Lavinia.curecheriu@uaic.ro 578 Teza Ciuchi Ioana Text

“ALEXANDRU IOAN CUZA” UNIVERSITY DEPARTAMENT OF PHYSICS Perovskite systems with ferroelectric/antiferroelectric character A thesis submitted by IOANA VERONICA CIUCHI in partial fulfilment of the requirements for the title of Doctor of Science in Physics SCIENTIFIC COORDINATOR : PROF. UNIV. DR. LILIANA MITOȘERIU Iași 2017 2 3 Table of Contents Abstract Acknowledgements Thesis organization CHAPTER 1 Introduction and theoretical considerations 1.1 Introduction 1.2 Brief history 1.3 Definition of Ferroelectricity and Antiferroelectricity 1.4 Basic properties of ferroelectric and antiferroelectric materials 1.4.1 Perovskite Structure 1.4.2 Domain switching: Ferroelectric versus antiferroelectric behaviour a) Domain switching in ferroelectric materials b) Domain switching in antiferroelectric materials 1.4.3 Tunabili ty a) Tunability in ferroelectrics b) Tunability in antiferroelectrics 1.4.4 Energy storage properties 1.5 Landau-Ginzburg-Devonshire Theory of Ferroelectricity 1.6 Landau Theory of Antiferroelectrics CHAPTER 2 Short literature review on PZT and La doped PZT systems 2.1 Introduction 2.2 General properties of PZT solid solutions 2.2.1 PbZrO 3 2.2.2 PbTiO 3 2.2.3 Phase diagram of the PZT solid solution 2.3 Properties of PLZT solid solutions 2.4 The High Zr content side of PZT and PLZT systems 2.5 Systems investigated for this thesis and motivation for the proposed research topic 4 CHAPTER 3 Characterization: Principles and Techniques 3.1 Introduction 3.2 Preparation of PLZT ceramics 3.2.1 Preparation of dense PLZT Pellets 3.3 Characterization techniques 3.3.1 Microstructural characterization a) Scanning Electron Microscopy (TEM) b) Transmission Electron Microscopy (TEM) 3.3.2 Structural Characterisation a) X-Ray Diffraction and High Resolution X-Ray diffraction a) Raman Spectroscopy 3.3.3 Electrical characterization a) Impedance spectroscopy b) Polarization-Field Hysteresis loop measurements c) Piezoelectric characterization CHAPTER 4. Microstructural and Structural Characterization 4.1 Introduction 4.2 Phase purity 4.2.1 Phase purity of calcined PLZT powders 4.2.2 Phase purity of PLZT sintered ceramics 4.3 Microstructure 4.4 Structural characterization 4.1.1 Phase characterization by XRD 4.1.2 Crystalline structural characterization by HXRD 4.5 Domain structure and local characterization by TEM 4.6 RAMAN analysis 4.7 Conclusions 5 CHAPTER 5 Room Temperature Electrical Properties 5.1 Introduction 5.2 Dielectric Properties 5.3 Piezoelectric properties 5.4 Conclusions CHAPTER 6. Study of antiferroelectric- to-ferroelectric switching 6.1 Introduction 6.2 Polarization vs. electric field study 6.2.1 Influence of the La3+ composition on ferroelectric/antiferroelectric properties of PLZT x/90/10 ceramics 6.2.2 Frequency dependence of AFE- to-FE switching of PLZT ceramics 6.3 Study of AFE- to-FE switching by XRD 6.3.1 Dynamic in situ XRD study 6.3.2 Ex situ High Resolution X-ray Diffraction (HXRD) study 6.4 Energy storage properties 6.5 Conclusions CHAPTER 7 Study of temperature-induced phase transitions in PLZT x/90/10 ceramics 7.1 Introduction 7.2 Dielectric study 7.2.1 Influence of La addition on the phase transitions of PLZT x/90/10 ceramics 7.2.2 Influence of frequency on the phase transition of PLZT x/90/10 ceramics 7.2.3 Analysis of FE/AFE-PE phase transition with Curie Weiss law 7.2.4 Analysis of FE/AFE-PE phase transition with modified Curie Weiss law 7.2.5 Influence of poling on the phase transitions sequence of PLZT x/90/10 ceramics 6 7.3 In situ XRD temperature study 7.3.1 Phase transitions in virgin PLZT x/90/10 ceramics 7.3.2 Phase transitions in poled PLZT 3/90/10 ceramic 7.4 Raman Spectroscopy Study 7.5 Revised phase diagram of the PLZT x/90/10 solid solution 7.6 Conclusions General Conclusions Appendix: Publication and conference participation Lists 7 Abstract The thesis entitled “Perovskite systems with ferroelectric/antiferroelectric character” is focus ed on the study of antiferroelectric properties in perovskite ceramic materials. The large majority of studies related to antiferroelectric phenomena have focus ed on systems with FE- AFE boundaries owing to room temperature electric field AFE- to-FE induced transition and their relevance to applications. Materials investigated in this study comprises solid solutions for which the room-temperature state is known beforehand to be either FE or antiferroelectric (AFE). The compositions were chosen from the phase diagram of lanthanum doped lead- zirconate titanate (PLZT x/90/10) system, with Lanthanum composition across the border between FE and AFE states, having the formula: Pb1-xLax(Zr 0.9Ti0.1)1-x/4 x/4O3 (with La3+ content x=0.020, 0.030, 0.031, 0.032, 0.033, 0.035, 0.038 and 0.040) . The local and macroscopic structure and properties of PLZT x/90/10 ceramics with selected compositions across the FE-AFE phase boundary were studied in en effort of understanding the effect of compositional fluctuation on the stability of AFE and FE phases. The room temperature structure and the stability region of FE and AFE phases has been established from high resolution synchrotron X-ray powder diffraction, Raman and transmission electron microscopy measurements. According to these studies, the FE- to-AFE border is very sensitive to the La addition and shows a broad compositional dependence structure and properties across it. From the detailed structural analysis , it was found that the room temperature state of PLZT compositions with x<0.020 is rhombohedral R3c, while ones with x>0.033 are AFE with orthorhombic Pbam structure. In-between, the compositions with 0.025≤x ≤0.033 show a coexistence of the AFE/FE phase. However, the results of Raman investigations suggest ed that the ground state of these compositions at low er temperature has a lower symmetry and this structure may be locally present even at room temperature in a small amounts and probably located in nanoregions . The use of complementary dielectric, in situ temperature XRD and Raman techniques to investigate the temperature phases stability of PLZT ceramics allow ed us to establish more precisely the phase diagram of PLZT x/90/10 in the compositional range around the AFE/FE phase boundary (0 1 and AFE phase is stabilized for <1 . Figure 1.2 Network of corner -linked octahedral where the A sites (yellow ball) is inside an octahedral cage of Oxygens The large m ajority of FE materials like Barium Titanate (BaTiO 3), Lead Titanate (PbTiO 3), Lead Zirconate Titanate (PZT), Lead Lanthanum Zirconate Titanate (PLZT) have the perovskite type structure. Barium titanate BaTiO 3 is the first laboratory prepared FE and it is the prototype perovskite FE structure. Barium ions (A ions), occupy the corner sites, titanate ions (B ions) are located in the centers of the c ube (the oxygen octahedral) and oxygen anions are on the face - 18 centers. The BO 6 building block led to the finding of a series of FE crystals with similar structure, such as KNbO 3, KTaO 3, LiNbO 3, LiTaO 3, PbZrO 3 and PbTiO 3. [43]. Pb(Zr,Ti)O 3 is the solid solution between AFE Lead Zirconate (PbZrO 3) and FE Lead Titanate (PbTiO 3). Lead zirconate - titanates and have the formula Pb(Zr 1-xTix)O3 (PZT). Crystalline structure and physical properties of PZT strongly depend on (x) lead titanate content. Titanium -rich compositions transform into tetragonal perovskite structures whereas the phase transformation in Zr -rich compositions is more complex. At low Ti content (x > 95%) and at room temperature the structure is orthorhombic and for x < 90% the structure is rhombohedral [ 45]. However, scientists started to pay more attention to the MPB in simple -structured pure FE compounds such as FE oxides. The morphotropic phase boundaries are present in many solid solutions and of particular interest are compositions PZT, relaxor FEs such as Pb(Zn 1/3Nb2/3)O3 – PT and Lead Magnesium niobate -lead titanate (1 - x)PbMg 1/3Nb2/3O3-xPbTiO 3). For example, PZT is a perov skite FE with a MPB between the tetragonal and rhombohedral FE phases in the temperature -composition phase diagram. High resolution x -ray powder diffraction measurements on homogeneous PZT sample s have shown that there is a monoclinic phase exists between the well -known tetragonal and rhombohedral phases [46-48]. AFEs have centrosymmetric structures which can be described by unit cells with oppositely directed dipoles generated by ionic displacements from a higher -symmetry reference structure. AFE class of materials include certain niobate perovskites, vanadates and complex perovskite oxides which satisfies the structural and energetic criteria for antiferroelectricity. Among them, PbZrO 3 is the most studied AFE oxide. It has Pbam orthorhombic distorted pero vskite structure [49]. Isostructural with PbZrO 3, AFE PbHfO 3 was discovered soon after the identification of antiferroelectricity in PbZrO 3 [20]. NaNbO 3 and AgNbO 3 are both perovskite AFEs with orthorhombic space group Pbcm with eight formula units per uni t cell [50-51]. Some d ouble perovskites may also show antiferroelectricity [ 20]. 1.4.2 Domain switching: Ferroelectric versus antiferroelectric behaviour a) Domain switching in ferroelectric materials The regions of the crystal with uniformly oriented spontaneous polarization are called FE domains [ 53-54]. In an as -processed ceramic, each grain consist of distinct types of uniform domains oriented in a preferential direction but polarizations of different grains could have different orientations governed by the local crystallographic symmetry. In the absen ce of electric 19 field, domains are oriented with equal probability along one of the crystallographic equivalent directions , resulting in the lack of piezoelectric effect. Thus the grains in a virgin polycrystalline sample are not uniformly polarized and the initial macroscopic polarization is zero ( point A from Fig.1.3 ). When an electric field is applied, domains start to orient along the applied field. This gives rise to an increase in polarization and a nonlinear behaviour describing the hysteresis curve. The study of the hysteresis loops, namely current -electric field (I -E), polarization -electric field (P - E) and strain -electric field (S -E), is one of the most important tools to investigate the behaviour and to assess the properties of FE/ferroelastic/ AFE materials [ 38]. The e lectrical parameters of interest for FEs include : coercive field (E c), spontaneous polarization (P s) and remanent polarization (P r). Spontaneous polarization (P s) is the polarization at maximum saturation field minus the induced contribution, remanent polarization (P r) is the polarization that persists at zero field, and coercive field (E c) is the field required to reverse the remanent polarization back to zero. A typical polarization reversal for FEs is sown in Fig. 1.3 When elec tric field is applie d on a virgin FE ceramic, at small values of the alternating electric field, the polarization increases linearly with the field amplitude. In this region, still no domain switching takes place. As the field is increased, the polarizatio n of domains with unfavourable direction of polarization will start to orient along the directions of the field. The polarization response in this region is high and strongly nonlinear. Further increasing of the fie ld will bring the system into the saturat ion of the polarization (point B). Subsequently, by decreasing the field, a remnant polarization P r remains when the E field is reduced down to zero (Point C). However when the field is applied in the opposite direction the polarization decreases down to z ero at a certain value of electric field named the coercive field Ec (point D). A new alignment of dipoles and saturation (point E) takes place if the field in the negative direction is further applied. When the field strength is then reduced to zero the polarization reverses to complete the loop [36, 38 ]. During the domain switching a macroscopic deformation and unit cell distortion take s place. For example the spontaneous polarization in PbTiO 3 lies along the c -axis of the tetragonal unit cell and the crystal distortion is usually described in terms of the shifts of O and Ti ions re lative to Pb. Under the application of an electric field along the spontaneous polarization direction, the Ti displacement is shifted along the field direction while c -axis of the lattice elongates and the a - axis of the lattice shrinks. This is the reason for the high performance piezoelectric ity in such ceramics [ 55]. It is worth mentioning that the domain switching takes place only on polar materials in the direction of polar axis (easy polarization axis). The polarization (order parameter) is aligned along definite 20 Figure 1.3 Schematic illustrations of the polarization switching: (A -C) the initial poling, (C-E) the electrical reversal, and (F -A) the electrical cycling. Under the application of an electric field, the B cations displacement is shifted along the electric field dir ection, giving rise to the lattice distortion. (The rectangles with blue arrows represent schematically the repartition of the two polarization states in the material ( e.g. in the cermic grains) at different fields . crystallographic directions ( e.g. the [111] directions for the rhombohedral phase). When the number of allowed directions is large, domain switching takes place quasi -continuous and the polarization is enhanced. The tetragonal structure show s a sizeable elongation along [001] and a large spont aneous polarisation in the same direction. There are six equivalent polar axes in the +ECPs -PrPrP E -EC Before poling 𝑬 Under electric field 𝑬 Under electric field After poling A B C D E F 21 tetragonal p hase corresponding to [100], [100], [010], [010], [001], and [001 ] directions of the cubic paraelectric state. A rhombohedral FE structure shows distortion an d polarisation are along [111] directions, giving rise to eight pos sible domain states: [111], [ -1-11], [1 -1-1], [11 -1], [1 -1- 1 ], [ 1 -1-1], [-1-1-1 ], and [ -1-1 1 ]. There are fourteen possible poling directions on monoclinic structure over a very wide t emperature range, which may in part explain the ceramic pi ezoelectric behaviour near the MPB boundary [ 56]. E I Figure 1.4 Current vs. field during domain switching in FEs (“current hysteresis”) Domain switching in FE ceramics is a complex process that is not very well understood. In order to induce piezoelectricity in the FE ceramic, those may be brought into a polar state by applying a strong electric field (10 -100 kV/cm) for a long time, usu ally at elevated temperatures , then cooled down to the room temperature under field . During this process (called poling ), the domains are realigned as close as possible to the field direction , thus making the ceramic become uniaxial piezoelectric. The poli ng process is essential for the technological applications of FE ceramics in sensors, actuators, and transducers that exploit the piezoelectric effect [ 57]. As clearly pointed out by J.F Scott in “Ferroelectric go bananas” [ 58], it is possible to observe a hysteresis even without any ferroelectricity. In order to assign the FE behaviour to a material the current curve versus electric field must be considered in order distinguish FE switching from artefacts. A schematic representation of a typical I(E) curve , representative for a Electric Field (E) Current (I) 22 FE, is shown in Fig.1.4 . The linear field ramp is positive through the ascent and the same, but inverted value through descent of the ac signal. During the FE switching a peak of this current is observed at a certain field , both in the rising part of the positive fields and in the falling part of the negative fields and corresponds to the coercive field. In non -FE materials which show FE-like hysteresis, the leakage current simply increases with the electric field and no FE switching peaks would appear in the transient current response. b) Domain switching in antiferroelectric materials Application of an electric field or pressure to the AFE can stabilize the parallel dipole arrangement and thus , lower its free energy. The dipole switching in AFE s under the application of electric field is a combination of strain behaviour and preferred orientation of AFE domains. A step-by step sequence is shown in Fig. 1.5 . Initially, in the virgin state of the AFE ceramics, the dipoles are rando mly oriented. After exposure to a low electric field value 𝐸<𝐸𝐴𝐹, the AFE domains are preferentially oriented but with its c axes still perpendicular to the field direction. In this stage the unit cell is deformed along the field direction: the size of the c axis is increased along longitudinal direction while unit cell size is decreased along the transverse direction. Once an E field applied is large enough 𝐸=𝐸𝐴𝐹 to induce the AFE -FE transition, the switching from non - polar state to polar state tak es place with a sudden increas e of polarization. At this step the oriented AFE dipoles change to oriented FE dipoles along the electric field direction and the structure of the unit cell is changed to a new FE structure. In comparison to sample exposed to an electric field, the primitive unit cell shows a slight decrease in the longitudinal dimension and a large increa se in the transverse one. If higher field is applied , the domains will continue to switch as close as possible to the field direction and the refore , the oriented FE will become poled , with piezoelectric character . After the field removal , there are two possible situation s, depending how far the system is from the AFE -FE phase boundary: 1) the FE phase return s to the oriented AFE state ( reversib le AFE to FE field induced transition ) or 2) the FE phase remain s in the poled state (irreversible field assisted AFE -to-FE transition ). However in both situations, the ceramic does not return to its virgin state, unless it is heated above its Curie temperature [3,31,59 ]. In conclusion, AFE materials can be divided into two subcategories, according to the stability of the ir induced FE state. If the induced FE state transforms back to the AFE phase after the field removal, it is named reversible AFE. I f the induced FE state persists after the applied field is removed, it is called 23 irreversible AFE. However, the induced FE state is not stable, it can return to AFE state at high temperatures [ 60], by applying a sufficiently high pressure [ 61] or by applyi ng an electric field of reverse polarity [ 62]. Usually the reversibility is recognized from the hysteresis loop shape: a double hysteresis loop is displayed for a reversible AFE -to-FE field induced transition or a single hysteresis loop for the AFE-to-FE field assisted irreve rsible transition. Figure 1.6 shows the typical P-E response for an AFE when the applied electric field strength is high enough to induce the FE state. The polarization shows a linear response (similar to a linear dielectric ) at low ele ctric field s, then the AFE -to-FE transit ion is induced at a critical field value E AF accompanied by a sudden increase in polarization. The as -induced FE phase turn s to the AFE state in the first quadrant during the field reversal at a lower field value, E FA, describing a hysteresis loop. Similarly, AFE -to-FE and Figure 1.5 Schematic diagram of the e volution of AFE -FE phase switching during application and rel easing of adequate electr ic field. Figure adapted from [ 59] FE–to-AFE transitions take place during the third quadrant on reversing electric field and subsequently increasing electric field [63 -67]. Random AFE, virgin stateOriented AFE, EOriented FE, E 1 2 Poled FE, E3 Piezoelectric effect E Poled FEOriented AFE OrReversible AFE to FE field induced transition Irreversible AFE to FE field assisted transition456 24 Figure 1.6 Representative electric -field-induced polarization hysteresis loop of AFE materials. FE-AFEFE-AFE AFE-FEE I AFE-FE Figure 1.7 Current -field dependence of AFEs A representative current –field (I –E) hysteretic curve of AFE ceramics is illustrated in Fig. 1.7. The high increase in current is caused by the switching of the dip lles. As can be seen in Fig. 1.7 four obvious peaks are observed in the I(E) curve during electrical loading. There is a sharp increase in current when the applied field just reaches the threshold values for the phase transition between AFE and FE state and a broad peak durin g decreasing of electric field just below the critical one that induced dipole re-orient ation back to the original anti -parallel state. Similarly, EAF EFAEFA Electric Field (E) Polarisation (P) EAFElectric Field (E) Current (I) 25 two peaks are present in the fourth quadrant during the field reversal with negative amplitude [ 3, 31]. 1.4.3 Tunability a) Tunability in ferroelectrics FEs show electric field tunable dielectric properties which have attracted extensive attention in recent years due to their potential applications for microwave devices such as tunable filters, phased array antennas, delay lines and phase shifters [ 68-69].Tunability, n, is defined as the ratio of its permittivity at zero applied electric field to its permittivity at a specific non -zero field [ 69]: 𝑛(𝐸)=𝜀(0) 𝜖(𝐸), (1.4.2 ) Relative tunability can be defined as 𝑛𝑟(𝐸)=𝜀(0)−𝜀(𝐸) 𝜖(0), (1.4.3 ) A typical “butterfly” permittivity -field dependence specific to FE materials is shown in Fig. 14, where the maxim a in permittivity appear at the coercive field ±E c. Figure 1.8 Typical field dependence of the dielectric permittivity of a tunable ferroelectric ceramic 26 b) Tunability in antiferroelectrics Similar with FEs, the AFE materials show field -dependen t dielec tric permittivity. However the (E) loop is quite different. As depicted in Fig. 1.9, permittivity changed in a similar way as for FEs during the half quarter of the electric cycle: as the DC electric -field increasing from 0 to maximum values, dielectric constant of AFEs increases sharply, and rea ches the first peak value at A and then decreases gradually. When electric field decrease s from its maximum value down to 0, a new dielectric maximum occurs at a lower field value. The f irst and second peak are related to the AFE -to-FE phase switching and FE -to-AFE respectively. A similar behaviour is noticed during the application of negative electric field. Therefore, four peaks are observed in the field dependence of permittivity during the field cycling [70] (two “butterfl y” loops ). There is another situation which is worth to be mentioned in this section. Fig. 1.10 show the DC field dependence of permittivity and its corresponding P(E) loop in of an AFE materials with irreversible AFE to FE field induced transition field-assisted irreversible AFE -to-FE phase transition ( during reversal of electric field the FE phase remain polarized and it does not return to AFE state) . As shown in Fig. 1.10 during i ncreasin g the bias electric field from 0, the dielect ric constant and the dielectr ic loss increase. Wh en the electric field reached the value of the AFE -to- FE switching field (about 30 kV/cm ), the dielectric constant drops while the FE polarization increases. When the same is further cycled the abrupt changes in dielectric cons tant and dielectric loss cannot be observed again and the materials respond to the electric field like a FE [71]. 27 Figure 1.9 DC field dependence of pemittivity in AFEs with reverssibile AFE to FE field induced transition [3] Figure 1.10 a) DC field dependence of permittivity and b) its corresponding P(E) loop in AFEs with irreversible AFE to FE field induced transition [71] 1.4.4 Energy storage properties Parallel plate AFE capacitors are the key components for high energy storage density , which recently becom e increasingly important especially in pulsed power circuit applications, such as hybrid electric vehicles (HEVs), medical devices, spacecraft, and electrical weap on systems . In b) a) 28 particular, AFEs offer enhanced energy density, lower hysteresis losses, good b reakdown strength and high dielectric permittivity [ 72]. Losed Energy (Jloss) Stored Energy(J)P (C/cm2) E (kV/cm)a) b) Lost energy (Jloss) Stored Energy (J) E (kV/cm)P (C/cm 2) Figure. 1.1 1 P(E) hysteresis loop and energy storage characteristics for a) FE and b) AFE ceramic s Fig. 1.1 1 shows the hysteresis loop and effective energy density of FEs and AFEs, respectively . The energy storage density, energy storage loss and energy storage efficiency can be determined from P(E) characteristics. The energy values can be calculated according to the definition of recoverable energy density of the FE capacitor, ( i.e. at the withdrawal of applied field) [ 71].: 𝐽=∫𝐸𝑚𝑎𝑥𝑑𝑃𝑃𝑚𝑎𝑥 𝑃𝑟 (Emax≡ applied electric field and P≡ polariza tion), (1.4.5) And the energy loss density: 𝐽𝑙𝑜𝑠𝑠=∫𝐸𝑚𝑎𝑥𝑑𝑃𝑃𝑚𝑎𝑥 0 and energy efficiency η=J/(J+J loss) (1.4.6) 1.5 Landau -Ginzburg -Devonshire Theory o f Ferroelectric ity Ginzburg is the first who developed a phenomenological theory for ferroelectricity [ 73]. He used the Landau theory of second -order phase transitions [ 74-75] for his formulations and applied a similar treatment as Devonshire [ 76-78]. The main variable of FE state at equilibrium are: the temperature (T), the spontaneous el ectric polarization (P), the electric field (E), the strain and the stress. The main FE characteristics which are going to be discussed in the following are polarization reversal (switching), and disappearance of spontaneous polarization above the FE phase transition temperature T c. The free energy F of a n homogeneous FE crystal can be generally expressed as a function of ten variables (three components of polarization, six components of the stress tensor, and 29 temperature). If it is assumed that the easy p olarization axis has the same direction as the applied electric field and the order parameter in the Landau theory has the same transformation properties as the polarization vector P, the Gibbs free energy density G in the Landau -Ginzburg polynomial expans ion in the uniaxial case can be expressed as: 𝐺=𝐹−𝐸𝑃=𝐹0+𝛼 2𝑃2+𝛽 4𝑃4+𝛾 6𝑃6−𝐸𝑃 , (1.5.1) where 𝐹0 is the free energy density of the paraelectric phase when 𝐸=0 ,𝛼 coefficient is pressure and temperature dependent while β, and γ are temperature -independent or less temperature , but pressure dependent. The equilibrium configuration is determined by finding the minima of F : 𝜕𝐹 𝜕𝑃=0, (1.5.2) and 𝜕2𝐹 𝜕𝑃2>0. (1.5.3) Thus the expression of electric field as a function of polarisation can be expressed as: 𝐸=𝛼𝑃+𝛽𝑃3+𝛾𝑃4, (1.5.4) The linear dielectric susceptibility above the transition temperature can be obtained by differentiating this equation with respect to P and then setting P = 0: 𝜕2𝐹 𝜕𝑃2=1 𝜒=𝑇−𝑇0 𝜀0𝐶, (1.5.5) 𝜒=𝑃 𝜀0𝐸=1 𝛼. (1.5.6) Where 𝛼 expression around the Curie point is 𝛼=1 𝜒=1 𝜖0𝐶(𝑇−𝑇0), 𝛽<0,𝛾=0, (1.5.7) The equation (1. 5.7) represent s the Landau -Ginzburg Devonsire of second -order ferro - electric phase transitions. The transition from paraelectric to FE phase takes place at temperature T0. The spontaneous polarization in the FE phase is given by: 𝑃𝑠=𝑃(𝐸=0)=√𝑇0−𝑇 𝛾𝐶 𝑇<𝑇0., (1.5.8) 30 Where , 𝑇>0 is the Curie temperature, and 𝐶>0 is the Curie Weiss constant. The Curie Weiss temperature T 0, where α changes sign , is close but not exactly coincident with the Curie temperature T C. The equation (1.3. 7) captures the Curie -Weiss behaviour observed in t he majority of the dielectrics in the paraelectric state, for 𝑇>𝑇0. The first and second derivatives of equation 1.5.4 are continuous and discontinuous, respectively , with temperature. In both situations, the difference of entropy (the first derivative o f free energy) is zero at T 0. These last two expressions suggest that Ps vanishes at T = T 0 and consequently , the dielectric susceptibility 𝜒 diverges. In the case of P S ≠ 0, β>0 corresponds to a second -order phase transition, while β< 0 corresponds to a first-order phase transition. For a second -order transition which occurs at T = T 0, the free energy will evolve continuously when decreasing temperature from the first curve with a single minimum corresponding to P = 0 (paraelectric sta te) in Fig 1 .12 to the F(P) dependence with two symmetric minima at finite polarizations P = ±P 0, corresponding to the FE state. The signs ± indicate that the polarization can be displayed in both direction s along the symmetry axis and corresponds to the two energetically equivalent polar states at zero field. In the paraelectric state, P s = 0 which means that α is a positive value when the system is in a stable paraelectric state. Figure 1.12: Second order phase transition. (a) Free energy (F) as a function of the polarization (P) at T > T 0, T = T 0, and T < T 0; (b) Spontaneous polarization P 0 (T) as a function of temperature (c) The susceptibility χ and its inverse, at the equilibrium condition P0(T) The variation of the F, P and 𝜒 associated with this second order phase transition as described by Landau -Devonshire theory are displayed in Fig. 1 .12. The free energy as a function of P for three typical temperatures is plotted in Fig. 1 .12 (a). The double wells in the energy P T TP0 T0 T0a) b) c)T0 T0 T0 PF 31 diagram at temperatures lower than T 0 (=T C) correspond to two stable FE states. As temperature goes below the critical temperature, the extreme at E = 0 becomes a local unstable maximum, while two minima emerge at ±E 0. The order parameter grows continuously from zero as the temperature is decreased from the critical point. It can be observed in Fig. 1 .12 (c) that as the temperature is raised, the bulk polarisation decreases continuously and vanishes at a temperature T0. If the Landau parameters meet the condition: β<0 and γ>0, the phase transition is of first order. 𝛼=1 𝜀0𝐶(𝑇−𝑇0),𝛽<0, 𝛾>0 . (1.5.9) The procedure for finding the spontaneous polarization and the linear dielectric susceptibility is conceptually the same as before, but now the quartic and sixth order terms cannot be neglect ed. Figure 1.13: First order phase transition. a) Free energy as a function of the polarization at T > TC, T = TC, and T = T 0 < TC; (b) Spontaneous polarization P as a function of temperature (c) The susceptibility χ and its inverse PP T TF TCT0 TC T0P0 a)b) c)T0 T0 TC 32 The temperature at which the transition takes place is, by definition, the Curie temperature TC which exceeds T0 . For temperature lower than this, the polarization at equilibrium is: 𝑃𝑠=(𝑃(𝐸=0))=±√1 2𝛾(⌈𝛽⌉+√𝛽2−4 𝐶(𝑇−𝑇0)𝛾) , 𝑇<𝑇0, (1.5.1 0) The contribution of the polarization to the dielectric polarizability in the paraelectric and FE phases is calculated from th e electric field expression (1.5 .4) as follows: 1 𝜒=[𝜕2𝐹 𝜕𝑃2]=𝑇−𝑇0 𝐶, 𝑇>𝑇𝐶, (1.5.1 1) 1 𝜒=[𝜕2𝐹 𝜕𝑃2]=8(𝑇−𝑇𝐶) 𝐶+3𝛽2 4𝛾, 𝑇<𝑇𝐶, (1.5.1 2) In this case the dielectric stiffness (inverse of the linear susceptibility) that does not vanish at T0, shows a finite jump in both the susceptibility and the spontaneous polarization at the transition. The spontaneous polarization changes suddenly by the magnitude (for E= 0 ). At any temperature between TC and T0 the unpolarized phase exists as a local min imum of the free energy. At T = Tc the three minima are energetically degenerate. As a consequence, the system’s behaviour at T = Tc will depend on whether is approaching Tc from lower or higher temperatures. More clearly, the system will be in one of the two finite polarization minima if it is heated from an initial low temperature Ti < TC, whereas it will be in a paraelectric state ( P = 0) if the initial temperature is high (Ti > TC). Fig 1.13. shows the dependences of free energy F on the parameters P and χ and χ-1 on temperature for the first order phase transition. The FE state is stable for T < T C, corresponding to the double wells. When temperature increases to T 1 (T1 > T C), the paraelectric state becomes mo re stable and the phase transition occurs at T = T C. At this temperature, the polarized state shows the same energy as the paraelectric state (the middle well). However, the reverse phase transition from paraelectric to FE phase takes place at a different te mperature T2 (T2 < T C). Therefore, thermal hysteresis occurs between heating and cooling for the first order phase transitions. This is a the key characteristics of this type of phase transitions. During heating, polarization decreases abruptly and discontinuously to zero at TC (Figure 1.13(b)) and the reciprocal of susceptibility χ-1 changes abruptly at TC. 33 The electric field can be calculated from the free energy with the relation : 𝐸=𝜕𝐺 𝜕𝑃=𝛼𝑃+𝛽𝑃3+𝛾𝑃5, (1.5.1 3) Figure 1.14. Schematic hysteresis in an idealized ferroelectric The equation (1.5.16 ) indicates a nonlinear dependence of the polarization P on electric field E. On the graphical representation it leads to a FE hysteresis loop. An ideal hysteresis loop is shown in Fig. 1.14. In a FE state (correspond to the condition 𝑇<𝑇𝐶) there are (at least) two minima of the free energy, corresponding to spontaneous polarizations of different spatial orientations. There is a barrier between these minima and a certain value of electric field is requested to switch the polarization. The Landau -Devonshire theory described previously predicts a FE hysteresis as shown schematically in Figure 1.14, for an ideal situation where all the dipoles have to be overturned together to switch from one polarization orientation to the other. 1.6 Landau Theory of Antiferroelectrics P E 34 Similarly to FEs, an AFE material transforms to a paraelectric phase at the Curie temperature TC. This critical temperature is related to a structural phase transition between two non-polar phases , with a dielectric anomaly at the high temperature side. In 1951, the macroscopic phenomenological theory of AFEs was firstly proposed by C. Kittel [ 14], together with the description of some intrinsic characteristics. Appling a similar formalism, as Landau for FEs, it was found that the susceptibility of AFEs will be continuous and nearly constant vs. temperature until the Curie point is reached , where a small discontinuity in the temperature coefficient may be present. The magnitude of the anomaly depends of the nature of the transition , but is much lower than for FEs . The AFE may obey the Curie Weiss law , but this is not an indicative of AFE cha racter [14]. This theory is limited and quite intriguing, because it does not explain some important phenomena observed in AFE materials , as for example the mechanism s to drive the relative spatial positions of the two sub lattices and the cell doubling du ring AFE phase transitions. Hatt et al. suggested a Landau -Ginzburg model for AFE phase transitions based on microscopic symmetry [79] Recently, Pierre Tolédano et al . extended Landau theoretical model to AFEs and demonstrated that there are symmetry criteria defining AFE transitions [ 81]. However the Kittel theory remain s the basis for describing antiferroelectricity. According with his model the AFEs have two equivalent lattices, Pa and Pb, which could be polarized independently and have an interaction between them. The sublattices polarizations Pa and Pb in the Kittel model can be regarded in terms of the molecular dipole moments oriented antiparallel to their adjacent dipoles, which results in a zero net polarization. In this case, in t he AFE state the formula units containing a positive dipole -moment component in the a direction form one sublattice with polarization Pa, and the formula units containing a negative dipole -moment component in the a direction form the other sublattice with polarization Pb. The free energy of an AFE system can be written as: ∆𝐺=𝑓(𝑃𝑎2+𝑃𝑏2)+𝑔𝑃𝑎𝑃𝑏+ℎ(𝑃𝑎4+𝑃𝑏4), (1.6.1) where f, g, h are phenomenological coefficients. He suggested a second -order transition from the paraelectric state to the AFE state, by truncating the free energy at the fourth order. The difference ∆𝐺 between AFE and FE state is small and therefore , an external electric field can induce a phase transition from AFE -to-FE state. When the field strength becomes sufficiently large, the polarization in the direction opposite to the field abruptly switches its orientation to become parallel to the field, resulting in a FE state. If 𝑔>0, the transition will favour the antiparallel orientation of 𝑃𝑎 and 𝑃𝑏, making the low -temperatu re phase to be AFE. On the other 35 hand, if 𝑔<0, the transition will favour the parallel orientation of 𝑃𝑎 and 𝑃𝑏, and the transition will lead to a FE state. However, Kittel model [ 14] is not realistic because 𝑃𝑎 and 𝑃𝑏 are assumed at the same location in space (or anywhere in the space) and it does not consider the symmetry change during the AFE-to-FE field induced transition. Tolédano et al. [81] developed a Landau model able to account more aspects of the AFE states, including local dipole orientation and crystalline structure changing , without the need of assum ing sublattices. In the following , the main characteristics of AFEs described by th is approach are summarised . The emergence of polar sites constitute s a necessary condition but it is not sufficient to explain the existence of a PE–AFE transition. In order to preserve the site symmetries at the macroscopic level and to allow a subsequent stabilization of a FE phase under applied electric field two condi tions are required: Condition 1 : At the PE –AFE transition, a set of crystallographic sites undergo a symmetry lowering that results in the emergence of polar sites and give rise to a local polarization . Condition 2 : The AFE space -group has a symmorphic polar subgroup coinciding with the local symmetry of emerging polar sites . Tolédano et al. suggested a definition of PA –AFE transitions, which does not imply the stabilization of a FE phase under applied field or a double hysteresis loop, and extends the current characteristics of AFEs to the larger class of materials in which a polar field -induced phase emerges from a non -polar phase: PE–AFE transitions are structural transitions between non -polar phases where the symmetry of crystallographic polar sites emerging at the local scale coincides with the symmetry of a polar symmorphic subgroup of the AFE space -group, allowing the emergence of an electric field induced polar phase at the macroscopic scale. The dielectric properties of AFE transitions are derived from the Landau potential: 𝜙(𝜂,𝑃,𝑇)=𝜙0(𝑇)+𝛼 2𝜂2+𝛽 4𝜂4+𝛾 6𝜂6+𝑃2 2𝜒0+𝛿 2𝜂2𝑃2−𝐸𝑃 (1.6.2) Where 𝛼=𝑎(𝑇−𝑇𝐶) (1.6.3) 36 The model involves one symmetry breaking parameter 𝜂 and the polarization P is a field - induced order -parameter. For "proper" AFE transitions η can be correlated with local dipole distribution either in a continuous formalism, as a polarization wave amplitude, or a combination of local dipoles belonging to antiparallel arrays of emerging polar sites. For "improper" AFE transitions η represents a structural (displacive or ordering) mechanism which typifies the lowering of symmetry at the transition, the emergence of an antiparallel polarization wave amplitude being an induced secondary effect of the preceding primary mechanism. The equations of st ate can be obtained if φ is minimized with respect to η and P: 𝜂(𝛼+𝛽𝜂2+𝛾𝜂4+𝛿𝑃2)=0 (1.6.4) 𝑃=(1+𝛿𝜒0𝜂2)=𝜀0𝜒0𝐸 (1.6.5) When 𝐸=0, the last two equations yield two possible stable phases: the P E phase (𝜂= 0,𝑃=0) and the AFE phase ( 𝜂≠0,𝑃=0). If 𝐸≠0 two phases are stabilized: a FE phase ( 𝜂= 0,𝑃≠0) and a phase in which the AFE has a non -zero total polarization ( 𝑃≠0) (weak FE of ferrielectric order may be present). Temperature -field T -E phase diagram is shown in Fig. 1.15. For 𝑇≥𝑇𝐶 the paraelectric phase is stable at 𝐸=0, and transform into a FE phase for 𝐸≠0. For 𝑇𝐶>𝑇>𝑇0 , the AFE phase [𝜂=±(−𝛼 𝛽)1/2 ,𝑃=0] which is stable at 𝐸=0, transform in a ferrielectric ( FI) or FE phase when 𝐸≠0. The equilibrium value of 𝜂 and 𝑃 of the induced FI or are given by 𝜂=±[(−𝛼−𝛿𝑃2)/𝛽]1/2 (1.6.6) where 𝑃 is a real root of the Cardan equation: 𝛿2 𝛽𝑃3−(1 𝜒0−𝛼𝛿 𝛽)𝑃+𝐸=0 (1.6.7) With the increasing of electric field the FI phase transforms across the second -order transition curve into a FE phase: 𝐸=±1 𝜒0(−𝛼 𝛿)1/2 (1.6.8) For 𝑇<𝑇0, the transition of the FI into the FE phase is of first order and it cross the region of the coexistence of the FI and FE phase. As the electric field increases, the FI phase tr ansforms discontinuously into the FE phase and the limit of stability of the FI (in the figure indicated with 𝐸𝑐1) correspond s to: 37 3𝛿3 𝛽𝑃4+𝛿𝑃2(4𝛿𝛼 𝛽−1 𝜒0)+𝛼(𝛿𝛼 𝛽−1 𝜒0)=0 (1.6.9) When the electric field decreases, the FE phase show s a limit of stability on the 𝐸𝑐2 curve. The meeting point of 𝐸𝑐1 and 𝐸𝑐2 provides the values of 𝑇0 and 𝐸0. Figure 1.15. Theoretical temperature –electric field (T –E) phase diagram associated with the free -energy given by Eq. (1. 6.2) for 𝛽 > 0 and 𝛼 >0. Hatched and hatched -dotted curves represent, respectively, second -order transition and limit of stability curves. The thermodynamic paths for T > T C, TC > T > T 0 and T < T 0 are described in the text. The dielectric susceptibility at the P E-AFE transition can be deduced from following equation . For a second order transition (𝛽>0) and below 𝑇𝐶 : 𝜒(𝑇)=𝜒0 1+𝛿𝛼𝜒 0𝑇𝑐−𝑇 𝛽 (1.6.10) The dependence of susceptibility on temperature depends on the sign of the 𝛿 parameter (Fig. 1.16 a)). Fig. 1.16 b)) shows the temperature dependence of a first order transition ( 𝛽<0) which occur when 𝑇1>𝑇𝐶. The dielectric permittivity upward (𝛿<0) or downward (𝛿>0) showing 38 discontinuity which reflect the attractive coupling between η and P requir ed for compensating the repulsive interactions between parallel dipoles. Figure 1.16. Temperature dependence of the dielectric sus ceptibility as given by Eq. (1.6.10 ) across a second -order (a) and first -order (b) transition The macroscopic hysteresis describes the dipole moments associated the AFE state. Antiparallel dipole ordering is energetically more stable than the parallel dipole ordering. However, if an external electric field is applied to domains, parallel to the Pa direction, this field will interact with the dipoles, causing Pa to increase and Pb to decrease in magnitude. In this way the parallel dipole arrangement is stabilized and thus , its free energy is lowered. At sufficiently large field strength EAF, the dipoles with Pa antiparallel to the external field will switch , so that Pa becomes parallel a nd the phase transition from AFE to FE phase is induced. This results in a state in which the dipole components are parallel along the a direction and antipa rallel along the b direction. It is a high field FE state with the polarization along a and with a cell size doubled that of the Kittel model [ 80]. When the field decreases to a critical value, E AF, the minimum energy of FE ordering restores back to the in itial magnitude and as a consequence the material recover its AFE phase. 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Guennou , Theory of antiferroelectric phase transitions , Phys. Rev. B 94, 014107 (2016) . 46 CHAPTER 2 Short literature review on PZT and La doped PZT systems 2.1 Introduction Lead zirconate-titanate Pb(Zr 1-xTix)O3 (PZT) is a solid solution between AFE Lead Zirconate (PbZrO 3) and FE Lead Titanate (PbTiO 3). PZT is the most widely used, most performing and cost effective available piezoelectric material. The studies on this systems start ed in 1952 when its FE behavior was reported [1]. Since then intensive studies regarding the comprehension of its properties, improvement of the processing technology and enlargement of fields of application were done [2 -16]. Due to their excellent dielectric, pyroelectric, piezoelectric, and electro optic properties, PZT-based materials have a variety of applications in high energy capacitors, nonvolatile memories, ultrasonic sensors, infrared detectors, electrooptic devices, and step -down multilayer piezoelectric transformers for AC -DC converter appl ications [3 -16]. Despite of 62 years of intensive studies on this system , it still presents a high interest in current scientific research. The main issues are related to the PbZr 1−x TixO3 in the Zr -rich zone , due to the FE -AFE transition and to the chemic al compositions close to the Zr/Ti ratio 48/52 because of the presence of the morphotropic phase boundary (MPB) between the Zr -rich rhombohedral phase and the Ti -rich tetragonal phase [9 -12,16 -18]. The main advantage of PZT is the possibility of controlling the electrophysical properties and temperature dependence of their properties with the aid of dopant additions [4,7-8, 10 -22]. Various modifications have been made with rare earth elements in order to improve the piezoelectric properties [23]. Lanthanum doped PZT (PLZT) materials developed in 1967 by Haertling et al. are one among the most promising compositions, showing very good optical and electrooptical properties [24 -26]. In order to understand and describe PLZT-based materials, PZT system should be first considered, starting from the end members PbZrO 3 and PbTiO 3 to the effect of different dopants and microstructures. 47 2.2 General properties of PZT solid solutions 2.2.1 PbZrO 3 Lead zirconate PbZrO 3, (denoted as PZ) is a prototype AFE material and the important end member of PZT crystalline solution series, with aCurie temperature of 233oC. At higher temperatures, it has the cubic perovskite structure with antipolar sublattices. During cooling a first- order phase transition occurs and the structure changes from cubic m3m to antipolar orthorhombic mmm with an eightfold increase of the number of atoms per unit cell. The phase transition is accompanied by a strong Curie–Weiss like anomaly of the dielectric constant (maximum value or permittivity up to 6,000) above the transition [27,28]. In contrast with FEs (for example, BaTiO 3), whose dielectric anomaly is related to the transition to a non-centrosymmetric low-temperature phase, the responsible for the critical divergence of the dielectric permittivity of AFEs is the softening of the lattice mode [29-30]. At room temperature PZ has an AFE orthorhombic (O) distorted perovskite structure with Pbam symmetry, as confirmed by diffraction experiments and predicted by first principles calculations. The ground Pbam state is close in energy to a rhombohedral (R) R3c phase [27, 28]. This small energy difference is a key signature of an AFE- to-FE field-induced transition [31]. The structure of PZ is very stable and it is very difficult to switch it into a polar phase by means of an electric field. The critical field for PZ at room temperature is greater than its dielectric breakdown strength, and a double hysteresis loop may only be seen at temperatures close but below the Curie point of 233°C where the AFE phase is less stable. The high temperature field-induced AFE- to- FE transition is possible because to the vicinity of the FE rhombohedral intermediate phase which had previously been observed just below Curie temperature [32]. 2.2.2 PbTiO 3 Crystalline structure and physical properties of PZT strongly depend on (x) lead titanate content (PbTiO 3, PT). PT system is a FE system with a strongly distorted tetragonal structure at room temperature. It undergoes a temperature induced phase transition at 490°C to a paraelectric cubic perovskite structure . Donor dopants like Nb+5, Bi+3 and La+3 decrease the tetragonal distortion and allow an easier displacement of domain walls, and thus an increase of polarizability. 48 A high stability of tetragonal phase was found in PT system when reducing grain size in nanoscale region, even don to around 7nm [3]. 2.2.3 Phase diagram of the PZT solid solution Lead titanate (PbTiO 3) and lead zirconate (PbZrO 3) form PZT solid solutions (PZT) over the whole composition range. PZT has also a perovskite structure ABO 3 with titanium and zirconium atoms occupying randomly B-sites, and lead atoms situated at the corners of the unit cell (A-sites) with oxygen atoms located at the surface centres. Both the lead and oxygen ions have radii of about 0.14 nm. Together they form a face-centred-cubic array, having a lattice parameter of nearly 0.4 nm, where octahedrally co-ordinated titanium or zirconium ions are located at the centre of the unit cell. This structure allows the existence of a large number and variety of stoichiometric compounds. Depending on concentration and different physical conditions it can give rise to changes in the crystal symmetry allowing the existence of tetragonal, orthorhombic, rhombohedral or monoclinic symmetry with FE or AFE behavior [34-38]. Temperature, pressure, and electric field may all induce structural, solid-solid phase transformations in PZT materials. By changing the Zr/Ti ratio materials with extremely different properties have been developed. Fig. 2.1 shows the phase diagram of PZT in the temperature range between 0°C and 500°C. For temperatures below the Curie temperature a spontaneous polarization driven by FE phases with tetragonal, rhombohedral or monoclinic structures as indicated in the phase diagram, are stable. When x is in the range between 0.48 and 1.0, the symmetry of Pb(Zr 1−xTix)O3 is orthorhombic AFE (Ao) up to T C, when it changes in tetragonal AFE A T. With the addition of titanium, the structure becomes rhombohedral FE (FR). This rhombohedral phase is divided into another two phases, i.e. a high temperature FR(HT) with rhombohedral symmetry space group R3m and a low temperature FR(LT) phase with a R3c tilted structure. The antiferrodistortive (or tilt) transition from the FR(LT) to the FR(HT) is accompanied with the loss of oxygen octahedron tilt angle and the corresponding superstructure. Composition s with a conc entration higher than x=48%, have tetragonal symmetry with P4mm space group and FE character (FT) , which turns into a cubic paraelectric one above the Curie temperature TC [1, 33 – 38]. High performance s of the functional properties were reported in PZT composition s near the morphotropic phase boundary (MPB). 49 Fig. 2.1 Phase diagram of PZT according to Jaffe et al. [1, 33] with the newly discovered monoclinic phase according to Noheda et al. [34-38]. P C means the paraelectric cubic phase, F T the FE tetragonal phase, F R the FE rhombohedral phase, F M the monoclinic phase, A T the AFE tetragonal phase, and A O an orthorhombic AFE phase The term morphotropic-nomenclature introduced by Goldschmidt is used to denote an abrupt structur al change in a solid solution with variation in composition [39 ]. The morphotropic phase boundary in PZT separates the tetragonal and rhombohedral phases and is nearly independent of temperature (Fig. 2.1) . It was observed experimentally that for the composition in the range of morphotrophic phase boundary (MPB) all the material properties and in particular, the piezoelectric constants are higher by comparison with all the other compositions. Six different orientations of the polarization axis are possible with a tetragonal structure and eight with a rhombohedric structure, the coexistence of the two phases was assumed at the MPB. This was the largely accepted view of the PZT phase diagram that persisted until the late 1990s, when new data revealed the existence of low symmetry phases which principally occur close to the AFE-FE and morphotropic phase boundaries. Thus, it has been demonstrated that the origin of the maximum piezoelectric performances at MPB is due to the presence of a low symmetry (monoclinic phase) , with twenty one possible directions of reorientation of the polarization axis. Intensive studies were Tc 50 performed on various composition near MPB in order to elucidate the scattered properties reported for similar compositions, . Cm space group was assigned by Noheda et al. for low temperature monoclinic phase [34 -38]. but Raghini et al. observed superlattice reflections by electron- diffraction studies, in the low temperature phase ,due to antiphase rotation of the oxygen octahedral, which cannot be accounted for in terms of the Cm space group [40 ]. D. M. Hatch et al. assigned a Cc space-group for the low temperature monoclinic phase [41 ], while D. Woodward et al. proposed intermediate phases close to the Zr-rich side compositions of PZT solid solutions : Pm (a subgroup of R3m and Pbam), and Pc (a subgroup of Pm, Pbam and R3c) [42 ]. Recently Glazer et al. have postulated a new diagram for PZT system [43 ], which was strongly contradicted by Pandey et al. , who demonstrated that Cc space group model represents the true structure of the ground-state phase while the high-temperature FE phase has got the Cm space group and not the R3m or R3m + Cm [44]. 2.3 Properties of PLZT solid solutions Lanthanum (La3+) is one of the most commonly used dopants which replaces Pb in the A site of the perovskite structure of PZT. Lead lanthanum zirconate titanate (Pb 1-xLax)(Zr yTiz)1-x/4O3 (PLZT x/y/z ) solid solutions were obtained from parent lead zirconate titanate PZT compounds by substitution of La for Pb. The PLZT solid solution was initially developed by Haertling, as a promising new material [ 25-26]. Indeed, PLZT system embraces all the compositional aspects and functional properties as dielectric, piezoelectric, pyroelectric, ferroelectric, and electrooptic (also being transparent, if properly densified) . Depending on the chemical composition, various FE, AFE or PE phases with slightly different dielectric properties and crystal structures may be formed. Haertling et al. [25,26] performed a great research in order to understand the structural properties of th is material. They produced and characterized hundreds of hot pressed specimens and defined the phase diagram of the PLZT system. The main observed effects are an increase of the mobility of the domain wall boundaries, which generates a soft behavior of the material, and a progressive reduction of the perovskite cell distortion, which brings to a reduction of the Curie temperature , when La content increases. Increasing the amount of the Lanthanum concentration in the PZT system , the Curie temperature decreases below room temperature and the behavior of the material goes through a FE, superparaelectric, electrostrictive and PE behavior . PLZT ceramics exhibit several important advantages, such as fast response speed, good temperature stability, high optical damage threshold, and relatively low electric field, which make them more applicable to the 51 modern laser technology. It has been shown that PLZT has a dispersive behavior for both rhombohedral and tetragonal phases, as well as an increase of the stability range of the AFE orthorhombic phase in the Zr-rich side of the phase diagram [24-26 , 45]. Figure 2.2 shows the PLZT phase diagram Several areas on the diagram are color coded for easy identification: (1) the FE tetragonal and rhombohedral phases are shown in orange, (2) the orthorhombic AFE phase in purple, (3) the cubic (nonferroelectric) PE phases in white, (4) the morphotropic phase boundary (MPB) in magenta, (5) the pyroelectric application areas near PbTiO 3 in blue, (6) the economically important MPB compositions that embrace almost all of the transducer applications in green, (7) the compositional area for AFE- to-FE, enforced-phase devices in gray, and (8) specific compositions in these regions in yellow [24]. Fig. 2.2 Phase diagram of PLZT system after Haertling et al. [24] Figure 2.2 shows that the effect of adding lanthanum to the PZT system is to: (1) maintain extensive solid solution throughout the system and (2) decrease the stability of the FE phases in favor of the PE and AFE phases, as indicated by the red line, which shows the reduction of T C with increasing lanthanum addition. At the 65/35 ratio of PZ/PT, a concentration of 9.0% 52 lanthanum (designated as 9/65/35) is sufficient to reduce the temperature of the stable FE polarization to slightly below room temperature, resulting in a material that is nonferroelectric and cubic in its virgin state. The cross-hatched area existing along the FE– PE phase boundary denotes a region of diffuse, metastable relaxor phases that can be electrically turned into a FE phase. Materials within this region exhibit a quadratic strain and electrooptic behavior [24 ].The solubility of lanthanum in the PZT lattice is a function of Zr/Ti ratio. The temperature-compositional phase diagram for the PLZT system from Fig. 2.2 was proposed in 1971 [26 ] and modified in 1974 [46 ]. The compositional dependence of the solubility limit is indicated by the dashed line adjacent to the fixed-phase region (double cross-hatched area). For the two end-member compositions, PZ and PT, these limits are 4 and 32 at.%, respectively. The morphotropic phase boundary that is located at the point that corresponds approximately to the composition Zr/Ti = 53/47 [15] in the Y –T diagram of PbZ 1−yTiyO3 is displaced towards the solid solutions with higher percentages of Zr as the La concentration increases. The morphotropic boundary of PLZT is observed in the vicinity of the 65/35 Zr/Ti composition [24 ] and is a superposition of FE rhombohedral, FE tetragonal and AFE tetragonal phases. In particular, PLZT with x≥6% attracted a high interest for investigati ng the nature of the so -called relaxor ferroelectric (FE) state. In case of trivalent ions (La3+), the excess charge must be compensated by vacancies in the perovskite lattice. Such vacancies can be formed either in the Pb sublattice (A-site vacancies) or/and in the (Ti, Zr) sublattice (B -site vacancies), in that case the compositions have the general formula Pb 1−3x/2 Lax(Zr yTi1−y)O3 and Pb1−xLax(Zr yTi1−y)1−1x/4 O3, respectively [47]. The properties of Pb 1−3x/2 Lax (Zr 1−yTiy)O3 depend on the position of a particular solid solution with respect to the hysteresis FE-AFE region in the ‘T i content –temperature’ phase diagram [47-48]. 2.4 The High Zr content side of PZT and PLZT systems The investigation of AFE materials became one of the most interesting recent topics, because of the field or mechanical stress induced AFE- toFE phase switching. . AFEs have recently received renewed attention also due to their high-electric performance and property of bi-tunability [9-12, 18, 29, 30, 32]. The AFE materials are commonly characterized by either a low dielectric permittivity with a field induced transition to a ferroelectric (FE) state at a critical electric field (named as threshold field), as in the case of the PbZrO 3 system. The AFE-FE phase transition could occur spontaneously due to several factors, such as, a change in the stress configuration promoted by external mechanical driving fields, an increase in the amplitude of the applied ac electric field and/or by thermal changes [18]. It is known that in FEs the domain structure is 53 characterized by a large spontaneous polarization, whereas classical AFEs are frequently centro- symmetric with zero polarization. Due to their relatively poor electrical strength (low breakdown fields), most PZ-based orthorhombic AFE ceramics are broken down before a critical switching field can be applied. As a consequence, the electric-field-induced AFE- to-FE phase transition of those AFE bulk ceramics were observed only at temperatures close to the FE rhombohedral transition. Therefore, in the search for new AFE materials with room temperature electric-field- induced AFE– FE phase transition, the scientific community focused on the development of materials with compositions near the AFE/FE phase boundary. AFE/FE phase boundaries have been reported to exist in PZT and PLZT in the Zr side of the phase diagram [1, 24 -26, 46] . In the PZT solid solution, an AFE/FE phase boundary was reported at Zr:Ti ratio approximately 96/4 (abbreviated PZT 96/4). The coexistence of (and competition between) AFE and FE phases is well-known in this region where the dominant phase is highly influenced by subtle changes in the Zr/Ti ratio and materials processing parameters [49 ]. At this boundary, switching from the AFE to the FE state using an applied field or stress induces large effective strains or charges [18 ]. Examination of the crystallographic features of Pb Zr1−xTixO3 samples with 0.07≤x≤0.20 by transmission electron microscopy have evidenced that the increase in antiferroelectricity in PZT leads to the fascinating sequence of the phase changes from the pure FE phase to the pure AFE phase via the two phase state and two ferrielectric phases [49 ] PZ or PZT compositions are often doped with Sn4+, Nb5+ for Zr4+/Ti4+ or La3+ ions for Pb2+ to stabilize the AFE state relative to the FE state over a relative wide temperature range. In this way the phase transition is broadened and the free energy difference between the AFE and FE phases is reduced [51-54 ]. According to Heartling, the Lanthanum doping of PZT 90/10 with La content from 2 at % to 4 at. % transits the system from FE (rhombohedral R3c) for composition with 2 at. % of La to AFE (Orthorhombic Pbam) for composition with 4 at. % of La (see Fig. 2.3 a)-b)). The phase diagram of La-doped Zr-rich compositions has been less addressed, although such compositions are important for applications in uncooled pyroelectric detectors, digital displacement transducers, high energy density capacitors, etc.[24-26, 46] However, some Pb 1- xLax(Zr 0.9Ti0.1)1-x/4O3 (PLZT x/90/10) compositions have been investigated and important information has been achieved, especially concerning structure –property relationship. The location of the AFE/FE phase boundary is very sensitive to the addition of La and i t is shifted towa rd higher concentrations of Ti in the solid solutions with increase of the La content [55-64]. La doping in PZT, suppress es the stability of the FE state and stabilize s the AFE state [ 55-64]. In general, the AFE phase is favored if tolerance factor t is low. PZ has the lowest t=0.98 in the PZT solid solution , which increases with the addition of PT. La substitution on the A site decreases t, 54 since La3+ (1.36 Å) is smaller than Pb2+ (1.49 Å), thereby inducing an AFE phase in, e.g., PLZT 4/90/10. This effect has been attributed to a disruption of the long-range dipolar interaction. The increase of lanthanum amount diminishes the energy barrier that separates the free energy minima corresponding to FE and AFE states. Fig. 2.3 PLZT phase diagram, after Haertling and Land a) across entire compositon Zr/Ti ratio.[3] Points indicate the compositions fabricated, PLZT 100x/90/10 (100x 0, 2, 4, and 10). AFE antiferroelectric, FERh ferroelectric rhombohedral, FETet ferroelectric tetragonal, SFE slim-loop ferroelectric, and PECubic paraelectric cubic and b) detail of the AFE/FE phase boundary a) b) 55 Therefore, competition between these phases manifests itself greatly. Wonderling et al. [65] reported that that substitution of La for Pb in the orthorhombic PbZrO 3 leads to both a change in atom coordinates and change in unit cell parameters toward the cubic structure. A similar behaviour may be expected for the high Zr side of PLZT. In fact, at high La contents (> 4 at.%), in PLZT 90/10 the long range dipolar disruption is sufficient to additionally destabilize the AFE state resulting in the onset of “relaxor” behavior [63]. In order to give a good comprehension of the structure of PLZT 90/10 the parent PZT 90/10 must be considered. The crystalline structure of PbZr 0.9Ti0.1O3 was studied by C. Michel et al. [66] with modest precision. PZT 90/10 is rhombohedral FE at room temperature, but neutron diffraction revealed superstructure peaks, which were interpreted also as a sign of orthorhombic symmetry [66]. Raman scattering of PZT 90/10 ceramics and single crystals revealed a soft mode of ETO symmetry with minimum frequency at 200 K, below the para -FE phase tr ansition temperature T C, as well as a new mode near 50 cm−1, both related probably with the doubling of the unit cell at 380 K [58, 67 ]. The presence of satellite reflexions correlated with either ferroelectric comme nsurate phase either to AFE incommensura te phase makes this system more interesting: Woodward considers that the appearance of superlattice reflections arises from rotations (tilting) of the octahedra [ 42, 68 ]. The structure of PLZT x/90/10 is even more controversial. As the La is incorporated i n PZT 90/10 system, the symmetry and polar character of PLZT x/90/10 change s: for small amounts of La (x=2) the macroscopic symmetry remains rhombohedral, but for x=4 the tetragonal symmetry or orthorombic is developed [57, 69]. However the reports on the structure of AFE PLZT are quite confusing. In the ea rly phase diagram of Haertling, the antiferroelectric composition were considered to have orthorhombic structure isomorphous with orthorhombic AFE PbZrO 3 (space group P2cb (JCPDS 75 -1607)) at that time [24-26]. Later, in 2002, J. Knudsen et al. report ed a tetragonal symm etry rather than orthorhombic [57, 70]. The confusion arise s due to the presence of superstructure reflexion which were correlated either with Pb displacements coupled t o an a− b− b− (a ≈ b) tilt system either with rotation of octahedra in antispace [ 57,70] . Furthermore, the competition between the FE and AFE interactions bring in to a conversion from the FE incommensurate phase to the AFE incommensurate phase in the FE/AFE region of the phase diagram of some materials [49]. Similary with La-substituted PZT 95/5 [49], an AFE incommensurate phase has been reported to exist between the rhombohedral FE and AFE phases of PLZT 90/10 [50, 57, 70 ]. Knudsen et al. report the appearance of an incommensurate AFE modulation only for PLZT (PLZT 4/90/10) (via electron diffraction) [ 57,70 ] while R . Villaurrutia report ed that the satellite reflections in diffraction patterns obtained for La -doped PZT with a Zr/Ti ratio of 90/10 and La content in the 56 range 2-4 are due to the presence of incommensurate AFE phase [59]. MacLaren studied the AFE structure of PLZT and demonstrated that this incommensurate phase represents a bridging phase between AFE and FE polarisation orderings in the lead zirconate titanate system (scanning transmission electron microscopy studies) [62 ]. However, the local symmetry, seems to be monoclinic. [57, 70] 0 1 2 3 420406080100120140160180200220240260 20406080100120140160180200220240260 (o C) La at. % Tm , dielectric data, Pelaiz et al. Tm, dielectric data, Knudsen et al. R3c-R3m, dielectric data, Pelaiz et al. TFE-AFE, ferroelectric and TEM data, Pelaiz et al. TAFE-FE, ferroelectric and TEM data, Pelaiz et al. Fig 2.4. Phase diagram of PLZT constructed from the data reported by A. Pelaiz – Barranco et al. [71, 72] and Knu dsen et al. [57] E. Breval et al. [55] has determined the phase diagram of PLZT based on dielectric properties at varying fields and temperature and proposed a phase diagram as a function of substituted (La), substituted titanium (Ti), electric field and temperature (Fig. 2 .5). Two orthorhombic PLZT phases (ORI and ORII) and three rhombohedral phases (RI (also called LRI), HRI, and RII) were identified. The basic PLZT structure with small substitutions at zero field and at room temperature is the orthorhombic ORI phase. With increasing temperature, the general phase sequence at zero field is ORI-ORII-LRI-HRI-RII-paraelectric phase. With increasing La substitution, the ORI phase disappears. With increasing field, the general phase sequence at room 57 temperature is ORI-ORII- RI-RII. With increasing Ti substitution, the ORI and ORII phases become less predominant and may disappear [55 ]. Fig. 2.5 Phase diagram of 2/95/5 and 4/5/95 as function of temperature and field [55 ] Recently, Pelaiz-Barranco et al. have shown that PLZT 2/90/10 and PLZT 3/90/10 exhibit different behaviors, from classical ferroelectricity, to the superposition of FE and AFE character on heating [71]. By dielectric, calorimetric, ferroelectric, and structural studies, they concluded that the phase sequence with the increase of the temperature, for the PLZT 2/90/10 composition, 58 can be considered as (i) a coexistence of the AFE and FE states at room temperature with no stabilization of the AFE phase, (ii) an AFE-FE phase transition around 90oC, and (iii) a transition to the paraelectric phase at T m. For the PLZT 3/90/10 composition, these phase sequences can be considered as (i) a coexistence of the AFE and FE states at room temperature with no stabilization of the AFE phase,(ii) a stabilization of the AFE state for temperatures around 90oC, and (iii) the transition to the paraelectric state at T m. However, these results are ambiguous since in an earlier paper [72] they stated that there is a sequence of FE to AFE (~90oC), AFE to FE (~170oC) and FE to PE (~190oC) phase transitions for PLZT 3/90/10 on heating from room temperature up to 250oC. E. Buixaderas studied the lattice dynamics of PLZT x/90/10 with x =0, 2, 4, 10 La at. % of La by far infrared, Raman and terahertz spectroscopies in a broad range of temperature (of 20– 800 K). Infrared active soft phonons anomalies were found near the PE-FE phase transition. Additionally they observed some Raman active phonons anomalies ~200 K below T C, due to another phase transition to a FE state with doubled unit cell. Samples with higher La content 4≤x≤10 display non-classical phonon softening. [58, 67 ] To date, there is no a structural refinement data to confirm any of these hypothesis and only antiphase rotations have been found. PZT 90/10 has been ascribed an a−a−a− octahedral tilt system, according to the Glazer notation, consistent with the rhombohedral FE distortion ( R3c) while the tilt system(s) in La-doped PZT 90/10 remain unknown. However, recently, antiferroelectric PLZT 90/10 was consider ed to have a structure similar with the prototypic orthorhombic PbZrO 3 rather than tetragonal [73-74]. The properties of some PLZT x/90/10 are studied using relative permittivity ( εr) versus temperature plots, and polarization versus field loops. A decrease in T C from 255 to 198°C occurs (cooling curve) as La concentration increases from 0 to 4 but without a significant change in the magnitude of the permittivity maximum (approximately 23,000) was reported. Further decreases and broadening of T C was observed for PLZT 4/90/10 and PLZT 10/90/10 accompanied by a reduction in relative permittivity (approximately 1000). For x=0 and x=2, a classic FE hysteresis loop was obtained whereas for x=4 and x=10 the samples show double hysteresis, characteristic of AFE phase [57, 63, 71 -72]. The potential of some PLZT x/90/10 to show electrocaloric properties has been evaluated by Pelaiz in a recent paper [75]. Their study was focused on composition close but far enough from AFE/FE phase boundary (PLZT 4/90/10, PLZT 5/90/10, PLZT 6/90/10). They observed an increase of the temperature of 0.06 K during the switching between AFE and FE for PLZT 4/90/10 [75]. 59 The AFE phase is capable of storing especially large amounts of electrostatic energy. Since the La is an AFE stabilizer, high energy storage can be found in PLZT x/90/10 compositions. An attempt to determine the energy storage of some PLZT composition was done by P. Wawrzała et al. [63]. For example the composition PLZT 3/90/10 shows energy storage of 0.0759 J/cm3. However the study of the energy storage was inadequately address ed since they studied only one composition across the AFE/FE border. Therefore the full potential applications as energy storage capacitors of PLZT x/90/10 have not been yet fully assessed. Knudsen et al. [57] also studied the properties of PLZT ceramics with a Zr to Ti ratio 90/10 and lanthanum content of 2 and 4 at.%. Their PLZT composition with x=2 at. % showed a P(E) FE behaviour with value of remanent polarization Pr=20 μC/cm2. Despite the crystallographic results which indicated that PLZT 4/90/10 is AFE, the loops did not show any AFE- to-FE induced transition for composition with x = 4 at.% because their maximum applied field for PLZT 4/90/10 of ~2.5 kV/cm was too low for causing the AFE-FE switching. It is worth mentioning that the composition studied by Knudsen and Buixaderas [57 -58, 67 ] were prepared creation of Pb vacancies on the A site of the PLZT perovskite cell [Pb 1-1.5xLax (Zr 0.9Ti0.1)O3], whereas the AFE PLZT samples with vacancies on B site prepared by Haertling had the formula Pb1−xLax(Zr 0.9Ti0.1)1−x/4O3 [24-26, 46 ]. Compared with B site vacancies, the A site vacancies in ABO 3 perovskite cell usually promotes a decrease in the remanent polarization, relaxation behavior, and affects the electrical performance of the ceramics (lower dielectric permittivity and higher dielectric losses) [61 ]. Therefore, some improvements in the FE/AFE properties are expected in ceramics with B site vacancy, with respect to the results published by Kndusen et al. [72]. The PLZT 90/10 materials are of fundamental scientific interest. Even if important information were obtained, the scientific results on these compositions are rather controversial then conclusive. The relationship between the local structure and macroscopic properties has not yet been solved. The crystallographic details of FE and AFE phases have not yet been clarified and the mechanism by which the FE to AFE transformation proceeds as function of composition is totally unknown. Furthermore , it is important to establish more precisely the range of coexistence of AFE/FE phases in order to better control their properties. The theory of AFE is under development and it still is, however, essential to understand the AFE -to-FE field -induced transfo rmation as well as possible from experimental point of view. Far enough from MPB of PLZT and from, the FE/AFE phase boundary of PZT 95/5, the PLZT 90/10 represent s an ideal system for the study of the properties of AFE materials. 60 2.5 Systems investigated for this thesis and motivation for the proposed research topic The fast increasing of modern equipment and electronics motivate the scientific community to be focused on the development of efficient materials for the generation, storage and distribution of electrical power. Suitable dielectric based solid-state capacitors with high enhanced energy storage density will play a key role in revolutionizing modern day electronic and electrical devices. Because of their high power density, high energy density, good discharge efficiency, low coercive field, low remnant polarization, and low dielectric losses, the AFEs become a strong competitor for future ceramic capacitors. These materials have also other important applications like pulse power circuits, displacement transducers, microactuators, pyroelectric security sensors, cooling devices, pulsed power generators [76]. The high recoverable energy density of AFE materials is due to the reversible electric field-induced AFE- to-FE phase transition main characteristic, since the energy stored in a cycle is proportional to the maximum polarization and applied electric field, while the losses are proportional to the remanent polarization (which in AFES is negl igible) . The FE-AFE phase transformation has been used to generate pulsed power by shock loading the material with explosives [ 18, 76 -78]. Their high recoverable energy storage density also allow s the miniaturization , which is highly requested nowadays. Moreover, AFEs possess fast charge – discharge speed and good fatigu e endurance owing to their unique field -induced switching between AFE and FE phase [ 18, 79 ]. Not only reversible but also irreversible field -assisted AFE – FE phase transformations are interesting for applications due to their enhanced toughening against fracture [ 78] and the possibility to be used in bistable optical information storage devices, based on differential light scattering [ 79]. During the AFE -to-FE phase transformation, a large strain response and a high volume expansion is allowed , which makes these materials very attractive for applications in high strain actuators and digital displacement transducers [ 77]. Therefore, the AFE materials possess the best combination of properties for the development of solid -state capacitors for future electronic applications among other areas. AFEperovskites are characterized by an antipolar structure close in energy to a related FE structure to wards which they can be driven by an applied electric field, strain, temperature or compositional changes. The phase bou ndary separating the AFE and FE phases has been the subject of considerable interest and intensive studies not only for their AFE -FE induced transition , but also because the compositions where more phases coexist are expected to present enhanced functional properties [ 57, 67, 80 -81]. Understanding the relationship between structure and prope rties can lead to both physical insight and design of new materia ls with enhanced material 61 properties. It was found also that the ground Pbam state is close in energy to a rhombohedral (R) R3c phase. This small energy difference is a key signature of an AFE material [78 ] and is also consistent with the existence of an intermediate FE R phase, in a very narrow temperature range just below TC, reported in some studies [79 ]; moreover, the transformation to the FE R phase at small Ti addition [82 ] and under strong electric fields point in the same direction. Recent papers clarified the origin of AFE and gave insight into the phase transition mechanism in PbZrO 3 [29- 30]. By lowering the free energy difference in favour of FE phase by adequate substitutional changes (like e. g. Ba, Sn, Ti, Nb) [ 83-84] the critical field necessary to induce the transition from the AFE to the FE phase is also lowered, opening the way to many different applications. Thus PZ-based AFE materi als (like e.g Zr-rich PZT) have been doped with different substituents in order to obtain an AFE/FE phase boundary at room temperature and to decrease the critical field for the FE phase induction [ 18, 85 -87]. Among them, Lanthanum doped lead zi rconate tit anate ceramics PLZT has a phase boundary separating the AFE phase with orthorhombic symmetry and FE tetragonal phase which has been the subject of considerable interest and intensive studies. The first phase diagram of PLZT established by Haertling [ 24-26, 46] shows a fine border between the FE and AFE phases, while in the phase diagram of PZT proposed by Jaffe et al. [1], there is a finite compositional region between FE and AFE phases. The phase diagram of PLZT at room temperature shows regions where FE a nd AFE phases are adjacent and this corresponds to closely similar free energies. The most important properties of PLZT are fast response, good temperature stability, excellent dielectric, ferroelectric and piezoelectric properties, which make them suitabl e for a large range of applications like large strain actuators and high charge storage devices, electrically controlled integrated structures based devices for high electrical field application [ 88]. Despite of its technological importance the AFE -FE indu ced phase transition in PLZT x/90/10 compositions has been less addressed. The associated functional properties of PLZT x/90/10 have been characterized by a limited number of authors only on few compositions (2/90/10, 3/90/10 and 4/90/10) across the FE-AFE phase boundary [ 64, 71, 75, 89 ]. Since the functional properties are also related to the various processing methods and sample characteristics (density, grain size, purity), a systematic investigation of dielectric, piezoelectric and ferroelectric propert ies for a larger range of compositions across the FE -AFE boundary is still necessary, as those properties are scientifically and technologically important for these compositions. Furthermore, as phase ’s superposition or low symmetry phase (monoclinic) has been reported in compositions across the boundary, an enhancement of functional properties is expected [ 90-91]. However, despite the evident advantages, there have only been few studies on properties of PLZT ceramics and the potential of these materials as sociated with its practical applications . 62 The current study aims at a better understanding of the antiferroelectricity of PLZT x/90/10 ceramics by a detailed comprehension of the structure dynamics, properties and phase transition induced by composition, electric field and temperature. A clearer insight into the evolution of the structure of PLZT across the FE/AFE phase boundary that is conditioned by a complex interplay of FE and AFE ordering, could potentially serve as a framework for the explanation of similar phenomena seen in other AFE systems. This may allow for the determination of the conditions under which the AFE behaviour of the specific systems might be altered or enhanced, thus allowing potential development of scenarios wherein the behaviour might be locally tailored for specific applications. A phase boundary between FE and AFE zone may allow develop a controlled amount of FE phase within the AFE matrix and vice versa, by inducing a local modification of the lattice dynamics. Such local structures could serve as potential tool for applications such as energy storage. If one is able to control the energy storage to be higher than conventional AFE ceramics, it would aid in reducing the size of the ceramics, thereby resulting in increased energy density per unit volume and further, in the miniaturisation of energy storage devices. The overall scientific goal of this study was to improve the knowledge concerning the role of La addition on dielectric properties and FE/AFE state of PZT 90/10 oxide perovskites ceramics . 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Ricote, Multifunctional Polycrystalline Ferroelectric Materials. Processing and Properties , Springer Series in Materials Science, 140, 782 (2011) 89. B. Araújo, D. A. Hall, M. E. Mendoza, J. A. Eiras, Effects of lanthanum modification on dielectric properties of Pb(Zr 0.90,Ti0.10)O3 ceramics: enhanced antiferroelectric stability , J. Mater. Sci., 43, 6087. (2008) 90. D.I. Woodward, J. Knudsen, I.M. Reaney, Review of crystal and domain structures in the PbZr xTi1−xO3 solid solution , Phys. Rev. B 72, 104110 . (2005) 91. R.S. Solanki, A. Senyshyn, D. Pandey, Space group symmetries of the phases of (Pb 0.94Sr0.06)(Zr xTi1−x)O3 across the antiferrodistortive phase transition in the composition range 0.620≤ x ≤ 0.940 , Phys. Rev. B 90, 214110 . (2014) 70 70 CHAPTER 3 Characterization: Principles and Techniques 3.1 Introduction Up to now, the fundamental theoretical aspects and general properties characteriz ing ferroelectric and antiferroelectric materials. The main recent literature results concerning PLZT and the unresolved fundamental issues were shown , in order to outline the importance of studying the systems with composition close and across the AFE -to-FE phase boundary. This chapter discusses the procedures used to prepare and the techniques and principles to characterize the PLZT ceramics . It includes the method of synthesis , powder X -ray diffraction, Impedance spectroscopy, Raman , Atomic Force Microscopy (AFM), Scan ning Electron Microscopy (SEM) and T ransmission Electron Microscopy (TEM) . Lastly , the procedure for the electrical measurements, the equipment employed for low and high field dielectric characterisation and temperature measurements are described. together with a short review of the principles of these method s. 3.2 Preparation of PLZT ceramics The dielectric, ferroelectric and electromechanical properties of oxide ceramics are considerably influenced by the synthesis method, chemical composition and sintering conditions [1-16]. Usually, the production process of piezoelectric ceramic materials is divided into th e following steps (Fig. 3.1): first ly the synthesis of perovskite phase is realised , then it follows the production of a green ceramic body by cold consolidation of these powders , followed by densification at high temperature (sintering). For the electrical and electromechanical characterisation, the ceramics are cut and polished and then electroded [15]. The usual method of preparation of powders is the traditional mixed oxide method [ 1, 15-16], which was also employed in the present work . Together with the composition and powder preparation, densification is a main 71 71 factor for obtaining good quality pore free ceramics. The conventional mixed oxides method requires a high calcination temperatu re (700 -900)°C of the starting materials (oxides) in order to promote the formation of perovskite phase . The high calcination temperature required by the solid – state reaction process leads to some disadvantages, such as large particle size, wide size distribution and a high degree of particle agglomeration. Thus, ball milling of the calcined material is a very critical step necessary in the processing in order to produce powders with high homogeneit y and low degree of agglomeration. Further, densification powder s into high dense ceramic s is critical to achiev e a high -quality ceramic product. Full density is rarely achieved with conventional sintering of ferroelectric ceramics , unless special techniqu es are employed to assist the sintering process during firing. An example of this is the use of an oxygen atmosphere for sintering lead -containing ceramics, such as PZT and PLZT. With air atmosphere only, relative densities of ~96% can be achieved, but using an oxygen atmosphere during sintering , this value can approach 99%. Another example is the use of PbO in excess to compensate the PbO loss (volatilization) during sintering , as well as for providing high densification rates via liquid -phase sintering. When both of these factors are used, bulk densities approaching 100% can be achieved. Typical sintering conditions for conventional PZT are 1250°C for 5 h with flowing oxygen [ 16- 17]. More dif ficult was t o obtain highly dense and pure perovskite phase in ceramics with high Zr content and for each composition various tests for optimisation of both calcination and sintering parameters have been realised . For example, high purity dense PLZT 90/1 0 ceramics have been produced by solid state method after calcination of precursors around 800°C for 1 h followed by sintering around 1200°C for 1 h in air or oxygen rich atmosphere [ 16-20]. In this work, PLZT x/90/10 ceramics with selected compositions across the FE -AFE phase boundary (with La3+ content x=0.020, 0.030, 0.031, 0.032, 0.033, 0.035, 0.038 and 0.040) according to the formula Pb 1-xLax(Zr 0.9Ti0.1)1-x/4 x/4O3) have been prepared by using powders synthesized by solid state reaction. The formulation was based on the assumption that La3+ would substitute Pb2+ on A site. In order to establish the best preparation conditions (calcination temperature and sintering temperature) four different compositio ns of PLZT were chosen as starting composition (two of which are out but very close to the FE -AFE border and two of which are in the middle of the border): 2.0/90/10, 2.5/90/10, 3.0/90/10 and 4.0/90/10. For these four compositions, two different calcinatio n temperatures and two different sintering temperatures were applied to obtain the sintered samples. 72 72 3.2.1 Synthesis of PLZT Powder The powders have been prepared from stoichiometric mixtures of PbO (purity 99.9%, Aldrich, Germany), ZrO 2 (purity 99%, Mel, Japan), TiO 2 (purity: 99%, Degussa) and La 2O3 (purity 99%, Aldrich, Germany). The starting oxides were weighted, wet mixed in water in a plastic jar for 48 h with zirconia milling media, dried in oven at 80°C and sieved at 200 µm. The fi rst approach in this work was focused on calcination temperature. Two different temperatures, 800 °C for 4h and 850°C for 4h were chosen, which are very close to other reports regarding the preparation of highly dense and pure perovskite phase PLZT 90/10 ac ross the boundary [ 19]. Thus the as prepared powder batch was divided in two. One part was calcined at 800 °C and the other one at 850 °C. During the calcination the powder was kept in a lidded alumina crucible. The calcined powders were crushed, sieved and then attrition milled at 400 rpm for 100 h with 10 wt . % poly(ethyleneglycol) (PEG, grade 10,000) as binder , then dried , reground and sieved again. The experimental proce dure of the powder preparation is shown in Fig . 3.1. Fig. 3.1 Flow diagram of the experimental procedure of the PLZT powder synthesis 73 73 3.2.2 Preparation of dense PLZT Pellets The powders previously obtained at two different calcination temperatures were used to prepar e the cold consolidated bodies, by using t wo different sintering temperatures: 1200 °C and 1250 °C. Fig. 3.2 Crucibles setup for sintering dense PLZT ceramics For each composition, disk shaped specimens (diameter ∼20 mm, thickness ∼2 mm) were uniaxially pressed (60 MPa) in a steel die followed by cold isostatic pressing (300 MPa) to increase the green density. The experimental procedure of the preparation of ceramics pellets is shown in Fig. 3.3. The most critical factor during sintering is the control of PbO vapour pressure to avoid the PbO weight lo ss and compositional change s. From the powder calcined at 800 °C for 4h two different sintered ceramics were prepared: one sintered at 1200°C for 2h a nd anot her one at 1250°C for 2 hours. The same procedure was applied for the powder calcined at 850 °C. The sintering was performed in air in lead -reach atmosphere maintained by using a PbZrO 3+5 wt. % ZrO 2 excess source in a closed Al 2O3 crucible (shown in Fig 3 .2). Thus, four categories of PLZT pellets were prepared , with t he compositions : 2.0/90/10, 2.5/90/10, 3.0/90/10 and 4.0/90/ 10 have been prepared , as following : I) calcin ation at 800 °C for 4h and sinter ing at 1200 °C for 2 h; II) calcin ation at 800 °C for 4h and sinter ing at 1250 °C for 2 h, III) calcin ation at 850 °C for 4h and sinter ing at 1200 °C for 2 h, and IV) calcin ation at 850°C for 4h and sinter ing at 1250 °C for 2 h .. According to the XRD patterns, the PZT powders prepared from this method were free of secondary phases and these ceramics have been used for the present study . ZrO plates PLZT disk Pack PbO rich atmosfere 74 74 Fig. 3.3 Flow diagram of the PLZT samples preparation from powder calcined at 800 °C for 4h In order to prepare the PLZT sample for electrical measurement, silver films were deposited by screen printingonto the polished faces follow ed by heat treatment at 7 50°C for h. 3.3 Characterization techniques 3.3.1 Microstructural characterization Understanding the relationship between micro structure and properties can lead to both 75 75 physical insight and finding enhanced materials properties. Electron microscopy techniques are widely used to analyse the surface micro structure s and compositional changes at a sub -micron level. With a magnification up to 200000x, high resolution down to 2nm together with the ability to generate localised chemical information these technique may give information about grain size, domains, defect and chemical homogeneity of powders and ceramics [ 21]. a) Scanning Electron Microscopy (S EM) The principle of such technique s is based on the interaction between electrons and sample: an incident electron beam focused onto the sample surface by a complex electromagnetic system is scanned across the sample’s surface , while emitted electrons are detec ted for each position by detector s and the intensity of the emitted electron s signal is displayed as brightness on a monitor and/or in a digital image file , creating in this way a scanning electron microscopy ( SEM ) image. The electrons interact with atoms in the sample, and produce several types of responses like: secondary electrons, backscattered electrons, Auger electrons, characteristic X -rays and photon of various energies. These are collected by specific detectors and multiple information concerning the samples can be extracted. Detected s econdary electrons and backscattered electrons give information about topography , phase and chemical composition of the investigated surface [21] together with the characteristic X -ray spectra of chemical elements (Energy -dispersive X -ray spectroscopy EDX) . For the present study, Scanning electron micrographs (SEM) were obtained on fresh fractured samples using a FEI Quanta 200 microscope. b) Transmission Electron Microscopy (TEM) Transmission electron microscopy (TEM) is one of the most precise technique for structur al characterization of materials at the atomic scale. With TEM technique many others characteristics of the sample, such as morphology, crystallization, stress , etc. can be studied . Compared with XRD , which analyses the crystal structure from the average Bragg diffractions from multiple grains of the sample, TEM is more precise, being able to distinguish between different crystal structures by extremely weak superlattice diffractions, which are invisibl e to XRD patterns . TEM uses a beam of highly focused electrons directed toward a thin sample which was polished down until becomes transparent for the electron beam (<200 nm). For the present work, a TEM JEOL JEM -2100 microscope (C s = 1.4 mm and point resolution 2.5 Å) operated at room temperature with an accelerating voltage of 200 kV was used. Cross sectional specimens (a 76 76 thickness of about 30 μm) were prepared from the as-processed pellets through standard 5 procedures including g rinding, cutting, dimpling, mechanical polishing and ion milling. The dimpled disks were thermally annealed above the Curie temperature to minimize residual stresses before Ar -ion milling to the point of electron transparency. 3.3.2 Structural Characterisation a) X-Ray Diffraction and High Resolution X -Ray diffraction In this study, several X-Ray Diffraction (XRD ) techniques, i.e. laboratory XRD and in situ high energy XRD with temperature chamber and in situ applied field were used to determine the room temperature structur es and to monitor the temperature and field -induced phase transitions of PLZT ceramics . These techniques are discussed in more details below. Generally, XRD analysis is used to determine the crystalline structure, phase composition and p hase purity of materials. Furthermore, using XRD data accurate lattice parameters can be obtained through Rietveld refinement [citare] . For ferroelectr ics, it has been demonstrated recently that XRD patterns are also sensitive to domain orientation and spontaneous strain [ 21-26]. The XRD technique is based on the diffraction of X-rays by crystalline plane of the sample . When an X -ray beam is incident on a surface of a crystalline material with an angle θ, a fraction of the beam is scattered by the atoms on the surface . The non -scattered part reaches deeper atoms in the crystal structure where a fraction is scattered. The constructive interference of the scattered X-rays represents the diffracted beam, which behaves as a specular reflection from a regular plane of atoms in the crystal [27-28]. The diffracted beam is generated according to the Bragg equation: Nλ = 2dsinθ where λ is the wavelength of the X -ray beam, d the crystal interplanar spacing, N is an integer that represents the orders of reflection, and θ the angle of incidence or reflection of the X - ray beam. 77 77 Fig. 3.4 Illustration of the X -Ray Diffraction setup showing the symmetry of the experiment as well as the Braggs law XRD analysis of the sintered PLZT compositions was performed at room temperature on the crushed samples with Bruker D8 Advance X -Ray diffractometer using Cu Kα1 radiation. The XRD patterns were recorded in the 2θ range from 10° to 80° with a step size of 0 .015° and a count time of 2s/step . Synchrotron radiation is an important tool for investigation of the properties of materials at atomic scale. Compared with laboratory XRD ’, Synchrotron XRD has several advantages like high flux, higher resolution, and rapid acquisition of high quality experimental data . High resolution XRD ( HXRD ) data show more features in the XRD patterns than laboratory detectors giving more accurate description of the crystal structure. Room temperature HXRD data were acquired using the Advanced Photon Source (APS) 11-BM-B beam -line at Argonne National Labs from Chicago (USA). Measurements were obtained using the following scan parameters: 78 78 – Start 2θ (°): -6.0 – End 2θ (°): 12.0 – Step Size (°): 0.001 – Time per Step (s): 0.15 – Scan Total Time (min): 30:00 – Wavelength (Å): 0.413895 In situ XRD is an innovative method used to determine the structural changes under different applied stimuli such as heat, pressure, and electric field. It is particular ly used to monitor the phase transition s during heating and cooling of ferroele ctrics . Recently, In situ XRD under electric fields technique have been demonstrate to be very useful in understanding the evolution of the structural changes due to field -induced strain, domain reor ientation or field -induced phase transition s [21-26]. In situ temperature and field technique s were applied to study the phase transition of PLZT ceramics induced by temperature and by electric field s. a) In situ field measurements A homemade time resolved XRD technique similar with that reported by Ref. [26] was used to measure the structural changes under in situ applied cyclic electric field. The XRD patterns were collected as a function of time during while on the sample was applied a triangular wave using at 0.03 Hz and amplitudes equal to or below the coercive field The XRD data are summed together , so that the results are presented as a function of the electric field sequences . The continuous acquisition of a selected pattern was possible within a l imited 2θ range for every 0.05 second during the application of a dynamic ac electric field up to the saturation polarization. For all the measurements, the applied electric fields were along the thickness direction of thin disk samples, under a triangular waveform with a loading rate of 0.8 mHz frequency. The data were collected in the -2 scan mode from the Platinum electrode flat face. X -ray diffraction scan parameters include a step size of 0.002° and a counting time per step of 1 s. b) In situ temperature measurements In-situ temperature measurements were realised utilizing PANalytical Empyrean XRD at the Analytical Instrumentation Facility at NC State University. The PLZT x/90/10 investigated ceramics were measured by using the following para meters: – Temperature Range: 25 – 450°C 79 79 – Ramp Rate: 2°C/minute – 2θ Range: 15 - 60° – Number of Scans: 82 b) Raman Spectroscopy The local structure of ferroelectric materials and properties at nanometric scale may differ significantly from the macroscopic ones; hence, the determination of the factors responsible for these differences is an important problem of the modern physics of ferroelectrics. Raman spectroscopy is very sensitive to local modifications like the short range order related to c hanges in the chemical bond length, angles and lattice dynamics of crystalline symmetry and is currently used for perovskite ferroelectrics. Therefore , this method is helpful to understand the significant differences in the physical properties of nanoscale ferroelectrics and their macroscopic analog ues. Raman spectroscopy is also used to study order -disorder phenomena and phase transitions in numerous solid solutions [29-34]. Basically, this technique is related to the inelastic light scattering in which th e change in the wavelength of scattered light is a result of the interaction of incident light with long wave length optical phonons of the scattering materials. The Raman spectrum provides information on the long wave length phonons with wave vectors |k| ≈ 0 [ 29-34]. For the present study, Raman spectra were obtained with a double monochromator (ISA/Jobin -Yvon/Horiba, Villeneuve d’Ascq, France) , as described in detail in in the paper of M. Deluca et a.l [34]. Excitation was provided by the 514.5 nm line of an Ar -ion. The maximum nominal power was adjusted to 100 mW, and the laser beam impinged on the sample either in a macro configuration (for low -temperature experiments) or by means of a 50× (NA = 0.55) microscope lens. Spectral fitting was performed with either Igor Pro 6 software package (WaveMetrics Inc., Portland, OR) or other multipeak fitting software (Labspec 4.02, Horiba/Jobin - Yvon), according to Gaussian –Lorentzian mixed modes . Low -temperature measurements were performed in a macro configuration ( the sample was glued onto a metallic support and attac hed to a He -gas cryostat cell , the instrument used was Cryotec, Daikin, Osaka, Japan ) while h igh- temperature measurements were carried out in a micro configuration by means of a Linkam TS1500 hot sta ge. 80 80 3.3.3 Electric al characterization Prior to the electrical measurements samples of two shapes were prepared: i) small (diameter of ~10 mm) and thin (thickness of ~0.5 mm) disks for P(E) measurement and ii) 20 mm diameter with thickness lower than 1/10 of diameter for the piezoelectric and dielectric measurements. a) Impedance spectroscopy The electrical measurement s were performed on disks samples via plan-parallel capacitor configuration with size of ~2 cm diameter and ~2 mm of thickness . The dielectric properties (real part of permittivity and dielectric losses) as a function of frequency were measured at room temperature by using a Hewlett Packard Imped ance Analyzer 4192A in the frequency range 100 Hz - 1 MHz. The measurement of the permittivity as a function of the temperature is among the most common tools to observe the phase transition in ferroelectric systems. and also reveals additional information on the dielectric properties of the material. For example, through the measurement of the imaginary component of the permittivity as a function of the frequency (and their variation with temperature) the conduction mechanisms or relaxation phenomena active may be deciphered. The measurements in the current study were conducted using an HP Impedance Analyzer 4192A . The samples were heated in aluminium furnace, and the temperature was controlled using an external thermocouple placed near the holder containing the sample. A bias of 1V was applied and several fixed frequenc ies in the range from 100 Hz to 1 Mz were used for measuring the dielectric properties versus temperature . The measurements were performed during heating from room temperature to 350 C with a rate of 1 C/min. I applied this technique to study dielectric properties (dielectric permittivity and conductivity) of a ferromagnetic material. The original results were published in the paper : M. Cernea, P. Galizia, I. V. Ciuchi , G. Aldica, V. Mihalache, L. Diamandescu, C. Galassi, CoFe 2O4 magnetic ceramic derived from gel and sintered by spark plasma sintering, J. Alloys Compd , (2016), (I.F. = 3.133 , A.I=0.551) . 81 81 b) Polarization -Field Hysteresis loop measurement s Polarization (P) –electric field (E) hysteresis loops have been obtained by using a modified Sawyer -Tower circuit ( Fig. 3.5 ) driven by a high voltage Trek amplifier, which provides sinusoidal signals of frequency 1 Hz and field strengths up to 6 kV/mm. A linear capacitor C ref is connected in series with the sample C pzt. The hysteresis loop was displayed by an oscilloscope in which the X channel measured the voltage across the R 2 and the Y channel measured the voltage induced across the capacitor C pzt. The quantity plotted on the horizontal axis is proportional to the field across the samples. Field was obtained by dividing the X channel reading by the thickness of the sample. If the capacitance of the capacitor C ref was known then the charge accumulat ed across the sample can be obtained by multiplying the measured voltage with the known capacitance of the capacitor C ref. Once the accumulated charge was determined , then the polarization was calculated by dividing the charge by the surface area of the el ectrode printed on the sample. During testing, the samples were immersed into silicone oil to ensure a n homogeneous tem perature and to prevent arcing [35]. Fig.3.5 Sawyer -Tower circuit used for polarization measurement c) Piezoelectric characterization Polycrystalline ceramics have a small spontaneous polarisation, since the polar axes of the individual grains are randomly oriented (macroscopically centrosymmetric state). In order to 82 82 induce a preferential orientation and loss of the inversion centre allo wing piezoelectricity, poling is necessary. In the conventional poling process, electric field is applied to the ceramics via the metal electrode. However, many factors influence the effectiveness of poling process. The drawback of conventional poling meth od is the occasional occurrence of fracture s and crack s due to the applied high electric voltage. Generally, the purpose of poling process is to induce maximum degree of domain alignment (in order to acquire the ∞mm symmetry) by implementing lowest electric field at temperature as close as possible to the room temperature. Experimentally, the procedure consists of placing the electrode ceramic in the device sample holder immersed in a bath of silicone oil. The process takes pla ce at temperature of 120° C (Fig. 3.6). The heating system ensures a constant temperature with a thermostat, while a magnetic bar keeps moving the bathroom ensuring the uniformity of the temperature. Once the set temperature is reached the voltage is appli ed for 40 minutes. For the present study, t he PLZT samples were subjected to electric fields from 3 to 5 kV/mm , depending on t he composition of the material. Fig. 3.6 Poling sample holder Due to the high voltages, the entire system is contained in a Faraday cage to avoid unwanted electrical manifestations. At the end of the 40 minutes, the oil bath is allowed to cool naturally down to the ambient temperature while keeping the applied voltag e. This part of the process takes approximately another 60 minutes. At this point the voltage is removed and the sample is poled. This process influences very much the material constants , like for example the dielectric constant . 83 83 The poling efficiency depe nds on the phase composition of the material and quality of sample (density, lack of cracks, etc) [36]. In order to find the piezoelectric constants, all the investigated PLZT samples were poled under a dc field of 3 kV/mm for 40 min at temperature of 120 ⁰C. The resonance and antiresonance frequencies fr and fa were measured using a Hewlett Packard Impedance Analyzer 4192A after 24 h from the poling process. The piezoelectric d33 constants were measured by a quasistatic Berlincourt -type d33 meter, while the electromechanical factors ( kp, k31), piezoelectric ( d33, d31) and voltage ( g31, g33) constants were determined by resonance and antiresona nce method according to the IEEE Standard of Piezoelectricity [ 36]. 84 84 Bibliography III [1] L. Pardo, J. Ricote, Multifunctional Polycrystalline Ferroelectric Materials. 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Deluca, H. Fukumura, N. Tonari, C. Capiani, N. Hasuike, K. Kisoda, C. Galassi and H. Harimac, Raman spectroscopic study of phase transitions in undoped morphotropic PbZr 1−xTixO3, J. Raman Spectrosc. 42, 488–495, (2011) [35] Y. Xu, Ferroelectric materials and their applications , North Holland Elsevier Sci. Publ., Amsterdam, (1991) [36] An American National Standard, IEEE Standard on Piezoelectricity, ANSI/IEEE Std 176 -1987. 87 87 CHAPTER 4. Microstructural and Structural Characterization 4.1 Introduction After the introductory chapters, in the following it will be presented the original results concerning the PLZT x/90/10 ceramics with La3+ compositions across the ferroelectric/antiferroelectric (FE/AFE) phase boundary. This chapter is dedicated to the study of the influence of the La3+ addition on the structure of PLZT x/90/10. Structural origins play a central role in understanding the physical properties of FEs and AFEs. The observed macroscopic FE/AFE properties are closely linked to the interplay between the variation in underlying polar order and the corresponding crystal structure. Depending on Zr/Ti ratio and dopant concentration, PLZT ceramics exhibit a variety of phases with different crystalline symmetries and different functional behaviours, such as ferroelectric, antiferroelectric , paraelectric and mixed phases. The FE/AFE nature of the phases presented in PLZT x/90/10 ceramics is not unanimously accepted but it was rather deduced from the shape of P-E hysteresis curves and from the aspect of the permittivity vs. temperature dependence. For example, a broad maximum permittiivty around T m was assigned to the transition from FE/AFE to PE state [1-23 ]. The available reports on the structure of the PLZT solid solutions on the Zr rich side of the PLZT phase diagram are controversial. The first phase diagram of PLZT established by Haertling [24 ] shows a fine border between the FE and AFE phases, while in the phase diagram of PZT proposed by Jaffe [25 -26], there is a finite compositional region between the FE and AFE phases. Precision measurements of crystallographic structure s at domain boundaries using Kiku chi diffraction methods confirm ed the close relationship of PLZT to PbZrO 3 (PZ) system [ 27-28]. Earlier, Jo na et al. established that t he structure of PZ is pseudo -tetragonal, being related to a peculiar relation between the orthorhombic a and b axes: b= 2a . Based on a combination of x -ray diffraction and neutron diffraction data they assigned an orthorhombic Pba2 space group with lattice parameters: ao, bo, co with ao =√2ac, bo =2√2 ac, co= 2a c where c denotes the lattice parameter of the cubic PE phase. This space group is not, however, centrosymmetric and therefore would allow a FE polarization. This structure is in contradiction with the theoretical Kittel studies which describe d PbZrO 3 as an AFE [29]. More recently, various subsequent temperature -dependent X -ray and 88 88 neutron single-crystal and powder diffraction studies determined the structure as centrosymmetric Pbam , and as proper AFE stable state [30]. The large majority of authors consider the Zr rich side of the PZT diagram as having a Pbam symmetry and stable AFE state. However, there is a study where a weak FE was reported in PZ [31] and the space group of the AFE orthorhombic phase in PZ is noncentrosymmetric, as previously suggested by Jona et al. [29]. However, only one experimental report evidenced a FE hysteresis loop with a saturation polarization of only 0.1 μC/cm2 at that time (value which is far below the previously estimated by Jona et al. ). More recently, Pintilie et al. [32] demonstrated unequivocally the existence of ferroelectricity in PZ films at 4.2 K with a remnant polarization of about 27 C/cm2, which is in a remarkable agreement with the value of 25 μC/cm2 estimated by Jona e t al. about 50 years ago. PZ films are in FE state only up to T= 16 K where a low temperature FE-AFE phase transition occurs. Their results were confirmed later by first principle calculations where the low temperature FE phase was estimated to have a rhombohedral structure with R3c space group, while the room temperature AFE phase was estimated to be a Pbam orthorhombic structure. Therefore, the room temperature crystalline structure of PZ is still considered to be as belonging to Pbam space group [32 , 33]. Lanthanum doping in lead zirconate has been found to increase the stability range of the AFE orthorhombic phase in the Zr -rich side of the phase diagram [ 34]. For example, PZT 94/6 exhibits a phase boundary between FE rhombohedral and AFE orthorhombic phases, which is very sensitive to the La3+ addition [ 24]. FE and AFE ordering coexistence was found in PLZT 6/80/20 [ 35]. The PZT 90/10 based composition was reported as being rhombohedral FE at room temperature, even neutron diffraction studies revealed signs of orthorhombic symmetry [2]. However, n eutron and electron diffraction data suggested that on the Zr -rich side the symmetry is lower than rhombohedral, and pointed out either a monoclinic Pc structure or a multiphase system. [12, 16, 36]. Studies of R. Villaurrutia and his coauthors suggested an incommensurate (AFE T/In) phase with not a unique crystalline symmetry [13, 14, 18, 27, 28] . TEM studies revealed a rhombohedral macroscopic symmetry of PLZT x/90/10 for small amount of La (x<2) and tetragonal together with an incommensurate s tructure for x >4 [ 2]. Ishchuk et al. [37] reported difficulties in the structure identification, while Refs. [ 4, 5] refined the composition PLZT 4/90/10 as orthorhombic instead of tetragonal. It has been observed experimentally from TEM investigation s that Lead zirconate titanate ceramics with a Zr:Ti ratio of 90:10 and doped with 2 -4% La display an incommensurate AFE phase. The presence of this phase make s the structure of PLZT 90/10 more interesting and more controversial [ 13, 14, 18, 2, 28 ]. It has been suggested that frustration from the competing FE and AFE ordering is responsible for the presence of such incommensurate structure. However, the exact cause of the incommensurate phase in these oxides is still under 89 89 debate. R. Villiarutina et al. exploited in detail the domain structures and the nanostructures contained in a composition range which demonstrates the incommensurate phases of interest: La- doped PZT with La:Zr:Ti ratios of 2:90:10, 3:90:10 and 4:90:10 [13, 14, 18, 27, 28 ]. Furthermore, the existence of an intermediate monoclinic Pc phase -a subgroup of both the orthorhombic and rhombohedral symmetries -was proposed recently by Woodward et al. [38]. In spite of many studies already carried out on the compositions La-PZT 90/10 the assignments of their phase symmetry is still under debate. It is even more important to establish more precisely the range of coexistence of AFE/FE phases in order to understand and to better control their properties. The stability of FE or AFE structures may be predicted using the Goldschmidt tolerance factor (t) as follows: 𝑡 =𝑅𝐴+𝑅𝑂 √2(𝑅𝐵+𝑅𝑂) where RA is the radius of large cation A (Pb2+), R B is the radius of small cation B (Zr4 or Ti4+) and R O is radius of anion (O2-). 0.00 0.01 0.02 0.03 0.040.98600.98650.98700.98750.9880 t La (at. %) Fig. 4. 1 . Goldsmith tolerance factor t for all PLZT x/90/10 investigated composition For t = 0.98 5-1.00 the cubic perovskite structure is favor ed, for t < 0.95 the AFE, while t > 90 90 1.00 is reported for FE materials. Generally, the AFE phase is favored for low tolerance factor. For example PZ has the lowest t~0.9821 in the PZT solid solution, which increases with the addition of Ti, since the Ti4+ ionic radius (0.605 A˚) is smaller than that of Zr4+ (0.72 A˚) [38, 39]. On the other hand, the appearance of AFE phase in PZT 90/10, (which is FE) may be expected since the La substitution on the A site decreases due to the La3+ radius (1.36 A˚) is smaller than Pb2+ (1.49 A˚). However, according with Ref. [40, 41] when doping with trivalent ions (La3+), for the stoichiometry Pb 1−xLax(Zr yTi1−y)1−1x/4 O, the charge balance is maintained by creation of vacancies of (Ti, Zr) on B site in the perovskite lattice. The calculated tolerance factor for all the PLZT x/90/10 investigated composition s are presented in Fig.4.1. The tolerance factor decreases linearly as the La content increases reaching the lowest value of t≈0.9859 for PLZT 4/90/10. All the composition s (including PZT 90/10) show a tolerance factor lower than 0.988, which indicate an AFE distortion. The PZT 90/10 based composition show the highest t value of 0.9878. However the PZT 90/10 is known to be rather FE than AFE [36 ]. Therefore a predominantly FE structure may be expected for low La addition and a predominantly AFE structure may be expected for higher La amount. Thus, La doping is expected to favor the formation of AFE phase. This is not surprising, since it is well known that the La addition induced AFE stabilization in the high Zr side of the phase diagram of PZT [25, 26] and PLZT. [1 -24]. Therefore, these results are in agreement with the phase diagram reported by Haertling and other publications which report AFE stabilization on PZT 90/10 by La doping [1 – 24]. In addition, when La is added to replace Pb ion on A site, it induces a lattice vacancy on the B site for electric charge mismatch and leads to disrupt the translational periodicity of the lattice. La substitution and the associated vacancy break the long range interaction between clusters. Above a critical amount of La substitution, decoupling is sufficiently strong that the long range FE interactions are destroyed and nanopolar domain clusters are established [35, 37]. According with D. I. Woodward et al. [38, 39 ] the composition with the tolerance factor in the range 0.985< t <1.06 are expected to have untilted structures. Perovskite compositions with 0.964< t<0.985 are usually tilted in antiphase and perovskites with t < 0.964 are expected to show in-phase and anti-phase tilting. As t continues to decrease, the stability of the perovskite phase decreases and eventually will not form. Since the all investigated PLZT x/90/10 show a tolerance factor higher then 0.985, an untitled FE or AFE structure may be expected. However the value of 0.9859 obtained for the composition PLZT 4/90/10 is close to the limit between antiphase tilted and untitled range and thus, a particular attention is necessary when the structure of this composition will be further identified. 91 91 Only a combination of macroscopic and local structure analyses can give a full picture of the macroscopic structure. All the accessible information concerning macroscopic and local structure s must therefore be discussed in context with theoretical structural model derived from the diffraction data. In the following , we will describe a detailed study which was undertaken to addre ss the existing controversies and to obtain new insights into the room temperature structure of the ground -state phase of PLZT in the La composition range of 0.02 ≤x ≤0.04 (around the FE- AFE boundary ) using synchrotron x -ray powder diffraction ( SXRPD), TEM, and Raman studies. Based on the literature review presented in the introduction of this chapter and the above mentioned assumptions, we establish ed the following objectives regarding the study of room temperature structure of the PLZT x/90/10 compositions: 1) To determine if high purity phase can be obtain ed and to choose the highest quality samples for further investigations ; 2) To determine the influence of La content on the FE and AFE phases evolution and to determine the range of phases coexistence; 3) To determine the structures of the phases content in the PLZT x/90/10 composition; 4) To determine whatever or not a monoclinic structure may be presented; 5) To check if an tilted structure is formed for La amount close or equal to 4 at %. 4.2 Phase purity In the following, a structural study aimed to unveil the suitable preparation conditions in order to obtain high purity PLZT phases and highly densified PLZT pellets will be presented . First, it will be shown the influence of the calcination temperature on the formation of perovskite phase, then the influence of the sintering temperature on the phase purity and densification of PLZT pellets. At the end, the best conditions will be chosen for the preparation of entire compositional range of PLZT x/90/10. The influence of compositional variation on the structure was investigated by HRXRD TEM and RAMAN. 4.2.1 Phase purity of calcined PLZT powders Details about the preparation were shown in the previous chapter of this thesis. Based on previous reports on preparation of some PLZT x/90/10 compositions [1 -23] the conditions for the calcination of PLZT powder were established at 800 C for 4h and at 850 C for 4h, respectively. 92 92 The XRD results for powder calcined at 800 and at 850 C for 4h are shown in Fig. 4.2 a)-b). As can be observed in the Fig. 4.3 a) , after calcining the mixed oxides at 800°C for 4 h the perovskite phase of PLZT is formed. However, a ll the XRD patterns show also the presence of very small intensity peaks of pyrochlore phase (marked with alpha). The existence of the pyrochlore phase is usually caused by off-stoichiometric precursor solution or the volatilization of PbO during the calcination process. There are some possible reasons for the formation of such secondary phase in PLZT solid solutions with compositions at the Zr-rich region: (i) the starting materials did not react and some secondary phases like PbO or pyrochlore results as consequence of incomplete reaction between oxides; (ii) the volatilization of PbO during the calcination process may cause deviations from the nominal composition and detrimental densification and functionality of the ceramics, and (iii) the metastable nature of the PLZT solid solutions with compositions at the Zr-rich region . Therefore, a higher calcination temperature may be need to be applied in this method. Fig. 4. 2 . Comparison of XRD patterns of PLZT calcined powders for four La concentrations (2.0, 2.5, 3.0, 4.0 at % ) calcined at a) 800C and b) 850 C for 4 h Since the PbO volatilization occurs at temperature above 900 C (melting point of 890°C)) we can exclude the second reason. Therefore, it is necessary to increase the calcination temperature in order to obtain a higher purity. Structural, electrical and thermal properties of the PLZT solid solution are extremely sensitive to the compositional variation. To study the PL ZT x/90/10 solid solution, it is essential to have the samples with exceptionally high quality and accuracy in terms of chemical composition and density. We thus increased the calcination temperature from 800 to 850oC for 4 h. The XRD results for powders calcined at 850 C for 4 h for the PLZT compositions 93 93 are shown in Fig. 1a). Diffraction patterns show a significant reduction of the PbO and pyrochlore phases in all the PLZT powders , while low intensity perovskite peaks are more visible , indicating a good reactivity between oxides at this temperature. XRD results show that the se PLZT powders are pyrochlore phase free, and diffraction pe aks match with the JCPDS database card of PLZT (No. PDF 77 -1194 ). 4.2.2 Phase purity of PLZT sintered ceramics The XRD patterns of the as-sintered PLZT disks obtained from powder calcined at 800 C, with four La content of 2.0, 2.5, 3.0, 4.0 at %, sintered a) at 1200 C for 2h and b) at 1250 C for 2h are show n in Fig. 2. a)-b), respectively. Fig. 4. 3 Comparison of XRD patterns of PLZT sintered ceramics obtained from powders calcined at 800 C with La content of 2.0, 2.5, 3.0, 4.0 at % sintered at: a) 1200 C and b) 1250 C. There is no evidence of any pyrochlore phase in all the sintered ceramics. However, the density was definitely far from the theoretical one, as it shown in Fig. 4.5 a). Although the pyrochlore phase was significantly reduced by sintering at 1200 C and 1250 C, the quality of the samples were not as good as to be used for the propose of these study since the relative density was lower than 90 %. Ceramics sintered at both 1200 C and 1250 C from powders calcined at 850 (Fig. 3 a)- b)) 94 94 are dense and free from pyrochlore phase and highly crystallized. As shown below, the relative densities of all these PLZT samples are higher than 95%. The results show that the dense PLZT ceramics pyrochlore-free can be obtained in a reproducible way by solid state reaction by using a calcination temperature of 850 C and a sintering temperature of 1250 C. Fig. 4. 4 Comparison of XRD patterns of PLZT sintered ceramics obtained from powder calcined at 850 C with four La concentrations (2.0, 2.5, 3.0, 4.0 at %) sintered at: a) 1200 C and b) 1250C for 2 h The densities of PLZT compositions 2/90/10, 2.5/90/10, 3/90/10 and 4/90/10 obtained after sintering, for two different calcination temperatures, are shown in Fig. 4 a). Usually, the density increases with increasing the sintering temperature. However, the density of PLZT ceramics is affected by the sintering condition in a more complex way. As it can be observed from Fig. 4 a) also the calcination temperature plays an important role in defining the final density of the ceramic body . As expected, the density is maximized (relative density above 95% for all the compositions) for ceramics prepared from powder calcined at 850 C and sintered at 1250 C (Fig. 4b)). The dependence of the experimental and relative density on the Lanthanum concentration shown in Fig. 4b) indicates a non-monotonous trend, with a distinct density minimum in the range of compositions x = (0.030, 0.035). A similar behaviour was reported earlier in a detailed study concerning the density and grain size evolution in PZT ceramics, where the non-monotonous evolution of density with composition was interpreted as related to the crystalline phase modification and in particular, to the phase coexistence, which seems to create specific properties of the calcined powders to provide a limited densification during sintering [42 ]. 95 95 In our compositions, the minimum densification level of 92% - 93.5% relative density for the La3+ content in the range (0.030, 0.035) can be due to the phase transformation from FE rhombohedral to AFE orthorhombic through a morphotropic region with phase coexistence. Therefore, it was decided to keep the calcination routine at 850 C for 4 h and the sintering temperature of 1250 C for 2 h to prepare all the PLZT ceramics for the further investigations . Fig. 4. 5 Lanthanum concentration dependences of the experimental (Archimede’s) absolute and relative densities of the PLZT ceramics: a) variation of PLZT density at various sintering temperatures for four representative compositions and b) the evolution of PLZT density for all the investigated compositions (calcination at 850oC, sintering at 1250 C). b) 96 96 4.3 Microstructure We studied the microstructure of these samples using scanning electron microscopy (SEM) in order to understand the composition dependence of grain morphology and sizes of the PLZT sintered disks (0.00≤ x ≤ 0.040). SEM images of cross section fracture of the PLZT sintered disks from Fig. 4.7, Fig 4.8 and Fig. 4.9 show dense, crack free morphology and small amount of pores in every composition.. It can be seen that the grain size is greatly influenced by the concentration of La3+. Fig. 4. 6 SEM micrographs of the fractured surfaces of PLZT 100x/90/10 compositions: (a) x=0.0 at. %, (b) x=2.0 at. %, (c) x=2.5 at. % and (d) x=3.0 at. %. The un-doped ceramic ( i.e., x = 0) is dense, well-crystallized and shows large faceted grains uniformly distributed, of diameters about 20 μm (Fig. 4.7 a)). As low amount of La is introduced a) b) c) d) 97 97 (i.e. x=0.020), the size of the grains drops significantly and the average grain size become small er of ~2.5 μm (Fig. 4.7 b). There is an inhomogeneous grain size distribution leading to a bimodal grain size distribution with few big grains of about 6-7 µm surrounded by many small grains (of about 2µm) observed for both low La amounts ( i.e. for x˂3.5) (Fig. 4.6 and Fig. 4.7 )). As x further increases to 3.8 and 4 .0, the microstructures look more homogeneous with smaller grains of about ~2µm, often with rounded aspect and uniformly distributed within the sample (Fig . 4.8). Thus, the grain growth in the doped samples is suppressed. For these compositions with finer grains, a lower level of density was also obtained. Fig. 4. 7 SEM micrographs of the fractured surfaces of PLZT 100x/90/10 compositions: (a ) x=3.1 at. %, (b) x=3.2 at. %, (c) x=3.3 at. % and (d) x=3.5 at. %, a) b) c) d) 98 98 For PZT-based ceramics, donor ( e.g. La3+, Nb5+, etc.) doping generally leads to an grain growth inhibition. A-site vacancies created by La3+ doping are supposed to be bound to the impurity ion (La3+) and the mass transportation is inhibited. As a consequence, the diffusion coefficient of vacancies decreases by La3+ doping and thus, the grain growth is inhibited [43-46]. It is worth mentionin g that the grain size predominantly decreas es with increasing the amount of AFE phase fraction. This may suggests that the larger grain size favor s the stability of AFE rhombohedral structure, while the smaller grain size favors the stability of the orthorhombic FE one. Fig. 4. 8 SEM micrographs of the fractured surfaces of PLZT x/90/10 compositions: (a) x=3.8 at. % and (b) x=4.0 at. %. 4.4 Structural characterization 4.4.1 Phase characterization by XRD Further, we have investigated the phase formation of the PLZT x/90/10 ceramics by XRD (Fig. 4.9a, with details from these patterns in the regions of interest in Fig. 4.9b). All the peaks could be indexed as belonging to the perovskite phase, which indicates that a pure phase is formed for all the studied ceramics. For the purposes of this thesis, Miller indices will generally be assigned to reflections on the basis of the pseudo-cubic unit cell. A composition-dependent crystalline structure with a gradual change of some specific peaks as a function of La content is observed. According to literature reports for similar compositions, the FE phase has R3c space a) b) 99 99 group while the AFE phase has a Pbam space group [2, 4, 5, 18, 13, 24 , 30]. In the following the XRD patterns of all the investigated compositions will be analysed according to the previous ly mention ed structures. The XRD diffraction pattern of the composition PLZT 0/90/10 as well as for 2/90/10, showed single (100), (200) peaks and split (111) reflections, indicating that their crystalline structure has a rhombohedral symmetry with R3c space group similar with the patterns no. PDF 77 –1194. These results agree with the phase diagram of PLZT proposed by [24 ] and other reports [2, 4, 5 ] for low La content (x≤ 0.0 2). The (101), (210), (110) and (211) main reflections of the 2.5/90/10 ferroelectric composition split, respectively, into (101) and (10-1), (210) and (20- 1), and (211) and (21-1) reflections, which indicate a rhombohedral distortion, a lower symmetry phase or that this composition has more than a single FE phase. As La3+ content increases from 2.5 to 4 at. % the splitted (111) peak emerge gradually in a single one, while the (100), (110), (200), (210), (211), (220) and (310) cubic reflections split into doublets typical to orthorhombic symmetry with pseudo cubic indices: (100) and (001), (110) and (101), (200) and (002), (210) and (201), (211) and (112), (220) and (202), (310) and (301), respectively. 10 20 30 40 50 60 70 80x=0,00x=4,00 x=3,80 x=3,50 x=3,30 x=3,20 x=3,10 x=3,00 x=2,50 O(301)O(310)O(212)O(221)O(202)O(220)O(112)O(211)O(102)O(201)O(210)O(002)O(200)O(111)O(101)O(110)O(001)Intensity(a.u.) 2O(100) x=2,00a) 37 3843 44 45b) x=4,00 x=3,80 x=3,50 x=3,30 x=3,20 x=3,10 x=3,00 x=2,50 O(002)O(200)O(111) x=0,00Intensity(a.u.) 2x=2,00 Fig. 4. 9 X-ray diffraction patterns (a) and details b) at room temperature for all the investigated PLZT x/90/10 compositions. Therefore, the substitution of Pb by La in the PZT 90/10 induces a transition from 100 100 rhombohedral to orthorhombic phase. The splitting of (111) and (200) peaks can be clearly seen in Fig. 4.9 b). Thus, t he results presented above confirm that the crystalline symmetry of compositions with x>0.02 chang ed gradually into orthorhombic. A small peak on the left side of the (111) indicate s that the rhombohedral reflection is still observed in AFE com positions up to 3.3 at. % of La. This suggests that a coexistence of r hombohedral and orthor hombic phases takes place in the range of composition (0.025≤ x<0.035) . Further more , while the intensities of (11 -1) and (001) rhombohedral peaks decrease , the intensity of the orthorhombic (002) reflexion gradually increase s for the same intermediate composition al range. Th is evolution indicate s a gradual modification of the amount of FE rhombohedral phase into an orthorhombic AFE phase, when increasing La addition. In addition, the XRD patterns for such intermediate compositions could not be fitted anymore with a single rhombohedral phase. Single orthorhombic phase is characteristic to the high La3+ amounts, x=0.035, x=0.038 and 0.040 which can be indexed very well w ith orthorhombic Pbam space group (PDF 70 -4844), as also suggested in the ea rly phase diagram of Haertling [24] and more recently, in structural studies by neutron diffraction [ 30]. The intermediate compositions, in the range x = (2.5, 3.5) , can be also fitted quite well with the same orthorhombic symmetry, but the signs of the rhombohedral phase can be still observed. Thus, a thorough investigation is necessary in the assessment of phase ’s superposition. In the following the characteristics of the (100) and (111) peaks will be presented in detail in order to give an insight into the evolution of the phase structure, as displayed in Fig. 4.11. Fig. 4. 10 Lorenzian fit peaks and peaks deconvolutions for PLZT composition 3/90/10 101 101 The intensity of the diffraction pattern of a particular phase in a mixture of phases depends on the concentration of that phase in the mixture. Thus, in order to quantify the FE and AFE phases evolution at the FE –AFE border, we have c oncentrated on the influence of La addition on intensities evolution of (100) and (111 ) reflections [4 7]. We used as example the study of Pandey et al. [47] for (Pb,Ba)ZrO 3 and we tried to determine the threshold x for the onset of the FE R phase starting from the AFE com positions. Diffraction profil es measured along the radial ( 2θ) direction in the pseudocubic (001) and (111) zones for the composition PLZT 3/90/10 are shown in Fig. 4 .10, where the Lorenzian functio n fits are superposed on the observed diffraction peaks. Depending on the nature of the phase presented in the investigated composition, FE or AFE, the (100) reflection appears as a singlet or a well split doublet in the XRD patterns, respectively. Similarly, the (111) reflection is a doublet and singlet in the XRD patterns of AFE and FE phases, respectively. Furthermore, the contribution of the (100) of the FE phase is nearly overlapped on the same peak of the AFE phase. The missing of the (11-1) reflection between 3.3 and 4 La3+ at. % content implies the disappearance of the FE phase in the compositional range. From individual fitting of every couple of reflections, we have determined the intensities of the (100), (001) and (111), (11-1) Lorenzian splitted peaks, for each investigated composition. As shown in the Fig. 4.10, an example of the profile fitting results for a PLZT 3/ 90/10 sample comprised of coexisting FE and AFE phases is given. Then we used theoretically calculated intensities to determine the ratios between intensities of the (100) and (001) AFE phase peaks , respectively between (111) and (11 -1) FE reflections. The results are plotted as function of La3+ content (Fig 4.11). The relative intensities reported in PDF 70 -4844 XRD pattern of the pure AFE phase and PDF 77 -1194 of pure FE phase are equal to 0.60 for (001) and (100) couple of reflections respectively and 0.35 for (11 -1) and (111) couple of reflections in the presenc e of pure AFE phase , respectively. It can be observed that the determined ratio of the intensities of (001) and (100), (11 - 1) and (111) reflections when the AFE and FE phases coexist is smaller than that for the pure AFE phase, or pure FE phase , respective ly. As x increases, the (11 -1)/(11 1) reflection ratio decreases from about 0.35 in the FE phase ( x≤0.020) to 0 in the AFE phase ( x≥0.035), indicating the disappearance of the FE phase. Contrariwise, the ratio of (001)/(100) increases from 0.00 in the FE phase ( x≤0.020) to 0.49 in the AFE phase ( x≥0.035). The monotonic variation of these ratios indicates the progressive transformation into AFE phase as La content increase s. Since the ratio values get for compositions x≥0.033 are close to the values determined for pure AFE phase, we can settle that the amount of AFE orthorhombic phase stabilizes on the expense of the 102 102 rhombohedral FE phase for higher La3+ amounts. This may be explained as follows: a small amount of La3+ cannot change dramatically the crystalline structure of PLZT ceramics, i.e. a rhombohedral phase still survives in the ceramics with x>0.025, but its amount is smaller. According with Fig. 4.11, a compositional range where both FE and AFE phase coexist (0.025≤ x<0.035) may be identified , but the amount of FE and AFE are reduced with respect to the pure FE phase or pure AFE phase. Therefore , the transition from rhombohedral FE towards orthorhombic AFE induced by the La3+ addition takes place progressively, t hrough a compositional range with coexistence of phases with similar free energies. Fig. 4. 11 The evolution with x of the (111)/(11 -1)R intensities ratio and of the (001)/(100) O intensities ratio The above presented structural data gave us an overview of the phase ’s evolution across the FE/AFE phase boundary. We can derive the main following conclusions: 1) PLZT x/90/10 with x = 0.0 25 for which ((001)/(001) intensities ratio is higher than 0) and x=0.033 for which ((111)/(11 -1) intensities ratio is 0) may be considered the threshold x values for the emergence of the AFE phase and for the disappearances of FE phase. 2) PLZT x/10/90 ceramics endure a series of crystalline symmetries with increasing La3+ content: rhombohedral FE phase (0≤x<0.020), coexistence of orthorhombic– 103 103 rhombohedral phases (0.025≤x<0.035), and orthorhombic AFE phase (x≥0.035). Further investigation will be performed for a quantitative evaluation of the relative amount of these phases and into assessment of the symmetry space group of all the investigated composition. More details about the structures will be presented in the following. 4.4.2 Crystalline structural characterization by HXRD Previously we had argued that phase composition of the PLZT system in the FE/AFE region should be a mixture of FE R3c and AFE Pbam phases. The widest phase coexistence region has been observed for (0.025≤x<0.035) were the FE R3c and AFE Pbam phases persist and then pure AFE is observed for x≥0.035. A systematic study aimed to deciphering the subtle changes in the local crystal structure, and understanding how these changes give rise to various physical properties was considered necessary. Using the high sensitivity of X-ray diffraction to fine octahedral tilt transitions, systematic composition-dependent superstructure or weak reflections in electron diffraction patterns can be identified and carefully indexed by the pseudocubic unit-cell to match specific structural changes in the lattice. In order to resolve the existing controversies about the structure of the ground-state phase of PLZT x/90/10 compositions, we used synchrotron x-ray powder diffraction (SXRPD) . This part of the study was realized in collaboration with Prof. J. L. Jones from North Carolina State University, Raleigh, North Carolina USA . The High Resolution X-ray diffraction patterns of the sintered PLZT ceramics the compositions with x=0.00, 0.020, 0.025, 0.030, 0.031, 0.032, 0.033, 0.305, 0.038 and 0.040 are shown in Fig. 4.12 . In agreement with previous XRD results, pure perovskite phase is confirmed for all the compositions. As previously discussed, the crystalline structure in the AFE state is characterized by orthorhombic distortions of the elementary crystal cell, whereas the one in the FE state undergoes rhombohedral distortions. To further confirm this, we focused our attention on (111), (200) and (211) x-ray diffraction characteristic peaks, which are the more relevant for the present study. In the presence of rhombohedral distortions the (200) line is a singlet and the (111) and (211) lines are a doublet whereas the (111) and (2 11) line are singlet while (200) line is a doublet for an orthorhombic distortion of the elementary cell. The variation of profiles of these lines with the change of La content in solid solutions is clearly seen in Fig. 4.13. The PLZT x/90/10 composition with x = 0.02 was earl ier regarded as rhombohedral phases on the basis of the singlet nature of the (200) and (211) profiles. According with the HXRD pattern of this composition, there are distinct changes in the diffraction profiles of this composition compared with the rhombohedral phase of 104 104 PZT 90/10 based one. As an increase in (111) peak intensity and a small intensity shoulder (marked with arrows) appears in the right side of (200 ) and (211 ) profiles together with large asymmetry of (200) which indicate the appearance of the orthorhombic phase in this composition. Thus a small amount of La3+ addition on the A site is sufficient to stabilize a small amount of AFE orthorhombic phase. Therefore the threshold x values for the onset of the AFE phase, (respectively for the decrease of FE phase amount) may be extended from the composition x=0.025 to the composition 0.02 of La. As previously, once the amount of La is further increased in the PZT 90/10, there are distinct changes in the diffraction profiles. The splitting of (111) perovskite peak is decreased while the splitting of (200 ) is increased. Since (200 ) is a singlet in rhombohedral and a doublet in orthorhombic symmetry, this is a clear indication for a two phase coexistence. In addition, the AFE phase becomes the majority phase , while the FE phase is the minority one. 5 10 15 20 25 30(301)(310)(220)(202)(211)(112) (201)(102) (200)(002)(111)(110)(101)(100)(001) x La 0.040 0.038 0.035 0.033 0.032 0.031 0.030 0.025 0.020 0.000Intensity (arb.unit) 2 (degrees) 105 105 Fig. 4 . 12 Powder HXRD of PLZT x/90/10 ceramics with compositions x=0.00, 0.020, 0.025, 0.030, 0.031, 0.032, 0.033, 0.035, 0.038, 0.040 Thus our results clearly show that the two phase region extends from x=0.0 20 to 0.0 33. With even higher La content (0.035≤x≤0.040), the FE type reflections disappear Fig. (4.13). Further, the superlattice reflection of FE phase also disappears and the (111) reflection of FE phase type reflection becomes a singlet whereas (200) is now a well splitted doublet, as can be seen from Fig. 4.13, suggesting the stabilization of AFE phase. Previously, the XRD data of PLZT x90/10 composition with x=0.0 32 was regarded as the onset composition for the FE phase. According to these results, the stabilization of AFE phase may be extended for the composition PLZT 3.5/90/10. Interesting, the composition with x=0.038 and x=0.04 show the AFE specific peaks, the characteristic splitting is 200 decreases wile (111 ) is broadened suggesting a modification of Pbam structure. 106 106 Fig. 4. 13 Evolution of the HXRD profiles of pseudocubi c (111), (200), and (211) reflect ions with compositions (x) for PLZT x/90/10 A quantitative confirmation for the above structural models is given by the results of Rietveld refinements for the entire composition al range. Thus, we have carried out a detailed Rietveld analysis on high resolution x-ray powder diffraction data of PLZT compositions across the FE/AFE phase boundary to settle the room temperature structure for all the compositions. The X-ray data were analyzed using the general structure analysis [48] program GSAS for Rietveld refinement. For the refinement, the initial values of the lattice parameters were obtained from our HXRD data by least squares method, whereas the values of the structural parameters were taken from literature [36, 49, 50]. The background was fitted with 9.75 10.00 11.2 11.6 13.8 14.4Intensity (arb.unit) 2 (degrees)0.040 0.038 0.035 0.033 0.032 0.031 0.030 0.025 0.020 0.000(211) (200) (111) x La 107 107 Intensity (arb.units)x=2.5 x=2.0x=0 5 10 15 20 25 30 2(deg.) 14.00 14.25(211) pc 11.4 11.7(200) pc 9.90 10.01(111) pcR3c Pbam2=1.23 Rwp=12.136 R3c Pbam2=1.24 Rwp=15.454 11.4 11.7(200) pc 14.00 14.25(211) pc 9.90 10.01(111) pc 5 10 15 20 25 30Pbam 14.00 14.25(211) pc 11.4 11.7(200) pc 9.8 10.0(111) pc 2(deg.)2=1.02 Rwp=11.16 R3c x=3.0d)c) * 2=0.9 Rwp=12.531 14.00 14.25(211) pc 11.4 11.7(200) pc 9.8 10.0(111) pc R3c *a) b) Fig. 4. 14 Observed (dots), calculated (continuous line), and difference (bottom line) profiles of PLZT (a) 0/90/10, (b) PLZT2/90/10, (c) PLZT 2.5/90/10 and (d) PLZT 3/90/10 obtained after Rietveld refinement using R3c and Pbam structural models. Vertical ticks b elow the peaks mark the position of the Bragg reflections. thirty sixth order polynomial, while the peak shapes were described by pseudo -Voigt profiles. In all the refinements, lattice parameters (a, b and c), positional coordinates (~x, y, z) and thermal parameters were varied , while o ccupancy parameters of all the ions were kept fixed during the refinement. The u se of generalized strain invariably led to improvement in the R factors, and thus, it was considered in all the refinements. The atomic coordina tes of Zr and Ti, Pb and La were only refined constrained at one position. 108 108 5 10 15 20 25 30 2(deg.)2=1.08 Rwp=18.653 14.00 14.25(211) pc 11.4 11.7(200) pc 9.90 10.01(111) pc R3c Pbamx=3.5 2=1.22 Rwp=13.023 14.00 14.25(211) pc 11.4 11.7(200) pc 9.90 10.01(111) pc R3c Pbamx=3.3 5 10 15 20 25 30 14.00 14.25(211) pc 2(deg.)2=1.07 Rwp=10.863 11.4 11.7(200) pc 9.90 10.01(111) pc R3c Pbam x=3.2* 2=1.22 Rwp=10.234 14.00 14.25(211) pc 11.4 11.7(200) pc 9.90 10.01(111) pc R3c Pbam **Intensity (arb.units)d)c) *a) b)x=3.1 Fig. 4. 15 Observed (dots), calculated (continuous line), and difference (bottom line) profiles of (a) PLZT 3. 1/90/10, (b) PLZT3.2/90/10, (c) PLZT 3.3/90/10 and (d) PLZT 3.5/90/10 obtained after Rietveld refinement using R3c and Pbam structural models. Vertical ticks below the peaks mark the position of the Bragg reflections. As shown earlier, the XRD profiles of FE phase show features similar to those expected for pure PZT 90/10, including the presence of a characteristic superlattice reflection on the left side of (111), which suggested oxygen octahedron tilt ed structure as for R3c. Accordingly, the rhombohedral R3c space group of PZT 90/10 was considered first in the Rietveld refinements for the FE phase. Fig. 4.14 a) depicts the observed, calculated and difference profiles for the refined structure. The experimental data are fitted almost perfectly based on a single R3c model for the composition PZT 90/10. This confirms that the PZT 90/10 is a single ferroelectric phase and belongs to the space group of R3c. It is reasonable to suppose that the intermediate FE phase existing in the ceramics is the same as PZT 90/10 based solid solution. In order to fully analyze 109 109 the nature of this phase(s), both the Pbam and R3c models are simultaneously applied in the subsequent refinements for the composition range. From Fig. 4.14 and Fig. 4.15 it can be seen that the R3c and Pbam space group yielded an acceptable fit with the diffraction pattern, as evidenced by the low value of Rp and Goodness of Fit ( GOF ) for the R3c model and examination of the from the magnified lower-angle region (see the inset). Therefore the mixture consisting of Pbam and R3c of symmetries are now confirmed by the theoretical analysis for intermediate composition. 5 10 15 20 25 302=1.47 Rwp=12.812 2(deg.)Pbam 14.00 14.25(211) pc 9.90 10.01(111) pc 11.4 11.7(200) pc 5 10 15 20 25 302=1.26 Rwp=15.094 2(deg.)Pbam 14.00 14.25(211) pc 11.4 11.7(200) pc 9.90 10.01(111) pcx=4.0 * *Intensity (arb.units)b) a) x=3.8 Fig. 4. 16 Observed (dots), calculated (continuous line), and differenc e (bottom line) profiles of (a) PLZT 3.8/90/10 and (b) PLZT 4/90/10 obtained after Rietveld refinement using Pbam structural model. Vertical ticks below the peaks mark the position of the Bragg reflections . Although the refinement of the structure of PZT with x꞊0.038 assuming as pure AFE Pbam phase gave acceptable ‘‘R’’ factors, the difference between observed and calculated profiles for pseudocubic reflection, like 200 and 211, is high (Fig. 4.16 a)). A higher value of Rwp (15.094) and GOF (1.26 ) together with the mismatch between the observed and calculated peak profiles may be observed. In addition, when compared with the composition PLZT 3.5/90/10 an anomalous broadening of the perovskite peaks is indicated . Similar results were obtaine d for the composition PLZT 4/9010 (Fig. 4.16 b) ). However, the rhombohedral space group R3c could be disregarded becaus e (111) superlattice reflection , which would be a characteristic R3c feature, is absent . The 110 110 assumption of a coexisting cubic phase in the refinement of the orthorhombic structure is however, physically unrealistic, since the AFE to PE cubic phase transition for this composition occurs at around 160 C. No additional reflections are observed and this indicates that the symmetry of the investigated materials may be lower on a local scale. Without any indication of a specific structure, it is difficult to explain these inconsistences. Whatever, it is obvious that the dominant phase is orthorhombic. 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0561562563564 390400410420430 Pbam volume PbamIIIIIR3c volume La (at. %) R3cI Fig. 4. 17 Variation of ferroelectric phase volume ( R3c) and antiferroelectric phase volume (Pbam) with La content for PLZT x/90/10 compositions (the unit size of the volume is A˚3 ) Therefore, Rietveld refinement analysis suggested that the FE phase has an a−a−a− tilt system (Glazer notation [52]). By analogy, it is proposed that the AFE orthorhombic phase has an a0b−b− tilt system. These results are in agreement with the tolerance factor determined at the beginning of this chapter. 111 111 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.00.000.250.500.751.00weight volume fraction (%) La (at. %) R3c Pbam Fig. 4. 18 Variation of rhombohedral R3c and orthorhombic Pbam phase fraction with increasing La amount for PLZT x/90/10. The continuous and discontinuous variations of the crystalline structures are better reflected by the volume variations as a function of La composition. Both the unit volumes of R3c and Pbam symmetries varies in a similar way as function of La content. Three regions can be identified: (I) a region where both volumes of orthorhombic and rhombohedral phases decrease when x increases, in the range 0≤ x≤0.03. This downward trend in lattice constants seems to stop at x = 3%, where a sudden increase is observed for the range of 0.03≤ x≤0.035 La composition: (II) region . For the composition with x≥0.038 the volume decreases again: (III) region. Since the radi us of La3+ (1.36 A˚) is smaller than of Pb2+ (1.49 A˚), the orthorhombic cell volume decreases with the introduction and increase of the La amount from x=0.020 up to x=0.030. In addition, the magnitude of antiparallel displacements of Pb2+ ions can be diminished since the radius for Pb2+ to displace is reduced by the contraction of the pseudocubic unit cells. This change of the lattice parameters may reflect a change in the crystalline structure, as it will be confirmed by the evolution of saturation polarization and piezoelectric constant as a function of La content. Discontinuities in the dielectric and piezoelectric properties of ferroelectric materials are indications of phase changes . It is worth to mention that Breval et al. [4, 5] repo rted two different possible orthorhombic structures at room temperature: one for La3+ content below 0.04 and another one for La3+ content higher or equal to 0.04. O rthorhombic lattice parameters for our compositions x=0.038 and 0.040, may be also an evidence of a different type of orthorhombic structure. In a recent spectroscopic 112 112 study [12, 16], the symmetry of PLZT 4/90/10 and PLZT 10/90/10 compositions was explained by either disordered rhombohedral phase, a monoclinic group or an orthorhombic one. Thus, systematic doping to replace the Pb3+ with smaller transition metal or RE dopants in the present case La3+ across the FE-AFE boundary not only may distort the cation spacing between the octahedra, which in turn could lead to rotations of the oxygen octahedra, but also may alter the long-range FE order, yielding to complex structures and ordering phenomena. Fig. 4. 19 Schematic illustration of the rotation (left) and in phase tilting (right) effect of TiO 6 octhaedra where Pb/La i ons are in the center of octahedra. Not only the lattice parameters but also the relative amount of FE/AFE and phases presented in the investigated compositions is expected to highly influence the FE and AFE properties. Variation of the rhombohedral and orthorhombic phase fraction with La concentration, in the composition range 0.000≤ x≤0.040, is shown i n Fig. 4.19. As expected, the fraction of rhombohedral phase decreases monotonically with increasing La concentration on the expense of the AFE phase one. The phase fractions of the rhombohedral and orthorhombic compositions are nearly equal for the range 0 .020≤ x≤0.025 of La composition. It is worth mentioning that in addition to the main perovskite peaks, the presence of many additional superstructure reflections may be observed. Three possible mechanisms may generate superlattice reflections in stoichiometric perovskites: cation ordering, octahedral tilts and 113 113 antiparallel cation displacements [51 ]. The cation ordering is ruled out for low amount of Ti in PZT [38]. One of the mechanism prese nt in different FE perovskites is the distortion from their cubic prototype symmetry by rotating or ‘tilting ’ of the corner-shared oxygen octahedra that form the ABO 3 perovskite structure. When an octahedron is tilted around one of the cubic [001] directions, the four adjacent octahedra in the plane normal to the tilt axis are constrained to tilt in opposite senses (Fig. 4.20). If the tilt is in the same sense as the central octahedron, the tilt is described as being ‘ in-phase’. If the tilt is in the opposite sense as the central octahedron, the tilt is defined as being ‘antiphase’ [38, 39]. Fig. 4. 20 Schematic illustration of the antiphase tilting (right) effect TiO 6 octhaedra obtained from the a ntiferroelectric Pbam phase for the composition PLZT 4/90/10 Pb/La Ions are in the center between octahedra while Zr/Ti are inside of octhaedra. Previously, we demonstrated that at room temperature, the ferroelectric phase of PLZT x/90/10 compositions has R3c structure. This structure contains a network of corner-sharing oxygen octahedra, where Pb2+/La3+ cations are inside the octahedra and Ti4+/Zr4+ cations fill in between the cavities with a−a−a − tilt system as per the Glazer notation [52] which gives rise to a rhombohedrally distorted cell along the (111) pseudocubic direction. The cations displace off their center of symmetry along (111) direction with the anion sublattice distorted and are rotated in antiphase around the rhombohedral axis. The a−a−a− tilt system give rise, according to the simple perovskite lattice, to two distinct types of weak superlattice reflections at 1⁄2{hkl}p and 1⁄2{0kl}p . The 1⁄2{hkl}p arises from rotations of the octahedra in antiphase while 1⁄2{0kl}p from cation displacement [52]. This results in a doubling of the unit cell along the c axis and, also, due to 114 114 successive rotations of the nearest-linked octahedra in the (001) p plane, doubling of the a and b unit-cell parameters [51 ]. The AFE phase has an orthorhombic structur e Pbam with a0b−b− tilted system. Fig. 4.20 and Fig. 4.21 depict the orthorhombic Pbam structure and the rhombohedral R3c structure drawn based on the results of refinement for the PLZT compositions . Fig. 4.20 is a schematic illustration of the inphase tilting of TiO 6 octhaedra while the Fig. 4.21 shows the antiphase tilting of AFE Pbam phase of PLZT. Despite the presence of some superlattice reflections which may be present in R3c and Pbam structure , the corresponding model used in Rieveld refinement is not enough complete to account for all the observed superlattice reflections. Therefore, these satellite reflections are convincing evidence that the investigated compositions are not purely rhombohedral and orthorombic. TEM studies performed on PLZT 2/90/10, 3/90/10, and 4/90/10 compositions revealed that such satellite reflexions may be interpreted as evidence of a long period ordered incommensurate phase [13, 14, 18, 27, 28]. This phase is formed by the disorder induced by La3+ doping which frustrates the formation of simple FE or AFE phases in this compositional range and promotes the formation of some long period ordered phases (2-3 nm) of unclear origin. The analysis of the ceramics within the whole composition range reveals that the appearance of the superlattice reflections is connected with the La content. For PZT 90/10 ceramics, no extra superlattice reflections are seen. Once the 2 % at. La amount is added on PZT 90/10, the supperlattice reflection may be observed and their intensity tend to increase with increasing La3+ up to 3.5 at. % concentration and then, it starts to decrease. This indicates that the replacement of Pb2+ ions by La3+ increases the driving force for an antiparallel shift of Pb2+ ions. The occurrence of these reflections is most probably driven by the disturbance of translational symmetry, similar with La doped AFE solid solutions [53 ]. It is worth mentioning that the multi peak structure of the reflection (211) in the inset of Fig. 4.14а)-d) may propose not only an orthorhombic but also a monoclinic distortion of the perovskite cells in the AFE state of the studied solid solutions. The structure of (211) peak definitely has a triplet hidden shape. Recent reports concerning the symmetry of compositions on the Zr rich side of the MPB [54] postulated that the crystal structure at ambient temperature was not rhombohedral, but monoclinic, since Rietveld refinement gave a better goodness of fit with Cm phase than for R3m or R3c symmetries . On the other side D. I. Woodward et al. [38 ] proposed the formation of two intermediate phases when PZT system transforms as a function of composition from FE rhombohedral R3c to orthorhombic AFE Pbam structures. At room temperature (ambient microscope) the phase sequence is 𝑅3𝑐 → 𝑃𝑐 → 𝑃𝑚 → 𝑃𝑏𝑎𝑚 while at lower temperature (20 K) the sequence is 115 115 R3c→ Pc→Pbam. Pm is a subgroup of R3m and Pbam , while Pc is a subgroup of Pm, Pbam and R3c. Pc and Pm phases develop only locally, at nanoscale level, and may not be observed by XRD. Similar with our results, the volume fraction of the intermediate phase increases as the Zr (in our study La) concentration increases ( x increases) until they are of sufficient size and number to stabilize the AFE Pbam cell [38 ]. Thus, the anomalies observed previously in our investigated PLZT x/90/10 compositions (Fig. 4.17 and Fig. 4.18) may be an evidence of such a phase transformation sequence. In addition, an enhancement of dielectric and piezoelectric properties and a region of temperature stability will be observed later from dielectric data for the same composition range, which supports the idea of a development of MPB. Recently, other symmetries were also proposed for PLZT 100x/90/10. The Cc monoclinic symmetry group was proposed by Solanki et al. [55 ]. According to Breval et al. , PLZT compositions with La x=0.02 and 0.04 have orthorhombic symmetry Pbam and C2mm groups respectively, at room temperature [4, 5 ]. From spectroscopy results published by ref. [16] the structure of PLZT (PLZT 4/90/10 and PLZT 10/90/10) compositions could be explained by either a disordered rhombohedral phase, a monoclinic group or an orthorhombic one. However, the presence of additional reflections at the AFE–FE phase boundary of the PZT phase diagram may give rise to multiple interpretations and the structure of PLZT 90/10 compositions still remain controversial. Additional investigations are required to understand the origin of superstructural reflections. In order to clarify the origin and the characteristics of the superlattice reflections in the phases of PLZT compositions, electron diffraction combined with TEM dark-field experiments were performed . 4.5 Domain structure and local characterization by TEM Knowledge about the way as La addition influences the perovskite tilt systems of PZT 90/10 is important to clarify the structur al symmetry of the investigated compositions. Small structural distortions which give rise to weak superlattice reflections may be difficult to observe by XRD . Electron diffraction has the advantage to be sensitive to superlattice reflections which 116 116 Fig. 4. 21 (a) Bright Field TEM image of PLZT 3.2/90/10 specimen showing a multi -domain grain with alternating AFE/FE confi guration . Domains 1, 3, 5 present both the incommensurate spots 1/n(a*+b*) with n=8-9 in the <001> pc ZA and ½(ooo) superlattice reflections in the <011> pc while for domains 2, 4 both features are missing. (b) Bright Field and (c) Dark Field TEM images showing the nanostructure streaking inside the AFE domains perpendicular to the direction of the satellite spot s. arise due to weak or short-range effects. Furthermore, in order to understand the different phase transformations associated with these compositions, a detailed knowledge of the domain structure and the nanostructure of the phases is an indispensable prerequisite. Thus, Transmission Electron Microscopy (TEM) technique was employed in order to investigate in detail the changes in domain and crystal structure across AFE/FE phase coexistence region for several key compositions PLZT x/90/10 compositions with x=0.025, 0.030, 0.032, 0.033, 0.035, 0.040. This part of the study was realized in collaboration with Alexandra Neagu and Prof Cheuk -Wai Tai from Department of Materials and Environmental Chemistry, Stockholm University, Stockholm, Suède. Previously, XRD results have suggested for 0 ⩽x<0.020 a r hombohedral FE phase, for 0.020 ⩽x <0.038 a phase coexistence region of orthorhombic and rhombohedral phases and for x⩾0.038 an orthorhombic 117 117 AFE phase. TEM analysis has revealed that the samples have more complex crystallographic features in the all investigated compositional interval. Mainly, bright field TEM images evidence that the all investigated compositions have a mixture of FE and ferroelastic domains typical of the ferroelectric PZT 90/10 and AFE PZ structures respectively [12, 14, 18, 28, 38, 51, 56-59]. Ferroelectric and ferroelastic domains show ½˂ooo˃ superlatice spots which indicate an R3c and Pbam antiphase tilt structures, in agreement with the previous XRD results. In addition, incommensurate spots along fundamental reflection of AFE phase are observed indicating that the AFE Pbam phase of PLZT x/90/10 is similar with one of the PZ system [12, 14, 18, 28, 38, 51, 56-59]. Fig. 4.22- a) is a cross-section TEM image of a representative PLZT 3.2/90/10 composition . From the previous studies , this composition is expected to reside at the FE-AFE phase boundary. The PLZT 3.2/90/10 ceramic displays a domain structure with a checkerboard-like pattern (Fig. 4.22-a)), which is typical for AFE state. Higher magnification reveals a modulated fine structure with regular fringes at a periodicity around 2 nm within the AFE domains. In addition, regular FE domain stripes were found in the PLZT 3.2/90/10 ceramic within the same grain (Fig. 4.22-a)). Domains denoted with odd numbers (1, 3, 5) are AFE, while domains denoted with even numbers (2, 4) are FE. The presence of satellites spots in the DP pattern of AFE domains is always accompanied by the presence of stripes in the BF and DF images, perpendicular to the direction of the spots as seen in Fig. 4.22-b) and (c). There are regions with the cell-doubled orthorhombic phase, as shown in Fig. 4.22 a) (in the left side). This structure is linked to the macroscopic AFE behavior. Similar patterns could, however, be obtained for the other PLZT x/90/10 investigated compositions. This indicates that these compositions present a two-phase mixture, where antiparallel cation displacements (most likely AFE clusters) are found in the fundamental FE matrix. Based on these results for the investigated samples, it can be concluded that La-dopant impurities induce a competition between the AFE and FE ordering due to the disruption of long- range dipolar interactions. a) Antiphase tilting of FE phase Previously, from XRD analysis , it was established that PLZT x/90/10 has a FE phase with R3c symmetry. Usually, R3c rhombohedral phase of PZT is characterized by cation displacements along ˂111˃p combined with antiphase octahedral rotations, denoted as a−a−a− [52]. This leads to a doubled unit cell in all three directions. This structure has previously revealed the existence of superlattice 1/2(hkl) and 1/2(hk0) (where h; k; l=2n+1). As suggested before, the ½(111) 118 118 superlattice spots are resulted from the oxygen octahedra tilting while the ½(011)-type superlattice spots are mainly attributed to the antiparallel Pb2+ displacements [38, 39 ]. Bright-field imaging revealed normal micron sized domain patterns at room temperature in all the investigated compositions which were characteristic of ferroelectricity (Fig. 4.22a)), domains denoted with 2 and 4). Diffraction patterns obtained from PLZT x/90/10 samples on FE zones and indexed with zone axes ˂011˃p are shown in Fig. 4.23. The strong reflections are indexed according to the simple perovskite lattice. The pattern of FE phase for all the investigated compositions exhibited superstructure reflections in electron-diffraction patterns at 1⁄2˂ hkl˃ positions consistent with rotations of the octahedra in antiphase around the ˂111˃p axis. The pattern is consistent with ferroelectric R3c space group. The observed superlattice reflection at 2θ=45o is ½ (311) which is allowed for the a−a−a− tilted structure. These observations are in agreement with previous XRD results and with those reported in the literature for similar compositions [2 ]. Fig. 4. 22 . Diffraction patterns obtained from PLZT x/90/10 samples indexed with zone axes ˂011˃ P which evidenced the superstructure reflections arising from antiphase tilting of R3c structure for PLZT composition a) 2.5/90/10, b) 3/90/10, c) 3.2 90/10, d) 3 .3/90/10, e) 3.5/90/10 and f) 4/90/10 . a)b) c) d) e) f)010 10-1000 000 000 000000010 10-1010 10-1000 010 10-1 10-1010010 10-1 119 119 b) Antiphase tilting of antiferroelectric phase Usually, the AFE phase on high Zr side of the phase diagram of the PZT is similar to the PZ structure with Pbam symmetry [12, 14, 18, 28, 38, 51, 56-59 ]. The AFE phase has quadrupled unit cells. The quadrupling arises from the antipolar order of the A-site cations along ˂110˃p directions. In addition, ½(ooo) superstructure reflections are observed, consistent with the a0b−b− tilted system associated with the antiphase rotation of octhahedra around the ˂110˃p cristalographic plane within PZ-like structure [12, 14, 18, 28, 38, 51, 56-59]. All the investigated samples exhibit a typical lamellar ferroelastic domain structure (100– 200 nm) consistent with previous studies of orthorhombic (AFEo) phase of PZ structured ceramics [12, 14, 18, 28, 38, 51, 56-59]. This indicates that the tilting of the FE octahedra has disappeared and the system Have in his composition only AFE. This is attributed to the effect of excessive La3+, which has destroyed the polar interactions between the Ti3+ ions, and thus any advantage gained by a change in symmetry order is frustrated due to the high concentration of nonpolar dopant cation [60 ]. Fig. 4. 23 Diffraction patterns obtained from PLZT x/90/10 samples indexed with zone axes ˂010˃ P which evidenced the superstructure refle ctions arising from antiphase tilting of Pbam structure for PLZT composition a) 2.5/90/10, b) 3/90/10, c) 3.2 90/10, d) 3.3/90/10, e) 3.5/90/10 and f) 4/90/10 a) b) c) d) e) f)000 001100 000100 001000100 001 000100 000 001100 001100 000 001 120 120 c) Incomenssurate phase As the La3+ content increases from x˃0.02 toward x=0.04, the diffraction data become highly complex. In addition to the appearance of circled 1/2 ˂011˃ spots, there are additional reflections along the ˂011˃ direction arrowed satellite spots appearing next to the main reflections. The weak satellite spots around the principal reflections indicate long period ordering of the crystal structure and are due to the modulation of the basic perovskite structure. These features are evidence of the presence of incommensurability in this composition [28 ]. A similar incommensurate phase was reported for PLZT x/95/5 and is believed to arise due to a competition between "broken" ferroelectric dipoles and AFE sublattice [58]. Similar with other reports [1, 14, 18, 27, 58 ] these reflections may be associated with AFE incommensurate phase. Electron transmission electron microscopy (TEM) investigations of PLZT x/90/10 allow as observing how this phase ch ange as function of La content. Fig. 4. 24 . Selected area dif fraction patterns obtained from PLZT x/90/10 samples indexed with zone axes ˂001˃ which evidenced the1/4(110) superattice refle ctions arising from AFE incommensurate phase for PLZT composition a) 2.5/90/10, b) 3/90/10, c) 3.2 90/10, d) 3.3/90/10, e) 3.5/90/10 and f) 4/90/10 ˂001˃pc010100a) b) c) d) e) f) 121 121 Diffraction patterns which show nanosized domains of an incommensurate modulated structure taken from all the investigated compositions are shown in Fig. 4.25 . The TEM study revealed that all the investigated compositions contained grains that in the SAED patterns showed satellite spots along <110>p directions with a periodicity of 8-9 (110) spacing’s and a modulation wavelength of ~24.4 Å for a pseudo-cubic unit cell of 4.15 Å. These satellite spots correspond with a long-period ordered incommensurate AFE structure. The presence of incommensurate reflections is always associated with the presence of AFE domains. For lower La -content (x=0.025, 0.030) grains that do not contain the incommensurate AFE phase at all have been found , suggesting chemical inhomoge neity. On the other hand, for higher La -content (x > 0.030) all the investigated grains presented the incommensurate modulated spots. For these compositions , most grains presented a multi -domain configuration with alternating AFE -FE domains . The incommensurate modulations spots along the <110>p shown in the SADPs can be expressed as: ha*+kb*+lc*±1/n(a*+ b*) where h, k, l, are integers and n can be directly measured from the diffraction pattern. The examination at higher magnifications reveals two sets of incommensurate modulation stripes along orthogonal ˂110˃ planes (Fig. 4.25-b )) for sample PLZT 3/90/10. Selected area electron diffraction along the ˂001˃ zone axis shows the corresponding two sets of satellite spots surrounding fundamental reflections (inset in Fig. 4.25-b )). In this figure, 1/4(110) superlattice reflections can be seen along two dimensions. The two-dimensionality may arise due to neighboring AFE domains which have the same (001) orientation but are rotated by 90 C with respect to each other. Similar structures have also been noted in Sm-doped BiFeO 3 [60]. Therefore, TEM investigation evidence that the room-temperature structure of PLZT x/90/10 AFE phase is similar with one of PZ, which is an incommensurate orthorhombic AFE [59]. It is often reported that this incommensuratephase a bridging phase which forms in materials with compositions at the phase boundary between FE and AFE [18 ]. These results are in agreement with the recent ones reported in ref. [18], in which it was shown that lanthanum doping of zirconium rich lead zirconate titanate gives rise to incommensurate, long-period AFE structures. According with their study, the incommensurate phase is FE [18]. The Pb atom displacements of about 28 pm and peak polarization values of about 60 μC cm−2 are in agreement with to reported values of polarization of the commensurate AFE PZ phase [29 ]. Therefore, La modification is 122 122 found to suppress the long-range AFE order by breaking apart the micrometer sized AFE domain into an ensemble of nanodomains. It is worth mentioning that, in addition to the previous mentioning structure s, domains with different type of structure were observed. The pattern of such structure is shown in Fig. 4.26. Fig. 4. 25 Selected -area electron diffraction obtained from composition PLZT 3.5/90/10 According with D. I. Woodward et al. [38, 39] this structure may be interpreted as being in phase tilting. This phase was observed for composition with x≥0.033 only in few domains. This structure is in high agreement with the assumptions of an in phase tilting expect when La content approach es 4 at.% at. Furthermore, an increase of the piezoelectric coefficients and remanent polarization will be later observed for the same compositional range, which indicate that the additional phase may be FE. Therefore the mismatch obtained from Rietveld refinement, especially for the composition with x≥0.035, and the anomalies obtained from lattice parameters may be explained with the presence of an additional phase with the main in phase tilt ed structure. In conclusion, PLZT x/90/10 compositions have a complex nanoscale domain mixture showing grains with alternating FE/AFE domains with features which correspond to R3c and Pbam structure . A higher amount of R3c phase was found in compositions with low La amount, while the Pbam phase is most prevalent for higher L a content. The FE phase has a-a-a- tilt structure. The primary feature of AFE phase is the presence of the 1/2˂ooo˃ superstructure reflections, consistent with the a0b-b- tilt system with are typically associated with antiphase tilting of oxygen octahedra and confirms that the superlattice structure of AFE phase is identical to one of PZ. Therefore, the local symmetry of FE/AFE phase s revealed by electron diffraction is consistent 123 123 with the overall symmetry previously obtained from x-ray diffraction analysis. Similar as reported in other woks [13, 14, 18, 27, 28], the La doping frustrates the formation of simple FE or AFE phases and promotes long period ordered phases (2-3 nm) with ordering along <110> with unit cells incommensurate with the primitive cubic unit cell. 4.6 RAMAN analysis 4.6.1 Introduction Raman spectroscopy is a powerful technique, which gives information about the vibrational modes of the atoms and molecules, being sensitive to the chemical composition, ionic bond and the local crystalline structure. This technique has been extensively used in the study of the phase transition in ferroelectric oxides [61, 65 ]. The subtle changes in bond distance, bond angle and crystal symmetry during the phase transitions can be detected by the evolution of different modes and changes in the mode frequency, peak intensity and width . It is known that the Raman spectr um of pure lead zirconate presents three main frequency regions, as reported in refs. [66, 67 ]: -a low frequency region from 0 to 180 cm−1, which corresponds to the lattice vibrations and bonds between Pb cation and other ions in the unit cell. This frequency range contains lines of high intensity and involves phonons corresponding to the lowest-frequency optic and acoustic branches activated by the antiferrodistortive transition, leading to the AFE structure [66]. In tetragonal samples the E(TO1) and A1(TO1) modes at frequencies similar to those modes in PbTiO 3 (∼80 and 150 cm−1, respectively) were observed [67 ]. -a medium frequency region between 200 and 450 cm−1, (Slater type) which shows internal modes related to certain polyatomic groups of the material. The modes are originated from the Zr and Ti atoms and reflect the Pb/La(Zr,Ti)O 3 vibrations; the vibration OZr/TiO is represented by the modes between 175 and 300 cm-1; the modes between 300 and 400 cm-1 are related to the torsion and distortion of (Zr,Ti)O. In tetragonal samples three main bands correspond to the E(TO2) mode, the B1 + E(TOs) peak stemming from the cubic silent F2u mode, and the A1(TO2) mode were observed [67 ]. -a high frequency which ranges between 400 –800 cm-1 is related the frequencies vibrations due the stretching of (Zr,Ti)O: the weak longitudinal optic modes E(LO2) and A1(LO2) modes 124 124 are located between 420– 450 cm−1, around 500 cm−1 contains the Zr-O elongation, In the band between 500 and 650 cm−1 contains the E(TO3) and A1(TO3) modes which are related to a mixture of vibrations of oxygen and B-site atoms. Above ∼700 cm−1, the E(LO3) and A1(LO3) modes with a possible contribution of the nonpolar oxygen breathing mode were observed [67 ]. All these distortion and elongation bands are present in the spectra corresponding to the PZT doped samples with a shift in frequencies. FE phases can be distinguished by the number of Raman-active modes. As example, for a full B-site occupational disorder and an effective, average Zr/Ti atom, the PbTiO 3-like (Z = 1) tetragonal FE phase has 3A1 + 4E + B1 Raman-active modes; the (Z = 1) rhombohedral FE phase has 3A1 + 4E Raman-active modes plus the silent A2; and the (Z = 1) monoclinic FE phase has 7A, + 5A [67]. Results and discussions Raman spectroscopy realised in collaboration with dr. Marco Deluca from Univ. of Leuben, Austria, was employed to complement the XRD and TEM investigations. In the following, a systematic study of PLZT x/90/10 by using Raman spectroscopy in order to evaluate the coexistence of FE and AFE phases and the ir evolution as a function of La content at room temperature will be present ed. First, the assignment of the Raman modes observed in the spectra at room temperature is discuss ed. The labelling of the Raman modes was done according to that proposed for tetragonal lead lanthanum zirconium titanate (PLZT) and PbTiO 3 [67 ]. The assignment of the AFE bands w as done on the basisof the reported ones for P Z single crystals [68]. Fig.4.27 and Fig. 4.28 show the Raman spectra for the studied compositions corrected for the Bose–Einstein temperature factor at various temperatures in the frequency range 25– 800 cm-1. This frequency region covers mainly the lattice modes of the material. Consistent with the results of x-ray diffraction and TEM, reported in the previous sections, PLZT x/90/10 ceramics display different active Raman modes depending on the La content. As shown in Fig. 4.27- c) the observed spectral features are numbered consecutively from 1 to 17, starting from the low-wavenumber region. The modes in the low frequency range (below 200 cm-1) are more focused because most active Raman modes for FE rhombohedral and AFE orthorombic were previously found in this range (4.28 -c)). Peaks 5, 6, 10 and 11 are found in all the FE phases of PZT-based materials, while peaks 1-4, were found predominantly i n the AFEs . The first mode, at 25 cm-1, is absent in AFE systems without La such as PbZrO 3, PZT-95/5, and PZTN-95/5/l. This mode has been reported previously in AFE PLZT 95/5 composition and it was interpreted as due to an intermediate high 125 125 temperature FE phase [69]. It is called the soft mode of AFE materials. This central mode is found to be responsible for the strong dielectric anomaly presented at Curie temperature (T C). The origin of the central mode arises from lattice disorder caused by strongly anharmonic hopping of Pb ions [70]. Thus, the presence of this mode in our compositions is plausible, since the La replaces the Pb on the A site and induces disorder in the lattice. 0100 200 300 400 500 600 700 800Intensity (a.u.) Frequency (cm-1) 4 3,8 3,5 3,3 3,2 3,1 3 2,5 24 Ka) 0100 200 300 400 500 600 700 800b)75 KIntensity (a.u.) Frequency (cm-1) 4 3,8 3,5 3,3 3,2 3,1 3 2,5 2 0100 200 300 400 500 600 700 800c) 71716 1514 1312 11109865 43 2100 K 4 3.8 3.5 3.3 3.2 3.1 3 2.5 2Intensity (a.u.) Frequency ( cm-1)A,+2A,,A,+A,, A,1 0100 200 300 400 500 600 700 800d)300 KIntensity (a.u.) Frequency (cm-1) 4 3,8 3,5 3,3 3,2 3,1 3 2,5 2 126 126 Fig. 4. 26 . Compositional dependence of the Raman spectra of PZT ceramics at 4 K, 75 K, 100 K and 300 K, respectively corrected from the Bose-Einstein factor. The proposed labels of the vibrational modes are in agreement with ones of the PbTiO 3. Raman features are numbered consecutively from 1 to 17, starting from low wavenumbers (see text). In addition, the appearance of this mode in our PLZT x/90/10 composition at this temperature may indicate that this phase is shifted to lower temperature or that another FE phase is present . It implies that the PLZT x/90/10 compositions may have, in addition, a local FE structure at room temperature, other than rhombohedral R3c and orthorhombic Pbam ones. In fact, in addition to Pbam and R3c XRD specific peaks, another shoulder was observed on the right side of (111). This shoulder may be related to this additional FE phase. The modes 2, 3, 4 and 6 cm-l can be assigned for cation (ZrO 2) lattice modes [67 ]. These are polar modes of AFE phase and are typically reported in the spectrum of PZ [71], which indicate that the AFE phase of PLZT x/90/10 is incommensurately modulated. The mode 3 may be present in both FE and AFE compositions. The mode 5 at 80 cm−1 at room temperature is presented only in FE samples (PLZT x/90/10 with x≤2). According with E. Buixaderas et al. [12, 16] this mode is connected with tilt of the oxygen octahedral and should be the antiferrodistortive soft mode of the FE phase. The band at ~60 (mode 3) and ~80 cm−1 (mode 5) frequency were assigned to the E(TO1) mode in tetragonal samples [67]. The mode 3 called AFD is expected to be the main component of the soft mode of the FE transition at T C. The modes no. 7 (at about 136 cm-1) originate from the off-center lead displacements, and are responsible for the spontaneous polarization Ps and the spontaneous strain presented in FE compositions. The mode no. 7 is linked with the transition from AFE (orthorhombic) to FE (rhombohedral) phase transition at 234°C in PZ single crystal [72 ]. This peak has a dominant A1(TO) character and is split in two components in the PLZT x/90/10 with x≤0.02. Modes 8 and 9 (204 and 250 cm-1) have been associated with Zr– O bonding; the bands corresponding to the modes 10, 11 and 12 (at 285, 340 and 390 cm-1) have been assigned to ZrO 3 torsions, while modes 13 and 14 (at 501 and 532 cm-1) are due to the Zr–O stretching. The assignments of these bands was realised by considering the reports for PbZrO 3 single crystals [68 ]. The mode 10 is called the ‘silent ’ mode and can become active owing to the local cubic symmetry breaking due to strong dynamic fluctuations [73 -75]. The mode 11 lies at a similar frequency to those calculated for the BO 6 rotation vibrations. Raman mode near 240 cm−1 can be related to the monoclinic or rhombohedral phases [76]. The small peak at 680 cm−1 observed for the same compositions (x≤2.5) might be connected with the localized breathing mode of the oxygen 127 127 octahedra [64, 77]. Above ∼700 cm−1, we observe the E(LO3) and A1(LO3) modes with a possible contribution of the nonpolar oxygen breathing mode [64 ]. 025 50 75100 125 150 175 200a) Intensity (a.u.) Frequency (cm-1) 4 3.8 3.5 3.3 3.2 3.1 3 2.5 24 K 025 50 75100 125 150 175 200b)75 KIntensity (a.u.) Frequency (cm-1) 4 3,8 3,5 3,3 3,2 3,1 3 2,5 2 025 50 75100 125 150 175 200d)300 KIntensity (a.u.) Frequency (cm-1) 4 3,8 3,5 3,3 3,2 3,1 3 2,5 2 025 50 75100 125 150 175 2007100 K 4 3.8 3.5 3.3 3.2 3.1 3 2.5 2Intensity (a.u.) Frequency (cm-1)654 321c) 128 128 Fig. 4. 27 . Compositional dependence of the Raman spectra in the low frequency band of PZT ceramics at 4 K, 75 K, 100 K and 300 K corrected from the Bose-Einstein factor. The proposed labels of the se modes was realised according to ones reported for PbTiO 3 [67] This mode is not a La-driven A-site vibration as its frequency is too high, but it is likely related to the presence of Ti and Zr vacancies promoted by the La addition. These vacancies affect locally the vibrations of the surrounding oxygens. Since the Raman peaks that represent the FE and AFE phases in PLZT ceramics were identified, we can now proceed with the fitting of the PLZT compositions at 100 K ( x from 2 to 4). The exact frequency of these three modes in different compositions after deconvolution is plotted in Fig. 4 .29. Fig. 4. 28 . The evolution of the frequency for the characteristic Raman modes as a Function of La content in PLZT x/90/10 (100K) . 129 129 The peak location and intensity is a function are functions of the types of bonds present in the material and this may help in understanding better the structure of PLZT x/90/10 compositions. The spectral modifications associated to the crossover from the rhombohedral to orthorhombic phase are clearly visible. All the spectra broadly follow a similar non-monotonically trends, as observed previously in XRD analysis. The main differences between the samples appear especially in the low-frequency part, below 200 cm−1. There is a distinct difference in pattern when increasing the La addition from x=0.02 to x=0.04, as following: – The spectrum of PLZT 2/90/10 and PLZT 2.5/90/10 at low temperature (100 K) displays three bands for frequencies lower than 200 cm−1 at 52, 61 and 128 cm−1, typical of the rhombohedral phase. The presence of rhombohedral symmetry is indicated by the coexistence of peaks 1, 3, 5 (the so called ‘triple mode’, by the small peak 8 at 490 cm-1 [69], and by the absence of the overdamped soft mode [78 ]. The small peak at 680 cm−1 observed for the same compositions ( x≤2.5) is connected with the localized breathing mode of the oxygen octahedra while for x ≥3 no trace s of this peak are seen [16]. It may be of interest to notice that the spectra for this compositions show a splitting of all E-symmetry modes into A+ A species (indicated in Fig. 4.27- c)). This kind of splitting is linked with monoclinic phase of PZT with compositions at the MPB. This indicates that the ground state structure of these samples may be of lower symmetry [65, 78 ]. – Upon increasing La content from 0.025 to 0.035 , several changes in the low temperature Raman spectra (100 K ) are observed . The most prominent feature in the Raman spectra is the loss of strength of the FE soft mode near 80 cm−1, as expected if the amount of FE phase is reduced. The disappearance of m odes at 430, 530 cm−1 and 675 cm−1 are another indicati on of a change in phase. Mode s 3 and 7 characteristic to the FE phase progressively disappear , suggesting that the rhombohedral symmetry is gradually lost . A minimum in the soft mode at 80 cm-1 and the AFD mode signaled that the onset of FE phase takes at 3. 5%. In the meantime the presence of several modes at low frequencies (less than 150 cm−1) which are the characteristics of the AFE order may be noticed . The appearance of the AFE phase is likely at 3% rather than at 2.5% La addition , as the AFE mode at 115 cm-1 and the emergence of the mo des around 60 cm-1 and 125 cm-1 appear for this composition . However, on zooming the higher freq uency spectra , we may notice a small shoulder in composition s with x≤ 2.5 corresponding to the mode 2 , which indicate s that a small amount of AFE phase may develop in this samples. The mode 1 (25 cm-1) could be hidden in the spectra for compositions with x≤2 due to high damping of the vibrations . This is in agreement with 130 130 previously discussed HXRD results and Rietveld analysis. A mixed AFE–FE structure is clearly indicated in the 0.25 ≤x≤3.5 composition range. -The Raman spectrum characteristic to the 0.038 and 0.040 compositions is intriguing. These samples still preserve the features of an AFE phase, but it shows a strong enhancement of the intensity of the FE phase specific bands. In particular, the intensity of modes 3 and 11 are enhanced. Such changes can indicate a modification in the structure. It worth mentioning that that also XRD and TEM results suggested a different structure for these compositions, but not completely resolved yet. To better understand the complexity of the structure of PLZT x/90/10 , in addition to the room temperature Raman spectra, the evolution of all the specific bands as function of La content was monitored. As shown in Fig. 4.27 a)-c) and Fig. 4.28 a)- c), a noticeable change in the Raman lines is observed mainly at lower wavenumbers (below 200 cm-1) when increasing of temperature from 4 K to 300 K. An enhancement of the Raman signal below 100 cm−1 is found at low temperatures, in contrast to the behaviour at 300 K. PLZT 2.0/90/10 and PLZT 2.5/90/10 compositions show a high intensity band at 35 cm-1, which decreases in intensity and splits in two bands at 39 and 53 cm-1 with increasing temperature from 4 K to 100 K. At the same time, for 0.03≤ x≤0.04 composition range, the band with frequency of 61 cm−1 at 4 K decreases in intensity and split in two peaks at 57 and 70 cm-1 at 100 K. The 207 cm−1 band grows in intensity; the 245 cm−1 jumps to lower frequency (240 cm−1) and increases in intensity, the splitting and intensity of the modes at 490 and 515 cm-1 are increased. In addition, the intensity of typical AFE mode at 112 cm-1 is enhanced with increasing temperature. The discontinuities both in the number of bands and in the band frequencies and the changes in the band’s intensities point to a phase transition from a different one to an orthorhombic structure, as temperature increases from 4 K to room temperature . This may indicate a diffuse transition from FE to AFE phase which finished at around roo m temprature during heating. Such a behaviour is in agreement with the results of the A. Pelaiz- Barranco et al. for the composition PLZT 3/90/10 that show a transition from FE to AFE when increasing temperature from 25 to 80oC. It is worth mentioning that from 4 K to 300 K, the features of the bands follow a similar trend as those observed as function of composition with titanium content from 0.6 to 0.40 for PZT, which were interpreted as due to the monoclinic-rhombohedral phase transition [79 , 32]. The intensity of AFD mode peak decreases from temperature 7 to 300 K. At room temperature (300 K) this peak is evident only for compositions with x ≤0.025. The Raman intensity of the AFD mode, decreases from the rhombohedral to the orthorombic side. This means that either the tilted phase disappears in samples with titanium of 0.35 % or the proportion of the tilted phase diminishes towards the orthorombic side in case of phase coexistence [67] . 131 131 However, HXRD and TEM previously presented results indicate that R3c like structure is the dominant FE phase. We may think that R3m phase local clusters can form at room temperature. This transition will be further discussed in more details. Only some representative data are shown here in order to understand the complexity of the phase transition of PLZT x/90/10 as function of x at room temperature. The phase coexistence may occur over a relatively wide temperature region. Therefore, either the FE or the AFE phase could be stabilized at room temperature in each composition. It is useful here to carry out a more detailed analysis of Raman spectra, in order to gain information on the structur al change as a function of La composition. Fig. 4.30 reports the peak position and intensity, the AFD mode at ~ 60 cm-1 for the PLZT x/90/10 compositions. This is an additional Raman band appearing in this frequency range ( 65 cm−1) and is associated with the antiphase tilt vibrations of the oxygen octahedra [67 ]. In the composition range from 2 to 4 at % of La, changes in the amplitude of AFD mode were observed: with increasing La content from 2 to 2.5 at % the peak amplitude begins to increase, then in the composition range 2.5≤ x≤3.at. % rapidly decreases , then for 3.3≤ x≤3.8 at. % starts to increase and then decreases again for x≥3.8. We can observe that the frequency corresponding to the AFD mode progressively shifts towards higher value, showing anomalies in the middle composition al range. Such kind of behaviour clearly indicates the existence of more than one phase transitions, in relationship with the tilt of the perovskite structure . The changes in the intensity of AFD Raman peak slope are detected in correspondence of the anomalies obtained previously on the evolution of the lattice of R3c phase from Rietveld refinement, which confirmed the structur al changes. 2,0 2,4 2,8 3,2 3,6 4,00,150,200,250,300,350,400,450,50 525456586062 Frequency (cm-1) Amplitude (a.u.) La (at. %) 132 132 Fig. 4. 29 . Position (characteristic frequency) and integrated intensity (amplitude), of the ~60 cm-1 peak of the Raman spectrum of the PLZT x/90/10 compositions as a function of La addition . The spectra obtained for x≤0.025 and x≥0.030 are similar to those reported by other authors [12, 16]. Similarly, they support the premise that the local symmetry in PZT and PLZT near the AFE phase boundary is lower than currently described in the accepted phase diagrams. A second- order transition from monoclinic (MC) to orthorhombic phase occurring at low temperature was reported previously on Pb(In 1⁄2Nb1⁄2)O3-Pb(Mg 1⁄3Nb2⁄3)O3-PbTiO 3 single crystals [80 ]. This transition is connected with antiphase tilts of oxygen octahedral, in analogy with the observations of PZT system [80 ]. The MC phase is a quasi-orthorhombic, space group Pm bridging phase mediates the structural transition from rhombohedral to tetragonal phase [80 ]. A similar situation may be found in our PLZT x/90/10 compositions: PLZT x/90/10 c ompositions with low La (x≤2.5) have lower -symmetry structure , while high La contents are compatible with symmetry phase higher than orthorhombic but very similar (most probabl e monoclinic Pc or Pm). A MPB between low-symmetry and high symmetry phases is likely as La content increase between 0.02 and 0.04. Furthermore , there are indications of a systematic splitting of the E -symmetry modes into monoclinic A -A doublets similar with one in the morphotropic samples. However, Raman scattering does not give a straightforward evidence for ma croscopic monoclinicit y in PZT [67]. Further studies are necessary to sustain th ese results. In addition , La concentrations can contribute to creat e point defects: La creates two types of defects, the substitute ion itself and vacancies which can in principle be located at the A and/or B sites, although it has been reported that they are mainly formed on the A sites. T hese defects may influence the formation of FE domains . For example, the vacancies Ti /Zr in the lattice can affect locally the vibrations of the surrounding oxygen octahedra at high frequencies [81]. Therefore, a systematic study of the Raman spectra of PLZT ceramics in a wide compositional range allowed us to analyze fine spectral features reflecting the subtle structural changes across the FE-AFE phase boundary. The AFD mode was detected for the whole compositional range (x = 0.00– 0.04) and this confirms the presence of frozen oxygen octahedral tilts in both rhombohedral and orthorombic PLZT ceramics at low temperatures of 100 K. In addition, we noticed an evidence for a monoclinic symmetry, such as the splitting of the E symmetry modes into A -A pairs for composition with x≤0.025. It appears that the Raman spectra can be interpreted in agreement with the assignments of the structure derived from HXRD analysis. 133 133 4.7 Conclusions The structure of PLZT x/90/10 ceramics with compositions 0 .0≤x≤0.04 have been investigated by a combination of techniques: XRD, HXRD, TEM and detailed Raman spectroscopi es. The combined results gave a more complete picture of the phase structure and composition evolution around the FE/AFE phase boundary in PLZT 90/10. and allow ed to describe the macroscopic to local structure evolution across the FE-AFE phase boundary of PLZT x/90/10 and to postulate the plausible origin for superior electromechanical properties at this phase boundary. Increasing La content from 0 to 4 % in the PZT 90/10 system produces a crossover from FE long-range order toward AFE state, although short-range order effect was also observed for high x≥0.025 La content. A broad overview of the crystal chemistry and domain structure of La doped PZT is described below: (i) Dense PLZT ceramics with ure perovskite phase can be produced by conventional solid state reaction; ii) All the compositions with x≤0.025 have macroscopic rhombohedral R3c symmetry state and exhibit superstructure with orientational and translational domains characteristic of an antiphase tilted FE perovskite. However, these compositions develop locally AFE clusters with lower symmetry FE phase, as reveal by TEM and Raman spectroscopy; (iii) A phase coexistence region with the rhombohedral R3c and orthorhombic Pbam structure, is established for the range of compositions 0.025≤ x≤0.035 ; (iv) The orthorhombic Pbam structure which has a quadrupled unit cell, similar to PZ was stabilized at higher values of x≥0.035 , since the tolerance factor decreased; (v) Development of local phase with distortion of Pbam structure like in phase tilting may take p lace in compositions with x≥0.038; (vi) In addition, on cooling PLZT x/90/10 compositions, a diffuse tilting structural phase transition was observed b y in situ Raman spectroscopy which starts at low temperature (~4 K) and extends to room temperature, which indicates that this compositions may develop locally other phases at room temperature. Both FE and AFE states coexist in all the PLZT x/90/10-based solid solutions and the AFE state initiates even at small substitution of La, i.e. x=0.02 by the development of loca l AFE domains, which gradually increase with increasing La addition. This intricate phase mixture is 134 134 perceived to be easily perturbed by external fields, temperature or stresses, and thus may lead to interesting electromechanical characteristics. The structural changes are expected to impact greatly on the dielectric and FE behavior of PLZT x/90/10. The existence of such phase superposition may be the reason for the unusual macroscopic characteristics. Therefore, tailoring the phase coexistence, especially by controlling the amount of AFE phase would constitute a very useful method to produce PLZT-based solid solutions with enhanced energy storage properties. The results of the microstructure, the data from structural characterization and TEM results were included in five original papers: 1. I.V. Ciuchi, F. Craciun, L. Mitoseriu, C. Galassi, Preparation and properties of La doped PZT 90/10 ceramics across the ferroelectric-antiferroelectric phase boundary , J. Alloys Compd. 646 (2015) 16-22. (I.F.=3.014 , A.I=0.558) 2. F. Craciun, F. Cordero, I.V. Ciuchi, L. Mitoseriu, C. Galassi, Refining the phase diagram of Pb 1-xLax(Zr 0.9Ti0.1)1-x/4O3 ceramics by structural, dielectric, and anelastic spectroscopy investigations , J. Appl. Phys. 117(18) (2015) 184103 (1-8). (I.F. =2.101 , A.I=0.68) 3. I.V. Ciuchi, L. Mitoseriu, C. Galassi, Antiferroelectric to Ferroelectric Crossover and Energy Storage Properties of (Pb 1-xLax)(Zr 0.90Ti0.10)1-x/4O3 (0.02 ≤x ≤0.04) Ceramics , J. Am. Ceram. Soc. 99(7) (2016) 2382-2387. (I.F.=2.841 , A.I=0.70) 4. A. Neagu, I.V. Ciuchi , L. Mitoseriu, C. Galassi, C. -W. Tai, Study of ferroelectric – antiferroelectric phase coexistence in La -doped PZT ceramics, European Microscopy Congress 2016: Proceedings, Wiley -VCH Verlag GmbH & Co. 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All.Com. 582, 680–687 (2014) 142 CHAPTER 5 Room Temperature Electrical Properties 5.1 Introduction The La3+ content not only affects the structure and microstructure evolution of the PLZT x/90/10 ceramics, but it is also expected to play an important role on the electrical and electromechanical properties. In fact, a high interest to find compositions within the range of crystalline phase superposition is related to the possibility to find enhanced functional properties [1-4]. The related functional properties of PLZT x/90/10 ha ve been investigated by a limited number of authors and only for few compositions (2/90/10, 3/90/10 and 4/90/10) across the FE- AFE phase boundary [5-6]. High values for d33~ coefficient of 100 pm/V and room temperature stable piezoelectric properties were reported in some AFE materials with irreversible field induced AFE -to-FE transition [7]. However , the piezoelectric properties of PLZT x/90/10 composition have not been explored yet. It was reported that the coupling between AFE and FE order parameters sometimes associated with intrinsic defects is responsible for remarkable FE properties in PLZT x/90/10 materials [8 -10]. Therefore a detailed investigation of the dielectric properties and piezoelectric properties with composition range across FE/AFE border is necessary. The aim of this chapter is to show the low field electric al and piezoelectric properties of PLZT x/90/10 ceramics determined by using Impedance Spectroscopy and piezoelectric measurements, in a wide range of frequency at room temperature . 143 5.2 Dielectric Properties Figures 5.1 a)- c) compare the variation of real part of permittivity and dielectric losses with frequency at room temperature, while the composition dependence of permittivity and losses at a few selected frequencies is presented in Fig. 5.2 a)- b). The dielectric permittivity slightly decreases and the loss gradually grows with increasing frequency in the AFE phase. The real part of permittivity does not show any particular dispersion phenomena in the investigated frequency range and its value shows a strong dependence on the La3+ content. However, careful analysis of the dielectric loss data reveal two weak dielectric relaxations. The first one is signalled by a small peak in the dielectric losses as a function of frequency (Fig. 5.1 b)). It occurs at lower frequencies , it is very weak and broad and it is centred on 10 kHz. The second dielectric relaxation is indicated by the small increase of dielectric losses in high frequency region. Similar relaxation phenomena were previously verified in NaNbO 3 and PbZrO 3 AFE ceramics [ 11-12]. Therefore , weak frequency dispersion exist in all the investigated compositions. It is worth mentioning that frequency dispersion is uncommon for AFE materials since th ey have non polar character. Dielectric relaxation in the microwave region were often reported for ‘normal ’ FE and for relaxors and it has been attributed to domain walls [13 ] and to nanometer polar regions [14 ], Respectively. This suggests that, despite of AFE dominating character, nanopolar regions with short-range order interaction still survive in the samples with high La content ( i.e. x≥0.030) [ 11, 15-16], as it was proved by HXRD, TEM and Raman investigations which have shown that small amounts of FE phase still exists for samples with high La content. In addition, the substitution of La for Pb yields the composition disorder at the A site, donates electrons, and this composition al disorder results in the formation of local lattice polarization clusters or nanodomain distribution. Therefore, the existence of nanodomains in the FE phase may be responsible for the appearance of frequency dispersion in PLZT x/90/10 solid solutions. When compared with PZT 90/10 based composition (Fig. 5.1 c)), the PLZT ceramics show lower frequency dispersion. The real part of permittivity of PZT 90/10 show a variation of 56% within the investigated frequency range while the real part of permittivity of PLZT x/90/10 decreases only slightly with increase of frequency (with about 10- 16%). Similar values and frequency dependence were reported for close related PLZT compositions [17 ]. 144 Fig. 5. 1 Room temperature dielectric properties: frequency dependence of the real part of permittivity a) and b) dielectric losses for PLZT x/90/10 ceramics while c) shows the same properties for the undoped PZT 90/10- ceramic. All the PLZT compositions show good dielectric characteristics with high dielectric permittivity (higher then 375) and low losses ( tan δ below 3%) in the overall investigate d frequency range, which make these solid solutions good candidates for capacitor applications. The 300350400450500550 1021031041051060.050.100.150.200.250.300.350.400.450.500.550.60c) Dielectric losses Frequency (Hz)PZT 90/10Real part of permittivity 145 incorporation of small quantity of La3+ to FE PZT 90/10 generates a soft FE behaviour and yields larger dielectric and piezoelectric coefficients. As shown in Fig. 5.2 a), the permittivity vs. La3+ addition presents a non-monotonous variation, with a noticeable maximum of about 1000 at 1 kHz, around the composition x=0.030. In the orthorhombic AFE range (x˃3), the permittivity is reduced, but still preserves values as high as 700-800 (Fig. 5.2 a)). Fig. 5.2 Room temperature dielectric properties: a) composition – dependence of permittivity at a few selected frequencies and b) composition – dependence of dielectric losses at a few selected frequencies for PLZT x/90/10 ceramics One of the main differences between the properties of FE and AFE materials is that the dielectric losses are markedly higher in FEs [18 ]. This is indeed confirmed for PLZT x/90/10 compositions. According with XRD studies for compositions with x≤3, the amount FE phase is higher than 25% while for x≥3.1 the amount of AFE phase content is higher than 75%. For PLZT composition with high amount of FE phase ( x≥3), higher losses were obtained, with a maximum for compositions around 3 at. % of La (Fig. 5.2 b)). However, for PLZT 3/90/10 the FE phase is only marginally stable over the AFE phase. The maximum permittivity and dielectric losses in the compositional range of about x=0.030 can be explain ed with the fact that the mixture of rhombohedral FE to orthorhombic AFE phase allows multiple possible directions of polarisation 146 (and minimum flat free energy), which bring in a high response of the system to the driving forces at room temperature. Enhancement of electromechanical properties was reported previously in many solid solutions with MPB, as PZT [19 ], lead magnesium niobate-lead titanate solid solution (PMN- PT) [20 ], bismuth sodium titanate -potassium niobate (BNBT ) [21] and sodium niobate and lithium niobate ( LNN) [ 22]. In addition, the grain size and density play an important role in the electromechanical properties. It is of great significance to remark that this increase of permittivity takes place in the same compositional range where a minimum density and small ceramic grain size were found (Chapter 4). Since the decreasing of porosity [ 23] and grain size normally causes a degradation of material constants as result of a kind of “dilution ” effect [ 24], it means that even higher values of permittivity values can still be achieved around this composition, by improving the sintering strategy for a better densification. In the present thesis, we choose to maintain the same sintering parameters for all the compositions, in order to describe the role of La3+ addition only. 5.3 Piezoelectric properties In the PLZT x/90/10 composition with x between the 0 and 3 at. % , the AFE phase can be transformed into the FE one by electric poling and it is stable below the depolarization temperature Td. The piezoelectric constants ( d33, d31), electromechanical coupling constants ( kp, k31), and piezoelectric voltage constants ( g31, g33) as a function of La content for the PLZT x/90/10 samples are shown in Fig. 5.3. For the compositions with La content in the range 0≤x≤0.3, both small and large -signal d33 increase with increasing amount of La. For higher La3+ addition, a sudden decrease of d33 (ten times reduction) is observed until the composition x=0.033, followed by a weak increase for x>0.033. The other constants : piezoelectric constant d 31, electromechanical coupling constants kp and k31 and piezoelectric voltage constant g31 and g33 show an analog ue type of variation when increasing La3+ addition, with an abrupt decrease after the maximum value at x=0.030 and a stabilisation of values for x>0.033 (Fig. 5.3). The maximum of small -signal d33 value reaches 110 pC/N at the composition PLZT 3/90/10. This is a quite high value when compared with PZT 90/10 based composition. As earlier proposed by Haertling in the early phase diagram [ 25], the electromechanical properties of PZT ceramics are sensitive to the La3+ addition, namely, piezoelectric constants should increase (due to the fact that La creates softening mechanism) with the addition of La3+ in FE PZT ceramics. This trend is confirme d in our study only for composition up to x=0.030, due to the fact that La3+ addition also induces a rhombohedral FE – orthorhombic 147 AFE transformation by breaking the long range interaction between FE active octahedra containing B site cations. It is worth mentioning that this composition is at the border between FE and AFE phase. According with crystallographic studies presented previously it shows 75% of volume of AFE phase and only 25% of volume of FE phase. The piezoelectricity presented in this sample may be explained with field induced AFE- to-FE irreversible transformation. A consequence of this transformation is that the samples remain in poled state, similarly as reported in literature [7]. However, the permittivity of the same composition is highest and this result cannot be explain ed with irreversible field induced AFE- to-FE transition, since prior and during the low field dielectric measurements, no high field was applied. Fig. 5. 3 Dependence of the room temperature piezoelectric coefficients upon the La concentration for PLZT x/90/10 ceramics. Another possible explanation of the enhancement of piezoelectric and dielectric properties for this composition may be the presence of lower symmetry phase. Low symmetry phases such as monoclinic in the Pc, Pm, Cc , space groups have been postulated recently on the basis of 0.050.100.150.200.25 05101520253035 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5020406080100k kp -k31 d(10-12m/V )g(10-3Vm/N ) -g31 g33 La mol % -d31 d33 148 electron diffraction patterns, located close to the AFE-FE phase boundary [26 -28]. Moreover, maximization of piezoelectric constants and permittivity have been reported for compositions with MPB [19 -22] which were explained by a FE softening mechanism due to the co-existence of phases or by the presence of a low-symmetry phase which allows multiple possible spontaneous polarization directions [28 ]. In addition, previously XRD and R aman results suggested that l ower symmetry phases may be p resented in this compositions. Therefore , according wt h these results, a MPB which lies between the compositions x=(0.025, 0.035) separating rhombohedral FE and orthorhombic AFE states may occur . This is in agreement with data reported by Ishchuk et al. who also proposed a MPB for similar compositions [ 29] The observed enhanced dielectric and piezoelectric properties for PLZT x/90/10 around x=0 .030 might suggest an exact location of the MPB in this system . A reduction of piezoelectric constants take place for composition with x>0.030. This agrees well with the result of FE measurements and crystallographic results that there is a stabilization of AFE phase for x ≥3.3. An effect of the FE -AFE transformation is a loss of spontaneous polarization and thus, of the piezoelectric activity. How ever, for compositions with x>3.3 the piezoelectric constants show a slight increase. This can be interpreted a gain as an evidence of a large compositional range with coexistence of FE and AFE phases. A similar trend was observed previously on the same co mpositions by Raman study. The intensity of Raman soft mode is enhanced for the compositions x≥3.8, which indica tes an increase of the FE amount or the formation of an orthorhombic structure with lower symmetry. 5.4 Conclusions In summary, the electric and piezoelectric properties of PLZT x/90/10 ceramics were investigated at room temperature covering a composition interval in vicinity of the FE-to-AFE phase boundary. Dielectric and piezoelectric result gives evidences of a large compositional range with AFE– FE phase coexistence , in agreement with the results of the structural analysis. The room temperature permittivity has a maximum for x = 0.03, while the piezoelectric constants show anomalies in the range of compositions x = (0.03, 0.035) which suggest that functional properties in the composition range, may be associated with the development of a MPB. This compositional range may be the subject of further study since an enhancement of dielectric and piezoelectric properties, which can be still increased if a better densification will be achieved. With the results presented above a on original paper regading the structure-properties relationship was performed: I.V. Ciuchi, F. Craciun, L. Mitoseriu, C. Galassi, Preparation 149 and properties of La doped PZT 90/10 ceramics across the ferroelectric-antiferroelectric phase boundary, J. Alloys Compd. 646 (2015) 16-22. (I.F.=3.014 , A.I=0.558) 150  Bibliography V 1. Xu, X. Dai, J. Li, D. 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Sobolev, The influence of the coexistence of ferroelectric and antiferroelectric states on the lead lanthanum zirconate titanate crystal structure , J. Phys.: Condens. Matter., 17, L177– L182 (2005) 153 CHAPTER 6. Study of antiferroelectric- to-ferroelectric switching 6.1 Introduction As already mentioned, AFE perovskites are characterized by an antipolar structure close in energy to a related FE polar structure. In their virgin state at room temperature, AFEs do not exhibit macroscopic polarization. When an electric field of an appropriate amplitude EAF is applied, the antiparallel dipoles will align in a parallel arrangement and a first-order AFE- to-FE phase transition takes place [1] . This characteristic is usually found in compositions with the FE/AFE phase boundary close to the room temperature. Depending on how far the composition is from the FE/AFE boundary, the AFE phases may be recovered (reversible phase transition) or not (irreversible phase transition) after removal of electric field . A double hysteresis P(E) loop with a specific E FA field at which the FE phase transform s back to AFE or a single hysteresis loop is displayed in P(E) behavior after the AFE – to-FE transition f or the reversible and irreversible field induced AFE-to-FE transition, respectively [2-7]. Usually, PbZrO 3-based AFE ceramics possess high switching fields of the order of magnitude as the electrical breakdown strength. In order to reduce the switching field, the scientists tried to lower the free energy difference in favour of FE phase by chemical substitution of PbZrO 3 (e.g., with Ba, Sn, Ti, La). New AFE materials have been developed with a moderate EAF and thus, the way for many AFE-based different applications was opened. Earlier studies [8-13] revealed that field induced AFE- to-FE transition takes place only for composition at FE/AFE phase boundary, but it was not clear how the amount of FE and AFE phases influence the AFE- to-FE phase switching. In Chapter 4 we demonstrated a coexistence of AFE and FE phases in PLZT x/90/10 ceramics with La3+ compositions across FE/AFE phase boundary. The amount of AFE phase’s increases sequentially while FE rhombohedral one is reduced as La3+ content increases. Combining FE and AFE phases in a single composit ion could be a method to lower and to tailor the switching field EAF for the AFE- to-FE transition, at resonable applied fields. Furthermore, for compositions with higher La3+ content (x ≥ 3.3), the AFE phase was stabilized. This indicates that those compositions may switch from an AFE- to-FE state under high electric field, similar as reported by Pelaiz et al. [14]. Therefore, such ceramics are optimum for a detailed study of the AFE- to-FE transition. 154 In this chapter, the field-induced AFE- to-FE transition n PLZT x/90/10 will be analysed in details. The transition was first investigated through the measurement of field dependent polarization P(E) (at a fixed frequency of 1 Hz) as a function of La3+ content and amplitude of the applied electric fields. The influence of the applied voltage and of the La3+ addition on the phase transitions dynamics will be discussed . Studies of frequency dependent ferroelectric hysteresis P(E) are often reported for FE materials. This method gives insights into the kinetic phenomena and mechanisms involved in the polarization switching process [15-17]. The remanent polarization is enhanced in ferroelectric materials when a low frequency field (~10-2 Hz) is applied [15-17]. From far as know there is only one reports regarding the frequency dependence of ferroelectric hysteresis of AFE materials [18] . In this chapter, the frequency dependence of the remanent polarization, saturation polarization, coercive field and P(E) loop curvature will be investigated for three representative compositions PLZT 3/90/10, PLZT 3.2/90/10 and PLZT 4/90/10 in low-radiofrequency range, from 3 mHz up to 1Hz. This part of the study is aim ed to clarify the influence of frequency in AFE domains reorientation and its role on the reversible/irreversible nature of the AFE- to-FE phase transition It is worth to remind here that the AFE- to-FE phase transition involves structural phase transformation , unit cell volume increment as well as domain switching and texture formation [19- 23]. D ue to the proximity of the FE/AFE phase boundary, a structural transition from orthorhombic Pbam to rhombohedral R3c is expected in PLZT x/90/10 AFE compositions during the AFE- to-FE phase switching . Direct observations of the microstructur al evolution during the AFE- to-FE field induced transition are very rarely reported and the available studies are limited to few patterns close to the transition [3, 19, 20, 22, 24, 25]. Therefore, is of high scientific importance to understand the evolution of the crystalline structures at low fields, before, during and after the AFE- to-FE phase transition. Hence, another objective of this chapter is to study the structur al evolution of the AFE and field-induced FE phases by using direct in situ x-ray diffraction measurements as a function of positive/negative applied field sequences. Two cases have been chosen for this study : (i) the reversible field induced transition in PLZT 4/90/10 and (ii) irreversible field assisted transition in PLZT 3/90/10 composition. Further, the structure of the high field induced FE phase will be investigated on poled PLZT 3/90/10 powders by high resolution x-ray-diffraction (HRXRD) technique. In comparison with linear dielectrics and FEs, as a result of the reversible AFE- to-FE phase transformation, the “storage energy area” is approximately a square instead of a triangle. This AFE specific property make AFE materials promising for application as high charge storage devices, since the ene rgy stored in a cycle is proportional to the maximum polarization and applied electric field, 155 while the losses are proportional to the remanent polarization, which is negligible. In the Chapter 4 we demonstrated that the AFE phase was stabilized for PLZT x /90/10 ceramics with La3+ content x ≥ 3.3.This suggest that those ceramics could show high-energy storage properties. Few reports on AFE properties and energy storage properties of PLZT x/90/10 compositions ha ve been previously reported [26, 27] . The variation of ratio between FE and AFE phases by adequate substitution of Pb2+ by La4+ may also help to ta ilor the energy storage density. Furthermore , an enhancement of energy storage properties could be expected for compositions range in which the AFE and FE phases coexist , since maximum values of the dielectric and piezoelectric properties were reported (Chapter 4) in the same compositional range . Therefore, at the end of this chapter, t he energy storage , the lost energy and the efficiency properties determined from polarization switching properties of PLZT x/90/10 ceramics will be discussed in detail. 6.2 Polarization vs. electric field study 6.2.1 Influence of the La3+ composition on ferroelectric/antiferroelectric properties of PLZT x/90/10 ceramics We check ed first the ability of PLZT x/90/10 ceramics to switch from AFE- to-FE during the application of an electric field sequence. The transition between the AFE and FE phases induced by electric field was investigated by measuring the dynamic P(E) hysteresis loops. In addition, to exclude any ambiguity regarding the origin of the polarization, the corresponding field dependent current curves I(E) were simultaneously collected, as they are known to be more sensitive to the polarization state [28]. As mentioned, there are two important processes to be analysed in this study: an irreversible and a reversible AFE- to-FE phase switching. The irreversible AFE- to-FE transition may be also called field assisted switching and it takes place when the AFE composition is sufficiently close to the AFE/FE phase boundary. This kind of transition is metastable, with the free energy difference being lower than the reverse switching threshold field energy [6]. In this case, the AFE state can be recovered by thermal annealing [29-33]. The reversible AFE- to-FE transition is also called field forced switching and it occurs in compositions outside the region of metastability. Hence, for exemplifying these processes , three representative compositions across the FE/AFE phase boundary were chosen: PLZT 3/90/10, PLZT 3.2/90/10 and PLZT 4/90/10. The virgin and second hysteresis P(E and I(E) curves of these three compositions obtained during the applications of maximum field ) 156 Fig. 6. 1 Room temperature P(E) hysteresis and I(E) loops of La3+ modified PZT x/ 90/10 during first and second loading (frequency of 1 Hz) from virgin sample at room temperatures. The hysteresis loops and current curves from (a) and (b) were obtained for the composition PLZT 3/90/10, (c) and (d) for PLZT 3.2/90/10 and (e) and (f) for PLZT 4/90/10, respectively. Here, the EAF-virgin and EAF represents the forward switching fields required to induce AFE- to-FE phase transition, the EFA is the threshold field for the recovery of the AFE phase and the Ec represents the coercive field of FE phase -80-60-40-20 020406080-40-30-20-10010203040 -80-60-40-20 020406080-1.5-1.0-0.50.00.51.01.5 -80-60-40-20 020406080-40-30-20-10010203040 -80-60-40-20 020406080-0.20-0.15-0.10-0.050.000.050.100.150.20 -80-60-40-20 020406080-40-30-20-10010203040 -80-60-40-20 020406080-0.15-0.10-0.050.000.050.100.15EAF EAFEFAEFAEC Virgin-first run The second run EAF-virginEAF-virgin EAF-virgin E (kV/cm) E (kV/cm) E (kV/cm) E (kV/cm)E (kV/cm) E (kV/cm)P (C/cm2) P (C/cm2) f) e)d) c)b) a) I (mA) I (mA) I (mA)P (C/cm2) 157 amplitudes ( Emax≥EAF) of the sine wave input were measured and the results are shown in Fig. 6. 1 a)-f). The forward phase switching field EAF-virgin, EAF and the backward switching field EFA was estimated from the intersections of the extrapolated steepest parts of the hysteresis loop with the horizontal axis. Each composition was set in a virgin state prior to the measurements by heating at 600 ⁰C for 4h. As it can be observed, at low electric fields Emax≤EAF , the loops are practically straight lines (they show a small slope dP/dE), indicating that these ceramics are in their AFE state at room temperature. As the electric field amplitude increases, once the EAF-Virgin field is reached, the field induced A FE-to-FE switching takes place by a sudden increase in polarization and extension of the P(E) loop. The PLZT 3/90/10 ceramic maintain a remanent polarization Pr of ~30 µC/cm and it shows a single P(E) loop during the second field cycle, suggesting a FE-like behaviour. Therefore, similar with other systems with compositions across the FE/AFE boundary [6, 31, 34] the composition PLZT 3/90/10 presents an irreversible transformation (it does not transform back to AFE state when field strength decreases below EAF). After the first quarter of the field cycle a remanent polarization Pr ~20 µC/cm2 is observed for PLZT 3.2/90/10 ceramic when the electric field returns to zero ( Fig. 6. 1 c)), suggesting the occurrence of an irreversible field-assisted AFE- to-FE transformation. The second loop of this intermediate composition is quite intriguing. It show a constricted loop (triple ” hysteresis loop) which suggests a mixed AFE/FE behaviour under high electric field (Emax~60 kV/cm). However, during the fourth quarter of the field cycle, when the electric field reaches zero value, the polarization suddenly goes to a very low value of Pr~10µC/cm2, indicating a FE- to-AFE transformation. Hence, during the first quadrant of the second cycle this composition behaves like an AFE ( i.e. it show a reversible AFE- to-FE phase transition). Therefore an irreversible field-assisted and reversible field induced AFE- to- FE transformation may be found in this composition under different sequences of the application of high field, thus suggesting that this composition is metastable . At low electric field E≤E AF-virgin the PLZT 4/90/10 show a similar behaviour like the other compositions. The first and the second P(E) loops of PLZT 4/90/10 evidences a typically AFE hysteresis loop, which switches to FE at EAF =60 kV/cm with increasing electric field and it switches back to the AFE state at EFA= 27 kV/cm during the field decreasing( Fig. 6. 1 e)). Therefore the AFE- to-FE transition in this composition is reversible. It is widely accepted that the reversible character of the field induced AFE- to-FE transition is demonstrated by double P(E) hysteresis loop typical for AFE s (and it was exactly what we go t during the final field cycle for PLZT 4/90/10). It is important to notice here that the threshold value of EAF-virgin =68 kV/cm and EAF-virgin =60 kV/cm determined from minor cycles for PLZT 4/90/10 and from the virgin loop of PLZT 3.2/90/10 are higher t han those corresponding to the second cycle ( 60 kV/cm and 50 kV/cm , respectively ), this suggesting that the 158 AFE phase is destabilized during the first field cycle and therefore, a lower field EAF is further needed to induce the AFE- to-FE transition. In addition, even for the AFE ceramics, the polarization does not return to zero when the electric field decreases to zero, but it keeps a small remanent value ( Pr ~2.5 µC/cm2). Thus, the strict reversibility of the phase transition and recovery of the initial AFE state after the field removal is called into question for these compositions. These findings are often report ed in AFE materials but they have not yet found a justification yet. Some reports attributed such phenomenon to the retention of the FE phase in AFE PLZT composition [35]. It is well known that XRD is sensitive not only to the change in the structure but also to changes in domain polarization [36, 37]. Therefore, we considered that a systematic in situ field microstructural study of AFE- to-FE transition may help to bring light into this question. The associated I(E) curves for PLZT 3/90/10, PLZT 3.2/90/10 and PLZT 4/90/10 obtained during polarization switching (from Fig. 6. 1 a), c) and e)) are shown in Fig. 6. 1 b), d and f). A strong positive polarization current peak in correspondence to EAF-virgin= 24 kV/cm is observed for PLZT 3/90/10 during the first half of the first cycle. A smaller current peak whose origin is probably related to kinetic effects (incomplete switching) will be described in the next section. During the third part of the first cycle, a negative peak is noticed when the applied field reached Ec= 10 kV/cm which can be ascribed to the ferroelectric domains switching. Only two peaks (one positive and one negative) are found during the second cycle in correspondence of Ec =10 kV/cm , which indicate that this composition behaves like a FE material and hence, in this case the AFE- to-FE has an irreversible character . The I(E) curve of PLZT 3.2/90/10 composition shows one positive sharp polarization current peak during the first half of the first electric cycle which corresponds to the forward switching field EAF-virgin=60 kV/cm. There is no negative peak during decreasing of the electric field which should correspond to FE– to-AFE switching field. This indicates that the PLZT 3.2/90/10 system remains in the FE state and hence, it is confirm ed that the field induced AFE– to-FE transition is irreversible during the first quadrant of the electric cycle. However, a peak is observed during the application of an electric field with reverse polarity at backward switching field EFA ~2 kV/cm and this suggests a FE- to-AFE field induced transition phase [38] rather than a possible recovery of the i nitial AFE phase. This transition will be discussed more in detail further. Upon increasing the electric field amplitude up to maximum value 60 kV/cm and then decreasing down to zero, two current peaks appears in correspondence to EAF~40 kV/cm and EFA~ 0 kV/cm switching fields. This behaviour indicates a reversible field induced AFE- to-FE transitions, in agreement with the corresponding P(E) loop (Fig. 6.1c) . Four current peaks can be distinguished in the I(E) curve of PLZT 4/90/10 ceramic during the second electrical cycle, which conforms that this composition exhibits a typical AFE n ature. The first positive peak corresponds to the capacitive charging of the ceramic sample up to the 159 phase transition EAF=60 kV/cm field, while the second peak corresponds to the charge necessary to establish the FE- to-AFE transition during the field decrease ( EF A=20 kV/cm). The reversible and irreversible AFE- to-FE phase transiti ons can be clearly distinguished in the corresponding switching current curve s, which confirm that the previous presented polarization study reflect the intrinsic nature of polarization and of the phase transitions. Therefore, our findings indicate that the FE polarization is stabilized in specimens with low La3+ amount ( x≤3.2), while AFE polarization phase only exists in PLZT specimens with a higher addition of La3+ (x=4). It is well known that the spontaneous polarization in a ABO 3-type perovskit e FE mostly emerge from the displacement of the B-site cations within the BO 6 octahedra [39] . The substitution of Pb2+ by La3+ into the PZT perovskite structure was compensated by forming of B-site vacancies. This would reduce the number of FE dipoles per unit volume and has a detrimental effect on the ferroelectric order, while the antiferroelectric character is favoured. It becomes interesting to further investigate how the electric field amplitude affects the high field induced polarization in al l the compositional range of PLZT x/90/10 across the AFE/FE border. For such a study, we have registered systematically the first polarization versus electric field dependences P(E) (virgin loops) of each PLZT x/90/10 ceramics by consecutively applying various field amplitudes ( Emax) of the sine wave input of low frequency of 1 Hz. The results are shown in Fig. 6. 2 a)-i). Booth, La3+ content and the amplitude of the applied electric field strongly influence the P(E) behaviour of PLZT x/90/10 ceramics. The evolution of macroscopic P(E) hysteresis with La3+ composition is a clear demonstration of the FE- to-AFE crossover induced at room temperature by La3+ modification in PLZT x/90/10 compositions, as reported previously in Chapter 5 according to the crystallographic analysis. At electric fields below EAF (20 kV/cm ≤ Emax ≤ 55 kV/cm), the measured P(E) loops of 2/90/10, 2.5/90/10, 3.0/90/10, 3.1/90/10 3.2/90/10, 3.3/90/10, 3.5/90/10, 3.8/90/10, 4/90/10 compositions [ Fig. 6. 2 a)-i)] are practically straight lines ( P(E) loops with small areas), which evidence the AFE dominan t behaviour at room temperature in these ceramics. As the electric field Emax increases (Emax≥EAF), the field induced AFE- to-FE transition takes place, and a high polarization is developed. The compositions 2.0/90/10, 2.5/90/10, 3.0/90/10, 3.1/90/10 show a typical FE macroscopic hysteresis loop at E max higher than EAF (1, 20, 25 and 35 kV/cm, respectively) [ Fig. 6. 2 a)-d)]. It is worth to mention that the P(E) loop apparently looks like a typical FE one, but this behaviour is reached only if the applied electric field is higher than a certain threshold field EAF below which the P(E) curves do not open. 160 Fig. 6. 2 Room temperature dynamic P(E) hysteresis loops of La3+-modified PZT x/ 90/10 ceramics at different magnitude s of the applied field at frequency of 1 Hz. (a), (b), (c), (d), (e), (f), (g), (h) and (i) represent 2.0 %, 2.5%, 3.0 %, 3.1 %, 3.2 %, 3.3 %, 3.5 %, 3.8 % and 4.0 % La3+ compositions , respectively. The shaded area represents the highest recoverable energy density for each composition, Wre=∫ EmaxdPPmax Pr. (Oy axis is in µC/cm2 and Ox axis is in kV/cm ). The P(E) hysteresis loops were obtained on samples already exposed to electric field E≥E AF The FE predominant behaviour indicate that more adjacent dipolar units cannot switch back to their original antiparallel directions after being subjected to a high electric field. Therefore, the field induced transition into the FE state is irreversible. Thus, the fields used in the P(E) experiments are enough to induce a FE-like behaviour instead of the AFE character in compositions with phase superposition, since their free energies are very close. However, after the AFE- to-FE transition takes place, the Pr value keeps increasing as the maximum of the applied electric field Emax is increasing, suggesting that FE poling process occurs. According to previous crystallographic results (Chapter 4) , the FE phase amount in these compositions decrease from 60 vol. % for PLZT 2/90/10 to 20 vol. % for PLZT 3.1/90/10. The FE 161 like single loop at high electric field suggests that a small amount of FE phase significantly influences the P(E) characteristics of the 2.5/90/10, 3.0/90/10, and 3.1/90/10 compositions. This behaviour is quite analogue with the results reported by other authors [6, 40, 41]. Interestingly, a superposition of both FE and AFE switching character (triple hysteresis loops with variable amounts of the FE/AFE contributions) is noticed in the composition dependence of P(E) with La3+ amount in the range 3.2≤x≤3.3 [ Fig. 6. 2(e) and (f)]. A remanent polarization Pr~20 µC/cm2 is observed when the electric field returns to zero, suggesting the occurrence of an irreversible field-assisted transformation from AFE- to-FE in these pellets. Consequently, in all the investigated PLZT 90/10 ceramics with La3+ content in the range 2.0 ˂ x ≤ 3.3,where coexistence of FE and AFE phases was observed, an irreversible field-assisted AFE- to-FE transformation takes place. The appearance of pinched hysteresis loops clearly shows that a metastable FE phase can coexist with the AFE phase in the poled samples. Similar as reported in another study [42], the appearance of the metastable FE phase may be explained with two possible scenarios . A fi rst one is that the metastable FE phase may coexist with the AFE phase in the virgin PLZT x/90/10 ceramics . The second one is that only the AFE phase exists in the virgin PLZT ceramics , the metastable FE phase is induced from the AFE phase during the application of electric field , and only a partial amount from induced FE phase will return to the AFE phase after removing the electric field. From the crystallographic study presented in the Chapter 4 , it is obvious that both FE and AFE phases coexisted at room temperature in this compositional range. However , the AFE state is domina nt , , since no FE switching occurs prior to the AFE -to-FE induced transition. It is worth mentioning that for composition with low La3+ amount (x ≥2.5) the AFE phase amount is higher than 50 %. Hence, we may propose that the amount of FE phase presented in the PLZT compositions highly modulate the AFE-to-FE phase transition induced by electric field. Therefore, as suggested by Ma et al. [42] for (1-x)(Bi 1/2Na1/2)TiO 3-xBaTiO 3 ceramics the FE/AFE boundary of PLZT x/90/10 may be a macroscopic AFE phase region. The FE phase is presented only locally and highly contribute s to the AFE -to-FE phase switching and the n the induced FE phase remains stabl e after removing the field in PLZT x/90/10 samples (similar ly as in (Pb,La)(Zr,Sn,Ti)O 3 and NayBizTi1- x-O3(1-x)-xBaTiO 3 [40, 43] . The PLZT 3.5/90/10, 3.8/90/10, and 4/90/10 ceramics revealed well – developed double hysteresis loops [Figs. 6.2 (g) –(i)], which indicate that under high electric field EAF (~65 kV/cm) a reversible field -induced AFE -to-FE phase switching takes place. This is a signature of AFE single -phase behavior and agrees very well with the crystalographic results presented in the Chapter 4. These findings confirm the results reported by Pelaiz -Barranco for similar compositions [14]. 162 Fig. 6. 3 The dependences of the forward switching field E AF necessary to induce the AFE -to-FE phase transition as a function of La3+ composition in PLZT x/90/10 ceramics. The dependences of the forward switching field EAF as a function of the La3+ composition in PLZT x/90/10 ceramics is shown in Fig. 6.3. As it can be observed the threshold field EAF necessary to induce an AFE- to-FE switching is progressively higher as the addition of the La3+ increases. The switching characteristics of the different compositions can be understood by using the phenomenological model proposed by Yang & Payne [6], in which a thermodynamic barrier separates the AFE and FE states. The observed monotonic increasing in the switching field EAF as the La3+ content increases implies that the difference of free energies between FE and AFE phases increases , and the stability of AFE phase is higher for PLZT 3.8/90/10 and PLZT 4/90/10 ceramics . Therefore, the La3+ addition shifts the AFE- to-FE and FE-to-AFE transition to higher fields in PLZT x/90/10 compositions. Similar results were published for (Pb,La)(Zr,Sn,Ti)O 3 [40] and Pb0.99Nb0.02[(Zr 0.57Sn0.43)1-yTiy]0.98O3 [44] ceramics with compositions close to the AFE/FE phase boundary. Similar to our results, the switching field EAF and EFA were observed to increase as the Ti content decreases in (Pb,La)(Zr,Sn,Ti)O 3 compositions[40] . The curve displayed in Fig. 6.3 can be considered as an example of a field-forced transformation illustrating forward field strengths EAF where EAF can be related to the barrier between AFE and FE states. Irreversible/reversible character of the transitions between AFE and FE states is illustrated in Fig. 6.3. The induced FE state is stable for PLZT x/90/10 with x<3.2, and the AFE- to-FE is 2.0 2.5 3.0 3.5 4.005101520253035404550556065 FEAFE FE+AFE FERe FEIrEAF (kV/cm) La (at. %)AFEFE 163 irreversible on switching. Therefore, for compositions in the AFE and FE coexistence range, the AFE and FE phases are energetically in competition, and the threshold field forces the transformation over the barrier into the FE state. The barrier between FE and AFE states does not vanish and hence, the reverse transformation back to the AFE state when the electric field decreases does not occur for composition with x≥3.5. Therefore, the typical AFE P(E) hysteresis loops and the low values of remanent polarization found for x=3.5 , confirm that all these compositions are basically AFE at room temperature. Fig. 6.4 The dependences of P s, Pr, Ps-Pr as a function of the La3+ composition in PLZT x/90/10 ceramics. The maxim um values of the polarization Ps of revealed from the P(E) cycles, the remanent polarization Pr, and the difference between them Ps-Pr as a function of La3+ compositions are plotted in Fig. 6.4 . The coercive field was Ec ~10 kV/cm for all the studied compositions. The Pr increase slightly in the low range of La3+ content 2.0≤x≤3.0 and i t show s a maximum value of 35.5 µC/cm2 for PLZT 3/90/10 composition. While x further increases to 3.5, the Pr suddenly reduces to 1.5 µC/cm2. The remanence observed for these boundary compositions indicates that some fraction of FE phase is retained in the PLZT compositions at zero field. This trend is similar to the change of the composition dependence of piezoelectric properties reported in Chapter 5 . It is important to note that the trend of Ps polarization show anomalies which agrees well with the anomalies observed for the variation of FE phase volume ( R3c) and AFE phase volume ( Pbam ) with La content for the same PLZT x/90/10 compositions (Fig. 4.18 from Chapter 4) . Therefore , the results presented above 164 demonstrate that the hysteresis characteristics, i.e. the AFE character and EAF values are enhanced with the increasing the amount of Pb2+ replaced by La3+ ions. The observed macroscopic AFE characteristics in PLZT with La3+ content x ≥ 3.5 are the result of the La3+ addition induced crystalline shift from FE rhombohedral to AFE orthorhombic phase boundary. Larger difference between Ps and Pr at the same applied field favour higher energy storage density values. A high difference between Ps and Pr is revealed for AFE composition with La content high or equal to 3.5 at %. Therefore, an improvement of energy storage properties is expected in this range of compositions. 6.2.2 Frequency dependence of AFE- to-FE switching of PLZT ceramics The measurement of P(E) loop at different amplitude of electric field E and different frequency ( f) give some additional information concerning the domain switching kinetics in FEs. The frequency dispersion enabled discrimination of different types of contributions including electric - field-induced lattice strains and domain wall motion [45]. Although it has b een generally known that the coercive voltage increases with increasing the applied voltage frequency in FEs, the mechanism that governs the switching has not been exactly described. It is commonly accepted that when using high frequencies, a slow polarization reversal start or proceed, but it may be incomplete. Therefore, for such cases of DWs with slow kinetics, polarization switching needs apparent higher fields for switching, as ex perimentally reported. From far as I know the role of field cycling frequency on the AFE hysteresis and AFE -to-FE switching of AFE materials was not reported at all. I f we compare the P(E) loops of PLZT 3.2/90/10 from Fig. 6.1 c) with P(E) loops of the sam e composition from Fig. 6.2 e) we can noti ce significant differences: the P(E) loop from Fig. 6.1 c) evidence a mixed reversible and irreversible character of AFE -to-FE field induced transition while P(E) loop from Fig. 6.2 e) evidence only irreversible character for the AFE -to-FE field induced transition. The major difference in the two experiments was that P(E) loops from Fig. 6.1 report the first cycles measured on virgin samples, while ones from the Fig. 6.2 refer to loops r ecorded after the samples were set in a kind of “dynamical equilibrium”, i.e. after being subjected to several number of field cycles. This indicates that the field history strongly influence s the stability of FE or AFE phase and consequently, the shape of the P(E) response in the present PLZT compositions . This also explains multiple contradictory reports concerning quite different shapes of the P(E) loops in such AFE -FE compounds. T herefore, we considered that it worth to investigate the frequency ( f) dependence of PLZT x/90/10 ceramics in order to clarify these rate -dependent hysteresis properties. 165 In the following, the role of the field amplitude and frequency on the polarization switching process will be presented, for three selected representative AFE compositions: PLZT 3/90/10, PLZT 3.2/90/10 and PLZT 4/90/10. Fig. 6. 5, Fig. 6.6 and Fig. 6. 7 show the evolution of the hysteresis loops at various frequencies and fixed amplitude of the electric field E, illustrating a strong frequency- dependence of the shapes of the loops. The amplitude of the electric field E was chosen in order to ensure the occurrence of the AFE- to-FE field induced/assisted transition (E=45 kV/cm for PLZT 3/90/10, E=60 kV/cm for PLZT 3.2/90/10 and E=72 kV/cm for PLZT4/90/10). The data from Fig. 6.5 was obtained from refreshed samples. The refreshment was performed in order to set to their virgin state by heating above their Curie temperatures (600 ⁰C for 4 h) and then slowly cooled down in the absence of the electric field, in order to bring the dipoles in their low-energy equilibrium state. The samples PLZT 3.2/90/10 and PLZT 4/90/10 were often broken during the application of electric field E>E AF ( the electrical break down field is close in value with the E AF switching field) and those were even more sensitive during application of electric field of very low frequency. Hence the data from Fig. 6.6 and Fig. 6.7 were obtained from samples which were previously exposed to electric field E>E AF of frequency of 1 Hz. Fig. 6.5a)-b) show the hysteresis for PLZT 3/90/10 obtained at different frequencies while Fig. 6.5c) shows the corresponding I(E) current loops corresponding to the data from Fig. 6.5 a). As it can be seen in Fig 6.5 a), once the field was higher than 2 7 kV/cm applied at the frequency of 1 Hz, the composition is switched and saturates at a field of 45 kV/cm. After the first cycle, the system maintained typical square loop with well saturated remanent polarization Pr ~30 µC/cm2, as previously found (Fig. 6.1 a)). The full switching in the first quarter of the applied field corresponds to the high I(E) peak from the Fig. 6.1 c) (black). However, it is evident that the polarisation switching starts at lower field, of about 22 kV/cm as shown by the anomaly denoted with triangle in the Fig. 6.5 a) and corresponding to the small black current peak (indicated by the arrow in Fig. 6.5 c)). Usually, such types of anomalies are explained as due either to incomplete switching ( i.e. the system needs more time than allowed by the frequency cycling f=1Hz) or it might be due to the influence of point charge defects which pin the DW movement placed in a local minimum energy, which are further de-pinned after the sample is subjected to more field cycles [46, 47] . When decreasing the frequency from 1 Hz to 1/300 Hz, the coercive electric field Ec decreases from 11 kV/cm to 7.7 kV/cm, while Pr is slightly enhanced from 29 to 33 µC/cm2 (the enhancement of polarization is visible 166 -50-40-30-20-10 01020304050-60-50-40-30-20-100102030405060a) 5 34 2 1P (C/cm2) E (kV/cm)1 1 Hz, Virgin 2 1 Hz 3 1/60 Hz 4 1/300 Hz 5 1 Hz 1 1/300 Hz, Virgin 2 1/300 Hz 3 1 Hz -50-40-30-20-10 01020304050-60-50-40-30-20-100102030405060b) 3 2P (C/cm 2) E (kV/cm)1 -50-40-30-20-10 01020304050-2,0-1,5-1,0-0,50,00,51,01,52,0 -30 0 30-0,0010-0,0008-0,0006-0,0004-0,00020,00000,00020,00040,00060,00080,0010 4 1/300 Hz -30 0 30-0,04-0,020,000,020,04 3 1/60 Hzc) 5 3421I (mA) E (kV/cm)1 1 Hz, Virgin 2 1 Hz 3 1/60 Hz 4 1/300 Hz 5 1 Hz Fig. 6. 5 a) Field-dependent polarization P(E) loops with the virgin loop obtained for frequency of 1 Hz and , b) Field-dependent hysteresis loops with the virgin loop obtained for frequency of 1/300s Hz and c) current curves obtained in the same condition as a) for PLZT 3/90/10 bulk ceramic. 167 only on the first half of the cycle). When the same amplitude of the electric field at f=1Hz is applied in sequence, the Ec increases to 10 kV/cm and Pr decreases to 30 µC/cm2, but does not return to the initial value of Pr= 30 µC/cm2 and Ec=11 kV/cm obtained at 1Hz, suggesting that these changes are not completely reversible. The current curves from Fig. 6.5 c) obtained in the same conditions as P(E) curves from Fig. 6.5a) confirm th at the polarization results is related to FE domain switching . During the application of an electrical cycle of f=1Hz it can be observed that there is one current peak in correspondence of AFE- to-FE field assisted transition at EAF~25 kV/cm and a negative current peak in the third quadrant at Ec~10 kV/cm. Then, during the second electric cycle of f=1Hz there is a positive current peak in the positive applied field region and a negative current peak in the third quadrant in the correspondence to critical field Ec~10 kV/cm, similar as in classic FEs . When the field of lower frequency is applied in sequence the value of the switching current drammatically decreases. For f=1/60 Hz the switching current peak has the maximum value at 5×10-2 mA while for f=1/300 Hz it has the value of 5×10-4 mA. The small current peak obtained at f=1Hz at E=22 kV/cm may be interpreted as being due to incomplete switching during the first field quadrant, as mentioned before. We have further investigated the behaviour of the virgin sample under the same field amplitude at ultralow frequency of 1/300 Hz (Fig. 6.5 b)). In these conditions an anomalous hysteresis is obtained for the virgin state. The AFE- to-FE irreversible phase transition take place at EAF= 21 kV/cm, which is slightly lower than EAF = 25 kV/cm obtained previously for f=1 Hz . The FE hysteresis shows a very high saturation polarisation Ps ~55 µC/cm2. A higher field value is necessary to achieve the saturation and thus, a higher Ps value may be achieved. In order to understand the effect of frequency, we kept unchanged the electric field amplitude. With increasing the frequency from 1/300 Hz to 1 Hz the polarization Ps decrease but still preserves a high value of 40 µC/cm2. Therefore, with the decrease of frequency from 1 Hz down to 1/300 Hz, the polarization hysteresis loops show an enhancement of the switchable polarization and decrease in EFA switching field. These results are similar to ferroelectric properties reported for morphotropic phase boundary systems where two ferroelectric phases coexist in equilibrium , as in PbTiO 3-BiScO 3 [48, 49] . It is worth nothing that the PLZT 3/90/10 shows 75 % volume fraction of the AFE Pbam phase and only 2 5% volume fraction of the FE R3c phase . The typical ferroelectric P(E) loop obtained for this composition after its field assisted AFE -to-FE transition indicates that AFE -to-FE switching field become very small in this situation, due to reduction of the energy barriers separating the two phases in the energy landscape. Similar results were reported for (1-x)PbZrO 3-(x)Bi(Mg 1/2Ti1/2)O3 AFE system with x=0.15, but the ferroelectr ic hysteresis P(E) loop of this composition was obtained at typical frequency of 1 Hz [50]. 168 -60 -40 -20 0 20 40 60-50-40-30-20-1001020304050 3 2P (C/cm 2) E (kV/cm)1 1 Hz 2 1/60 Hz 3 1/300 Hz1a) -60 -40 -20 0 20 40 60-0,009-0,006-0,0030,0000,0030,0060,009b) -60 -30 0 30 60-1.20×10-40.001.20×10-43 300 s, 0.0033 Hz 1 1Hz 2 1/60 Hz 3 1/300 HzI (mA) E (kV/cm) Fig. 6. 6 a) Field-dependent hysteresis P(E) loops and b) current curves I(E) at various frequency and fixed amplitude of electric field 60 kV/cm for PLZT 3.2/90/10 bulk ceramic. The data were acquired starting from 1 to 3, respectively . Frequency dependent polarization P(E) and current I(E)curves of the AFE PLZT 3.2/90/10 ceramic cycled consecutively at 1 Hz, 0.016 Hz, and 0.0033 Hz with the same maximum field strength of 58 kV/cm, are shown in Fig. 6.6 . a)-b). Similar with the composition PLZT 3/90/10, the shape of the hysteresis loop change dramatically with decrease of frequency. The hysteresis loops at low frequency (0.016 Hz and 0.0033 Hz ) for the PLZT 3.2/90/10 material shows two distinct variations as compared with the curve obtained at frequency of 1 Hz ( Fig. 6. 1 c)). The first one is an obvious enhancement in the switchable polarization from Pr ~ 22 µC/cm2 for f=1 Hz to Pr~36 µC/cm2 for f=1/300 Hz and the other is a decrease in the AFE- to-FE switching field from EAF ~42 kV/cm to EAF ~37 kV/cm within the first quadrant. After the AFE- to-FE phase transition, during decreasing the field, a high remanent polarization of Pr~30 µC/cm2 is observed , probably due to the residual domain reorientation of the FE phase. However, the Pr decreases slowly when the field reverses the polarity . As a consequence, the AFE- to-FE transition take place when the sample was still showing a remanent value of Pr ~5 µC/cm2 at negative fields of -40 kV/cm. It is worth nothing that t he amplitude of EAF switching field from the first quadrant decreas es with decreasing of cycling field frequency while during the third quadrant, the EAF switching field increases when decreasing the frequency (from -33 kV/cm at f=1 Hz to -37 kV/cm for f=1/300 Hz) . Based on these data we may hypothesize that during the application of a periodic electric field of adequate amplitude for a long time (300 s) the PLZT 3/90/10 sample switch es from an initial AFE 1 169 state to a FE1 state with high polarization Ps~ 44 µC/cm2. Then, during the field reversal the polarization decreases to a low value Pr~30 µC/cm2 which suggest that the FE phase was destabilized. Hence, during FE- to-AFE phase switching the system returns to a AFE 2+ FE 1 state, but with a low polarization of FE1. Successively, during the application of an electric field of negative polarity, it switch es from AFE 2+ FE 1 to a FE2 phase characterised by a lower polarization Ps ~20 µC/cm2 and then, during the decreasing of electric field, is switch es back to the initial AFE 1 state. Another important feature of the PLZT 3.2/90/10 quasi-static hysteresis (low frequency) is that the P(E) loop is highly asymmetric : Pr=30 µC/cm2 and Ps~45 µC/cm2 from the first quadrant are different than Pr~6 µC/cm2 and Ps~33 µC/cm2, respectively, obtained during the third quadrant. The asymmetry may arise from the fact that, during the field reversal, the system does not have enough time to revert to the same initial AFE 1 state and thus, when the negative EAF threshold field was approaching, the system was found in a different state energy. The current curves form Fig. 6.6 b) obtained in the same conditions as the P(E) hysteresis loops confirm that the observed anomalies in the hysteresis curves are not related to extrinsic effects. Well defined peaks were observed in the correspondence of AFE- to-FE and FE- to-AFE phase transitions, which clearly indicate that domain switching and phase transitions occurred during the measurements of P(E) hysteresis loops. In addition, the current value dramatically decreased (from 0.005 mA to 1.2×10-4 mA) with frequency decrease (from 1 Hz to 1 /300 Hz) suggesting that the domains switch with very low electrical losses. The hysteresis loop for the PLZT 4/90/10 material cycled in a continuos sequence at 1/0.03, 1/0.2, 1/0.3, 1, 1/60, 1/180, 1/300 Hz , exhibits an even more interesting variation. As it can be seen from Fig. 6.7 a), there is a higher enhancement in the switchable polarization (P s~55 µC/cm2 and Pr~18 µC/ cm2 at f=1/300 Hz) and a stronger decrease in the EAF field during the first half of the cycle for the P(E) loops obtained at frequency lower than 1 Hz. Such frequency dependence indicates that the AFE- to-FE phase transition becomes diffuse. Clearly, these loops can be divided into three types: a pinched FE-like loop without saturation but with dominant FE character for high frequency (1/0.03 Hz), AFE-like hysteresis for frequency of 1 Hz with well saturated polarization, and an asymmetric loop with booth AFE and FE contributions, high remanent polarisation and without saturation for intermediate frequency. Similar with the composition PLZT 3.2/90/10, the loops obtained for PLZT 4/90/10 ceramics for low frequency (f˂1Hz) are asymmetric. However, the asymmetry is lower for this last composition, since an enhancement of remanent polarization was also observed for the negative half of the loops. This may be explained by considering that during the field reversal, the system has almost zero remanent polarization (close to the initial AFE 1 state). Similar as for the composition 170 PLZT 3.2/90/10, the current value dramatically decreases ( i.e. for the frequency lower or equal with 1/60 Hz, the current value was lower than the resolution limit of the oscilloscope of ~10-5 A and it was no more possible to record it) with the decrease of frequency. This fact indubitably suggests that leakage currents do not occur and the variation of polarization is related to the intrinsic properties of PLZT 4/90/10 sample [28, 47]. -80 -60 -40 -20 020 40 60 80-60-40-200204060a) 1/0.03 Hz, 1/0.2 Hz 1/0.3 Hz 1 Hz 1/60 Hz 1/180 Hz 1/300 Hzdecreasing fP (C/cm 2) E (kV/cm) -80 -60 -40 -20 020 40 60 80-2,0-1,5-1,0-0,50,00,51,01,52,0b) 1/0.03 Hz 0.2 Hz 0.3 Hz 1 Hz 1/60 Hz 1/180 Hz 1/300 HzI (mA) E (kV/cm)decreasing f Fig. 6. 7 a) Field-dependent hysteresis loops and b) current curves at various representative frequencies and at fixed field amplitude for PZT PLZT 4//90/10 bulk ceramic. The data were acquired in a continuos sequence, starting from high to low frequencies. It is worth to mention that an enhancement and a high value of Pr together with very low value of the current were previously observed only in FE systems with MPB (morphotropic phase boundary) where lower symmetry phases and polarization rotation were present [47] . A similar change of FE hysteresis ( i.e. decreasing of coercive field and increasing of Ps with frequency decrease) was reported as function of composition for the (Bi 1/2Na1/2)TiO 3–(Bi 1/2K1/2)TiO 3 system while it transit from rhombohedral to monoclinic state [51]. No similar behaviour was reported in AFE materials until now. However, we may hypothesize a similar transition (from AFE orthorhombic to FE monoclinic) in the PLZT 3/90/10, PLZT 3.2/90/10 and PLZT 4/90/10 under high field conditions. A structural study is necessary to further confirm this hypothesis. In conclusion, according to the variations of the PLZT loops induced by various frequencies , the samples cycled at f<1 Hz show an increasing extent of polarization effects. Based on the 171 experimental data we may assume that the AFE PLZT materials undergo multiple AFE- to-FE phase transitions under high electric field depending on the field frequenc y, since at low frequencies there was a significant remanent polarization. This latter behaviour nicely illustrates that phase modifiction phenomena may be induced at various field amplitudes and frequencies ( i.e. phenomena with various kinetic) between energy close polymorphic phases. These phenomena are very interesting and not [previously reported. The assumption of a metastable character or two FE phases successively induced by various electric fields may be considered to understand the nature of the anomalous hysteresis in the PLZT x/90/10 system. Structural dynamic in situ studies under similar conditions of the applied field (amplitude and frequency) are necessary to explain and to sustain these results. In situ synchrotron experimental investigation is in progress; a detail ed analysis will be performed in the near future. 6.3 Study of AFE- to-FE switching by XRD 6.3.1 Dynamic in situ XRD study In the following, a study regarding the influence of electric cycling on the structure evolution in reversible and irreversible field induced/assisted of PLZT 4/90/10 and PLZT 3/90/10 will be shown . These compositions have been selected considering their location with respect to the FE/AFE phase boundary: (i) PLZT 4/90/10: single AFE phase (Pbam orthorhombic symmetry in the virgin state), and presenting only AFE macroscopic character, as shown by the P(E) data (Fig. 6.1 e)). (ii) PLZT 3/90/10: located at the AFE/ FE boundary (with mixed Pbam orthorhombic (75 wt.%) and R3c rhombohedral phases (25 wt.%) in the virgin state), showing macroscopic FE character (Fig. 6.1a)). The ground Pbam state is close in energy to a rhombohedral R3c phase [52]. Since the investigated composition are close to the AFE/FE boundary of PLZT x/90/10, a structural transition from Pbam to R3c is expected in AFE composition during the application of electric field of adequate amplitude. Such a structural transition should show significant changes on the (111) and (200) peseudocubic peaks since the (111) reflection is split in the rhombohedral state and unsplit in the orthorhombic patterns, while the (200) reflection shows a single peak in rhombohedral structure and it appears split in two (200)/(002) peaks in the orthorhombic phase. Therefore, in order to gain a more precise idea about the structural modification at the AFE- to-FE transition, the evolution of the (111) and (200) pseudocubic reflections was monitored during the application of electric field E≥EAF (Fig. 172 6.8 and Fig. 6.9). Figure 6.8 a), b, and c) show the diffraction patterns of PLZT 4/90/10 ceramic, subjected to an alternative electric fields of triangular wave form during the first (virgin cycle) and second electrical cycle. Fig. 6. 8 Contour plots of diffraction intensities for a) and b) 002 pc/200 pc, and c) 111 pc reflections for the PLZT 4/90/10 composition during a complete triangular field cycle with amplitude ±74 kV/cm and frequency 0.8 mHz during a) and c) first and b) second elec trical cycle. The diffraction profiles from d) show the (002) pc/(200) pc reflections before application of electric field (black), at the coercive field of -74 kV/cm (red) and after application of electric field (green). The subscript pc indicates reflections indexed with the pseudocubic primitive cell. A schematic drawing to describe the experimental sequence of applied electric field is included in the left hand side of the figure. EAFE -FE and E FE-AFE represent the forward switching field for inducing the AFE -to-FE transition and the backward switching field of the FE -to-AFE recovery, respectively. 173 The diffraction pattern between 37°< 2θ <38° corresponds to the (111) lattice planes, while the diffraction pattern in the range 43°< 2θ < 44° corresponds to the (002) and (200) lattice planes. Initially, at zero field, the (200) diffraction splits into two peaks wile (111) is a singlet, in agreement with orthorhombic symmetry as shown in Chapter 4 for this composition in the virgin state. Upon increasing the electric field (Fig. 6.8 a) and c)), the appearance of (111) and (200) diffraction profiles remains stable until reaching E AF-Virgin =68 kV/cm, when the (111) and (200) diffraction lines profiles change suddenly: the (111) peak shift to lower 2θ while the (002) peak disappears and (200) increases in intensity. This indicate that the AFE- to-FE domain switching was quickly induced at this threshold field. The shift towards lower 2θ of the {111}pc diffraction line profile indicates that the preferred orientation induce non-uniform strains in the lattice. The resultant profile of (111) peak appears symmetric and is not splitted as expected for a rhombohedral structure. Furthermore, { 200} pc reflection develops a distinct shoulder on the lower 2θ side (Fig. 6.8 (d)). This suggest that, in addition to the domain switching process, the PLZT 4/90/10 composition is subjected to a modification towards a lower symmetry new phase under high electric fields E≥ E AF-Virgin =68 kV/cm. The present data are not sufficient to determine exactly the structure of th is field-induced new phase, but it is obvious that the high field phase is a FE one, according to the aspect of P(E) loops (Fig. 6.1e)). When the field decreases from 74 kV/cm, the (111) and (200 ) peak s shift back to the initial position at EAF=28 KV/cm, when the ceramic has recovered its orthorhombic structure characteristic to the virgin state. The recovery of (111) and (200) position of the initial AFE orthorhombic state (Fig. 6.8 b)) during the field reversal suggests that the field-induced AFE- to-FE phase transition is reversible. However, when the electric field returns down to zero , it is noticed that the intensity (200) AFE peak is significantly reduced (the integrated intensity is about one fourth from one of the virgin state), while the intensity of (002) is enhanced . This suggests that the sample experience d also some type of irreversible process during the first field cycling. This effect is further maintained during the second cycle as show n in Fig. 6.8 b) for the ( 200) reflecti on. Furthermore , from Fig. 6.1 e) it can also be found that the EAF-Virgin =60 kV/cm from the second cycle is lower than the EAF=68 kV/cm from first cycle. This is in agreement with the polarization results from Fig. 6.1e). Therefore, the AFE-to-FE transition is not completely reversible. Similar as in tetragonal perovskite phases the intensity interchange between (002) and (200) indicates 90ș domain reorientation [19, 53, 54] . In a previous work, T. Lu et al. [19] evidenced that the change in (002) intensity of AFE phase after field -induced phase transition is related to the strain -driven preferred orientation. The preferred orientation has a quite considerable effect on the polarization of AFE sample s since it dest abilize the 174 Fig. 6. 9 Contour plots of diffraction intensities for a) and b) 111 pc/-111 pc, c) and d) 200 pc/002 pc reflections for the PLZT 3/90/10 composition as a function of triangular bipolar electric field of ±45 kV/cm amplitude and frequency of 0.8 mHz during a) and c) first, and b) and d) second electrical cycle. The subscript pc indicates reflections indexed w ith the pseudocubic primitive cell. A schematic diagram to describe the experimental sequence of applied electric field is included on the left side of the figure. E AFE -FE and E FE-AFE represent the forward switching field for inducing the AFE -to-FE transit ion and the reverse switching field of the FE -to-AFE recovery, respectively. AFE phase and bring in lowering of threshold field EAF. The evolution of the features (111) and (200) pseudocubic peaks as a function of the applied electric field is shown in Fig. 6.9 a)-d) for the composition PLZT 3/90/10. In the virgin state this ceramic belongs to the range of coexistence of AFE and FE phases (Chapter 4). The doublet nature of (200) and the singlet nature of (111) peaks at 175 zero field suggests that the dominant phase is orthorhombic. This is in agreement with the previous Rietveld calculations (Chapter 4) showing a phase superposition with 75% orthorhombic phase and 25% rhombohedral state. Similarly with the composition PLZT 4/90/10, the present structural results directly verify the FE structure induced by an electric field at EAFE-FE virgin = 24 kV/cm and it also indicate a fast kinetics of the AFE- to-FE transition. As expected for an irreversible AFE- to-FE field induced transition, when the electric field reaches zero at the end of the second quarter of the cycle, the (111) and (200) peaks features of high field FE phase are preserved. Therefore, the XRD results clearly illustrate the irreversibility between AFE and FE states in this composition . During reversing the polarity of the electric field in the third quarter of the cycle, a small recovery of the (111) and (200) AFE specific peaks position and features, is observed for an amplitude of electric field of ~6 kV/cm. This indicates an unusual field assisted FE- to-AFE phase transition and suggest that the FE phase induced by the electric field can be restored back to the AFE phase when a field with opposite polarity is applied. The recovery of AFE phase takes place within a limited field range (4-7 kV/mm). Further increase of the field amplitude brings the system back in to the FE state, as reveled by the (111) and (200) peak position for the field value of ~9 kV/mm during the third quarter of the cycle. The FE- to-AFE field induced transition is non-typical, because it is expected that electric fields favor the FE phase with parallel electric dipoles instead the AFE one. However, the P(E) loop (Fig. 6.1a)) does not show any evidence which can confirm the recovery of AFE and again the AFE- to-FE induced transition. Hence, in order to confirm the appearance of the electric field-induced FE- to-AFE phase transition we performed in situ XRD measurement during the second electric cycle by applying an electric field of the same amplitude and reverse polarity. As observed in the Fig. 3 b) and d), the (111) and (200) AFE phase signature XRD peaks appear again in a limited field range (4-7 kV/cm) during the first and the third quarter of the second cycle showing the highest amplitude at the same coercive field electric EFE-AFE ~6 kV/cm. This experiment confirms the occurrence of AFE phase out of a FE phase. The fact that the hysteresis loop of 3/90/10 measured during the first and second cycle does not show any evidence of the FE-to-AFE field induced transition may be explained by the fact that the frequency of 1 Hz is too fast to induce the AFE- to-FE transition since the frequency of the hysteresis loop (1Hz) is higher than the frequency of the field used for the in situ XRD measurements (0.00083 Hz). The induced AFE state from FE was previously observed in a NaNbO 3-based lead-free ceramic and in polycrystalline Pb 0.99{Nb 0.02[(Zr 0.57Sn0.43)0.92Ti0.08]0.98}O3 ceramics [38, 55]. We further investigated the FE- to-AFE phase transition more in detail through the analysis of the FWHM of (111) peak in the vicinity of the FE- to-AFE transition and the results are shown in Fig. 6.9. It can be observed that a small increase of FWHM of (111) peak is already present for field values 176 in the range of -3≤E≤0 kV/cm. A larger value of FWHM of (111) peak can be associated by the appearance of a new phase, and hence, the results from Fig. 6.9 indicate that the FE- to-AFE phase transition is induced but only in a very small amount. However, the FWHM of (111) diffraction peak increase during the application of electric field of reversed polarity and it shows a maximum value for E~6 kV/cm. The high value of FWHM of (111) may be related to coexistence of FE and AFE phase s, which means that the transformation FE- to-AFE phase is not completed. The figure indicates that the AFE- to-FE phase transition takes place in a large range of the electric field. Therefore, the XRD data evidences that the field induced FE-to-AFE phase transition takes place with a slow kinetics. These observations indicate also that the transition from FE- to-AFE is diffuse. Thus, the key factor to determine the field-induced FE- to-AFE phase transition is to expose the poled ceramic to a field with reverse d polarity of adequate amplitude for a long time. The explanation for the induced AFE phase at the coercive field with a reversed polarity was attributed to volume contraction/reduction which happen during application of electric field of reverse polarity on poled sample [38] . Fig. 6. 10 Electric field dependence of FWHM of the (111) diffraction peak during FE -to-AFE and AFE -to-FE induced transitions for the composition PLZT 3/90/10 . Further increase of the electric field cause a decrease of FWHM of (111) while the (002) peak disappears. These changes are consistent with a high field FE phase, i.e. the previously induced AFE phase was converted into the FE phase for a field higher than 10 kV/cm. It is worth mentioning that 177 the transition happen rather slowly as the applied electric field level increases. When compared with the AFE- to-FE transition from the first quadrant it can be observed that the new AFE- to-FE phase transition takes place at a larger field and its threshold value E AFE- FE ~10 kV/cm is significantly lower than the E AFE-FE virgin = 24 kV/cm threshold field. Furthermore, this field amplitude is almost the same with the coercive field E c of the induced FE phase (Fig. 6.9 (b), second cycle). It should be noted that it is impossible from the present data to determine the space group of the field-induced FE phase. H owever, the multi peak feature of the (002) reflection indicates that rather than R3c, Cm and Cc are the most probable candidates. This should be confirmed by a detailed crystallographic study in the future. 6.3.2 Ex situ High Resolution X-ray Diffraction (HXRD) study We will present here the results of a high resolution structural analysis performed at room temperature (HXRD) on PLZT 3/90/10 poled powders. We have chosen to investigate this composition because it is clear from its hysteresis loop and piezoelectric properties that it retains at room temperature the high field induced FE phase after unloading the electric field. We have followed a similar experimental approach as K.V. Lalitha et al. [48] . The HRXRD pattern was recorded after subjecting the PLZT 3/90/10 ceramics to a strong electric field EAF ~30 kV/cm (for 30 min), followed by grinding into powder for the high resolution powder diffraction experiment. It was supposed that the grinding procedure itself did not modify the polarisation state of the ceramic grains. Figure 6.11 shows a comparison of the HXRD profiles of 110 pc, 111 pc, 200 pc, 211 pc, 220 pc and 222 pc main perovskite reflections of the PLZT 3/90/10 samples, in the virgin state and after poling at ∼30 kV/cm. The 110 pc is a doublet for both rhombohedral and a orthorhombic distortions. For a rhombohedral structure, 111 pc is a doublet whereas for orthorhombic structure, the 111 pc is a singlet. The pseudocubic 200 pc and 220 pc reflections form doublets (002 and 200, 202 and 220) while 222 pc is a singlet for pure orthorhombic phase. For a rhombohedral structure 200 pc is a singlet while 220 pc and 222 pc are doublets. The 211 pc is a triplet for rhombohedral phase and a doublet for the orthorhombic one. As observed in Fig. 6.11 the 110 pc diffraction profile of poled PLZT 3/90/10 shows an additional peak,the 200 pc pseudocubic peak is a triplet, while 220 pc and 211 pc pseudocubic peaks split into four peaks, which suggests that the structure of poled sample is neither orthorombic nor rhombohedral or a combination of both. The complex character of these peaks may indicate the presence of a low symmetry monoclinic phase. The splitting of the XRD peaks of poled ceramic from Fig.6.11 can be understood in terms of the monoclinic phase with Cm space group discovered by Noheda et al. [56] 178 It may, however, be mentioned that for the monoclinic phase, the pseudocubic (220)pc reflection is usually observed to have three peaks but higher resolution XRD data may reveals the presence of an additional fourth peak, as also obtained for poled PLZT 3/90/10. However, a small intensity 002 pc peak is observed in Fig. 6.11, which indicates that the orthorhombic phase may coexists with the monoclinic phase (maybe because some AFE state was recovered during the mechanical crushing) . 14.0 14.2(211)pc 2 virgin poledo 16.2 16.3 16.4(220)pc 2 virgin poled o ooIntensity(a.u.) 8.10 8.16(110)pc 2 virgin poled 9.84 9.96 10.08(111)pc 2 virgin poled 19.9 20.0 20.1(222)pc 2 virgin poled 11.4 11.6(200)pc 2 virgin poledoo Fig. 6. 11 Selected pseudocubic Bragg profile of electrically poled and of virgin PLZT 3/90/10 powders 179 5 10 15 20 25 30Intensity(a.u.) o2Cm Pbam2=0.99 Rwp=11.925 Fig. 6. 12 The observed (dots), calculated (continuous line) and difference (bottom) HXRD profiles after Rietveld refinement of the structure of poled PLZT 3/90/10 powder using Cm and Pbam space groups. The Bragg positions are shown in the bottom inset by vertical lines: the top one corresponds to Cm and the bottom one corresponds to Pbam structure. Rietveld Refinements of the high resolution synchrotron XRD data have shown a Cm monoclinic phase, as discovered by Noheda et al. for PbZr 1-xTix [56] together with the Pbam structure, which can explain the induced FE phase diffraction pattern from Fig. 6.12 and Fig. 6.13. In order to provide a better visual clarity, only some of the representative pseudocubic Bragg peaks are shown on a magnified scale. Figure 6.13 depicts the XRD profiles of the pseudocubic 110 pc, 111 pc, and 200 pc, 211 pc, 220 pc and 222 pc reflections along with the Rietveld fits while full pattern refinements is shown in Fig. 6.12. The Cm and Pbam structural models could nicely explain all the features of the diffraction pattern for the poled samples and an excellent overall fit was obtained. Rietveld analysis suggests that the monoclinic is the major phase with volume fraction of 75.1%. The monoclinic structural parameters of the poled specimen are: a m=5.800, b m=5.825, c m=4.141 and α=91.4904. The cell volume of AFE Pbam has marginally increased by 0.012 ( A˚)3 after poling. The weight volume fraction of the Pbam is 24.9%. The enhancement of polarization observed 180 previously (Fig. 6.5, Fig. 6.6 and Fig. 6.7) can be explained within the framework of the polarization rotation model. The monoclinic Cm phase allows continuous rotation of the polarization vector within the pseudocubic 110 pc plane. In particular, the polarization vector may become parallel to [111]pc, [001]pc, or [110]pc direction, leading to a change of the structure like rhombohedral (R3m), tetragonal (P4mm), or orthorhombic (Amm2) phases, respectively [48] . These structural results validate the enhancement of polarization data whereas the application of a field of a very low frequency would favour a low symmetry FE phase which corresponds exactly to the observed monoclinic distortion and field induces the monoclinic phase. In addition anomalies observed in the P(E) loop may be explained with different structures induced by the orientation in different planes of the Cm structure, as previously mentioned. 2Intensity(a.u.) 19.9 20.0 20.1(222) pc 14.04 14.16(211) pc 11.5 11.6(200) pc 9.90 10.01(111) pc 8.10 8.19(110) pc 16.12 16.38(220) pc o Fig. 6. 13 The observed (dots), calculated (continuous line) and difference (bottom) HXRD profiles of the enlarged views for the 110 pc, 111 pc, 200 pc, 211 pc, 220 pc and 222 pc peaks obtained after Rietveld refinement of the structure of poled PLZT 3/90/10 powder using Cm and Pbam space groups 181 These data have revealed that poling of PLZT 90/10 induces a monoclinic distortion. This field-induced new phase is stable even after the field removal. I n Chapter 3 we have shown that PLZT 3/90/10 exhibits an anomalous large piezoelectric response when poled at room temperature. These anomalous features remained unclear so far from the structural standpoint. The presence of a field- induced monoclinic phase in PLZT 3/90/10 can explain the superior piezoelectric coefficient of this composition (d 33 ~100 pC/N) as compared with that of PZT 90/10 (d 33= 65 pC/N). In conclusion, the present structural detailed analysis provided an explanation for the large piezoelectric coefficient of PLZT x/90/10, the field induced phase transition behaviour and high polarization values at room temperature and the anomalies observed in the P(E) loops under high field conditions and low frequency. Further studies of the poled samples are in progress, and they will be reported in a subsequent publication. The present results support the existence of a field-induced kinetically stabilized monoclinic phase at the orthorhombic-rhombohedral phase boundary crossover. More extensive studies are necessary to confirm the nature of the monoclinic phase. Although a satisfactory fit was found using Cm and Pbam space groups, the monoclinic distortion is only slightly different with respect to the rhombohedral structure and thus, the relative importance of considering the monoclinic structure with respect to the rhombohedral structure need to be tested. For similar compositions, other monoclinic structures such as Pm, Pc and Cc were proposed [57, 58]. Therefore,further tests are necessary in order to exclude any ambiguity regarding the crystalline structure of the poled PLZT 3/90/10 ceramic. 6.4 Energy storage properties To design a suitable energy storage AFE material, at least three requirements have to be satisfied simultaneously: high breakdown field, low coercitive field, large saturated polarization, and small remanent polarization. According to the P(E) loops reported in earlier sections, it is obvious that all the PLZT x/90/10 with x≥3.5 satisfy these conditions. Previously, in Chapter 4 , it was demonstrated that the FE and AFE phases coexist in a large range of compositions (2.5≤x≤3.3). All the ceramic samples with compositions in the range of phase superposition show good dielectric properties (ԑ’≥700 and tan δ below 3%) in a broad frequency range. In addition, the compositions with higher La3+ content (x≥ 3.5 ) show reversible field-induced AFE- to-FE transition. Moreover, PLZT-based AFEs have a high Curie temperature (~170oC) [59, 60] which means that the AFE phase may be stable in a wide temperature range. This indicates that these ceramics could have potential for 182 applications in high energy-storage capacitors. Although a lot of works concerning energy storage properties of AFE ceramics have been published, the influence of the amount of FE phase on the energy storage properties of the FE and AFE composite materials was not yet investigated. There are several reports on the energy storage properties of some Pb 1-xLax(Zr 0.9Ti0.1)1-x/4O3 (PLZT x/90/10) compositions and important insight on this system has been achieved [14, 26], but a systematic study on the influence of La3+ content on the energy storage properties has not been reported so far . Therefore, it is now interesting to investigate how the coexistence of AFE and FE phases influence the energy storage properties in PLZT 90/10 ceramics, in order to better control their properties and to settle their potential application in energy storage devices. The energy-storage density, was evaluated from the electric-field induced polarization P(E) hysteresis loops of the PLZT x/90/10 ceramics. Only the dynamical FE loops from Fig. 6.2, obtained during application of different amplitudes of sinusoidal signals (1≤ Emax≤65 kV/cm) at the frequency of 1 Hz are considered for this study. The recoverable energy , loss energy and efficiency values have been calculated using to the definition of recoverable ener gy density of the FE capacitor, ( i.e. at the withdrawal of applied field Emax.): 𝑊𝑟𝑒=∫ 𝐸𝑚𝑎𝑥𝑑𝑃𝑃𝑚𝑎𝑥 𝑃𝑟 (Emax≡ applied electric field and P≡ polarization), the energy loss density 𝑊𝑙𝑜𝑠𝑠 =∫ 𝐸𝑚𝑎𝑥𝑑𝑃𝑃𝑚𝑎𝑥 0 and energy efficiency η=W re/(W re+W loss) [7, 12, 13, 61 -63]. Fig. 6. 14 a) Energy storage density Wre and b) efficiency η at room temperature for PLZT 90/10 samples with various amounts of La3+ under different applied electric field Emax. 183 Fig. 6.14 shows the dependence Wre (Fig. 6.14 a)) and efficiency η (Fig. 6.15 b)) on external field obtained for PLZT x/90/10 ceramics with different La contents. As expected, the La3+ addition induce an enhancement of the energy storage density. Higher energy storage, Wre ~0.819 J/cm3, was obtained for the samples with low La content (x≤2.5) during application of electric field of Emax =30 kV/cm. The Wre of PLZT 3/90/10 is ~1.30 J/cm3 at 40 kV/cm applied field. The PLZT ceramics 3.1/90/10, 3.2/90/10 and 3.3/90/10 have Wre equal to ~1.38 J/cm3, ~1.25 J/cm3, ~1.23 J/cm3, respectively, at Emax of 45 kV/cm. The AFE ceramic 3.5/90/10, 3.8/90/10 and 4/90/10 have Wre of ~1.85, ~1.27, and ~0.58 respectively at Emax of 65 kV/cm. These results evidence that with increasing La3+ addition, the maximum value of Wre is higher at a higher value of the applied electric field. This is related to the AFE nature of the samples and AFE -to-FE field induced transition. As La3+ increases, the amount of AFE phase is also increased and a higher EAF is necessary to induce the AFE -to-FE transition. Fig. 6. 15 Electric field -dependence of the recoverable energy storage density Wre (a) and of the lost energy Wloss (b) of the La3+ modified PZT 90/10 ceramics with La3+ at. % content varying from 2 to 4 Contrary from Wre values, η of the PLZT x/90/10 ceramics is observed to decrease as the amplitude of applied field increases. As can be seen in Fig. 6.14 b) at lower field ( Emax≤15 kV/cm), η increases as La content increases and the PLZT compositions with higher x (x≥3) values show high η (~95%). When the electric field Emax increas es from 15 kV/cm toward higher values, η displays lower values for the same compositions , but not lower than 40% . The dependence of Wre and Wloss on electric field and La3+ content is complex. A more clear dependence of the Wre and Wloss as function of electric field Emax (y in logarithmic scale) is illustrated in Fig. 6.15 . As La3+ addition increases, the highest recoverable and loss energy are shifted at higher value of electric field and two borders 184 defined by high and low energy values are displayed. The lower limit is determined by the energy values of the AFE phase, while the upper border delimit ated the energy values of FE phase and FE field-induced transition for AFE ceramics . As La3+ content increases, the difference between these limits become larger and the values of energy are shifted to higher values of the electric field and energy. Concomitantly with the AFE-to-FE field induced/assisted transition (at Emax≥ E F) there is as step-like increase in energy [64] and the values of Wre and Wloss of AFE ceramics are shifted from the lower limit to the higher limit. The PLZT 2.0/9 0/10 ceramic is predominantly FE. This is the reason why i t does not show step like increase in energy at higher electric field. This is the composition which Wre values start to draw the superior limit for low value of electric field. The PLZT compositions with La3+ content in the range 2.5≤x≤3.5 evidence a step -like increase in Wre and Wloss wich is related to AFE -to-FE field induced transition. It is worth to no te that, the energy value where the step -like increase takes place on the superior limit seems join the value Wre of the previous FE composition and it displaces Wre at higher value as the field Emax is further increases. The samples with high La3+ content (x≥3.5), show similar values of Wre and Wloss for Emax lower than EAF , while for Emax values higher than EAF, the energy Wre shifts to higher values. When compared with lower La3+ content compositions, the ceramics with x> 3.5 show a slight redu ction in the highest Wre energy and an enhancement in the inferior border energy, the difference between the limits being reduced. In spite of the fact that a small steep increase in energy is displayed, (Fig. 6.15 a)) (due to the transition from AFE -to-FE takes) and the PLZT 4/90/10 compositions have its highest Wre value nearer to AFE lower limit rather than to FE upper border. Therefore, the limits start to closeness each other for the ceramic s with x> 3.5. PLZT ceramics with La3+ addition in the compositional range 3.1 ≤ x ≤ 3.5 have a similar η value in a large range of moderate electric field 21 -42 kV/cm. Therefore we may argue that the ceramics with amount of fraction rhombohedral phase in the range 0.21≤ f ≤0.8 % and orthorhom bic phase amount in the range 0.79≤ f ≤0.91 % are more preferable for energy storage devices due their lower electric field AFE -to-FE induced transition. The highest Wre value of about 1.8 J/cm3 obtained for PLZT 3.5/90/10 ceramic is quite high value of energy density in comparison with other latter reports for AFE ceramics [12, 30, 65, 66] . Thus, the stabilization of AFE phase cause s an increment of the energy storage capacity . However, accordin g to the crystallographic study, the highest amount of AFE phase is obtained for PLZT 3.8/90/10 and PLZT 4/90/10 ceramics. The compositions PLZT 3.5/90/10 still ha s a low amount of FE phase of about 0.08 %. It is obvious that a low amount of FE phase enhan ce the energy storage properties. This aspect is worth to be studie d more in detail in the future. Enhancement of energy storage in AFE’s caused by La3+ addition was reported in other studies 185 [30, 64] . In the paper published by J. Parui et al. [64], it was demonstrated that La3+-modified PZ thin films are able to induce enhancement of charge storage by La3+ induced crystallographic orientation. The low dielectric polarization and low break -down strength limit the energy density of AFE ceramics to 2 J /cm3 [24, 30, 61, 65, 67, 68] . The highest energy W re of 6.4 J/cm3 up to now has been obtained on PLZST -based AFE composition at very high field of 400 kV/cm and with an efficiency of 67 % [10]. Thin film AFE materials show higher values , but the low thickness may be a limitation in the practical requirements [11, 63] . While the system evolves from a FE state to an AFE state, as a function of La3+ composition and from an AFE state to a FE state as a function of electric field, the recovered energy was maximized for compositions exhibiting coexistence of the AFE and FE phase and whose perce nt weight amount of FE phase is higher than 0.1 % ( PLZT x/9 0/10 compositions with x=2.50, 3.0, 3.10. 3 .20). In addition, the energy efficiency of the same compositions is about 60%. The highest energy density of 1.85 J/cm3 is obtained for AFE PLZT 3.5/90/10 ceramic with a high efficiency of 65% at the electric field of 65 kV/cm. The value of energy density (0.85≤ Wre≤1.85) together with energy efficiency in the range (41≤ η≤65 %) obtained for a relatively low applied field (30≤ Emax≤65 kV/cm), point out that the investigated PLZT ceramics show high potent ial for pulse power capacitors applications. Therefore , the FE/AFE phase coexistence ha s a great influence for the energy enhancement due to the interaction between the response of the FE phase component and of the field -induced AFE -to-FE phase transition. The main conclusion of this study is that, the energy storage may be tuned by adjusting the La3+ addition in 90/10 PLZT ceramics. Furthermore, it can be found an optimum composition wich can display both high energy storage and high efficiency, at reasonable available fields, imposed by a given application. 6.5 Conclusions To understand the origin and features related to the phase transition between FE and AFE states is of considerable interest for solid state physics. However, until now this phenomenon has not been amply covered in the literature. In this Chapter we have carried out an in-depth electrical and structural characterization of AFE- to-FE field induced switching of high quality ceramic materials of PLZT x/90/10 ceramics with La3+ composition at the FE/AFE phase boundary. The presented findings are important to the fundamental studies of field-induced phase transitions in AFE s. From the above presented results the main following remarks may be mentioned: 186 I. The study of P(E) loops indicated that the AFE phase can be induced at ambient temperature through a compositional modulation of PLZT x/90/10 system. The ground state is AFE at ambient temperature for PLZT virgin compositions at the FE/AFE border. The AFE PLZT ceramics shows the ability to switch between an antipol ar (AFE) and a polar (FE) state under high electric field , similar with PbZrO 3-based ceramics. The polarization results demonstrate that the stability of the induced FE phase for ceramics in the FE/AFE compositional border can be tailored through the compositional substitution of Pb2+ by La3+. The recovery of the AFE state during electrical field reversal is highly influenced by the strength of the AFE order. The high field induced FE phase remains and the AFE phase is not recovered upon removal of the applied field for PLZT x/90/10 ceramics wit h x≤3.1 (the AFE order is weakened) . The two competing phases are nearly equally stable for x=3.2 and x=3.3 and the electric field -induced AFE -to-FE transition becomes irreversible. With still higher La3+ content, i.e. x≥3.5, the as processed ceramics are dominantly AFE at room temperature . For these compositions, the AFE phase is stable and is almost completely recovered during unloading of the applied electric field (the AFE order is str ong). In addition, these findings confirm the crystallographic resul ts from Chapter 4, where the c oexistence of AFE and FE phases or a pure AFE phase for virgin compositions away from the FE/AFE region with a high excess i n La3+ was found. More specific, there is a broad phase boundary between FE and AFE phases at ambient temperature in these compositions. EAF field increases as the amount of AFE phase increases in PLZT x/90/10 compositions. Therefore, the La3+ chemical modifier helps to stabili ze the AFE phase and hence to manipulate the AFE -to-FE phase transition in PLZT x/90/10 based ceramics . II. In situ XRD and polarization studies provides significant insights in understanding the AFE- to-FE field induced transition of AFE materials. It is widely accepted that typical AFE double hysteresis loops indicate the reversibility of AFE- to-FE field-induced transition. In this work we demonstrated that the AFE- to-FE phase transition of composition with x=0.04 which show double hysteresis loops is not fully reversible. Both investigated composition with (x= 0.03 and 0.04) turn to significant preferred orientation in the final AFE state after the exposure to threshold field necessary to induce AFE- to- FE phase switching. An electric field induced structure develops in both compositions during the AFE- to-FE phase switching . On the other side it is widely accepted that the AFEs which show single ferroelectric loops are characterized by an irreversible AFE- to-FE field assisted transition ( i.e. the AFE state is not recovered during decreasing of electric field or after field removal). More important, we demonstrated that electric fields can induce an AFE phase out of a FE phase if previously poled in the composition x=0.03. This transition is unusual since it has been widely accepted that electric fields favor the FE phase with parallel electric dipoles over the AFE 187 phase. In addition, the current results demonstrate that the FE- to-AFE switching is diffuse and it takes place in a broad field range. This kind of transition has rarely been reported previously. III. A remarkable result of this study is the very high remnant polarization Pr~58 µC/cm2 of PLZT 3/90/10, PLZT 3.2/90/10 and PLZT 4/90/10 compositions, achieved during the application of electric field E≥E AF of very low frequency f~1/300 Hz. There are not similar previous reports on AFE hysteresis loops for any AFE material. To unveil the mechanisms involved in these polarization responses, it is necessary to better understand the domain switching mechanism in similar conditions. In particular, it is needful to reveal whether nucleation and growth of FE domains together with the change of the structure contribute to the switching process. Further, the structure of high field induced FE phase in PLZT 3/90/10 was investigated . Rietveld study of ex situ HXRD data obtained for powders from poled PLZT 3/90/10 composition stated th at the system is monoclinic. Therefore, the crystallographic results confirm without any doubt the enhancement of FE polarization on a longer timescale. The structural changes of the phases involved in the phase -change functional response deserve a further deeper study. In view of these findings, the early hypothesis of the R3c space group representing the structure of the high fi eld induced FE phase of PLZT samples seems to be not verified , at least in the AFE/FE boundary compositional region. The high values of polarizat ion and the induced monoclinic FE phase may rather suggest the development of a high field induced complex morphotropic phase boundary. IV) In the previous chapter, the structural analysis of PLZT x/90/10 virgin powders failed to provide a relationship between the anomalous piezoresponse of x = 0.30 and structural factor(s). Ex situ study of electric-field induced structural changes revealed that the composition exhibiting the highest piezoelectric response is the one which also exhibits significantly enhanced polarizability of the lattices of both coexisting (monoclinic and orthorhombic) phases. The enhanced piezoelectric properties may arise from the irreversible transition from orthorhombic to monoclinic , since after the electric al poling a significant fraction of the monoclinic phase remains in the poled PLZT 3/90/10. V) A detailed study regarding the composition and electric field-dependent energy storage properties in PLZT ceramics across gthe FE-AFE boundary has been undertaken. The highest recoverable energy density of 1 .8 J cm−3 and a high efficiency (η~60) have been achieved on 3.5% of La modification (located in orthorhombic and rhombohedral phase coexistence region) under the applied electric field of 65 kV/cm. This behavior has been explained on the basis of competing FE and AFE phases during the field-induced transition. Therefore, these results suggests that these AFE’s have application potential for use as energy storage. These results highlighted also the possibility tha t 188 high energy-storage density could be obtained in other similar system which show FE and AFE phase’s coexistence. The results of energy storage properties were published in the dedicated paper: I.V. Ciuchi , L. Mitoseriu, C. Galassi, Antiferroelectric to Ferroelectric Crossover and Energy Storage Properties of (Pb 1-xLax)(Zr 0.90Ti0.10)1-x/4O3 (0.02 ≤x ≤0.04) Ceramics , J. Am. Ceram. Soc. 99(7) (2016) 2382-2387. 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App. Phys. 119 ,(12) (2016) 124106- 124106. [68] X. Hao, J. Zhou, S. An, Effects of PbO Content on the Dielectric Properties and Energy Storage Performance of (Pb 0.97La0.02)(Zr 0.97Ti0.03)O3 Antiferroelectric Thin Films, J. Am. Ceram. Soc. 94(6) (2011) 1647-1650. 195 CHAPTER 7 Study of temperature -induced phase transitions in PLZT x/90/10 ceramics 7.1 Introduction The functional properties of perovskite s are intimately related to the complex displacive phase transitions that readily occur. The stability balance among the ferroelectric (FE), ferrielectric, paraelectric (PE) and antiferroelectric (AFE) states have been one of the most intriguing topic of ferroelectric ity [1-4]. The first information about the structural transition which occur s in PLZT x/90/10 arise from the early phase diagram of G. H. Heartling [3, 4] . Recent studies reveal more details about the thermal of stability of FE/AFE state in PLZT x/90/10 [5-7]. In the majority of publications reporting properties of PLZT x/90/10 -based solid solutions, the researchers hold to an accepted point of view that the following sequence of the structural phase transformations described in [5-7] takes p lace during heating in PLZT x/90/10 compositions: a low temperature FE phase to high temperature FE phase and then to PE phase (FELT-FEHT-PE) for compositions which have a predominant FE phase (x≤2) and AFE -PE for compositions predominantly AFE (x≥3). In t hese early studies it was proposed a scenario concerning the general tendencies in the formation of different phase states in few solid solutions across the FE/AFE. More recently E. Buixaderas et al. [8, 9] and A. Pelaiz -Barranco et al. [10] show ed new evidences concerning the phase transition s sequence of PLZT x/90/10 which are different than the previously accepted one. More precisely , E. Buixaderas et al. have shown the presence of phonon anomalies in Raman spectrum of ~200 K below TC, which was assigned to another phase transition to a FE state with double unit cell [8, 9] . A. Pelaiz -Barranco et al. evidence d that the phase transitions sequence for PLZT 3/90/10 on heating from room temperature up to 250oC is: FE-to-AFE (~90oC), AFE -to-FE (~170oC) and FE -to-PE (~190oC) [10]. However , this study is mainly based on polarization P(E) study under the effect of temperature. Th ese results (obtained during heating and under the effect of electric field) may be different with respect to ones of the virgin sample s, as we demonstrated in the previous chapter, where we found for this composition irreversible AFE -to- FE phase switching during application of electric field of adequate amplitude. Therefore , the 196 previous results on the phase transitions of PLZT x/90/10 under the effect of temperature have to be completed by more accurate investigations . The aim of this chapter is to show a more complete picture of the phase transitions for PLZT x/90/10 ceramic swith La -content in the range from 0 to 0.4 a t.%. In our study , the solid - solid phase transitions of PLZT compositions are characterized by using the combined results of dielectric measurements, in situ temperature XRD and Raman spectroscopy. These combined techniques are expected to contribute in assessing the temperature stability of long - and short - range structures of PLZT x/90/10 and to solve the previous controversi al aspects concerning the phase transition sequence. Chapter 4 have shown that PLZT x/90/10 ceramics develop strong piezoelectricity through electric field -induced phase transitions during the poling process. Many poled AFE compositions exhibit an AFE state above the so -called thermal depolarization temperature [11, 12] . A similar behaviour may be expected in the presen t PLZT compositions .. It is important to mention that all the data reported in this chapter were recorded during increasing temperature sequence s. 7.2 Dielectric study In the following , it will be presented and discussed the dielectric results regarding th e influence of La content on the phase transition s of PLZT x/90/10 and then, the influence of the frequency on the phase transition s. Further , the Curie Weiss and modified Curie Weiss laws were applied to analyse the FE/AFE -to-PE phase transition s. In addi tion, the influence of poling on the phase transition sequence for some representative compositions (PLZT 3/90/10, PLZT 3.1/90/10 and PLZT 4/90/10 ) will be shown . 7.2.1 Influence of La addition on the phase transition s of PLZT x/90/10 ceramics The r esults of real part of permittivity ( ε’), imaginary part of permittivity ( ε”) and dielectric losses ( tan δ ) vs. temperature, as obtained by measur ing the capacity during heating from 25 to 350oCC at frequency 100 kHz of are shown in Fig. 7. 1 a)-c). Several interesting features are visible in the temperature dependencies of ε’and tan δ of PLZT x/90/10 with different La content. For each composition, the largest main peak observed in higher temperature range, is associated to the temperature of the AFE/FE -PE phase transition ( Tm), according to previous reports for similar 197 materials [5, 7]. It is known that the increase of La content in PLZT has two important effects: it reduces TC and induces a relaxor behaviour [13, 14] . Fig. 7. 2 shows the changes in εm and Tm as a function of La addition obtained at frequency of 100 kHz. Indeed, doping PZT 90/10 with lanthanum has a significant effect on the location and value of permittivity maximum εm. By increasing the La amount from 0 to 3 at. a %, ca dramatic decrease in Tm from 260 to 190oC is noticed . The variation of Tm for La additions from 2.5 % to 4% is relatively small (within 5oC). The obtained value for the ε m for the PLT 90/10 composition was found to be around 15 x103. 50 100 150 200 250 300 3500,010,1 (d) on heating @ 100 kHzT2tan  Temperature (oC)Tm T1 Fig. 7. 1 a) Real part of permittivity ( ε’), b) imaginary part of permittivity ( ε”), c) detail from a) and d) dielectric losses ( tan δ ) of PLZT x/90/10 compositions measured during heating at a fixed frequency of 100 kHz. 100100010000 50 100 150 200 250 300 350101001000(b)' on heating @ 100 kHz(a)" Temperature (oC) 0 2,00 2,50 3,00 3,10 3,20 3,30 3,50 3,80 4,00 20 30 40 50 60 70200300400500' Temperature (oC)(c) 0 2,00 x 198 A further increment in the La content resulted in a continued suppression of the main dielectric peak (related to the AFE/FE -PE phase transition) and shift in the permittivity maximum to lower value s (⁓5x103) as shown in Fig. 7. 1 and Fig. 7. 2. This may be explain ed with the fact that the long range FE interactions are weakened as the La content increases. We demonstrated in a previous Chapter th at for compositions with La content above x≥0.0250, the AFE state begins to become more stable than the FE state. Tm, now associated with the formation of the AFE rather than the FE state, does not shift down with additional increment of the La content (as shown in Fig. 7. 2) but is rather stabilized. This suggest s that the interactions within the AFE sublattices are less affected by the La content. In addition, the short range AFE coupling is less affected by defects than the long range dipolar interactions responsible for the FE state [14, 15] . Two peaks or inflections were observed in the imaginary part of the dielectric permittivit y and dielectric losses, in temperature range lower th an the FE/AFE -PE temperature, for PZT 90/10 and PLZT 2/90/10. The signature of these last two transitions is more visible in the loss peak, which may be explained as due to a change in the mobility and amplitude of charge and polar relaxations, which are influenced by the octahedral tilting [16]. Both inflection points are marked by arrows in the Fig. 7. 1c). The anomaly indicated with T 1 near 80°C is a sign for the structural change between R3m and R3c of the FE phase and is related to the gradual increasing of the oxygen octahedra tilt [17, 18] . The other dielectric anomaly T 2 looks like a very small step observed in temperature range between the tilt transition T 1 and T m and it was interpreted in literature as due to the coupling between tilt and polar modes [19]. A similar transition has been previously reported for Zr -rich P bZr 1−xTixO3 (PZT) compositions [20]. From the analysis of dielectric curves of Fig. 7. 1a)-c), it can be observed that T 1 increases, while T 2 decreases with increasing amount of La. For compositions with x≥3% these transition are not anymore observed and this effect is a consequence of the AFE phase stabilization induced by La modification. PLZT x/90/10 compositions with x>0.030 show only the a nomaly at Tm (Fig. 7. 1). As previously shown by the structural analysis , the FE phase persist until relatively high x , but the short range AFE coupling is stronger because of the disruption of the long range dipol ar FE interactions . This make s the AFE behaviour as being dominant in these compositions . In the high temperature paraelectric region, ε’ again increase s when te mperature increase s. Similarly is observed for ε” and tan δ vs. temperature p lots. This phenomenon is caused by the dielectric relaxation of the space charge polarization and it is associated with ac conduction effects [21]. It seems that t he space charge polarization is enhanced with increasing the amount of lanthanum , as indicated by the increase of loses in the same temperature range. This phenomen on is caused by the p -type conductive carriers and the vacancies of lead or oxygen [22, 23] . 199 A careful anal ysis has reve aled an additional anomaly in the tan δ (T) or ε”(T) plots with a corresponding anomaly in ε’ at temperature of 35C. It is evident from Fig. 7. 1 a)-c) that this maximum shift s below room temperature for all the compositions x≥3. The corresponding temperature is found to decrease with increasing La content. The origin of this anomaly a t 300 K is not yet clear. Similar with other reports in literature [24], this can be a signature of a nonpolar antiferrodistortive (AFD) transition (doubling of the unit cell). For example, such AFD transition is accompanied by a small dielectric peak in t he permittivity in high Zr -content of PZT [24] and by a broad feature in the dielectric loss and of the anelastic spectra of PZT with composition at morphotropic phase boundary (MPB) (x ≤0.054) [25]. Fig. 7. 2 Maximum permittivity εm and its corresponding temperature Tm as a function of La composition obtained at 100 kHz. This anomaly was assigned to the gradual appearance of the oxygen octahedra tilt at MPB of PZT [16, 26] . As result from our data, this peak is rather broad and th e broadening observed only in the temperature dependence of tan δ may also occur due to extrinsic reasons, such as defects/impurities dynamics. However, the temperature corresponding to the anomaly in ε’(T ) is independent of the measur ing frequency and it can rather be regarded as being of intrinsic origin and linked with the AFD transition. It is well known that FE states show significantly higher losses than AFEs [27]. Another possible interpretation of these results is that this comp osition may have 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5190200210220230240250260270 051015202530 Tm La at %@ 100 Hz m (o C)m*103 m 200 a broad transition from FE -to-AFE state. In fact, a broad dielectric peak was observed near 350 K, which has previously been considered to be a monoclinic to orthorombic transformation in lead – free (K 0.5Na0.5)NbO 3-5%LiNbO 3 solution s [28] and in Pb(In 1/2Nb1/2)O3-Pb(Mg 1/3Nb2/3)O3-PbTiO 3 single crystals [29]. NaNbO 3 present s successive phase transitions during cooling from non -polar (PE) to AFE and finally reaching a FE state [30]. Similar with another AFE material (Pb 0.97La0.02(Zr 0.75Sn0.25xTix)O3) [31] the transition may be interpreted as a phase transformation from FE -to-AFE. In fact this maximum is more evident for PLZT x/90/10 composition with x≥3 , which ha s a dominant AFE character . In addition, this anomaly can be less evident for PZT 90/10 and PLZT 2/90/10 if the transition take s place between two FE states of very similar energies. It is worth to remind that in Chapter 4 the Raman study show ed evidences for a transition from monoclinic to rhombohedra/orthorhombic in a similar temperat ure range. Therefore , this anomaly also may be linked to this structural transition. As the La increases, the phase transition temperatures from FE -to-AFE gradually shift to lower temperature s. It results that the increase of La content in PLZT x/90/10 cer amics improve s the AFE properties. This is in agreement with crystallographic studies on similar PLZT x/90/10 compositions [10]. The phase stability and sequence of phase transition s of PLZT x/90/10 compositions should be similar with ones of PLZT x/95/5 system. As reported by Dai et al. for PLZT x/95/5 , the low – temperature AFE state is a co mmensurate orthorhombic and the high -temperature an incommensurate orthorhombic. At temperatures between the PE and incommensurate AFE states, a FE phase with slim P(E) loop (relaxor -like) was found at lower La contents (x ≤ 4 at.%) , while a multicell cubic region is shown at higher La contents (x>3 at.%) [14]. In many PZ crystals and ceramics (high -Zr-content lead zirconate titanate compositions) a transie nt FE phase of rhombohedral symmetry was observed in a narrow temperature interval ( ∼10 K) close to AFE -PE phase transition [32-34]. Its existence seems to be related to the defect concentration and stoichiometry [32-35]. In the Pb 0.99Nb0.02 [(Zr 0.57Sn0.43)1-yTiy]0.98O3 (PNZST 43/8/2) ceramics a shoulder observed in the temperature dependence of ε’ near 75°C represents the transition from the FE -to-AFE phase. At lower y -values, this transition moves to lower temperatures, due to a significant increase from the AFE phase to the FE phase [36, 37] . Despite of the similarity of our PLZT x/90/10 system with PLZT x/95/5, PNZST 43/8/2 and high -Zr-content lead zirconate titanate compositions , there is no evidence of the above presented phases transition sequence in our dielectric data. Investigation of the temperature stability by other techniques, like Raman and in situ temperature XRD, are necessary to further clarify th ese aspects. 201 7.2.2 Influence of frequency on the phase transition of PLZT x/90/10 ceramics In orde r to give more insight in the phase transition of th ese compositions, the influence of frequency on the dielectric properties as a function of temperature have been investigated in detail. The real part of the dielectric permittivity and dielectric losses for the PLZT x/90/10 as function of temperature at several frequencies (0.1, 1, 10, and 100 KHz) are shown in Fig. 7. 3, Fig. 7. 4, Fig. 7. 5, Fig. 7. 6 and Fig. 7. 7. For PLZT x/90/10 ceramics with 0≤x≤2 La content, the FE/AFE – PE phase transition temperature ( Tm) did not show any frequency dependence , which is typical of ‘normal’ FE -PE phase transitions. Fig. 7. 3 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PZT 90/10 and PLZT 2/90/10 ceramics. 50 100 150 200 250 30002000400060008000100001200014000160001800020000 80 85 90 95 100 105 110220240260280300320Dielectric permittivity Temperature (oC)Real part of permittivity (’) Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPZT 90/10 a) 50 100 150 200 250 3000,020,040,060,080,100,120,14b) 80 100 1200,0260,0280,0300,0320,0340,0360,0380,040Dielectric losses Temperature (oC) PZT 90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses ( tan ) Temperature ( oC) 50 100 150 200 250 30005000100001500020000250003000035000 50 1001000Dielectric permittivity Temperature (oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 2/90/10Real part of permittivity (’) Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 2/90/10c) 0 50 100 150 200 250 3000,000,020,040,06d) 100 1500,04PLZT 2/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses Temperature (oC) PLZT 2/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses ( tan ) Temperature ( oC) 202 In addition to this transition, different dielectric dispersion s were observed in various temperature ranges. Using the highest dielectric peak (which signal s the transition from FE/AFE – to-PE) as a reference , we may divide the temperature data in : high temperature range (temperature s higher then one corresponding to the FE/AFE -PE inflection ) and the low temperature range, for lower temperature s. Permittivity rise at low frequencies while the diele ctric losses decreases as frequency increase at high temperatures in the PE phase is ue to a finite conductivity [21]. This phenomen on almost disappears at higher frequencies, showing that it could be related to a low frequency relaxation process corresponding charged species with low mobility [23]. This behaviour is similar for all investigated PLZT x/90/10 compositions. Fig. 7. 4 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 2.5/90/10 and PLZT 3/90/10 ceramics. Fig. 7. 4 a)-d) show the temperature dependence of dielectric permittivity and dielectric losses for the composition PLZT 2.5/9010 ( Fig. 7. 4 a) and b)) and PLZT 3/90/10 ( Fig. 7. 4 c) and d)) obtained at several frequencies. Depending on the temperature range, the dielectric losses for 50 100 150 200 250 30005000100001500020000 100 15012001400 Dielectric permittivity Temperature (oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 2.5/90/10 Real part of permittivity (’) Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 2.5/90/10 0 50 100 150 200 250 3000,000,010,020,030,04PLZT 2.5/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses ( tan ) Temperature ( oC) 50 100 150 200 250 3000100020003000400050006000700080009000100001100012000Real part of permittivity () Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 3/90/10c) 50 100 150 200 250 300 3500,000,020,040,06d) PLZT 3/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses (tan ) Temperature ( oC) a) b) 203 PLZT 2.5/90/10 and PLZT 3/90/10 show different frequency depend ences. Only one peak related to the FE/ AFE -PE transition can be distinguished in Fig. a) and c). Similar with previous investigated compositions the dielectric permittivity peaks does not show frequency dependence. However , the peak corresponding to the s ame transition from dielectric losses show a slightly frequency dependence : it decreases in amplitude and it shift to higher temperature as frequency increase from 100 Hz to 100 kHz. . Fig. 7. 5 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 3.1/90/10 and PLZT 3.2/90/10 ceramics. The temperature range 25 -100C is dominated by a broad maximum which shift s to lower temperature with increas ing frequency . The amplitude of this maximum increase as frequency increases from 100 Hz to 100 kHz. In the temperature range from 100C up to temperature at which AFE/FE -PE transition occurs ( 180C) it can be observed a slight change of the slope of the dielectric l osses. In this temperature range the change of the amplitude dielectric losses of the 50 100 150 200 250 300010002000300040005000600070008000900010000a) Real part of permittivity (’) Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 3.1/90/10 0 50 100 150 200 250 3000,000,010,020,03b) PLZT 3.1/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses (tan ) Temperature ( oC) 50 100 150 200 250 300010002000300040005000600070008000c) 50 100 15070080090010001100120013001400 Dielectric permittivity Temperature (oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 3.2/90/10 Real part of permittivity (’) Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 3.2/90/10 0 50 100 150 200 250 3000,000,010,020,03d) PLZT 3.2/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses ( tan ) Temperature ( oC) 204 PLZT 2.5/90/10 is less affected by frequency while dielectric losses of PLZT 3/90/10 do no present any dielectric dispersion. Fig. 7. 6 Temperature dependence of the real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 3.3/90/10 and PLZT 3.5/90/10 ceramics. The dependence on temperature of real part of permittivity and dielectric losses of PLZT 3.1/90 /10, PLZT 3.2/90/10, PLZT 3.3/90/10, PLZT 3.590/10, PLZT 3.8/90/10 and PLZT 4/90/10 show a similar behaviour with change of frequency and those will be analysed together. The results are shown Fig. 7. 5 a)-d), Fig. 7. 6 a)-d), and Fig. 7. 7 a)-d). The maximum observed in the temperature dependence of real part of permittivity do es not show any change with increase of frequency from 100 Hz to 100 KHz. This tendency is similar for all the compositions. Contrary, the anomaly of dielectric losses corresponding to FE/AFE -PE phase transition is highly dependent on frequency. The maximum of the anomaly is shifted at higher temperature as frequency increases while its amplitude is decreased. In addition, the peak become s broader as La content increases. 50 100 150 200 250 300 35002000400060008000a) Real part of permittivity (’) Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 3.3/90/10 50 100 150 200 250 300 3500,000,010,020,030,04b) PLZT 3.3/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses ( tan ) Temperature ( oC) 50 100 150 200 250 300050010001500200025003000350040004500500055006000c) Real part of permittivity (’) Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 3.5/90/10 50 100 150 200 250 300 3500,000,010,020,030,040,05d) a)PLZT 3.5/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses ( tan ) Temperature ( oC) 205 The broad peak observed in the temperature dependence of dielectric permittivity is the main feature of FE with a diffuse phase transition [38-40]. Fig. 7. 7 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 3.8/90/10 and PLZT 4/90/10 ceramics. If for the previous compositions (x≤3) the conduction effect were observed only for temperature s corresponding to the paraelectric range, for composition with x≥3 the ac conduction phenomena are shifted at temperature lower (up to 100C), then the temperature at which FE/AFE -to-PE transition occurs. As it results from Fig. 7.5 b) and d), Fig. 7.6 b) and d ), and Fig. 7.7 b) and d) for temperature s higher then 100C the dielectric losses for PLZT x/90/10 compositions with x≥3 decrease with increas ing frequency and it merge together with the curve of dielectric losses obtained for the PE temperature range. For temperature range below 100C the dielectric losses increases as frequency increase from 100 Hz to 100 kHz. 50 100 150 200 250 30050010001500200025003000350040004500a) Real part of permittivity (’) Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 3.8/90/10 50 100 150 200 250 300 3500,000,020,040,060,080,10b) PLZT 3.8/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses ( tan ) Temperature ( oC) 50 100 150 200 250 300 350 40010001500200025003000350040004500c) Real part of permittivity (’) Temperature ( oC) 0,1 kHz 1 kHz 10 kHz 100 kHzPLZT 4/90/10 50 100 150 200 250 3000,000,020,040,060,08d) PLZT 4/90/10 0,1 kHz 1 kHz 10 kHz 100 kHzDielectric losses (tan ) Temperature ( oC) 206 7.2.3 Analysis of FE/AFE -PE phase transition with Curie Weiss law The permittivity of a first -order normal FE can be described in the paraelectric state by the Curie -Weiss law : εr=C 𝑇−TCW (7.1) where C is the Curie constant and Tcw is the Curie -Weiss temperature at which the deviation begins and the local order polarization is induced. These two parameters may provide considerable information of the FE transition, including the transition type (displacement or order/disorder) and the order of the phase transition (first or second order, is reflected by the Curie constant) in FEs. Like FE materials, many AFE s show a high dielectric constant peak near the polar -to-antipolar transition temperature driven by the softening of the lattice vibration [41]. A similar l aw was suggested by W. Cochran for the transition from AFE -to-PE state [42]. According to his theory the transition from PE -to-AFE state when cooling can be associated with a thermally -induced instability of a different mode whose wavelength is on the order of a lattice parameter. Hence the Curie Weiss law was successfully applied to study AFE materials [2, 14] . In the fol owing the Curie –Weiss l aw will be also applied to describe the characteristics of the FE/AFE -to-PE transition of the PLZT x/90/1 0 ceramics. The dielectric data measured at 100kHz reported in Fig. 7. 8 were fitted with Eq. (7.1) and the values of TCW and Curie constant were evaluated. Fig. 7. 8 a) shows a plot of the reciprocal dielectric constant as a function of temperature for PLZT x/90/10 compositions at 100 kHz , wher e the dashed lines are extrapolations using the Curie -Weiss relationship. With the increase of La addition, TCW of the ceramic samples distinctly shows a descending trend. As shown in Fig. 7. 8 b) the deviation from the Curie –Weiss, given by relation ∆T=(Tcw -Tm) where ∆𝑇 = 0 for a second order transition, while for the first order transition ∆𝑇≠ 0, increase s when the amount of lanthanum dopants increase s.. A small difference between Tcw and Tm as found for PLZT 2/90/10 and PLZT 2.5/90/10 indicates that the spontaneous polarization abruptly vanishe s at the temperature Tm and tail to zero at lower temperature s. It is well known that La induce dispersive behaviour and thus , deviations from the Curie –Weiss behaviour may be expected on higher La content compositions. Indeed , ∆T gradually increases as La content increases in the PZT 90/10 system. The value of ∆T=80 C for PLZT 4/90/10, impl ies that the diffuse phase transition (DPT) behaviour has strengthened with increasing the addition of La. The dielectric response and Curie temperature should be frequency independent for a second -order phase transition. In order t o understand better this behaviour, the frequency dependences of TCW 207 was determined. Fig. 7. 8 c) shows a comparison of the T CW as a function of composition for several frequenc ies. From these data, it is clearly observed that PLZT x/90/10 concentration s above x=0.3 present a dispersion of the determined values of TCW which increases with increasing frequency. The dielectric dispersion below Tm indicates a typical relaxor FE behaviour arising from the responses of polar microdomains within the relaxation spectrum. The structural disorder may arises as a consequence of the complex defect chemistry, structural defects (vacancies and Fig. 7. 8 (a) Reciprocal dielectric constant 1/’ as a function of temperature at 100 kHz for PLZT x/90/10 compositions (the dashed lines are fittings with the Curie -Weiss law); (b) Tm and TCW vs. La addition; (c) Frequency dependence of Curie Weiss temperatures mobile ions) [43] which may be associated with the substitution of La in PZT 90/10 lattice. It is reported that the decrease of the transition temperature with La doping in PZT system is in fact a 50 100 150 200 250 300 3500.00000.00050.00100.0015 50 100 150 200 250 300 350 4000,0000,0010,0020,0030,0040,0050,0060,007 Temperature ( oC)1/’ 0 2,00 Temperature ( oC)1/’ 0 2,00 2,5 3 3,1 3,2 3,3 3,5 3,8 4 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0120150180210240270 (o C) La at % Tm Tcw@ 100 kHz 0,0 0,5 1,0 1,5 2,0 2,5 3,0 3,5 4,0120140160180200220240260TCW (oC) La (at. %) 1kHz 10 kHz 10 kHz 100 kHz x b) a) c) 208 function of the concentration of the point defects and dilatation centers [15, 44] . This dispersion is similar to that previously found in PLZT 100x/95/5 by Handerek et al. [45] and to that reported in La-doped lead titanate (Pb,La)TiO 3 by Kuwabara et al. [46]. The decrease of TCW with La doping probably reflects also a weakening of the long -range coupling between dipoles. As result from structural analysis at room temperature presented previously, the long -range FE interactions may persist until relatively high La contents ( i.e. x≤3); however, the short -range AFE couplings become dominant due to the disruption of the long -range dipole interactions induce by La addition ( i.e. x≥3.1). The dielectric parameters were found to be frequency de pendent which may suggest that a second order phase transition rather than a dipole glass and/or relaxor polar phase take s place. Strong frequency dispersion around TCW suggest s a diffuse phase transition (DPT) for PLZT x/90/10 compositions with x≥3. This transition needs to be re -evaluated by an adequate model. 7.2.4 Analysis of FE/AFE -PE phase transition with modified Curie Weiss law PLZT x/90/10 ceramics with x≥3 exhibit b road εm with strong frequency dispersion abround the dielectric maxima (from tan δ vs. T plots) and a significant departure from the Curie – Weiss behaviour, which suggests a ‘‘diffuse’’ FE/AFE -PE phase transition for these compositions. The diffuse transition is generally explained in terms of a distribution of the local transition temper atures in different regions of the sample with respect to the average transition temperature TC, due to local statistical fluctuations in the composition [41, 43] . In the following the diffuse character of the phase transition from the PE phase into the or dered FE or AFE state will be discussed . In order to examine the diffusivity of the FE/AEF -PE phase transition of PLZT x/90/10 solid compositions, a modified Curie –Weiss formula has been used for analysing of temperature dependences of the real part of die lectric constant. log⁡((1 𝜀−1 𝜀𝑚𝑎𝑥)=𝛾log(𝑇−𝑇𝑚)−log𝐶′ (7.2) where C′ is the modified Curie –Weiss constant, and λ is the diffuseness exponent, which changes from 1 to 2 for normal ferroelectrics to relaxor ferroelectrics, respectively. In the case of γ = 1, a normal Curie –Weiss law is obtained and describes a normal FE -PE phase transition [47]. 209 The value of γ was employed to investigate the degree of diffusivity of the phase transition for the present PLZT x/90/10 compositions. The λ (diffusivity) value was determined from the slope of Fig. 7. 9 Log(1/ε -1/εm) vs. Log(T−T m) for two representative PLZT compositions log(1/ε r − 1/εm) versus log(T − T m) plot, as shown in Fig. 7. 9 for PLZT 2/90/10 and PLZT 4/90/10 under different frequencies. A good linear relationship is observed in all specimens. Fig. 7. 10 Diffuseness exponent λ and modified Curie Weiss constant C ’ vs. La content for PLZT x/90/10ceramics Figure 7.10 a)-b) shows the variation of γ and C with La amount in PLZT x/90/10 ceramics. Booth λ and C’ show a similar trend evolution with La content , with a tendency of increase when -0.4 0.0 0.4 0.8 1.2 1.6 2.0 2.4-6.5-6.0-5.5-5.0-4.5-4.0-3.5-3.0PLZT 2/90/10 log (T-T m)log(1/-1/m) 0,1 KHz 1 KHz 10 KHz 100 KHz Liniar fits -0.4 0.0 0.4 0.8 1.2 1.6 2.0 2.4 2.8-6.5-6.0-5.5-5.0-4.5-4.0-3.5-3.0-2.5PLZT 4/90/10 log (T-T m)log(1/-1/m) 0,1 KHz 1 KHz 10 KHz 100 KHz liniar fits 0.00.51.01.52.02.53.03.54.04.51.051.101.151.201.251.301.351.401.451.50 La ( mol %) 100 Hz 1 kHz 10 kHz 100 kHz 0.00.51.01.52.02.53.03.54.04.502468101214161820C’ *10 5 La ( mol %) 100 Hz 1 kHz 10 kHz 100 kHza) b) 210 the concentration of La increases from 0 to 3 at.%. For the samples with x≤3 λ has values between 1.10 and 1.40 , which indicate the diffuse character of the phase transition in the se compositions . The disorder in B site increases with increasing La addition in this compositional range. An anomaly of both λ and C’ parameters can be observed for compositions in the range of 33.5 at.% of La. This trend is similar to the anomaly of the unit cell parameters found in the same compositional range (Fig. 4.17 from Chapter 4). At this stage, we cannot propose an explanation for its origin. 7.2.5 Influence of poling on the phase transitions sequence of PLZT x/90/10 ceramics In order to exploit the high temperature application s of PLZT -based piezoelectric materi als, not only the Curie temperature but also the depoling temperature are of high importance. To further investigate the phase -transition behavior, the temperature dependence of the real part of permittivity ( ε’) and loss tangent ( tan δ ) for selected pole d PLZT compositions was measured (PLZT 3/90/10, PLZT 3.1 /90/10 and PLZT 4/90/10). Fig. 7. 11 shows the temperature – dependence of ε’ and tan δ (at selected frequency of 10 kHz) collected during heating from ceramic samples after their exposure to high poling field of 30 kV/cm at room temperature for 30 min. A similar peak as for virgin samples at Tm is observed in poled ceramics. Interesting, the FE/AFE -to- PE phase transition of poled PLZT is affected by electric field. The poled PLZT 3/90/10 display s a much lower permittivity ε’m~4300 at its temperature Tm of 182°C with respect to the virgin sample , which shows a permittivity ε’m~1100 0 at its Tm of 199°C ( Fig. 7. 4c) and d)). Th e poled PLZT 3.1/90/10 presents a permittivity of ε’m~10000 at its Tm of 200°C while in virgin condition it show s a slightly lower permittivity ε’m~9000 and Tm of 196°C. In addition to the peak corresponding to the FE/AFE -to-PE phase transition, another broad anomaly (a hump in the tan δ vs. T curve and a steep like increase in the ε’ vs. T dependence) can be observed at lower temperature , at ~100°C for PLZT 3/90/10, at ~70°C for PLZT 3/.190/10 and at ~35°C for PLZT 4/90/10 , respectively . The dielectric p roperties of the poled PLZT ceramics are similar to ones reported by Pokharel et al. for AFE Pb 1-xBaxZrO 3 (x = 0.05, 0.10) [48, 49] . They correlated such anomaly to the transition from the AFE orthorhombic phase to FE rhombohedral phase. Similarly, poled compositions with formula (Pb 0.97La0.02(Zr 0.65Sn0.22Ti0.13)O3 and Pb0.97La0.02Zr0.90Sn0.025Ti0.075)O3) show an AFE -FE transition at a temperature referred to as 211 depolarization temperature Td [50, 51] . A more detail ed study was performed recently rega rding the influence of poling process on the phase transition s, as we described in our recent paper [52]. 010002000300040005000 50 100 150 200 250 3000,010,020,030,040,050,06Tm Temperature (oC)Real part permittivity () Dieelectric losses (tan ) ’ TdPLZT 3/90/10 a) tan @ 10 kHz 01000200030004000500060007000800090001000011000 50 100 150 200 250 3000,030,040,050,060,070,080,090,10b) ’ TmPLZT 3.1/90/10 Td Dieelectric losses (tan )Real part permittivity () tan @ 10 kHz@ 10 kHz Temperature (oC) 01000200030004000500060007000 50 100 150 200 250 3000,010,020,030,04c) Dieelectric losses (tan )Real part permittivity () Temperature ( oC) ’Tm TdPLZT 4/90/10 tan @ 10 kHz Fig. 7. 11 Variation with temperature of the real part of dielectric permittivity ε’, loss tangent tan δ, of poled ceramics withy compositions: a) PLZT 3/90/10 sample, b) PLZT 3.1/90/10 and c) PLZT 4/90/10 measured during heating in the temperature range of 25 –300°C at the frequency of 10 kHz. 212 It can be seen that both FE -AFE and AFE -PE phase transitions occur in a wide temperature range, and the dielectric peaks at Td and Tm indicate that poled PLZT ceramics possess a diffus e behaviour of their phase transitions . In particular, the FE -AFE phase transition is widened, and the diffus e character is enhanced with the increase in the La content. Similar as for virgin samples, tan δ in poled PLZT are smaller than 0.03 below 200°C and sharply increase as result of increasing ac conductivity above this temperature. . According to A. Pelaiz – Barranco et al. the conduction behaviour at high temperatures on similar compositions (PLZT 2/90/10 and PLZT 4/90/10) is due to vacancy hopping processes and relaxations in oxygen vacancy -related dipoles [21]. The results presented above indicate that the phase transition s of such PLZT compositions are compl ex and they are accompanied by structural modifications . In order to correlate the results of the dielectric study concerning the phase transitions with information concerning the chang es in the structure induced by temperature and field, more sophisticated technique s are necessary. 7.3 In situ XRD temperature study 7.3.1 Phase transitions in virgin PLZT x/90/10 ceramics To clarify the field and temperature effects on the phase transformational sequence summarized in the previous section and for a better understanding of the PLZT structure -property relation, structural analysis was carried out by collecting X -ray patterns from virgin PLZT x/90/10 crushed ceramics as a function of temperature in the range of 25-300°C. Structural transitions in the material are detected through the observ ation of changes in the diffraction peak splitting. The diffraction profiles measured alo ng the radial (θ -2θ) direction, during the heating cycle, in the pseudocubic {100} pc and {200} pc zones are shown in Fig. 7. 11 and Fig. 7. 12 The pseudocubic diffraction line profiles {100} pc and {200} pc consist i n orthorhombic 100 and 001, 002 and 200 doublets and a residual rhombohedral 100 and 002 singlets. These lines are the most important characteristic s for the AFE phase. The distortion of these peaks is characteristic of the simultaneous presence of the abo ve-mentioned phases in the sample. As the the temperature corresponding to the dielectric peak Tm is approaching the aspect of 100 and 200 peaks chan ges continuously ( Fig. 7. 12and Fig. 7. 13). As temperature increases, t he couple of reflection 001 and 213 100 and 002 and 200 peaks increase in intensity, and those closeness itself quite significantly from each other by sequential shifting of 100 and 200 on the lower 2 ° side. These are indicati ons of the gradual increase of the unit cell volume when increas ing temperature. These changes are expected to influence the AFE/FE properties of PLZT 3/90/10 when the sample is subjected to thermal modifications. In fact, as reported by Ref. [35], the hysteresis behaviour of similar Fig. 7. 12 A contour plot of diffraction intensities as a function of temperature obtained from {001}pc for PLZT x/90/10 ceramics composition changes with increasing temperature from room temperature up to 130°C. According to their results, in this temperature r ange the shape of the hysteresis loops of PLZT 3/90/10 changes from a typical FE loop to a double hysteresis loop, which is a consequence of enhancement of AFE character with increasing temperature. However, unlike in Ref. [35], our 50100 150 200 250 300 35021,221,421,621,8 x=3,1x=2,5 100,011002100Intensity x=3,2PE AFE+FE 50100 150 200 250 300 35021,221,421,621,8 x=3,0 100,01100Intensity 50100 150 200 250 300 35021,221,421,621,8 100,01100Intensity 50100 150 200 250 300 35021,221,421,621,8d)c)b)2222 Temperature (oC)A 100,01100Intensitya) 50100 150 200 250 300 35021,221,421,621,8 x=3,3 100,01100Intensity 50100 150 200 250 300 35021,221,421,621,8 x=3,5 100,01100Intensity x=3,8 50100 150 200 250 300 35021,221,421,621,8 100,01100Intensity 50100 150 200 250 300 35021,221,421,621,8h)g)f)e)2222 Temperature (oC)PE AFE+FE x=4,0 100,0110021003100Intensity 214 dielectric data (Fig. 4 a)) do not show any evidence for a low temperature FE -AFE phase transition. XRD data evidences the AFE interactions are enhanced, as the temperature increases. These modifications indicate that a diffuse FE -AFE transition takes place, in agreement with ot her reports [51]. At the temperature of Tm the features of these peaks change significantly: the double {200} pc and {100} pc diffraction line profiles are transformed in to a singlet through the disappearance of the peak at higher scattering angle. The 200 and 100 Bragg peaks become more evident and this is consistent with the hypothesis of a cubic structure. Fig. 7. 13 . A contour plot of diffraction intensities as a function of temperature obtained for {200}pc of PLZT x/90/10 ceramics 50100 150 200 250 300 35043,443,643,844,044,244,4 x=3,1x=2,5 100,011002100Intensity x=3,2PE AFE+FE 50100 150 200 250 300 35043,443,643,844,044,244,4 x=3,0 100,011002100Intensity 50100 150 200 250 300 35043,443,643,844,044,244,4 100,01100Intensity 50100 150 200 250 300 35043,443,643,844,044,244,4d)c)b)2222 Temperature (oC)A 100,01100Intensitya) 50100 150 200 250 300 35043,443,643,844,044,244,4 x=3,3 100,0110021003100Intensity 50100 150 200 250 300 35043,443,643,844,044,244,4 x=3,5 100,0110021003100Intensity x=3,8 50100 150 200 250 300 35043,443,643,844,044,244,4 100,011002100Intensity 50100 150 200 250 300 35043,443,643,844,044,244,4h)g)f)e)2222 Temperature (oC)PE AFE+FE x=4,0 100,0110021003100Intensity 215 Thus the AFE -PE transition occurs simultaneously via the disappearance of the orthorhombic distortion. The temperature induced FE/AFE -PE phase transition detected by the structural analysis is consistent with the dielectric behaviour from Fig. 7. 1 and Fig. 7. 2. It is worth mentioning that there is no evidence of a FE -AFE phase transition in dielectric data. In addition, the previous presented XRD data did not show any proo f of a low temperature transition ( 35C) from a monoclinic to rhombohedral or orthorhombic as suggested by the dielectric data. Further study is necessary to clarify this aspect. A detailed structural analysis was 5,845,855,86 11,6911,7011,71 8,208,218,228,238,24 -40 04080120160561562563564a (Ao) b (Ao) c (Ao) VIVIII IIVolume (Ao)3 Temperature (oC)I a) 5,845,855,865,87 11,6611,6811,7011,72 8,208,218,228,238,24 306090120150560562564566b) a (Ao) b (Ao) c (Ao) IV III IIVolume (Ao)3 Temperature (oC)V Fig. 7. 14 The evolution of lattice parameters (a, b, c ) and of the unit cell volume with temperature for a) PLZT 3.5/90/10 and b) PLZT 4/90/10 ceramics 216 performed by Rietveld refinement on two representati ve samples by using the XRD patterns obtained below Tm. According with the structural study presented in Chapter 4, the majorit ary phase for PLZT 3 .5/90/10 and PLZT 4/90/10 is AFE with Pbam structure. Only the evolution of lattice parameters and unit cell volume with temperature for PLZT 3.5/90/10 and PLZT 4/90/10 ceramics is shown here ( Fig. 7. 14), where several anomalies are observed in the t emperature range 40 -165C. Usually, such anomalies of the structural parameters are connected with structural phase transitions. However, t he XRD pattern s corresponding to the anomalies did not show any characteristics in order to be identified with a nother symmetry than Pbam. Therefore , for the results presented only the orthorhombic phase Pbam was used to fit the data. From th ese data it is hard to assess the anomalies to a different symmetry, but it is clear that below Tm there are five distinct structures . The se anomalies were confirmed by the Raman analysis and their origin will be discussed below. 7.3.2 Phase transitions in poled PLZT 3/90/10 ceramic The anomalies observed in the dielectric properties near the piezoelectric depolarization temperature and Curie temperature indicate that a structural change occurs at this temperature. To clarify the field and temperature effects on the phase transformational sequence and for a better understanding of the PLZT structure -property relation, structu ral analysis was also carried out by collecting X -ray patterns as a function of temperature in the range 25 -300°C first on poled (during heating) and then on depoled (during cooling) PLZT 3/90/10 ceramics. We now compare and discuss the dielectric properti es and X -ray results, by referring in particular to Fig. 7. 15 and Fig. 7.16, where results of dielectric investigation and of X -ray analysis as a f unction of temperature for the PLZT 3/90/10 composition in the virgin and poled state, respectively, are shown . In addition, the variation of ε’, tan δ and 1/ ε’ as a function of temperature , measured for the virgin and poled ceramic at a fixed frequency of 10 kHz during the heating cycle from 25 -300°C has been included in to Fig. 7. 15 and Fig. 7. 16 for a better correlation between the dielectric and structural properties. The d iffraction profiles measured along the radial (θ -2θ) direction in the pseudocubic {111} pc and {200}pc zones were monitored during heating ( Fig. 7. 15 b)-c)) and cooling cycle s (Fig. 7. 16b)- c)) of the poled PLZT 3/90/10 ceramic. A coexis tence of rhombohedral and orthorhombic phases is observed in the temperature range 25 -75°C during heating of poled PLZT 3/90/10 pellet (Fig. 7. 15 b)-c)), in agreement with one suggested by the aspect of P(E) loops presented in Chapter 6 (where an irreversible AFE -to-FE field assisted transition was observed). 217 020004000600080001000012000 50 100 150 200 250 3000.000.020.040.06 0.00000.00020.00040.00060.00080.0010 50 100 150 200 250 30043.443.643.844.044.244.4AFE PEa) b) c)IntensityIntensity2 2111pc 510.0193333554778620050 100 150 200 250 30037.237.437.637.838.0 ș ș ș ș ș Temperature ( C)Temperature ( C)tan  Dieelectric losses (tan ) Temperature ( C) 500.01375225031254000 200pc’ ’ 1/Real part permittivity (’)Real part permittivity (’) @ 10 kHz TCWTm Fig. 7. 15 a) Variation with temperature of the real part of dielectric permittivity ε’, loss tangent tan δ , and reciprocal permittivity ( 1/ε’) of the virgin PLZT 3/90/10 sample, measured during the heating in the temperature range of 25–300°C at the frequency of 10 kHz. The pink line represents the fit with the Curie -Weiss relation in the high temperature range. The AFE/FE -PE transition tempera ture Tm, and Curie -Weiss temperature TCW are marked by arrows. b) Contour plots of {111} pc and c) {200} pc regions of diffraction patterns for PLZT 3/90/10 . Data from a) were obtained during heating of virgin PLZT 3/90/10 ceramic while the selected areas of X-ray patterns of b) {111} pc and c) {200} pc were obtained during cooling of poled PLZT 3/90/10 ceramic. 218 010002000300040005000 50 100 150 200 250 3000.000.020.040.06 0.0000.0010.0020.0030.0040.005 50 100 150 200 250 30043.443.643.844.044.244.4PE AFEFE + AFE ș c)b) 510.01933335547786200 ș FE 50 100 150 200 250 30037.237.437.637.838.0 IntensityIntensity2a) Temperature ( C)tan ’ ’TCWTdTm ș 2 Temperature ( C) ș Temperature ( C) ș 500.01375225031254000@ 10 kHz 111pc 200pc 1/Real part permittivity (’)Real part permittivity (’) Dieelectric losses (tan  ) Fig. 7. 16 a) Variation with temp erature of the real part of dielectric permittivity ε’, loss tangent tan δ , and reciprocal permittivity ( 1/ε’) of the PLZT 3/90/10 sample after applying a field of 30kV/cm field at frequency of 10 kHz , measured during the heating in the temperature range of 25–300°C. The red line is the fitting with the Curie -Weiss relation from high temperature dielectric data. The AFE/FE -PE transition temperature Tm, depolarization temperature Td and Curie -Weiss temperature TCW are marked by arrows. A contour plot of diffraction intensities as a function of temperature obtained from b) {111} pc and c) {200} pc. Data from a), b) and c) were obtained during heating of poled PLZT 3/90/10 ceramic 219 The pseudocubic diffraction line prof ile {200} pc consist of an orthorhombic 002 and 200 doublet and a residual rhombohedral 002 singlet. There are no obvious changes in the temperature range of 25 -75°C. In the vicinity of dielectric anomaly Td, the appearance of 111 and 200 peaks chan ges. In the temperature range 75 -135°C the 111 rhombohedral peak shift s towards lower 2θ values and its intensity decreases, while the intensity of 111 orthorhombic peak increases. In the same temperature range , the peak 002 became broader and nonsymmetric. Th ese characteristics suggest that the FE interactions are weakened while the AFE are enhanced, as the temperature increases. These changes indicate that a diffuse FE -AFE transition takes place, as also observed by other researchers [51]. Once the temperature of ~135°C is reached, the features of these peaks change significantly: the {200} pc diffraction line profile split into two, while {111} pc become a singlet, being consistent with the orthorhombic distortion. Therefore, the FE -AFE transition was completed at ~135°C . The AFE -PE transition is found at a higher temperature, of 180°C, as shown by the appearance of cubic 111 pc and 200 pc diffraction profiles and a simultaneous disappearance of the orthorhombic 200 peak distorsion. The microstructural evolution du ring heating and the detected phase transitions with temperature are consistent with the dielectric behaviour from Fig. 7. 15 a).On cooling from 300°C down to 180°C the XRD pattern s show a similar evolution with temperature as found for virgin PLZT 3/90/10 sample ( Fig. 7. 13). The PE -AFE transition occurs since the orthorhombic distortion of the Bragg peak becomes evident via the appearance of an additional peak at lower scattering angle ( Fig. 7. 16 b) and c)). From Fig. 7. 15 a) with Fig. 7. 15 b) and c) it can be seen that t he AFE -to-PE transition temperature depends on the heating /cooling sequence . The difference of ~14°C in the transition temperatures between the P E and AFE phases during heating (dielectric data) and cooling cycles (X RD data) indicate s that PE -AFE is of the first-order [53]. Upon further decreasing temperature from 180°C down to the room temperature, the appearan ce of 111, 200, and 002 diffraction peaks changes continuously (Fig. 7. 16 b) and c)) in a similar fashion as for the virgin composition ( Fig. 7. 12 b)). Since the unit cell volume decre ases when decreasing temperature t he 002 and 200 peaks becomes less strong and less broad, also separati ng quite significantly from each other by progressive shifting of 200 on the higher 2° side while t he 111 peak s hifts gradually to higher 2 ° angle It is clear from these data (Fig. 7. 15 and Fig. 7. 16) that the {200} diffraction line is split exactly at ~100°C, in correspondence to the dielectric anomaly corresponding to the FE-to-AFE phase transition during heating , but is absent for the cooling cycle data . Therefore, the recovery of the AFE phase is completed during the heating proces s. In conclusion, these trans itions occur at the AFE -FE and FE -P boundaries of an intermediate FE phase, similarly as found in earlier experiments for Pb0.97La0.02(Zr 0.63Ti0.13Sn0.24)O3 and Pb(Zr 1−xTix)O3 (0≤x≤0.2) FE/AFE 220 compositions [35, 54] . Thus, after heating in the temperature range 25-300°C, the poled PLZT 3/90/10 sample converts , as expected, to the initial state before poling ( AFE /FE phase coexistence). The TCW of poled PLZT 3/90/10 is located in the vicinity of the temperature corresponding to the limit where the recovery of AFE was completed, corresponding to 135°C. 7.4 Raman Spectroscopy Study In the Chapter 4 the Raman study of the phase transitions o f PLZT x/90/10 system below room temperature, has been presented. In particular, anomalous structures were found below room temperature, which could be interpreted by a transition between a monoclinic to rhombohedral/orthorhombic structures. The present study aims to improve the knowledge and to clarify the thermal stability for the intermediate phases in th e FE/AFE region. One can expect that the evolution of Raman modes should be directly connected to the anomalies observed in temperature dependence of dielectric constants and the XRD patterns obtained at variable temperature s. The Raman spectrum of PLZT x/ 90/10 solid solutions at the temperatures from 200 – 600 K are shown in Fig. 7. 17, Fig. 7. 18 and Fig. 7. 19. The assignment of Raman Modes was presented in Chapter 4. The symmetry indexation was not possible because of the polycrystalline character of the samples. In this section only some representative Raman modes will be discussed . In particular , the evolution of the Raman spectra with temperature will be observed, in order to establish the absence /appearance of discontinuous changes that would indicate a phase transition. When the temperature increas es from 200 to 600 K , a continuous modification of the Raman spectra is observed. As it can be seen in Fig. 7. 17, Fig. 7. 18 and Fig. 7. 19 the relative intensities of Raman modes decreas e when increasing temper ature and some modes disappear at characteristics temperatures. In the meantime, broadening, mode merging, and frequency shifting are observed in Fig. 7. 17, Fig. 7. 18 and Fig. 7. 19. The broadening of Raman modes with increasing temperature support the idea that higher structural disorder exists in the Ti -O bond of the TiO 6 octahedra w hen increasing temperature. This behaviour may be associated with the nucleation of nano polar domains within the AFE/FE matrix. These are characteristics which were usually reported in perovskite relaxor ferroelectrics [55] and are in agreement with our previous dielectric study , where a diffus e FE/AFE -PE transition was observed . Consistent with the results of XRD and dielectric study reported in the previous section s, the spectra of PLZT x/90/10 ceramics can be divided into several specific temperature ranges , which shift to lower or to higher temperatures 221 Fig. 7. 17 Raman spectra at different temperatures of PLZT a) 2/90/10, b) 2.5/90/10, c) 3/90/10 and d) PLZT 3.1/90/10 ceramics 0150 300 450 600 750 Wavenumber [cm-1]PLZT 2/90/10 1311652 IVIIIII 500480460440420400380360340320T [K]Raman Intensity250 I1 150 300 450 600 750IPLZT 2.5/90/10 11T [K] 420 440 460 480400380360340300 320250 Raman Intensity Wavenumber [cm-1]13 IV 0150 300 450 600 750PLZT 3/90/10 200 Raman Intensity Wavenumber [cm-1]1311 652 IVIIIII 480440420400380360340300T [K]Raman Intensity250 I1 0150 300 450 600 750200 300 Raman Intensity Wavenumber [cm-1]PLZT 3.1/90/10 1311652 IVIIIII 500480460440420400380360340320T [K] 250 I1a) b) c) d) 222 as a function on the La content . The most noticeable observations can be derived from the dramatic changes of Raman spectra around ~19 cm-1 (mode 1), ~35 cm-1 (mode 2), ~135 cm-1 (mod e 5), ~200 cm-1 (mode 6), ~340 cm-1 (mode 11) and ~500 cm-1 (mode 13). The first temperature range is located from low temperature -250 to 350 K ( -100 °C to ~70°C). The main features reflected in this temperature range are the vanishing of mode 1 and the disappearance of the splitting of all E -symmetry modes into A+A” species , which are connected with the existence of a monoclinic phase (indicated in Fig. 4.27 -from Chapter 4)) [56, 57] . Those are indicati ons of the tranformation of the monoclinic structure into a rhombohedral or orthorhombic one. This transition is better reflected for the composition PLZT 2.5/90/10 ( Fig. 7. 17 b)) where the splitting of the E -symmetry mode is better seen and the transition occurs at higher temperature 360 K (70°C). Therefore, the Raman results support the hypothesis of a phase transition from monoclinic to rhombohedral or orthorhombic in th e unpoled PLZT polycrystalline solid solutions between 25°C and 70°C , depending on the La content. It is worth mentioning that E. Buixaderas et al. observed anomalies in the evolution of Raman modes as a function of temperature below room temperature. They have associated this anomaly with a transition to an another FE state with doubled unit cell , but they also argue d that further studies are necessary to clarify the origin of this anomaly [9]. However, in a more recent paper [8], the same authors argue d that the local symmetry in PLZT X/90/10 is lower than currently described in the accepted phase diagrams. In the temperature range from about 320 -440 K (40°C to 160°C) the weak shoulders of mode 5 and mode 13 show a continuous decrease in intensity as the temperature increases. The mode 5 are characteristic of the FE order , being related the spontaneous polarization and the spontaneous tetragonal strain , while the mode 5 at 510 cm-1 (E(TO4) mode) is characteristic to the rhombohedral phase [37]. Interesti ng, these shoulders disappear in the neighbourhood of the expected depolarization temperature Td value obtained from the poling study. Other characteristics in the Raman spectra are the softening of the mode 2 around 25 cm-1 and the attenuation of the bands around 250, indicating a FE -AFE phase transition. Hence , it is likely that this transition is connected with the loss of FE order and the enhancement of AFE one. The occurrence of FE -to- AFE phase transition is also confirmed by previous studies on simi lar PLZT x/90/10 compositions [10, 58] . The main features observed in the temperature range from 380 -440 K (110°C to 160°C) are the softening followed by the disappearance of the shoulders 5 and the formation of one main band which shift downward in the low -frequency range at about 85 cm-1. The mode 5, at ~135 cm-1, is also a characteristics of AFE commensurate phase while the mode at about 85 cm-1 is typical for 223 AFE incommensurate phase of well -known PbZrO 3 [32, 59] and PbHfO 3 AFE materials [60]. Hence , the vanishing of mode 5 and the individual ization of mode at 85 cm-1 during heating may suggest a transition between two similar AFE structures, incommensurate -to-commensurate, as similar reported for PbHfO 3 [60] and Pb 0.99Nb0.02[(Zr0.57Sn0.43)1−yTiy]0.98 [37]. These new insights regarding the phase relationships, are inconsistent with previous TEM studies on the same compositions presented in Chapter 4 and with the results of I. MacLaren, et al. [61], who reported an incommensurate character of the room temperature AFE phase in PLZT x/90 /10 compositions. However , it should be stressed that TEM, as a technique for local structure analysis, cannot conclusively determine the overall phase structure of a polycrystalline ceramic specimen . In addition, the invasive sample preparation of ceramics for TEM analysis may alter the phase symmetry composition in systems with phase superposition, as is the case of our PLZT ceramics . In many investigations of PbZrO 3- based crystals and ceramics , a transition from A FE-to-FE phase of rhombohedral symmetry was observed in a narrow temperature region ( ∼10 K) below PE phase. The appearance of FE phase seems to be related to the defects concentration and stoichiometry [32-34, 59, 62] . XRD data from Fig. 7. 14 have revealed a change of the lattice parameters in a region below and very close to the temperature range of AFE/FE -PE phase transition , which may be connected to the presence of a high temperature FE phase. A careful inspection of the Raman mode s in the low frequency region (Fig. 7. 19 b)-d)) evidences that the mode at 27 cm-1 is in fact a composite one, which starts to bulge at 30°C and then gradually splits into two peaks (20 cm-1 and 38 cm-1) upon further heating. The low frequency mode (20 cm-1) keeps on shifting up with increasing temperature until disappe ars at temperature corresponding to Tm, while a new mode (38 cm-1) still persist s until high temperature range ( 500 K). The temperature corresponding to the splitting of the mode occurs is in the same range of temperature where anomal ies in the lattice parameters vs. temperature were observed ( Fig. 7. 14). In addition, recent studies of Pelaiz et al. [10] have revealed a slim ferroelectric loop (relaxor -like) just below Tm for PLZT x/90/10 compositions with La -contents (x≤4 at.%). Therefore , the change of the structural parameters and of Raman spectrum of PLZT x/90/10 composition s in a narrow temperature r ange just below Tm is related to the appearance of a high temperature FE phase , as similar ly found for AFE PbZrO 3 compositions [63]. Even if Raman modes are forbidden in the Pm3m space group, the Raman spectra of all PLZT x/90/10 ceramics still show three broad bands at high temperatures, in the cubic phase. The source of these bands presented in the PE cubic phase was studied for tetragonal PLZT x/40/60 [64]. The presence of Raman -active modes could be ascribed to off -centred cations and to the presence o f nanoscale chemical/structural inhomogeneities related to symmetry breakdown at nanoscopic 224 Fig. 7. 18 Raman spectra at different temperatures of PLZT x/90/10 ceramics 0150 300 450 600 750PLZT 3.2/90/10 Raman Intensity Wavenumber [cm-1]3001311652 IVIIIII 500480460440420400380360340320T [K] 250I1 0150 300 450 600 750 Raman Intensity Wavenumber [cm-1]PLZT 3.3/90/10 30013 11652 IVIIIII 500480460420400380360340320T [K] 250 I1 0150 300 450 600 750440 Raman Intensity Wavenumber [cm-1]PLZT 3.5/90/10 3001311 652 IVIIIII 500480460420400380360340320T [K] 250 I1 0150 300 450 600 750200 500 Raman Intensity Wavenumber [cm-1]PLZT 3.8/90/10 4403001311652 IVIIIII 480460420400380360340320T [K] 250I1a) b) c) d) 225 0150 300 450 600 750T [K] Raman Intensity Wavenumber [cm-1]PLZT 4/90/10 500440300 480420400380360340320250 1311652 1 IVIIIIII Fig. 7. 19 Raman spectra at different temperatures of PLZT x/90/10 ceramics 0 20 40 60 80520T [K] PLZT 4/90/10 Raman Intensity Wavenumber [cm-1]440 460420400380 * 010203040506070* 500 Raman Intensity Wavenumber [cm-1]PLZT 3.8/90/10 440 480460420400380T [K] 0 20 40 60 80* Raman Intensity Wavenumber [cm-1]PLZT 3.3/90/10 300 500480460420400380360340320T [K] 250a) b) c) d) 226 scales (polar nanoregions) [55, 64] . Therefore, the presence of Raman signal above T C for our PLZT x/90/10 compositions indicate s that the local symmetry of these compositions contrast significantly from the average cubic perovskite structure. 7.5 Revised phase diagram of the PLZT x/90/10 solid solution The first “composition−temperature” phase diagram for the PLZT x/90/10 system was published early in [3, 4] . A more recent revised phase diagram based on the results of a combined dielectric and anelastic spectroscopy study was proposed in one of our paper s [19]. With r espect to the early diagram proposed by Hearting [3, 4], in addition to the transition from F LT and F HT, a new border has been added on the low temperature La content range: an intermediate curve between TT and Tm lines de termined from the temperatures corresponding to the anelastic and dielectric anomalies , which have been attributed to the intermediate FE phase with tilt instability TIT (Fig. 7. 20). Fig. 7. 20 Phase diagram of PLZT x/90/10, based on structural results, dielectric and anelastic data as published in Ref. [19] The temperature dependent dielectric permittivity, piezoelectric response, and Raman spectroscopy, XRD analysis in a broad temperature range presented in the previous sections have 227 revealed the presence of some unknown phase transitions. We have carefully ident ified the temperature at which these transition s occur for each composition and we joined th ese data as a function of La content. Figure 7.21 represents our proposed revised phase diagram of the PLZT x/90/10 system , where previous results of Pelaiz et al. [7, 10] and Knudsen et al. [5] together with our experimental data are represented together in order to give a more complete pictures about the stability of phases. The data show excellent agreement between phase transition temperatures Fig. 7 .21 The temperature vs. La addition phase diagram for polycrystalline PLZT x/90/10 ceramics. determined by different methods . Several temperature regimes seemingly exist , corresponding to the FE phase with monoclinic structure (F M), a FE low temperature (F LT) region, FE high temperature (F HT) phase, a coexistence of FE low temperature (F LT) and AFE incommensurate (AFE IC), AFE incommensurate (AFE IC), AFE commensurate (AFE C), a high temperature FE phase (FIT/F/FC) and paraelectric (PE) phases, respectively. A border between monoclinic F M and F LT or between monoclinic FM and a coexistence of AFE C and F LT phases was drawn f or all PLZT x/90/10 composition in temperature range close to the room temperature. Depending on the La content, the phase diagram obtained fo r temperature higher then room temperature ( 25C) can be 0,0 0,5 1,0 1,5 2,0 2,5 3,0 3,5 4,020406080100120140160180200220240260280300 20406080100120140160180200220240260280300 AFEIC FLT + AFECFHT Temperature (o C)PECTemperature (o C) FLT FM TM-LT dielectric data TM-LT XRD data TM-LT, Raman data TLT-HT, dielectri data TLT-HT, Pelaiz et al. [7] TFE-AFE, dielectric data TFE-AFE, XRD data TFE-AFE, Pelaiz et al. [10] TCW, dielectric data TAFEc-AFEic, XRD data TAFEc-AFEic Raman data TIT/TF/TFC, dielectric data TIT/TF/TFC, XRD data TIT/TF/TFC, Raman data TIT/TF/TFC, Pelaiz et al. [10] Tm, dielectric data Tm, XRD data Tm, Pelaiz et al. [7] Tm, Knudesn et al. [5] La (at. %)FTT/F/FC AFEC 228 divided in two parts: a low La content range (x≤2) where the FE phase behaviour dominates , and a high La content range (x≥3) where the AFE behaviour is predominant . Below Tm with small La content (x<0.025), one can find a succession of phases similar to the parent phase PZT 90/10. A tilt instability border which separate s the low FE LT (R3c) and high FE HT (R3m) temperature phases, has been evidenced. In addition , a new transit ion, associated with the disappearance of disordered tilting associated with the long range order of the R3c phase, previously found in Zr -rich Pb(Zr,Ti)O 3, is confirmed in the rhombohedral PLZT x/90/10 compositions [65]. With the addition of lanthanum (x≥ 3), the phase sequence change significantly. Clearly, La modification enhances the stability of the AFE state over that of the FE phase during heating. At higher La contents, the temperature stability range of the incommensurate AFE state was stabilized w hile the FE LT was suppressed , since a transition between FE and AFE C was found. The temperature of this transition decreases with increases of La content ranging from 150C for PLZT 3/90/10 to 35C for PLZT 4/90/10. The FE -AFE phase transformation is diffuse , i.e. it takes place over a relatively wide temperature interval (as result ed from the dielectric and XRD data). A narrow temperature (from 110C to 150C) range with incommensurate structure is also shown between AFE IC and high temperature FE (F IT/F/FC) phases for PLZT x/90/10 composition with x≥3.2. On further heating AFE IC disappeare s and a FE F IT/F/FC state was stabilized. This region is shown in the phase diagram as F IT/F/FC. A “triple” point where FE rhombohedral R3c, FE rhombohedral R3m, AFE commensurate, AFE incommensurate, and another high temperature FE phases meet , is seen for the compositional range 3≤x≤3.3. The results presented above and the proposed phase diagram from Fig. 7.2 are similar with the results published by X. Dai et al. for Pb1-xLax(Zr 0.95Ti0.05)O3 (PLZT 100x/95/5) compositions [14]. 229 7.6 Conclusions The present study extends the understanding concerning the temperature stability of the phase symmetries and establishes the coexistence range of multiple phases in the PLZT x/90/10 system with La composition near the FE/AFE region. We have investigated t he influence of La3+ ion substitutions in the B -site of the crystal lattice of PLZT x/90/10 solid solutions on the temperature of the FE/AFE -PE phase transition s. The substitution with La3+ of Zr/Ti ions in the B -site perovskite position leads to the decrease of stability of the FE phase and decrease of the critical temperature for the FE/AFE -PE phase transition. PLZT x/90/10 ceramics with x≥3 are characterized by a highly diffuse phase transition from the PE to the FE or AFE polar ordered phase on cooling . The diffuseness of the phase transition increases considerably when the composition of the solid solution moves toward the compositions corresponding to the pure AFE phase. The l ack of previous studies concerniong the low-temperature pha se transformations in this system adds the difficulty to distinguish the structures below Curie temperature. Our study revealed some new important features of the phase transformational sequence of PLZT x/90/10 ceramics. The detailed structural study of PL ZT x/90/10 solutions in a broad range of temperature reveal ed a complex sequence of phase transitions. It was proposed a new phase diagram of the PLZT x/90/10 system based on dielectric, XRD results , and Raman spectroscopy data. New temperature phase bound aries separating monoclinic FM and R3c/Pbam phases, F LT and AFE C, AFE C and AFE IC, AFE IC and a high temperature FE phase F TT/F/FC were revealed. Therefore , with these combined method s, we added important knowledge to the clarification of phase succession in PLZT x/90/10. The phase boundary between monoclinic M and AFE C and F LT phases and between AFE C and FLT phases coexistence to single AFE C are diffuse, they are found at temperature very close each other and very close to the room temperature. This expl ains controversial previous results reported for this system, as well as the unusual electric -field-induced multiple phase transition s, complex poling and switching properties of PLZT x/90/10 compositions under low frequency loading (as demonstrated in Cha pter 6). All the observed phenomena and the proposed interpretations represent very useful information which contribute to a better understanding of the mechanisms involved in the temperature and field induced phase transitions in FE/AFE materials and prov ide ways to drive their functional properties towards the desired ones. 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The experimental results have revealed new findings which are important in understanding of the field-induced AFE- to-FE transition . The main results are summarized below:  The influence of the La addition on the microstructure of PLZT x/90/10 ceramics  Increasing the amount of La3+ from 2 at.% to 4 at.% result in a slight decrease of ceramic grain size from 5 μm to 1 μm, for the same calcination and sintering parameters (calcination at 850C for 4h and sintering at 1250  C for 2h) .  The influence of the La addition on crystalline structure of PLZT x/90/10 ceramics  Dense PLZT ceramics with p ure perovskite phase have been produced by conventional solid state reaction.  The complex structural analysis (XRD, HRXRD, TEM, Raman) have revealed that increasing La content from 0 to 4% in the PZT 90/10 system produces a crossover from FE long-range order towards AFE state : the compositions with x≤0.025 have a macroscopic rhombohedral R3c symmetry and develop locally AFE clusters with lower symmetry FE phase; the range of compositions 0.025≤x≤0.035 show a phase coexistence region with the rhombohedral R3c and orthorhombic Pbam structure; at higher values of x≥0.035 the orthorhombic Pbam structure to PbZrO 3 was stabilized.  In addition, Raman spectroscopy indicates that this compositions may develop locally other low symmetry phases at room temperature  The influence of the La addition on dielectric and piezoelectric properties of PLZT x/90/10 ceramics 236  Dielectric and piezoelectric studies confirm ed the results of the structural analysis showing a large compositional range with AFE–FE phase coexistence across the FE/AFE border.  The room temperature permittivity has a maximum for x = 0.03, while the piezoelectric constants show anomalies in the range of compositions x = (0.03, 0.035). These anomalies were interpreted as being related to the phase superposition or to the development of a morphotropic phase boundary.  Study of the AFE- to-FE field induced transition in PLZT x/90/10 system  Polarisation vs. electric field dependence confirms the structural calculations according to which these compositions transform from FE to AFE states, with a stabilization of the AFE phase at higher La3+ content.  The AFE PLZT x/90/10 ceramics shows the ability to switch between an antipolar AFE and a polar FE state under high electric field, similar ly with PbZrO 3-based ceramics.  The study of P(E) loops indicated that La3+ chemical m odifier helps to stabili ze the AFE phase and hence to manipulate the AFE-to-FE phase transition in PLZT x/90/10 based ceramics : the composition with x≤3.1 show s macroscopic FE behavior; the compositions with x=3.2 and x=3.3 show an irreversible AFE -to-FE transition ; the ceramics with La3+ content, x≥3.5, present a field induced AFE -to-FE reversible phase transition. The AFE -to-FE switching field increases as the amount of AFE phase increases in PLZT x/90/10 compositions .  In situ X-ray diffraction study indicates that PLZT 4/90/10 and PLZT 3/90/10 compositions, in addition to the AFE- FE phase switching, exhibit irreversible preferred orientation after experiencing the field-induced AFE- to-FE phase switching. An electric field-induced structure develops in both compositions, which has a reversible character during the field decrease in PLZT 4/90/10 and an irreversible character in PLZT 3/90/10.  Structural analysis of pre-poled PLZT 3/90/10 ceramics shows that it is possible to induce successive FE-to-AFE and AFE- to-FE transitions when fields with reversed polarity are applied in sequence. The field range required to induce the AFE phase is broad, and the phase transition is kinetically slow. This kind of transitions has been rarely reported before.  A remarkable result of this study is that the PLZT 3/90/10, PLZT 3.2/90/10 and PLZT 4/90/10 compositions show very high remnant p olarization Pr~58 µC/cm2 during the application of electric field E≥E AF of very low frequency f~1/300 Hz . These results are confirmed by Rietveld study of ex situ HXRD data obtained for powders from poled PLZT 3/90/10 composition , which stated that the system is monoclinic. 237  A study regarding the composition and electric field -dependent energy storage properties in PLZT ceramics across the FE-AFE boundary has been performed . A recoverable energy density of 1.8 J cm−3 and a high efficiency (η~60) have been achieved on 3.5% of La modification , which suggests that these AFE compositions are interesting for potential applications for energy storage.  Study of the influence of the La addition on phase transitions sequence in PLZT x/90/10 ceramics  The substitution with La3+ of Zr/Ti ions in the B-site perovskite positions leads to the decrease of stability of the FE phase and decrease of the critical temperature for the FE/AFE-PE phase transition. This transition for compositions with x≥3 is highly diffuse. The diffuseness of the phase transition increases considerably when the composition of the solid solution moves toward the compositions corresponding to the pure AFE phase.  It was proposed a new phase diagram of the PLZT x/90/10 system where new temperature phase boundaries are proposed: a border between monoclinic and low temperature FE or high temperature FE phases; a boundary between low temperature FE phase and AFE commensurate; in a limited temperature range there is a transition between AFE commensurate and AFE incommensurate state , a high temperature FE phase was found between AFE incommensurate and PE phase. Therefore, the knowledge concerning the phase transition sequence was completed for PLZT x/90/10 ceramics. 238 List of Figures Fig. 1. 1 An ideal cubic ABO 3 perovskite unit cell ................................................................... 16 Fig. 1. 2 Network of corner-linked octahedral where the A sites (yellow ball) is inside an octahedral cage of Oxygens ....................................................................................................................... 17 Fig. 1. 3 Schematic illustrations of the polarization switching: (A-C) the initial poling, (C-E) the electrical reversal, and (F-A) the electrical cycling. Under the application of an electric field, the B cations displacement is shifted along the electric field direction, giving rise to the lattice distortion. (The rectangles with blue arrows represent schematically the repartition of the two polarization states in the material ( e.g. in the cermic grains) at different fields. ............................................ 20 Fig. 1. 4 Current vs. field during domain switching in FEs (“current hysteresis”) .................... 21 Fig. 1. 5 Schematic diagram of the evolution of AFE-FE phase switching during application and releasing of adequate electric field. Figure adapted from [59 ] ................................................... 23 Fig. 1. 6 Representative electric-field-induced polarization hysteresis loop of AFE materials. .. 24 Fig. 1. 7 Current-field dependence of AFEs ............................................................................. 24 Fig. 1. 8 Typical field dependence of the dielectric permittivity of a tunable ferroelectric ceramic25 Fig. 1. 9 DC field dependence of pemittivity in AFEs with reverssibile AFE to FE field induced transition [3] ............................................................................................................................ 27 Fig. 1. 10 a) DC field dependence of permittivity and b) its corresponding P(E) loop in AFEs with irreversible AFE to FE field induced transition [71] ................................................................ 27 Fig. 1. 11 P(E) hysteresis loop and energy storage characteristics for a) FE and b) AFE ceramics28 Fig. 1. 12 Second order phase transition. (a) Free energy (F) as a function of the polarization (P) at T > T 0, T = T 0, and T < T 0; (b) Spontaneous polarization P 0 (T) as a function of temperature (c) The susceptibility χ and its inverse, at the equilibrium condition P 0(T) ............................................ 30 Fig. 1. 13 First order phase transition. a) Free energy as a function of the polarization ............. 31 Fig. 1. 14 Schematic hysteresis in an idealized ferroelectric ..................................................... 33 Fig. 1. 15 Theoretical temperature–electric field (T – E) phase diagram associated with the free- energy given by Eq. (1.6.2) for 𝛽 > 0 and 𝛼 >0. Hatched and hatched-dotted curves represent, respectively, second-order transition and limit of stability curves. The thermodynamic paths for T > TC, TC > T > T 0 and T < T 0 are described in the text. ................................................................ 37 Fig. 1. 16. Temperature dependence of the dielectric susceptibility as given by Eq. (1.6.10) across a second-order (a) and fir st-order (b) transition ........................................................................... 38 Fig. 1. 17 Typical electric field dependence P(E) of the polarization of an AFE material together with the energy diagrams explaining the double hysteresis loops .............................................. 39 239 Fig. 2. 1 Phase diagram of PZT according to Jaffe et al. [1, 33] with the newly discovered monoclinic phase according to Noheda et al. [34-38]. P C means the paraelectric cubic phase, F T the FE tetragonal phase, F R the FE rhombohedral phase, F M the monoclinic phase, A T the AFE tetragonal phase, and A O an orthorhombic AFE phase .............................................................. 49 Fig. 2. 2 Phase diagram of PLZT system after Haertling et al. [24] ........................................... 51 Fig. 2. 3 PLZT phase diagram, after Haertling and Land a) across entire compositon Zr/Ti ratio.[3] Points indicate the compositions fabricated, PLZT 100x/90/10 (100x 0, 2, 4, and 10). AFE antiferroelectric, FERh ferroelectric rhombohedral, FETet ferroelectric tetragonal, SFE slim- loop ferroelectric, and PECubic paraelectric cubic and b) detail of the AFE/FE phase boundary ................................................................................................................................................. 54 Fig. 2. 4 Phase diagram of PLZT constructed from the data reported by A. Pelaiz – Barranco et al. [71, 72] and Knudsen et al. [57] ......................................................................................... 56 Fig. 2. 5 Phase diagram of 2/95/5 and 4/5/95 as function of temperature and field [55] ............ 57 Fig. 3. 1 Flow diagram of the experimental procedure of the PLZT powder synthesis ............... 72 Fig. 3. 2 Crucibles setup for sintering dense PLZT ceramics .................................................... 73 Fig. 3. 3 Flow diagram of the PLZT samples preparation from powder calcined at 800°C for 4h ................................................................................................................................................. 74 Fig. 3. 4 Illustration of the X-Ray Diffraction setup showing the symmetry of the experiment as well as the Braggs law .............................................................................................................. 77 Fig. 3. 5 Sawyer-Tower circuit used for polarization measurement ........................................... 81 Fig. 3. 6 Poling sample holder .................................................................................................. 82 Fig. 4. 1 Goldsmith tolerance factor t for all PLZT x/90/10 investigated composition ............... 89 Fig. 4. 2. Comparison of XRD patterns of PLZT calcined powders for four La concentrations (2.0, 2.5, 3.0, 4.0 at % ) calcined at a) 800C and b) 850 C for 4 h ............................................ 92 Fig. 4. 3 Comparison of XRD patterns of PLZT sintered ceramics obtained from powders calcined at 800 C with La content of 2.0, 2.5, 3.0, 4.0 at % sintered at: a) 1200 C and b) 1250 C. ................................................................................................................................................. 93 Fig. 4. 4 Comparison of XRD patterns of PLZT sintered ceramics obtained from powder calcined at 850 C with four La concentrations (2.0, 2.5, 3.0, 4.0 at %) sintered at: a) 1200 C and b) 1250C for 2 h ..................................................................................................................... 94 Fig. 4. 5 Lanthanum concentration dependences of the experimental (Archimede’s) absolute and relative densities of the PLZT ceramics: a) variation of PLZT density at various sintering temperatures for four representative compositions and b) the evolution of PLZT density for all the investigated compositions (calcination at 850oC, sintering at 1250 C). ............................... 95 240 Fig. 5. 1 Room temperature dielectric properties: a) composition – dependence of permittivity at a few selected frequencies and b) composition – dependence of dielectric losses at a few selected frequencies for PLZT x/90/10 ceramics 145 Fig. 5. 2 Dependence of the room temperature piezoelectric coefficients upon the La concentration for PLZT x/90/10 ceramics. ............................................................................... 147 Fig. 6. 1 Room temperature P(E) hysteresis and I(E) loops of La3+ modified PZT x/ 90/10 during first and second loading (frequency of 1 Hz) from virgin sample at room temperatures. The hysteresis loops and current curves from (a) and (b) were obtained for the composition PLZT 3/90/10, (c) and (d) for PLZT 3.2/90/10 and (e) and (f) for PLZT 4/90/10, respectively. Here, the EAF-virgin and EAF represents the forward switching fields required to induce AFE- to-FE phase transition, the EFA is the threshold field for the recovery of the AFE phase and the Ec represents the coercive field of FE phase .................................................................................................. 156 Fig. 6. 2 Room temperature dynamic P(E) hysteresis loops of La3+-modified PZT x/ 90/10 ceramics at different magnitudes of the applied field at frequency of 1 Hz. (a), (b), (c), (d), (e), (f), (g), (h) and (i) represent 2.0 %, 2.5%, 3.0 %, 3.1 %, 3.2 %, 3.3 %, 3.5 %, 3.8 % and 4.0 % La3+ compositions, respectively. The shaded area represents the highest recoverable energy density for each composition, Wre = PrPmaxEmaxdP . (Oy axis is in µC/cm2 and Ox axis is in kV/cm). The P(E) hysteresis loops were obtained on samples already exposed to electric field E≥E AF ..................................................................................................................................... 160 Fig. 6. 3 The dependences of the forward switching field E AF necessary to induce the AFE- to-FE phase transition as a function of La3+ composition in PLZT x/90/10 ceramics. ........................ 162 Fig. 6. 4 The dependences of P s, Pr, Ps-Pr as a function of the La3+ composition in PLZT x/90/10 ceramics. ................................................................................................................................ . 163 Fig. 6. 5 a) Field-dependent polarization P(E) loops with the virgin loop obtained for frequency of 1 Hz and , b) Field-dependent hysteresis loops with the virgin loop obtained for frequency of 1/300s Hz and c) current curves obtained in the same condition as a) for PLZT 3/90/10 bulk ceramic. .................................................................................................................................. 166 Fig. 6. 6 a) Field-dependent hysteresis P(E) loops and b) current curves I(E) at various frequency and fixed amplitude of electric field 60 kV/cm for PLZT 3.2/90/10 bulk ceramic. The data were acquired starting from 1 to 3, respectively. .............................................................................. 168 Fig. 6. 7 a) Field-dependent hysteresis loops and b) current curves at various representative frequencies and at fixed field amplitude for PZT PLZT 4//90/10 bulk ceramic. The data were acquired in a continuos sequence, starting from high to low frequencies. ................................ . 170 241 Fig. 6. 8 Contour plots of diffraction intensities for a) and b) 002 pc/200 pc, and c) 111 pc reflections for the PLZT 4/90/10 composition during a complete triangular field cycle with amplitude ±74 kV/cm and frequency 0.8 mHz during a) and c) first and b) second electrical cycle. The diffraction profiles from d) show the (002) pc/(200) pc reflections before application of electric field (black), at the coercive field of -74 kV/cm (red) and after application of electric field (gree n). The subscript pc indicates reflections indexed with the pseudocubic primitive cell. A schematic drawing to describe the experimental sequence of applied electric field is included in the left hand side of the figure. E AFE-FE and E FE-AFE represent the forward switching field for inducing the AFE- to-FE transition and the backward switching field of the FE- to-AFE recovery, respectively. ............................................................................................................................ 172 Fig. 6. 9 Contour plots of diffraction intensities for a) and b) 111 pc/-111 pc, c) and d) 200 pc/002 pc reflections for the PLZT 3/90/10 composition as a function of triangular bipolar electric field of ±45 kV/cm amplitude and frequency of 0.8 mHz during a) and c) first, and b) and d) second electrical cycle. The subscript pc indicates reflections indexed with the pseudocubic primitive cell. A schematic diagram to describe the experimental sequence of applied electric field is included on the left side of the figure. E AFE-FE and E FE-AFE represent the forward switching field for inducing the AFE- to-FE transition and the reverse switching field of the FE- to-AFE recovery, respectively. ............................................................................................................. 174 Fig. 6. 10 Electric field dependence of FWHM of the (111) diffraction peak during FE- to-AFE and AFE- to-FE induced transitions for the composition PLZT 3/90/10. ................................... 176 Fig. 6. 11 Selected pseudocubic Bragg profile of electrically poled and of virgin PLZT 3/90/10 powders ................................................................................................................................... 178 Fig. 6. 12 The observed (dots), calculated (continuous line) and difference (bottom) HXRD profiles after Rietveld refinement of the structure of poled PLZT 3/90/10 powder using Cm and Pbam space groups. The Bragg positions are shown in the bottom inset by vertical lines: the top one corresponds to Cm and the bottom one corresponds to Pbam structure. ............................. 179 Fig. 6. 13 The observed (dots), calculated (continuous line) and difference (bottom) HXRD profiles of the enlarged views for the 110 pc, 111 pc, 200 pc, 211 pc, 220 pc and 222 pc peaks obtained after Rietveld refinement of the structure of poled PLZT 3/90/10 powder using Cm and Pbam space groups ............................................................................................................................ 180 Fig. 6. 14 a) Energy storage density Wre and b) efficiency η at room temperature for PLZT 90/10 samples with various amounts of La3+ under different applied electric field Emax. 182 242 Fig. 6. 15 Electric field-dependence of the recoverable energy storage density Wre (a) and of the lost energy Wloss (b) of the La3+ modified PZT 90/10 ceramics with La3+ at. % content varying from 2 to 4 .............................................................................................................................. 183 Fig. 7.1 a) Real part of permittivity ( ε’), b) imaginary part of permittivity ( ε”), c) detail from a) and d) dielectric losses ( tan δ ) of PLZT x/90/10 compositions measured during heating at a fixed frequency of 100 kHz. .............................................................................................................197 Fig. 7. 2 Maximum permittivity εm and its corresponding temperature Tm as a function of La composition obtained at 100 kHz. ............................................................................................ 199 Fig. 7.3 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PZT 90/10 and PLZT 2/90/10 ceramics. .............................................. 201 Fig. 7.4 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 2.5/90/10 and PLZT 3/90/10 ceramics. .....................................202 Fig. 7. 5 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 3.1/90/10 and PLZT 3.2/90/10 ceramics. ...................................203 Fig. 7. 6 Temperature dependence of the real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 3.3/90/10 and PLZT 3.5/90/10 ceramics. ............................... 204 Fig. 7. 7 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 3.8/90/10 and PLZT 4/90/10 ceramics. .....................................205 Fig. 7.8 (a) Reciprocal dielectric constant 1/’ as a function of temperature at 100 kHz for PLZT x/90/10 compositions (the dashed lines are fittings with the Curie-Weiss law); (b) Tm and TCW vs. La addition; (c) Frequency dependence of Curie Weiss temperatures ......................................207 Fig. 7. 9 Log(1/ε-1/ε m) vs. Log(T−T m) for two representative PLZT compositions ................... 209 Fig. 7. 10 Diffuseness exponent λ and modified Curie Weiss constant C’ vs. La content for PLZT x/90/10ceramics ...................................................................................................................... 209 Fig. 7. 11 Variation with temperature of the real part of dielectric permittivity ε’, loss tangent tan δ, of poled ceramics withy compositions: a) PLZT 3/90/10 sample, b) PLZT 3.1/90/10 and c) PLZT 4/90/10 measured during heating in the temperature range of 25– 300°C at the frequency of 10 kHz. ............................................................................................................................... 211 Fig. 7.12 A contour plot of diffraction intensities as a function of temperature obtained from {001}pc for PLZT x/90/10 ceramics ........................................................................................ 213 Fig. 7.13 . A contour plot of diffraction intensities as a function of temperature obtained for {200}pc of PLZT x/90/10 ceramics ......................................................................................... 214 Fig. 7.14 The evolution of lattice parameters (a, b, c) and of the unit cell volume with temperature for a) PLZT 3.5/90/10 and b) PLZT 4/90/10 ceramics ..........................................215 243 Fig. 7.1 a) Real part of permittivity ( ε’), b) imaginary part of permittivity ( ε”), c) detail from a) and d) dielectric losses ( tan δ ) of PLZT x/90/10 compositions measured during heating at a fixed frequency of 100 kHz. .............................................................................................................197 Fig. 7. 2 Maximum permittivity εm and its corresponding temperature Tm as a function of La composition obtained at 100 kHz. ............................................................................................ 199 Fig. 7.3 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PZT 90/10 and PLZT 2/90/10 ceramics. .............................................. 201 Fig. 7.4 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 2.5/90/10 and PLZT 3/90/10 ceramics. .....................................202 Fig. 7. 5 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 3.1/90/10 and PLZT 3.2/90/10 ceramics. ...................................203 Fig. 7. 6 Temperature dependence of the real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 3.3/90/10 and PLZT 3.5/90/10 ceramics. ............................... 204 Fig. 7. 7 Temperature dependence of real part of dielectric permittivity and dielectric losses at various frequencies for PLZT 3.8/90/10 and PLZT 4/90/10 ceramics. .....................................205 Fig. 7.8 (a) Reciprocal dielectric constant 1/’ as a function of temperature at 100 kHz for PLZT x/90/10 compositions (the dashed lines are fittings with the Curie-Weiss law); (b) Tm and TCW vs. La addition; (c) Frequency dependence of Curie Weiss temperatures ......................................207 Fig. 7. 9 Log(1/ε-1/ε m) vs. Log(T−T m) for two representative PLZT compositions ................... 209 Fig. 7. 10 Diffuseness exponent λ and modified Curie Weiss constant C’ vs. La content for PLZT x/90/10ceramics ...................................................................................................................... 209 Fig. 7. 11 Variation with temperature of the real part of dielectric permittivity ε’, loss tangent tan δ, of poled ceramics withy compositions: a) PLZT 3/90/10 sample, b) PLZT 3.1/90/10 and c) PLZT 4/90/10 measured during heating in the temperature range of 25– 300°C at the frequency of 10 kHz. ............................................................................................................................... 211 Fig. 7.12 A contour plot of diffraction intensities as a function of temperature obtained from {001}pc for PLZT x/90/10 ceramics ........................................................................................ 213 Fig. 7.13 . A contour plot of diffraction intensities as a function of temperature obtained for {200}pc of PLZT x/90/10 ceramics ......................................................................................... 214 Fig. 7.14 The evolution of lattice parameters (a, b, c) and of the unit cell volume with temperature for a) PLZT 3.5/90/10 and b) PLZT 4/90/10 ceramics ..........................................215 Fig. 7. 15 a) Variation with temperature of the real part of dielectric permittivity ε’, loss tangent tan δ , and reciprocal permittivity ( 1/ε’) of the virgin PLZT 3/90/10 sample, measured during the 244 List of abbreviations FE - ferroelectric AFE - antiferroelectric PE - paraelectric MPB - morphotropic phase boundary PZT - PbZr xTi1-xO3 PLZT – La-substituted PZT PLZTx/10/90 – Pb1-xLax(Zr 0.9Ti0.1)1-x/4O3 245 List of original publications 1. I.V. Ciuchi , F. Craciun, L. Mitoseriu, C. Galassi, Preparation and properties of La doped PZT 90/10 ceramics across the ferroelectric-antiferroelectric phase boundary , J. Alloys Compd. 646 (2015) 16-22. (I.F.=3.014 , A.I=0.558) Cited by 2 publications: [1] X. Wang, T. Yang, J. Shen, Y. Dong, Y. Liu, Phase transition and dielectric properties of (Pb,La)(Zr,Sn,Ti)O 3 ceramics at morphotropic phase boundary, J. Alloys & Compd. 673 (2016) 67-72. 2. F. Craciun, F. Cordero, I.V. Ciuchi , L. Mitoseriu, C. Galassi, Refining the phase diagram of Pb 1-xLax(Zr 0.9Ti0.1)1-x/4O3 ceramics by structural, dielectric, and anelastic spectroscopy investigations , J. Appl. Phys. 117(18) (2015) 184103 (1-8). (I.F. =2.101 , A.I=0.68) Cited by 1 : [1] A. Peláiz-Barranco, Y. González-Abreu, Y. Gagou, P. Saint-Grégoire, J.D.S. Guerra, Raman spectroscopy investigation on (Pb 1−xLax)(Zr 0.90Ti0.10)1−x/4O3 ceramic system, Vib. Spectrosc. 86 (2016) 124–127. 3. I.V. Ciuchi , L. Mitoseriu, C. Galassi, Antiferroelectric to Ferroelectric Crossover and Energy Storage Properties of (Pb 1-xLax)(Zr 0.90Ti0.10)1-x/4O3 (0.02 ≤x ≤0.04) Ceramics , J. Am. Ceram. Soc. 99(7) (2016) 2382-2387. (I.F.=2.841 , A.I=0.70) Cited by 8 : [1] R. Xu, Q. Zhu, J. Tian, Y. Feng, Z. Xu, Effect of Ba-dopant on dielectric and energy storage properties of PLZST antiferroelectric ceramics, Ceram. Int. 43(2) (2017) 2481-2485. [2] D. Zheng, R. Zuo, Enhanced energy storage properties in La(Mg 1/2Ti1/2)O3-modified BiFeO 3- BaTiO 3 lead-free relaxor ferroelectric ceramics within a wide temperature range, J. Eur. Ceram. Soc. 37(1) (2017) 413-418. [3] B. Luo, X. Wang, E. Tian, H. Song, H. Wang, L. Li, Enhanced Energy-Storage Density and High Efficiency of Lead-Free CaTiO 3–BiScO 3 Linear Dielectric Ceramics, ACS Appl. Mater. Interfaces. 9(23) (2017) 19963-19972. 246 [4] R. Xu, J.T. , Q. Zhu, T. Zhao, Y. Feng, X. Wei, Z. Xu, Effects of phase transition on discharge properties of PLZST antiferroelectric ceramics, J Am. Ceram. Soc. 100(8) (2017) 3618-3625. [5] B. Li, Q. Liu, X. Tang, T. Zhang, Y. Jiang, W. Li, J. Luo, High Energy Storage Density and Impedance Response of PLZT2/95/5 Antiferroelectric Ceramics, Materials 10(2) (2017) 143. [6] B. Li, Q.-X. Liu, X.-G. Tang, T.-F. Zhang, Y.-P. Jiang, W.-H. Li, J. Luo, High temperature dielectric anomaly and impedance analysis of (Pb 1−3x/2Lax)(Zr 0.95Ti0.05)O3 ceramics, J. Mater. Sci.: Mater. El. (2017) 1-10. [7] F. Li, J. Zhai, B. Shen, X. Liu, K. Yang, Y. Zhang, P. Li, B. Liu, H. Zeng, Influence of structural evolution on energy storage properties in Bi 0.5Na0.5TiO 3-SrTiO 3-NaNbO 3 lead-free ferroelectric ceramics, J. Appl. Phys. 121(5) (2017) 054103. [8] R. Xu, J. Tian, Q. Zhu, T. Zhao, Y. Feng, X. Wei, Z. Xu, Effects of phase transition on discharge properties of PLZST antiferroelectric ceramics, J Am. Ceram. Soc. 100(8) (2017) 3618-3625. 4. M. Cernea, P. Galizia, I. V. Ciuchi , G. Aldica, V. Mihalache, L. Diamandescu, C. Galassi, CoFe 2O4 magnetic ceramic derived from gel and sintered by spark plasma sintering, J. Alloys Compd , (2016), volum, numar (I.F. = 3.133 , A.I=0.551) Cited by 6 : [1] P. Galizia, C. Baldisserri, C. Capiani, C. Galassi, Multiple parallel twinning overgrowth in nanostructured dense cobalt ferrite, Mater. Des. 109 (2016) 19-26. [2] J. Jin, X. Sun, M. Wang, Z.L. Ding, Y.Q. Ma, The magnetization reversal in CoFe 2O4/CoFe 2 granular systems, J. Nanopart. Res. 18 (2016) 383. [3] R. Zhang, L. Sun, Z. Wang, W. Hao, E. Cao, Y. Zhang, Dielectric and magnetic properties of CoFe 2O4 prepared by sol-gel auto-combustion method, Mater. Res. Bull. (2017). [4] P. Galizia, M. Cernea, V. Mihalache, L. Diamandescu, G. Maizza, C. Galassi, Easy batch-scale production of cobalt ferrite nanopowders by two-step milling: Structural and magnetic characterization, Mater. Des. 130 (2017) 327-335. [5] P. Galizia, C.E. Ciomaga, L. Mitoseriu, C. Galassi, PZT-cobalt ferrite particulate composites: Densification and lead loss controlled by quite-fast sintering, J. Eur. Ceram. Soc. 37(1) (2017) 161-168. [6] P. Galizia, Production and morphological and microstructural characterization of bulk composites or thick films for the study of multiphysics interactions, PhD Thesis, Politecnico di Torino, Porto Institutional Repository, (2017). 247 5. A. Neagu, I.V. Ciuchi , L. Mitoseriu, C. Galassi, C. -W. Tai, Study of ferroelectric - antiferroelectric phase coexistence in La -doped PZT ceramics, European Microscopy Congress 2016: Proceedings, Wiley -VCH Verlag GmbH & Co. KGaA (2016) ( http://emc- proceedings.com/abstract/study- of-ferroelectric-antiferroelectric-phase-coexistence- in- la-doped-pzt-ceramics/ ) 6. I. V. Ciuchi , C. C. Chung, C. M. Fancher, J. Guerrier, J. S. Forrester, J. L. Jones, L. Mitoseriu and C. Galassi , “Field -induced antiferroelectric to ferroelectric transition in (Pb 1–xLax)(Zr 0.90Ti 0.10)1–x/4O3 investigated by in situ X -ray diffraction”, J. Eur. Ceram. Soc. , 2017, volum, numar daca au aparut sau doi (In Press) (I.F.=3.411 , A.I=0.697) 7. I.V. Ciuchi , C. M. Fancher, C. Capiani, J.L. Jones, L. Mitoseriu, C. Galassi, Field induced metastable ferroelectric phase in PLZT 3/90/10 ceramics, J. Eur. Ceram. Soc. (2017), (Review to be submitted) (I.F.=3.411 , A.I=0.697) 248 List of participations to international conferences 1. I. V. Ciuchi , C. Capiani, L. Mitoseriu and C. Galassi, Composition-dependent dielectric and energy-storage properties of Pb 1-xLax(Zr 0.9Ti0.1)1-x/4O3 ceramics , Ceramics for Energy, June 2017, Faenza, Italy, (Oral) 2. I. V. Ciuchi , L. Mitoseriu, and C. Galassi, PLZT x/90/10 ceramics for energy storage , Piezo 2017 -Electroceramics for End Users IX, Madrid (Spain), February 2017, (Oral) 3. A. Neagu, I. V. Ciuchi , L. Mitoseriu, C. Galassi, C.- W. Tai, Study of ferroelectric- antiferroelectric phase coexistence in La-doped PZT ceramics, The 16th European Microscopy Congress, Lyon, France, 2016 (Poster) (http://emc-proceedings.com/abstract/study- of- ferroelectric-antiferroelectric-phase-coexistence- in-la-doped-pzt-ceramics/) 4. I. V. Ciuchi , J.E. Guerrier, C. Chung, L. Mitoseriu, J.L. Jones and C.Galassi, Diffuse Phase transitions and Curie Weiss behavior of of Pb 1-xLax(Zr 0.9Ti0.1)1-x/4O3 (0.02 ≤ x ≤ 0.04) ceramics , 1st Research Triangle Nanotechnology Network Research Symposium, Raleigh, US, 2016 (Oral) 5. I.V. Ciuchi , F. Craciun, M. Deluca, L. Mitoseriu and C. Galassi , Phase transitions and Curie Weiss behaviour in La3+ doped PZT 90/10 ceramics with compositions across the antiferroelectric/ferroelectric phase boundary, 13th European Meeting on Ferroelectricity, Porto, Portugal, July 2015 ( Oral ) 6. R. Pullar, I. V. Ciuchi , P. Galizia and C. Galassi , Novel TiO 2-doped semiconducting hexagonal ferrites, 13th European Meeting on Ferroelectricity, Porto, Portugal, July 2015 ( Poster ) 7. I.V. Ciuchi , F. Craciun, L. Mitoseriu, C. Galassi, Temperature dependence of dielectric, piezoelectric and elastic properties of La-doped PZT 90/10 across ferroelectric- antiferroelectric boundary , International Confere4nce European Ceramic Society, Toledo, Spain, June 2015 (poster) 8. I. V. Ciuchi , L. Mitoseriu, C. Galassi, PLZT Based Antiferroelectric Plate Capacitors with High Energy Storage Density , Ceramics For Energy, Faenza, Italy, May 2015 (poster) 9. I.V. Ciuchi , P. Galizia, L. Mitoseriu, C. Galassi, Magnetic and dielectric properties of cobalt ferrite/titania composites, Magnet, Bologna, Italy, Febbruary 2015 (poster) 249 10. I.V. Ciuchi , F. Craciun, L. Mitoseriu, C. Galassi, Dielectric Properties of La3+ doped PZT Ceramics across the antiferroelectric/ferroelectric phase boundary, PIEZO -Electroceramics for End-Users VIII conference, Maribor, Slovenia, January 2015 (oral) 11. P. Galizia, C. Baldisserri, I. V., Ciuchi, C. Galassi, Investigation of new magneto-dielectric titania-cobalt ferrite composites, E-MRS Fall Meeting, Warsaw, September 2014 (Oral) 12. I. V. Ciuchi, F. Craciun, C. Galassi, L. Mitoseriu , Piezoelectric Properties of La3+ doped PZT Ceramics across the antiferroelectric/ ferroelectric phase boundary , European Conference on Application of Polar Dielectrics , July, Vilnius 2014 (Poster) 13. I. V. Ciuchi, M. Cernea, B. S. Vasile, R. Trusca, C. Galassi , Synthesis and dielectric properties of one-dimensional BNT-BTCe@SiO 2 core-shell heterostructures, Electroceramics XIV Conference, June 2014, Bucharest (Oral) 14. I.V. Ciuchi , F. Craciun, C. Galassi, L. Mitoseriu , Lanthanum Dependence of Piezoelectric Properties of PLZT Ceramics with Zr/Ti Ratio of 90/10, Electroceramics XIV Conference, June 2014, Bucharest (Poster) 15. R. Stanculescu, I. V. Ciuchi , I. Turcan, P. Galizia, C. Capiani, C. Galassi and L. Mitoseriu, Preparation and dielectric investigations of Ba 0.60Sr0.40TiO 3 ferroelectric ceramics with different degree of porosity, Electroceramics XIV Conference, June 2014, Bucharest (Oral) 16. I. V. Ciuchi , M. Cernea, C. Galassi, Piezoelectric and Dielectric Characterization of BNT 0.92 BT0.08 (5 mol% Ce doped), Coated with SiO 2, COST Action MP0904-Closing Conference, Genoa, January, 2014, (Poster) 17. R. Stanculescu, I. V. Ciuchi , I. Turcan, P. Galizia, C. Capiani, C. Galassi and L. Mitoseriu , Microstructural and dielectric investigations of Ba 0.60Sr0.40TiO 3 ferroelectric ceramics with different degree of porosity, COST Action MP0904-Closing Conference, Genoa, 2014, (Poster) 18. L. Stoleriu, L. Curecheriu, I. V. Ciuchi , C. Galassi, L. Mitoseriu , FORC Investigation of Ferroelectric-Antiferroelectric Crossover in PLZT (x/90/10) Ceramics , COST MP0904 Single- and multiphase ferroics and multiferroics with restricted geometries, (SIMUFER), January 2014 (Poster, Presenting Author) 19. R. E. Stanculescu, I. Turcan, I. V. Ciuchi, Investigation of the role of porosity on the functional properties of Ba 1-XSr xTiO 3 ceramics produced by using graphite forming agent ,, “The Tenth Students’ Meeting”, SM-2013 Processing and applications of ceramics , The Third Early Stage Researchers Workshop” COST MP0904 – SIMUFER, Novi Sad, Serbia, November, 2013 (oral) 250 20. I. V. Ciuchi, C. Galassi, L. Mitoseriu, Temperature dependence of the main piezoelectric paramaters in very soft, soft and hard piezoelectric ceramic discs, “The Tenth Students’ Meeting”, SM-2013 Processing and applications of ceramics, The Third Early Stage Researchers Workshop” COST MP0904 – SIMUFER, Novi Sad, Serbia, November, 2013 (oral) 251 Research stages and participations to training schools during the PhD activity 1. JECS TRUST Research Stage (15 February-15 May 2016) performed in collaboration with Prof. J. Jones Research Group, Department of Materials Science, in North Carolina State University, Raleigh, North Carolina, USA The main activity was: ▪ Investigation of the structure and phase transitions of PLZT (according to the formula Pb1-xLax(Zr0.9Ti0.1)1- has been carry on by using in situ XRD field measurements and in situ temperature XRD measurements ▪ Rietveld Refinement on high resolution Synchrotron XRD data and on traditional XRD 2. Research Stage (26 May-31 May 2014) performed in Department of Radiophysics, Vilnius University, 3 Universiteto St, LT-01513 Vilnius, Lithuania The main activity was: ▪ Magnetic permeability measurement of magneto-dielectric ceramic composites ▪ Learning about measuring magnetic permeability and dielectric permittivity with EIA Bullet 3. “One day satellite workshop-Inelastic scattering in ferroic materials” (28 th June 2015) The main thematic wastheoretical and practical demo courses regarding RAMAN spectroscopy in ceramic composites. The stage was performed at University of Porto, Portugal 4. International FIB-SEM/AFM (JECS Trust) school ( 29 September-03 October 2014) performed at Sabanci University Campus, Istanbul, Turkey The main thematic was:  Theoretical courses and practical deomos concering SEM, FIB and AFM 252  Basics and special application of this techniques on powder, bulk solid ceramic materials and ceramic based nanocomposite systems