properties.

2. Tuning surface properties To investigate the correlation between the surface properties and the bias field, as well as retention properties, the ferroelectric state of the PZT/SRO/STO samples is probed by measuring the out of plane (OP) PFM signal. Figure 1a shows that the as-grown samples with upward (dark contrast) is the preferred polarization direction. By selectively applying positive (+5 V) or negative (-5 V) voltage across the PFM tip, the polarization direction can be switched between downward (bright) and upward (dark) states with sharp domain boundary, which confirms the high-quality nature of the ferroelectric materials studied here. The results also indicate that the upward polarization state is the preferred state for the PZT/SRO/STO sample, and further suggests a bias field pointing from the bottom electrode toward the top surface of the PZT film.



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

To understand the driving force for this upward preferred polarization, a systematic study of the electrostatic properties of the SRO/PZT interface is necessary. Figure 1b shows a schematic drawing of the band alignment across the SRO/PZT interface. Electrostatic equilibrium condition requires that PZT layer and SRO layer should have the same Fermi energy and continuous vacuum level.

With the work functions of SRO layer (4.6 to 4.8 eV)52 as well as the

semiconductor electron affinity (~3.5 eV)1 and band gap (3.4 eV)1 of the PZT layer, the band alignment across the interface can be constructed, as shown in Figure 1b, in which 1 eV energy potential difference from the interface will be formed to make these two materials in electrostatic equilibrium.53 Consequently, the PZT side forms the potential well for electrons to accumulate, while leaving the positive screening charges at the SRO side. Thus, the separation between positive and negative charges will result in an internal field pointing from SRO to PZT to compensate the Fermi energy difference between SRO and PZT, which makes the upward state more energetically favorite. At this point, it is unclear how the electronic structure at the PZT/electrode interface will affect potentials which are measurable at the sample surface which can be studied by Kelvin probe force microscopy in correlation with domain switching events.54

Since the bias field is determined by the stacked charges, the straightforward ideas to control this field would be either inducing another charge-distribution between PZT layer and vacuum/air or directly controlling the charge density at the PZT/SRO interface. To test the validity of the first approach, small pieces of the identical as-grown sample were annealed at 380oC (heating rate 20oC/min, 30 min holding time, cooling rate 5oC/min) with ambient pressure of 10 mTorr. Figure 1c shows that after annealing, the polarization direction was changed from upward state (dark) to downward state (bright). This experiment clearly demonstrates that the as-grown


Page 6 of 35

Page 7 of 35

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


ferroelectric state can indeed be switched by defect engineering through thermal annealing which has been shown previously by other methods.21

The oxygen vacancies induced can be

accumulated both in the bulk and/or at the surface regions. However, the formation of oxygen vacancies in the bulk would result in an increase of the out-of-plane lattice constant in perovskite oxides,55,

56

which is not observed in our X-ray diffraction results, as shown in Figure S1.

Therefore, it is reasonable to deduce that the oxygen vacancies are mainly formed at the PZT surface while having a small gradient inside the film as shown in the inset of Figure 1d. , one oxygen

According to the chemical defect forming equation:

vacancy will be formed along with two free electrons. The free electrons will move toward the SRO/PZT interface due to the band alignments as shown in Figure 1b, which sets up a energy well for electron at the bottom of the conducting band at the interface. As a consequence, the separation of the positive charge oxygen vacancies and negative charge free electrons will induce an extra field (Eanneal) with opposite direction to the bias field, which becomes the driving force to switch the ferroelectric polarization direction from upwards to downwards. Assuming as small as 0.2% oxygen vacancies (σ = 4*1012 /cm2) formed at the top surface, the dielectric permittivity

ε of PZT is 20057 and the PZT thickness d is 70 nm, the induced potential difference between different sides of the PZT films,

(

) can be estimated to be 0.5 V. To experimentally

estimate the induced internal field, the difference of PFM hysteresis loops for as-grown and annealed samples are compared, as shown in Figure 1d. Notable displacement of the hysteresis loops along the voltage axis was observed after the annealing process. To address this issue, the bias voltage is introduced here as

, which is -0.11 V and +0.3 V for the as-

grown sample and annealed sample, respectively. This change of the bias voltage clearly shows that the annealing process can induce an extra field as large as 0.4 V, which would be strong 7 ACS Paragon Plus Environment


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

Page 8 of 35

enough to switch the polarization direction at higher temperature with suppressed coercive voltages.

Since the reduction reaction is the major source of oxygen vacancy forming, the charge density n is given by the charge neutrality condition: n ≈ 2[Vo2 + ] ∝ e −1 / kT P (O2 ) −1 / 6 . From this relationship, it is clear that the oxygen vacancy concentration is closely related with the annealing pressure and temperature. Thus, a systemic study of the annealing process as functions of pressure and temperature is necessary.

The first set of samples was annealed at different oxygen pressures ranging from 40 Torr to 0.01 mTorr at 380oC following the same sequence as previously described. To establish the relationship between the preferred polarization direction and the annealing pressure, OP PFM measurements were carried out, as shown in Figure 2a-2f. The results can be clearly classified into three categories: 1) for the annealing pressure at 40 Torr (Figure 2a) and 3 Torr (Figure 2b), the samples show mainly upward polarization; 2) for the annealing pressure at 10 mTorr (Figure 2d), 0.4 mTorr (Figure 2e) and 0.01 mTorr (Figure 2f), the samples mainly show downward polarization; 3) while the sample annealed at 100 mTorr shows a mixture of both upward and downward polarizations consistent with predicted domain structures.58

The corresponding

coercive and bias voltages are shown in Figure 2n. The bias voltages slightly increase from negative side to positive side when decreasing the annealing pressure. Note that for the sample annealed at 100 mTorr, PFM loops measured on domains pointing upwards (closed symbol) show an averaged negative bias voltage whereas domains pointing downward (open symbol) display a positive bias voltage on average. However, the observed trends are rather weak when


Page 9 of 35

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


considering the error bars. On the other hand, the coercive voltages almost monotonously increase as the pressure decreases due to the accumulation of oxygen vacancies inside the film, which can hinder the movement of domain wall.

To study the annealing temperature dependence, the annealing pressure is fixed at 100 mTorr and the temperature is varied from 280oC to 580oC. Figures 2g-2m show that when T < 300oC, the polarization is upward; when 300oC < T < 500oC, the polarization is the mixture of both upward and downward; when 500oC < T, the polarization becomes upward again. The coercive and bias voltages in Figure 2o show that the coercive voltages are mainly constant. For the bias voltage, domains pointing upwards show dominantly negative bias voltages and domains pointing downwards show close to zero or positive bias voltages. If only as-grown upward oriented domains are analyzed, the bias voltages increase with increasing annealing temperature up to 480oC followed by a decrease at higher temperatures. Following the oxygen defect forming relationship, with identical pressure, the higher temperature will lead to dramatically enhanced oxygen vacancy density, and as a consequence, larger annealing field (Eanneal) as shown in Figure 1d, which can be employed successfully to explain the increase of the bias voltage below 480oC. However, the results above 480oC are beyond the expectation. This discrepancy can be understood with two possible hypotheses. When the sample annealed at high temperature, because of its high volatility, Pb will be evaporated to form negatively charge Pb vacancies VPb”. As a result, the net positive charge density at the top surface could decrease as the temperature increases. Another possible mechanism is related to oxygen vacancies movement. At high temperature, the diffusion constant of the oxygen vacancies becomes larger, and more and more vacancies will move from the top surface to the bottom surface, i.e., the SRO/PZT interface, as



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

Page 10 of 35

has been reported in a similar interface of LSMO and SrRuO3. 59 Both scenarios can be used to successfully explain the suppression of the annealing field Eanneal (or decreasing of bias voltages), as well as the back switching of the ferroelectric polarization. Nevertheless, when looking at the change of bias voltage, it can be concluded that the sign of the bias voltage is always a good indicator for the internal fields determining the global (or local) polarization direction.

Having achieved the full control of the ferroelectric polarization state by using thermal annealing process, the ferroelectric domain retention, which is directly relevant to the memory applications, is studied next. Small domains against and in the direction of the as-grown polarization direction (upward) are switched and imaged as a function of time. Figure 3a shows the example of domains switched against the as-grown polarization direction. While domains switched in the asgrown polarization direction (negative voltages) are stable, domains switched against the asgrown domain orientation disappear in just over 100 min for an initial domain size of around 120-130 nm (Figure 3b). This difference in domain stability can be explained by the bias field discussed above, which drives the domains back to the preferred polarization direction to lower the total energy. The appearance of domain walls will increase the total energy and leads to further domain instability. Thus, a systematic study of the domain stability depending on the size of the domain is highly desired and shown in Figure 3c. To switch domains of different sizes, different pulse amplitudes and durations were used. In addition to switching domains by voltage pulses, small areas were scanned and poled with an applied dc voltage. Independent of the domain switching procedure, all domains with sizes up to 300 nm in diameter disappeared as function of time in the similar manner. In first order approximation, the experimental data for


Page 11 of 35

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


retention properties can be nicely fitted with an exponential decay function

with

D0 being the initial domain size, t is time, and τ the characteristic decay time. The decay time as function of D0 reveals an almost exponential relationship (Figure 3d), indicating the importance of domain size for retention. We speculate that the observed strong correlation between the domain size and the characteristic decay time could be explained based on the scenario of a possible non-penetrating nature of the switched domains,9 meaning the switched domains don’t reach the bottom electrode for stabilization. In that case, the length of the domain wall becomes a critical factor for the relaxation time. Note that an increase of temperature up to 125˚C did not affect the retention properties as shown in Figure S3 in the supporting information. Knowing the fact that as-grown (upward) state direction shows better retention performance comparing with the unfavorable downward state, the focus turns to study the influence of the annealing procedure on domain stability, as shown in Figure 4. Three represented samples with different preferred ferroelectric states have been selected for this analysis. In all cases, the domains switched upward with negative voltages pulses are stable (Figure 4a) regardless of the initial polarization directions. For domains switched downward (Figure 4b), the mixed domain state behaves more instable while the reversal of the polarization direction after annealing at low oxygen partial pressures did not affect the domain instability. This becomes even more evident when the sample with reversed domain orientation (0.01 mTorr, 380˚C) was investigated in more detail. Figure 4c shows the stability of domains of different sizes switched with positive voltages. Similarly, as the as-grown sample, the curves are fitted with a simple exponential decay function and the decay time is plotted as a function of initial domain size and compared with the as-grown sample (Figure 4d). It can be concluded that despite the change in bias voltages and domain orientation of the PZT film, the retention properties are unchanged and the domain instability is



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

Page 12 of 35

driven by the same processes. This can be explained by the fact that the oxygen vacancies introduced through annealing locate mainly in the surface-near region, while the charges at the PZT/SRO interface are unchanged. From the experiment, it can be concluded that those charges are the driving force for the domain instability and therefore, unchanged through the annealing process. This means the retention properties are not directly related to the measured bias voltages, and therefore, chemical doping at the surface through thermal annealing treatments are not sufficient for retention improvement.

3. Tuning interface properties In order to manipulate the charge density at the PZT/electrode interface to actively tune the retention properties, PZT was grown on La0.7Sr0.3MnO3 (LSMO) as bottom electrode, which is a p-type metal with the work function of 4.2 to 4.8 eV60 depending on the oxygen content. Although the work function of LSMO is compatible with that of the SRO, we observed a distinct different as-grown state as shown in the inset of Figure 5a.

These results can be easily

understood based on a recent proposed polar catastrophe model,61 in which a polar discontinuity between polar metal LSMO and PZT will induce a large internal field at the interface and lead to downward state as the preferred direction for MnO2 terminated LSMO layer as studied here.62 As shown in Figure 5a and 5b, the as-grown domain orientation and the domain stability are reversed compared with those of the PZT/SRO. When the effect of domain size on the decay time was investigated (Figure 5c), it was found that all domains decay with the same rate independent on their initial domain sizes. This becomes even clearer when the domain size as function of time is normalized by the initial domain size D0 (Figure 5d). All retention curves fall on the same curve. This shows that the physical mechanisms responsible for the domain


Page 13 of 35

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


instability are fundamentally different as compared to that of the PZT/SRO (Figure 3). However, the nature of the physical mechanism is unknown at this time. Here we speculate that while the domain instability in PZT/SRO is driven mainly by the domain walls (explaining the close correlation with the domain size) for the PZT/LSMO case the interface properties might play an essential role instead. This would explain that the domain size has no influence on the domain stability.

In order to combine the opposite retention characteristics of SRO and LSMO to improve the overall retention performance, PZT was grown on SRO which was covered with a thin (1 nm and 6 nm) LSMO layer with the stacking sequence as PZT/LSMO/SRO. The thickness of the LSMO layer is described by the parameter x. Below 4 nm LSMO thickness, the orientation of the global polarization direction is like an SRO alone (Figure 6a-6c). LSMO thicker or equal than 4 nm changes the polarization direction to the one of LSMO alone (Figure 6d-6e). We speculate that this might be due to switching of the termination layer of LSMO during the growth from LaSrO terminated to MnO2 terminated layer or through residual effects of the underlying SRO layer which is not fully screened by a ultrathin LSMO layer, since the LSMO has a critical thickness of around 2.4 nm, below which the thin films behave as insulators.63 When measuring the bias voltage of the ferroelectric hysteresis loops, a strong change from negative to positive Vbias was observed (Figure 6f) to be consistent with the observed as-grown polarization orientation for the different samples.

The measurement of the retention properties revealed that domains switched upwards with negative voltages remain stable independent of the LSMO thickness (Figure 7a) comparable



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

Page 14 of 35

with SRO electrodes. It is surprising that in this case, the domain stability does not reflect the stable polarization direction shown in Figure 6 which changes between a LSMO thickness of 3 and 4 nm. In Figure 7, the thickness of the LSMO layer is described by the parameter x. For a thickness of 3 nm LSMO domain growth was observed after long periods of time. For domains switched downwards with positive voltages (Figure 7b), it was found that with increasing LSMO thickness the domains become bigger under the same switching conditions and more stable as they are for LSMO electrodes alone. When comparing the samples with different LSMO thickness, the sample with 6 nm LSMO shows the best retention properties. The stability of the domains pointing downwards is improved compared to SRO alone while without the compromise of the domain stability at the opposite direction. Therefore, the combination of bottom electrodes seems to be a useful tool to improve and equalize the retention properties. At this point, the exact origin if the observed unusual domain stability behavior for mixed electrodes, especially in light of the stable polarization direction of the film is unknown. Detailed structural and chemical investigation of the interfaces and the film itself would need to be conducted which is subject to further studies. However, the results in Figure 7 show an unexpected trend which is different than the cases of the pure electrodes suggesting the creation of metastable material bahvior.

In order to study the retention properties of PZT/6 nm LSMO/SRO in more detail, temperaturedependent measurements were performed. While there is no effect of temperature on domain stability in PZT/SRO (Figure S3), the domain stability is strongly affected by temperature for the mixed electrode sample, as shown in Figure 8, in which only the normalized domain diameter is shown. The experiments were repeated multiple times for each temperature and the domain size


Page 15 of 35

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


scattered quite drastically in different regions for the same switching conditions. However, as for the case of pure LSMO electrodes (Figure 5), all curves for one temperature fell on a universal curve when the diameter is normalized by the initial diameter after switching which justifies this approach (Figure S4). For domains pointing upward, an increase in temperature results in a strong decrease in decay time constant, i.e., domains become more instable (Figure 8a). This means the increased stability of upwards-oriented domains in the samples with mixed electrodes is a result of restricted kinetics, which can be enhanced by temperature. For domains switched downward with positive voltages, an increase in temperature results in the growth of domains (Figure 8b). From the fitted decay times, two different activation energies for domain shrinkage and domain growth can be observed (Figure 8c). They can be determined to be 0.64+/-0.09 eV for domain shrinkage and 0.47+/-0.20 eV for domain growth. We note that other electrode materials or a different order in electrode stacking might lead to a different result. We speculate that the different domain kinetics observed for SRO and LSMO bottom electrodes could stem from the existence of surface domains (SRO) vs. switched-through domains (LSMO). In the case of surface domains, the domain wall length forms a strong driving force for domain stability, which naturally depends on the domain size. In the case of switched-through domains, the domain wall length does not depend on the domain size as well as the observed retention behavior. The exact This and the identification of physical mechanisms behind the domain stability and instability will be subject to future research.

4. Conclusion In summary, the influences of surface and interface properties on static and dynamic ferroelectric properties, such as domain structure, internal bias field and domain stability were investigated



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

Page 16 of 35

for epitaxially grown PZT thin films. This study not only provides the important knowledge for the probe-based information storage technology but also forms a generic guideline for the manipulation of ferroelectric properties since the samples need to be interfaced with electrodes for any electrical measurements. Annealing at low oxygen partial pressures modifies the concentration of oxygen vacancies at the top surface and subsequently changes the global polarization direction as well as Vbias. However, retention properties remain unaltered which suggests that the charges at the PZT/electrode interface determine the domain instabilities. This has been further confirmed by the retention results on samples with SRO and LSMO electrode, in which it was found that the bottom electrode strongly affects the bias field and retention properties. In both cases, the domains oriented against the as-grown polarization direction are unstable and vanished over time. While for SRO, the decay time is proportional to the size of the domain, for LSMO all domains show the same decay time independent of the domain size, indicating two very different physical mechanisms. A combination of LSMO and SRO was shown to improve the domain stability creating stable domains independent of their orientation. Temperature dependent measurements reveal a kinetic frustration of domain instability as origin of this observation. This is an important step to overcome some of the technological challenges associating with practical applications of ferroelectric materials in electronic devices.

5. Experimental Details Thin film synthesis: The PZT films, used in this study were grown on (001) –oriented SrTiO3 (STO) substrates with SrRuO3 (SRO) and (La0.7Sr0.3)MnO3 (LSMO) as bottom electrodes by pulsed laser deposition (PLD) method, as described previously.40-42 X-ray diffraction patterns (shown in S1 in the


Page 17 of 35

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


supporting information) indicate that the PZT films are each single-phase, highly oriented along the c axis, and essentially epitaxial. The thickness of SRO, LSMO and PZT is 50 nm, 50 nm and 70 nm respectively by controlling the growth time and is confirmed by X-ray reflectometry.

Piezoresponse force microscopy: Standard PFM imaging and poling was applied. While poled domain squares were imaged under standard imaging conditions (1 Vac voltage with the frequency of 30 kHz), the nanodomains were imaged at frequencies close to the contact resonance frequency64 around 350 kHz in order to decrease the imaging voltage to 0.04 Vac so as to reduce the change in domain size as the result of the applied imaging voltage. To trace the domain evolution, a series of 10-20 domains were switched in a row and the domain size was averaged. The domain size was determined by using the strong PFM phase contrast between domains of different orientations. To investigate the stability of domains in the same direction as the as-grown polarization direction, large areas were poled to the opposite direction of the as-grown polarization and experiments were performed in smaller sections in the middle of the poled area (see also section 2 in the supporting information for more details). PFM hysteresis loops were measured under standard imaging conditions (1Vac at 30 kHz) while sweeping a DC voltage applied simultaneously.



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

Figures


Page 18 of 35

Page 19 of 35

(c)

(a)

1 µm

2.3 eV

Ec EF Ev

(d)0.2 Piezoresponse (a.u.)

Ebias

3.5 eV

PZT

3.4 eV

SRO

1 eV

(b) 4.8 eV

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


As-grown Annealed

0.1

0.0

VO•• VO•• PZT

-0.1

eeee SRO -6

-3

0

3

6

Voltage (V)

Figure 1. OP PFM contrast of SRO/PZT thin films with as-grown (a) and after annealed at 10 mTorr and 380oC (c); (b) schematic drawing of band alignment across the SRO/PZT interface; (d) Piezoresponse hysteresis loops of as grown sample and sample annealed at 10 mTorr and 380oC. The inset in (d) shows a schematic drawing of oxygen vacancies distribution for the annealed sample.



(a)

(b)

40000 mTorr

(c)

3000 mTorr

(d)

(e)

9 mTorr

100 mTorr

(f)

0.4 mTorr

0.01 mTorr

1 µm

(g) 280

(h) oC

(i)

330

oC

380

(k)

(j) oC

430

oC

480

(l) oC

530oC

(m) 580oC

1 µm

(n)

(o)

0.6

open: down domains closed: up domains

2

0.4

0.6 0.4

0.2

0

0.0

-1 -2

Vc

+

Vc

-

-0.2

5

4

0

0.0 -0.2

-1

Vbias

10

0.2 Vc [V]

1

Vbias [V]

1 Vbias [V]

2

Vc [V]

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

Page 20 of 35

3

2

1

0

-1

-2

-0.4

-0.4

-0.6

-2 -0.6 250 300 350 400 450 500 550 600 Annealing Temperature [C]

10 10 10 10 10 10 10 Annealing Pressure [mTorr]

Figure 2: PFM images for PZT/SRO after annealing at 380˚C with different oxygen partial pressures ranging from 40 Torr to 0.01 mTorr (a-f) and at 100 mTorr oxygen pressure at different temperatures from 280˚C to 580˚C (g-m). Coercive voltages and bias voltages extracted from average PFM hysteresis loops as function of (n) annealing pressure with fixed temperature at 380 ˚C and (o) annealing temperature at fixed pressure of 100 mTorr. The color legends in Fig. (n) and (o) are identical.


Page 21 of 35

(a)

(b)

Time

Diameter D [nm]

160 140 120 100 80 10 V, 10 ms -10V, 10 µs

60 40 1

10

100

1000

Time t [min]

(d) 15 V, 10 ms 10 V, 10 ms 10 V, 1 ms

350 300

300 nm, 6 V 230 nm, 6 V 150 nm, 6 V

250 200 150 100

Decay time τ [min]

(c) Diameter D [nm]

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


1000

100

50 1

10

100

1000

100

Time t [min]

150 200 250 Diameter D0 [nm]

300

Figure 3: (a) Imaging of nano domains through the PFM phase image as function of time for PZT/SRO. The inlet shows the PFM image of two squares poled with positive (bright square) and negative voltages (small dark square). (b) Size of domains as function of time for positive and negative switching voltages applied. (c) Time-dependent domains stability for domains with different initial sizes switched with positive voltage through application of short pulses or poled by a scanning tip. Lines indicate fit to an exponential decay function. (d) Decay time from exponential fit in (c) as function of domain size.



Change of retention properties due to annealing (b)

(a)

160 120 80 -10V, 10ms As-grown (up) 40 100mTorr, 430C (up/down) 0.01mTorr, 380C (down) 0 1 10 100 1000

Diameter D [nm]

Diameter D [nm]

200

140 120 100 80 60 40 10V, 10ms As-grown (up) 20 100mTorr, 430C (up/down) 0 0.01mTorr, 380C (down) 1 10 100 Time t [min]

Time t [min]

(c)

(d)

180 160 140 120 100 80 60 40 20

as-grown (up) 0.01mTorr, 380C (down)

0.01mTorr, 380C

Decay time τ [min]

Diameter D [nm]

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

Page 22 of 35

1

15 V, 10 ms 10 V, 10 ms 10 V, 1 ms 10 Time t [min]

1000

100

100

100

150 200 250 Diameter D0[nm]

300

Figure 4: Size of nano-domains as function of time for (a) negative and (b) positive switching voltages for the as-grown PZT/SRO sample, the sample annealed at 100 mTorr at 450˚C showing a mixed domain state and the sample annealed at 0.01 mTorr at 400˚C showing a reversed polarization direction. (c) Time-dependent domains stability for domains with different initial sizes switched with positive voltages for the sample annealed at 0.01 mTorr at 400˚C. Lines indicate fit to an exponential decay function. (d) Decay time from exponential fit in (c) as function of domain size compared with Figure 1d.


Page 23 of 35

(b) 160

160

140

140 Diameter D [nm]

Diameter D [nm]

(a) 120 100 80 60 40 PZT/LSMO/STO, -8 V, 1 ms PZT/SRO/STO, -10 V, 100 µs

20 0

1

10 100 Time t [min]

120 100 80 60 40 PZT/LSMO/STO, 8 V, 1 ms PZT/SRO/STO, 10 V, 10 ms

20 0

1000

(c)

1

10 100 Time t [min]

1000

(d) 1.2 Diameter D [norm.]

350 Diameter D [nm]

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


300 250 200 150

-9V, 1ms -9V, 1ms -9V, 500µs -9V, 200µs -12V, 100µs

100 50 0 1

10

1.0 0.8 0.6 -9V, 1ms -9V, 1ms -9V, 500µs -9V, 200µs -12V, 100µs

0.4 0.2 0.0

100 1000 Time t [min]

1

10 100 Time t [min]

1000

Figure 5: Size of nano-domains as function of time for (a) negative and (b) positive switching voltages for PZT/LSMO compared to PZT/SRO. The inlet in (a) shows the PFM image of two squares poled with negative (dark) and positive voltages (bright) for PZT/LSMO. (c) Timedependent domains stability for domains with different initial sizes switched with negative voltage pulses in PZT/LSMO. (d) Data displayed in (c) after normalization of the domain size by the initial domain size.



(a)

x = 1 nm

(b)

x = 2 nm

(c)

(d)

x = 4 nm

(e)

x = 6 nm

(f) 2 1 Vc [V]

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

Page 24 of 35

0

x = 3 nm

Vc

+

Vc

-

Vbias

-1 -2 1 2 3 4 5 6 LSMO thickness x [nm]

Figure 6: PFM images after a double square poling experiment with positive and negative voltages to identify the as-grown polarization direction for PZT/ x nm LSMO/SRO sample with (a) x = 1 nm, (b) x = 2 nm, (c) x = 3 nm, (d) x = 4 nm, (e) x = 6 nm. (f) Coercive voltages extracted from average PFM hysteresis loops as function of LSMO thickness.


Page 25 of 35

(b) 300

Diameter D [nm]

Diameter D [nm]

(a)

250 200 150 100 -8V, 10 ms 50 0

LSMO thickness x = 6 nm x = 4 nm

1

10

x = 3 nm x = 2 nm x = 1 nm

100 Time t [min]

1000

200 180 160 140 120 100 80 60 40 20 0

10 V, 10 ms LSMO thickness x = 6 nm x = 4 nm x = 3 nm x = 2 nm x = 1 nm

1

10

100 Time t [min]

1000

Figure 7: Size of domains as function of time for (a) negative and (b) positive switching voltages for the PZT/x nm LSMO/SRO sample series for x ranging from 1 to 6 nm. (a)

(b)

(c) -2

2.2

0.8 0.6 0.4 0.2

-10V, 10ms 50°C 60°C 75°C 100°C 110°C 125°C

1

10 100 Time [min]

10V, 10ms RT 75°C 100°C 125°C

2.0 1.8 1.6

-3 -4

1.4 1.2

1000

-6 -7 -8

1.0 0.8

PZT/6nm LSMO/SRO/STO -10V, 10ms 10V, 10ms

-5 ln(1/τ)

1.0

Diameter D [norm.]

Diameter D [norm.]

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


-9 1

10 100 Time t [min]

1000

-10 0.0024 0.0026 0.0028 0.0030 0.0032 -1 1/T [K ]

Figure 8: Size of domains as function of time for (a) negative and (b) positive switching voltages for PZT/6 nm LSMO/SRO samples at different temperatures ranging from room temperature to 125 ˚C. (c) Arrhenius plot for the fitted decay times in (a) and (b).

Supporting Information The following files are available free of charge.



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

Additional information of sample structure and properties (PDF).

Corresponding Author *Corresponding authors: [email protected] (characterization) [email protected] (sample growth)

Author Contributions NB, RR, and PY designed the experiment. NB conducted the SPM-based characterization and PY performed the sample growth and X-ray analysis. All authors contributed to data interpretation and manuscript writing.

Acknowledgement A portion of this research was conducted at the Center for Nanophase Materials Sciences, which is a DOE Office of Science User Facility. P.Y. was supported by the National Basic Research Program of China (Grant No. 2015CB921700) and National Natural Science Foundation of China (Grant No. 11274194).

References (1)

Scott, J. F.; De Araujo, C. A. P. Ferroelectric Memories Science 1989, 246, 1400-1405.

(2)

Alexe, M.; Gruverman, A.; Harnagea, C.; Zakharov, N. D.; Pignolet, A.; Hesse, D.; Scott,

J. F. Switching Properties of Self-Assembled Ferroelectric Memory Cells Appl. Phys. Lett. 1999, 75, 1158-1160.


Page 26 of 35

Page 27 of 35

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


(3)

Paruch, P.; Tybell, T.; Triscone, J. M. Nanoscale Control of Ferroelectric Polarization

and Domain Size in Epitaxial Pb(Zr0.2Ti0.8)O3 Thin Films Appl. Phys. Lett. 2001, 79, 530-532. (4)

Kim, Y.; Cho, Y.; Hong, S.; Buhlmann, S.; Park, H.; Min, D. K.; Kim, S. H.; No, K. Tip

Traveling and Grain Boundary Effects in Domain Formation Using Piezoelectric Force Microscopy for Probe Storage Applications Appl. Phys. Lett. 2006, 89, 172909. (5)

Kim, Y.; Cho, Y.; Hong, S.; Buhlmann, S.; Park, H.; Min, D. K.; Kim, S. H.; No, K.

Correlation between Grain Size and Domain Size Distributions in Ferroelectric Media for Probe Storage Applications Appl. Phys. Lett. 2006, 89, 162907. (6)

Min, D. K.; Hong, S. Servo and Tracking Algorithm for a Probe Storage System Ieee

Transactions on Magnetics 2005, 41, 855-859. (7)

Park, H.; Jung, J.; Min, D. K.; Kim, S.; Hong, S.; Shin, H. Scanning Resistive Probe

Microscopy: Imaging Ferroelectric Domains Appl. Phys. Lett. 2004, 84, 1734-1736. (8)

Shin, H.; Hong, S.; Moon, J.; Jeon, J. U. Read/Write Mechanisms and Data Storage

System Using Atomic Force Microscopy and Mems Technology Ultramicroscopy 2002, 91, 103110. (9)

Woo, J.; Hong, S. B.; Min, D. K.; Shin, H. J.; No, K. Effect of Domain Structure on

Thermal Stability of Nanoscale Ferroelectric Domains Appl. Phys. Lett. 2002, 80, 4000-4002. (10)

Ko, H.; Ryu, K.; Park, H.; Park, C.; Jeon, D.; Kim, Y. K.; Jung, J.; Min, D. K.; Kim, Y.;

Lee, H. N.; Park, Y.; Shin, H.; Hong, S. High-Resolution Field Effect Sensing of Ferroelectric Charges Nano Lett. 2011, 11, 1428-1433. (11)

Hong, S.; Park, M. N., Resistive Probe Storage: Read/Write Mechanism. In Scanning

Probe Microscopy, Springer: 2007, pp 943-973.



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

(12)

Page 28 of 35

Wu, W. B.; Wong, K. H.; Choy, C. L. Interface-Oxygen-Loss-Controlled Voltage Offsets

in Epitaxial Pb(Zr0.52Ti0.48)O3 Thin-Film Capacitors with La0.7Sr0.3MnO3 Electrodes Appl. Phys. Lett. 2004, 85, 5013-5015. (13)

Warren, W. L.; Dimos, D.; Pike, G. E.; Tuttle, B. A.; Raymond, M. V.; Ramesh, R.;

Evans, J. T. Voltage Shifts and Imprint in Ferroelectric Capacitors Appl. Phys. Lett. 1995, 67, 866-868. (14)

Grossmann, M.; Lohse, O.; Bolten, D.; Boettger, U.; Schneller, T.; Waser, R. The

Interface Screening Model as Origin of Imprint in PbZrxTi1-xO3 Thin Films. I. Dopant, Illumination, and Bias Dependence J. Appl. Phys. 2002, 92, 2680-2687. (15)

Park, B. H.; Noh, T. W.; Lee, J.; Kim, C. Y.; Jo, W. Effects of Interface Charges on

Imprint of Epitaxial Bi4Ti3O12 Thin Films Appl. Phys. Lett. 1997, 70, 1101-1103. (16)

Pike, G. E.; Warren, W. L.; Dimos, D.; Tuttle, B. A.; Ramesh, R.; Lee, J.; Keramidas, V.

G.; Evans, J. T. Voltage Offsets in (Pb,La)(Zr,Ti)O3 Thin-Films Appl. Phys. Lett. 1995, 66, 484486. (17)

Kim, J.; Hong, S.; Buhlmann, S.; Kim, Y.; Park, M.; Kim, Y. K.; No, K. Effect of

Deposition Temperature of TiO2 on the Piezoelectric Property of PbTiO3 Film Grown by Pbo Gas Phase Reaction Sputtering J. Appl. Phys. 2010, 107, 104112. (18)

Park, M. Y.; Hong, S. B.; Kim, J. Y.; Kim, Y. S.; Buhlmann, S.; Kim, Y. K.; No, K. S.

Piezoresponse Force Microscopy Studies of PbTiO3 Thin Films Grown Via Layer-by-Layer Gas Phase Reaction Appl. Phys. Lett. 2009, 94, 092901. (19)

Hiboux, S.; Muralt, P. Origin of Voltage Offset and Built-in Polarization in in-Situ

Sputter Deposited PZT Thin Films Integrated Ferroelectrics 2001, 36, 83-92.


Page 29 of 35

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)

Wang, R. V.; Fong, D. D.; Jiang, F.; Highland, M. J.; Fuoss, P. H.; Thompson, C.;

Kolpak, A. M.; Eastman, J. A.; Streiffer, S. K.; Rappe, A. M.; Stephenson, G. B. Reversible Chemical Switching of a Ferroelectric Film Phys. Rev. Lett. 2009, 102, 047601. (21)

Highland, M. J.; Fister, T. T.; Richard, M.-I.; Fong, D. D.; Fuoss, P. H.; Thompson, C.;

Eastman, J. A.; Streiffer, S. K.; Stephenson, G. B. Polarization Switching without Domain Formation at the Intrinsic Coercive Field in Ultrathin Ferroelectric PbTiO3 Phys. Rev. Lett. 2010, 105, 167601. (22)

Feigl, L.; Janolin, P. E.; Yamada, T.; Iwanowska, M.; Sandu, C. S.; Setter, N. Post-

Deposition Control of Ferroelastic Stripe Domains and Internal Electric Field by Thermal Treatment Appl. Phys. Lett. 2015, 106, 032902. (23)

Guo, E.-J.; Roth, R.; Herklotz, A.; Hesse, D.; Doerr, K. Ferroelectric 180 Degrees

Domain Wall Motion Controlled by Biaxial Strain Adv. Mater. 2015, 27, 1615-1618. (24)

Balke, N.; Granzow, T.; Roedel, J. Degradation of Lead-Zirconate-Titanate Ceramics

under Different DC Loads J. Appl. Phys. 2009, 105, 104105. (25)

Balke, N.; Lupascu, D. C.; Granzow, T.; Roedel, J. Fatigue of Lead Zirconate Titanate

Ceramics. I: Unipolar and DC Loading J. Am. Ceram. Soc. 2007, 90, 1081-1087. (26)

Rodriguez, B. J.; Nemanich, R. J.; Kingon, A.; Gruverman, A.; Kalinin, S. V.; Terabe,

K.; Liu, X. Y.; Kitamura, K. Domain Growth Kinetics in Lithium Niobate Single Crystals Studied by Piezoresponse Force Microscopy Appl. Phys. Lett. 2005, 86, 012906. (27)

Pertsev, N. A.; Petraru, A.; Kohlstedt, H.; Waser, R.; Bdikin, I. K.; Kiselev, D.; Kholkin,

A. L. Dynamics of Ferroelectric Nanodomains in BaTiO3 Epitaxial Thin Films Via Piezoresponse Force Microscopy Nanotechnology 2008, 19, 375703.



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

(28)

Page 30 of 35

Kiselev, D. A.; Zhukov, R. N.; Ksenich, S. V.; Kozlova, A. P.; Bykov, A. S.;

Malinkovich, M. D.; Parkhomenko, Y. N. Investigation of the Ferroelectric Properties and Dynamics of Nanodomains in LiNbO3 Thin Films Grown on Si (100) Substrate by Scanning Probe Microscopy Techniques Thin Solid Films 2014, 556, 142-145. (29)

Chen, Y.-C.; Wang, G.-F.; Tai, H.-H.; Chen, J.-W.; Huang, Y.-C.; Yang, J.-C.; Chu, Y.-

H. Non-Volatile Domain Nucleation and Growth in Multiferroic BiFeO3 Films Nanotechnology 2011, 22, 254030. (30)

Woo, J.; Hong, S.; Setter, N.; Shin, H.; Jeon, J. U.; Pak, Y. E.; No, K. Quantitative

Analysis of the Bit Size Dependence on the Pulse Width and Pulse Voltage in Ferroelectric Memory Devices Using Atomic Force Microscopy J. Vac. Sci. Technol., B 2001, 19, 818-824. (31)

Morozovska, A. N.; Eliseev, E. A.; Kalinin, S. V. Domain Nucleation and Hysteresis

Loop Shape in Piezoresponse Force Spectroscopy Appl. Phys. Lett. 2006, 89, 192901. (32)

Morozovska, A. N.; Eliseev, E. A. Screening and Size Effects on the Nanodomain

Tailoring in Ferroelectrics Semiconductors Phys. Rev. B 2006, 73, 104440. (33)

Wang, B.; Woo, C. H. Atomic Force Microscopy-Induced Electric Field in Ferroelectric

Thin Films J. Appl. Phys. 2003, 94, 4053-4059. (34)

Lichtensteiger, C.; Fernandez-Pena, S.; Weymann, C.; Zubko, P.; Triscone, J.-M. Tuning

of the Depolarization Field and Nanodomain Structure in Ferroelectric Thin Films Nano Lett. 2014, 14, 4205-4211. (35)

Dahan, D.; Molotskii, M.; Rosenman, G.; Rosenwaks, Y. Ferroelectric Domain

Inversion: The Role of Humidity Appl. Phys. Lett. 2006, 89, 152902. (36)

Ievlev, A. V.; Morozovska, A. N.; Shur, V. Y.; Kalinin, S. V. Humidity Effects on Tip-

Induced Polarization Switching in Lithium Niobate Appl. Phys. Lett. 2014, 104, 092908.


Page 31 of 35

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


(37)

Gruverman, A.; Tokumoto, H.; Prakash, A. S.; Aggarwal, S.; Yang, B.; Wuttig, M.;

Ramesh, R.; Auciello, O.; Venkatesan, T. Nanoscale Imaging of Domain Dynamics and Retention in Ferroelectric Thin Films Appl. Phys. Lett. 1997, 71, 3492-3494. (38)

Auciello, O.; Gruverman, A.; Tokumoto, H.; Prakash, S. A.; Aggarwal, S.; Ramesh, R.

Nanoscale Scanning Force Imaging of Polarization Phenomena in Ferroelectric Thin Films MRS Bull. 1998, 23, 33-42. (39)

Chen, W. J.; Zheng, Y.; Xiong, W. M.; Feng, X.; Wang, B.; Wang, Y. Effect of

Mechanical Loads on Stability of Nanodomains in Ferroelectric Ultrathin Films: Towards Flexible Erasing of the Non-Volatile Memories Sci. Rep. 2014, 4, 5339. (40)

Nagarajan, V.; Prasertchoung, S.; Zhao, T.; Zheng, H.; Ouyang, J.; Ramesh, R.; Tian, W.;

Pan, X. Q.; Kim, D. M.; Eom, C. B.; Kohlstedt, H.; Waser, R. Size Effects in Ultrathin Epitaxial Ferroelectric Heterostructures Appl. Phys. Lett. 2004, 84, 5225-5227. (41)

Maksymovych, P.; Jesse, S.; Yu, P.; Ramesh, R.; Baddorf, A. P.; Kalinin, S. V.

Polarization Control of Electron Tunneling into Ferroelectric Surfaces Science 2009, 324, 14211425. (42)

Arenholz, E.; van der Laan, G.; Fraile-Rodriguez, A.; Yu, P.; He, Q.; Ramesh, R. Probing

Ferroelectricity in PbZr0.2Ti0.8O3 with Polarized Soft X Rays Phys. Rev. B 2010, 82, 140103. (43)

Gruverman, A.; Auciello, O.; Tokumoto, H. Scanning Force Microscopy for the Study of

Domain Structure in Ferroelectric Thin Films J. Vac. Sci. Technol., B 1996, 14, 602-605. (44)

Balke, N.; Bdikin, I.; Kalinin, S. V.; Kholkin, A. L. Electromechanical Imaging and

Spectroscopy of Ferroelectric and Piezoelectric Materials: State of the Art and Prospects for the Future J. Am. Ceram. Soc. 2009, 92, 1629-1647.



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

(45)

Page 32 of 35

Kalinin, S. V.; Morozovska, A. N.; Chen, L. Q.; Rodriguez, B. J. Local Polarization

Dynamics in Ferroelectric Materials Rep. Prog. Phys. 2010, 73, 056502. (46)

Roelofs, A.; Schneller, I.; Szot, K.; Waser, R. Piezoresponse Force Microscopy of Lead-

Titanate Nanograins Possibly Reaching the Limit of Ferroelectricity Appl. Phys. Lett. 2002, 81, 5231-5233. (47)

Ganpule, C. S.; Nagarajan, V.; Hill, B. K.; Roytburd, A. L.; Williams, E. D.; Ramesh, R.;

Alpay, S. P.; Roelofs, A.; Waser, R.; Eng, L. M. Imaging Three-Dimensional Polarization in Epitaxial Polydomain Ferroelectric Thin Films J. Appl. Phys. 2002, 91, 1477-1481. (48)

Kleemann, W. The Relaxor Enigma - Charge Disorder and Random Fields in

Ferroelectrics Journal of Materials Science 2006, 41, 129-136. (49)

Shvartsman, V. V.; Kholkin, A. L. Domain Structure of 0.8Pb(Mg1/3Nb2/3)O3-

0.2PbTiO3 Studied by Piezoresponse Force Microscopy Phys. Rev. B 2004, 69, 014102. (50)

Balke, N.; Gajek, M.; Tagantsev, A. K.; Martin, L. W.; Chu, Y. H.; Ramesh, R.; Kalinin,

S. V. Direct Observation of Capacitor Switching Using Planar Electrodes Adv. Funct. Mater. 2010, 20, 3466-3475. (51)

Dittmer, R.; Jo, W.; Rodel, J.; Kalinin, S.; Balke, N. Nanoscale Insight into Lead-Free

BNT-BT-xKNN Adv. Funct. Mater. 2012, 22, 4208-4215. (52)

Hartmann, A. J.; Neilson, M.; Lamb, R. N.; Watanabe, K.; Scott, J. F. Ruthenium Oxide

and Strontium Ruthenate Electrodes for Ferroelectric Thin-Films Capacitors Appl. Phys. A 2000, 70, 239-242. (53)

Scott, J. F. Ferroelectric Memories Springer: Berlin, 2000.


Page 33 of 35

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


(54)

Kim, Y.; Bae, C.; Ryu, K.; Ko, H.; Kim, Y. K.; Hong, S.; Shin, H. Origin of Surface

Potential Change During Ferroelectric Switching in Epitaxial PbTiO3 Thin Films Studied by Scanning Force Microscopy Appl. Phys. Lett. 2009, 94, 032907. (55)

Sun, J. R.; Yeung, C. F.; Zhao, K.; Zhou, L. Z.; Leung, C. H.; Wong, H. K.; Shen, B. G.

Strain-Dependent Vacuum Annealing Effects in La0.67Ca0.33MnO3-Delta Films Appl. Phys. Lett. 2000, 76, 1164-1166. (56)

Tarsa, E. J.; Hachfeld, E. A.; Quinlan, F. T.; Speck, J. S.; Eddy, M. Growth-Related

Stress and Surface Morphology in Homoepitaxial SrTiO3 Films Appl. Phys. Lett. 1996, 68, 490492. (57)

Xu, R.; Karthik, J.; Damodaran, A. R.; Martin, L. W. Stationary Domain Wall

Contribution to Enhanced Ferroelectric Susceptibility Nat. Commun. 2014, 5, 3120. (58)

Ma, D. C.; Zheng, Y.; Wang, B.; Woo, C. H. Domain Structures of Ferroelectric Thin

Film Controlled by Oxidizing Atmosphere Appl. Phys. Lett. 2011, 99, 142908. (59)

Borisevich, A. Y.; Lupini, A. R.; He, J.; Eliseev, E. A.; Morozovska, A. N.; Svechnikov,

G. S.; Yu, P.; Chu, Y.-H.; Ramesh, R.; Pantelides, S. T.; Kalinin, S. V.; Pennycook, S. J. Interface Dipole between Two Metallic Oxides Caused by Localized Oxygen Vacancies Phys. Rev. B 2012, 86, 140102. (60)

de Jong, M. P.; Dediu, V. A.; Taliani, C.; Salaneck, W. R. Electronic Structure of

La0.7Sr0.3MnO3 Thin Films for Hybrid Organic/Inorganic Spintronics Applications J. Appl. Phys. 2003, 94, 7292-7296. (61)

Nakagawa, N.; Hwang, H. Y.; Muller, D. A. Why Some Interfaces Cannot Be Sharp Nat.

Mater. 2006, 5, 204-209.



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

(62)

Page 34 of 35

Yu, P.; Luo, W.; Yi, D.; Zhang, J. X.; Rossell, M. D.; Yang, C. H.; You, L.; Singh-

Bhalla, G.; Yang, S. Y.; He, Q.; Ramasse, Q. M.; Erni, R.; Martin, L. W.; Chu, Y. H.; Pantelides, S. T.; Pennycook, S. J.; Ramesh, R. Interface Control of Bulk Ferroelectric Polarization P. Natl. Acad. Sci. USA 2012, 109, 9710-9715. (63)

Huijben, M.; Martin, L. W.; Chu, Y. H.; Holcomb, M. B.; Yu, P.; Rijnders, G.; Blank, D.

H. A.; Ramesh, R. Critical Thickness and Orbital Ordering in Ultrathin La0.7Sr0.3MnO3 Films Phys. Rev. B 2008, 78, 094413. (64)

Jesse, S.; Mirman, B.; Kalinin, S. V. Resonance Enhancement in Piezoresponse Force

Microscopy: Mapping Electromechanical Activity, Contact Stiffness, and Q Factor Appl. Phys. Lett. 2006, 89, 022906.

TOC graphic


Page 35 of 35

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



Manipulating Ferroelectrics through Changes in ... - ACS Publications

Recommend Documents