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Acta Materialia 57 (2009) 2328–2338

Substrate heterostructure effects on interface composition,microstructure development and functional properties of PZT thin films

Salah Habouti a, Abdelilah Lahmar a, Matthias Dietze a, Claus-Henning Solterbeck a,Vladimir Zaporojtchenko b, Mohammed Es-Souni a,*

a Institute for Materials & Surface Technology, University of Applied Sciences, Grenzstr. 3, 24149 Kiel, Germanyb Faculty of Engineering, Christian-Albrechts-University, Kiel, Germany

Received 21 October 2008; received in revised form 2 February 2009; accepted 3 February 2009Available online 4 March 2009

Abstract

Platinum- and (La0.8,Sr0.2)MnO3 (LSMO)-terminated silicon substrates were used for the liquid-phase deposition of Pb(Zr0.52,Ti0.48)O3(PZT) thin films. Different layer thicknesses ranging from 100 to 600 nm were processed by sequential coating. Characterization of thefilms involved X-ray diffraction, atomic force microscopy and X-ray photoelectron spectroscopy (XPS) combined with depth profilingto probe the interface composition. The films deposited on Pt exhibit an intermetallic layer, PtxPb, after annealing at 500 �C in air. Thisfilm has been used to establish the XPS signature of the intermetallic phase which consists of a negative shift of the peak position of Pt(4f)due the electron transfer from Pb to Pt. In all cases pure phase perovskite thin films were obtained after short annealing at 700 �C. XPSdepth profiling shows unambiguously the existence of an intermetallic layer, PtxPb, of approximately 10 nm at the interface between Ptand PZT, while an interdiffusion layer of�30 nm was observed between LSMO and PZT. The impacts of interfacial layers on microstruc-ture development and functional properties translate in the formation of specific textures, i.e. a pronounced (111)-texture on Pt due tolattice matching between (111)-PZT and (111)-PtxPb, and a random film orientation on LSMO, and a substantial thickness dependenceof the dielectric and ferroelectric properties, though specific behaviors were observed for the two different substrate heterostructures.� 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: PZT; XPS; Depth profiling; Interface

1. Introduction

Chemical solution deposition (CSD) is a well-estab-lished processing method for functional thin films of vari-ous stoichiometries and structures: it combines cost-effectiveness with a high degree of stoichiometry andmicrostructure control [1]. CSD, however, involves differ-ent specific steps which could all influence the final micro-structure and hence the properties of the thin film. Given acertain substrate heterostructure, the precursor solution,including starting reagents, chelating agents and solvents,the pyrolysis conditions, including pyrolysis atmosphereand temperature, and the final annealing conditions all

1359-6454/$36.00 � 2009 Acta Materialia Inc. Published by Elsevier Ltd. Alldoi:10.1016/j.actamat.2009.02.005

* Corresponding author. Tel.: +49 4312102660; fax: +49 4312102661.E-mail address: Mohammed.Es-souni@fh-kiel.de (M. Es-Souni).

play a role in determining the final microstructure [2,3].One class of functional thin films that is being extensivelystudied are perovskite-based ferroelectric thin films. Specialattention has been devoted to Pb(Zr,Ti)O3 (PZT) thin filmswhich can be used in a variety of devices, including ferro-electric memories, surface acoustic waves and microelectro-mechanical systems (MEMS) [4–6]. PZT films are usuallyprocessed on Pt-terminated substrates, mostly (111)-Pt/Ti/SiO2/Si, that can withstand the high reactivity of theprecursor solution components. The films grown on thesesubstrates have been shown to adapt two different fiber tex-tures, i.e. (100), (11 1), but unfortunately there is no pre-cise knowledge of the conditions which can drive theformation of the one or the other texture. It is agreed upon,however, that the pyrolysis temperature and atmosphereare critical issues in this respect. For deposition on Pt

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S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338 2329

bottom electrodes, the nucleation mechanisms have beenanalyzed in terms of heterogeneous nucleation either onPtPb intermetallic phases (generally Pt3Pb), which arebelieved to arise from an interfacial reaction between thethin film and the Pt electrode [7,8], or on a PbO buffer layerwhich allegedly forms in the early stages of pyrolysis [8].Nucleation on Pt3Pb is reported to lead to (111), whereasnucleation on PbO would lead to (100) preferred orienta-tions. In both cases nucleation is supposed to be facilitatedby lattice matching between the perovskite thin film andthe interfacial phase. Fè et al. [9] investigated early pyroly-sis products using attenuated reflection Fourier transforminfrared spectroscopy, and found that the (111)-textureforms when the acetate groups were not completelyburned-off due to the short pyrolysis time at 350 �C (insuf-ficient oxygen supply at the substrate–film interface, orpyrolysis under reducing conditions) and when a dryingstep at 200 �C was adapted where complete hydrolysis ofOR-radicals took place. Otherwise a mixed (111)/(1 00)-texture was formed. Fè et al. discussed these results interms of thermodynamic and kinetic considerations with-out taking into account the formation of template layers.However, in previous work [10] we have shown that the(111)-texture can be obtained at a lower pyrolysis temper-ature of 300 �C and longer pyrolysis time, whereas pyroly-sis under UV light resulted in (100)-texture, keepingotherwise the pyrolysis conditions the same.

In the present work we investigate the effects of two sub-strate heterostructures on the formation of interfacial lay-ers and how they affect the microstructure, dielectric andferroelectric properties of PZT thin films with morphotrop-ic stoichiometry. We first direct our attention to the analy-sis of the interfacial layer between platinum and PZT andwe establish its X-ray photoelectron spectroscopy (XPS)signature. Although the formation of an intermetallic layerin the early stages of pyrolysis has been reported [3,8], athorough XPS characterization of this layer is still lacking.

The second substrate heterostructure of interest is(La0.8,Sr0.2)MnO3 (LSMO)-terminated (111)-Pt/Ti/SiO2/Si substrate. LSMO has been used as a bottom electrodefor PZT-based capacitors to improve their resistance againstpolarization fatigue [11]. However, investigations of interfa-cial compositions between PZT and LSMO do not exist,apart from our own work on substrate effects on the proper-ties of PbTiO3 (PTO) thin films, where substantial interdiffu-sion between PTO and LSMO was shown [12].

Nucleation kinetics as well as the final microstructureobtained for PZT films, including second phases (e.g.pyrochlore), grain size and texture are well known todepend largely on the Zr/Zr + Ti ratio. In the absence ofa thorough thermodynamic analysis of this system it canonly be speculated that Zr/Zr + Ti ratio affects the activityof PbO in the solid solution and subsequently perovskiteformation [13]. Furthermore, quantitative analysis ofPZT films with increasing Zr/Zr + Ti ratios using XPShas been shown to present some challenges related to theoverall lower peak intensity obtained for Pb that probably

arises from the Pb ionization cross-section being affectedby Zr. Sugiyama et al. [14] were the first to report on prob-lems associated with XPS analysis of PZT. A survey of theliterature on XPS depth profiling of PZT on Pt shows thatresearchers often calculate stoichiometries of the PZT filmswith very underestimated Pb and overestimated Zr con-tents [15–18]. While our own results on PbTiO3 [19] hadshown that Scofield atomic sensitivity factors (ASFs) gavea very good match between calculated and nominal stoichi-ometry, erroneous results were initially obtained for thePZT films. Therefore, it appeared imperative to us first toperform a calibration using different Zr/Zr + Ti ratiosand consequently calibrate the Scofield ASFs.

2. Materials and methods

The procedure for precursor solution preparation can befound elsewhere [20,21]. The starting materials for the prep-aration of the precursor solution were Pb(II)–acetate trihy-drate (Pb(CH3COO)2�3H2O), zirconium-n-propoxide (Zr(OCH(CH3)2)4) and titanium iso-propoxide (Ti(OCH(CH3)2)4)which were mixed to yield a Pb1.1(Zr0.52Ti0.48)O3 stoichi-ometry. Excess lead is present in order to compensate forlead oxide loss at the annealing temperature chosen. Thestock solution was stabilized with 3 mol of acetylacetone,and diluted with 2-methoxyethanol to give a final concen-tration of 0.6 mol L�1. Finally, the precursor solutionwas filtered through a 0.2 lm syringe filter. Spin coatingwas performed at 5000 rpm for 30 s on two alternative sub-strate heterostructures consisting of bare or LSMO-termi-nated (11 1)Pt(150 nm)/Ti(10 nm)/SiO2/Si. The procedurefor the preparation of the LSMO buffer layer is describedelsewhere [12]. Pyrolysis was conducted in air, at 350 �Con a hot plate followed by an annealing step at 500 �Cfor 5 min. For XPS measurements one layer (1L) wasdeposited using a 0.25 mol L�1 concentrated sol, giving alayer thickness of �30 nm for PZT on Pt. For PZT onLSMO a thicker layer of �80 nm was used, since in previ-ous work [12] we have shown that La and Mn could bedetected on the surface of a PbTiO3 film of 30 nm thicknessin an PbTiO3/LSMO film heterostructure. Further, a cali-bration of the XPS signal was conducted with 1L thin filmshaving different Zr concentrations while keeping the Pbconcentration constant (see below for more details). Crys-tallization of the perovskite phase was conducted at700 �C for 10 min, after the two steps of pyrolysis at 350and 500 �C.

Phases and preferential orientations of the films werecharacterized by X-ray diffraction (XRD) using a four-circle diffractometer (Seifert 3000 PTS) utilizing Cu Ka radi-ation (0.15418 nm) at 40 kV and 40 mA. XPS characteriza-tion of the 1L PZT films was performed using an OmicronFull Lab System. For depth profiling it was necessary tofind sputtering conditions that would not lead to artifactsrelated to preferential sputtering, e.g. of oxygen, and thusto the reduction of Pb, seen in previous work [22–24]. Suit-able parameters were 9.7 � 10�5 mbar for argon pressure

2330 S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338

and 10 mA for ion current at 160 V. The sputtering ratecorresponding to these parameters was 1.7 nm min�1 (mea-sured with a profilometer). The roughness of the sputteredarea was measured by atomic force microscopy (AFM,Nanostation, SIS, Herzogenrath, Germany). Systematicshifts of the XPS spectra (due to surface charging effect)were corrected using the 1 s line of carbon at 284.5 eV.

For electrical measurements, Pt top electrodes of 0.6 mmdiameter were sputtered through a shadow mask and subse-quently annealed at 350 �C for 10 min to heal eventualdefects introduced during sputtering. Ferroelectric proper-ties were obtained using a ferroelectric testing system (TFAnalyzer 2000, aixACCT, Aachen, Germany). Small signaldielectric properties were determined in the frequency rangefrom 100 Hz to 1 MHz using a computer-controlled imped-ance analyzer (Agilent 4192A, impedance analyzer). Themodulation voltage amplitude was 25 mV.

The data obtained are subject to errors that includeinstrumental and peak-fitting/integration errors.

The error in the XPS composition is dominated by theuncertainty of the calibrated ASFs. The errors are com-puted from linear regressions to the data of Fig. 3, and yieldcomposition errors in the range of ±8% for Pb concentra-tion and ±4% for Ti and Zr concentrations; the error barsare shown in Figs. 5 and 6. The error on the binding energyis DE = ±0.1 eV.

The error in the reciprocal capacity is dominated by thespread in the pad area. Errors from the impedance measure-ment itself are smaller by more than an order of magnitude.The accuracy of the thickness is limited by the roughness ofthe surface, while the measurement of an average thicknessby ellipsometry and the preparation of a film with the desiredthickness are more accurate. The value taken here is half theminimum–maximum difference of the AFM topography.

3. Results and discussion

3.1. XRD

The XRD patterns of the 1L films after pyrolysis on a hotplate at 350 �C show no peaks other than those of the sub-strate. Upon annealing at 500 �C in air, an intense peak in

Fig. 1. XRD pattern of the PZT thin films deposited on platinized silicon (a andannealing at 500 �C where the intense peak belongs to the intermetallic PtxPb phto the nanocrystalline pyrochlore phase; (b) XRD patterns after annealing at 70on LSMO showing the random crystallization of the perovskite phase; the pe

the vicinity of the 111-Pt peak appears that can be attrib-uted to the intermetallic phase PtxPb, see Fig. 1a. The broadpeak around 30� (2H) is characteristic of the pyrochlorephase (a Pb-deficient phase of general formula Pb2TiZrO6).Incipient nucleation and growth of the perovskite phase canalso be seen in the weak 100 peak around 20�. This film wasused for the characterization of the XPS attributes of theintermetallic phase. Annealing at 700 �C under ambientatmosphere leads to the formation of strongly (111)-tex-tured films regardless of their thickness (Fig. 1b). The peakintensity ratios I111/(I111 + I100 + I110) all attain approxi-mately 99%. The (111)-texture predominates despite thenumerous coating sequences, indicating that following lay-ers simply adapt the orientation of the first layer formed,thus minimizing the energy barrier for nucleation andgrowth. The line characteristic of the interfacial intermetal-lic layer at 2H = 38.44� is no longer detectable. However,this is no proof of the absence of the intermetallic phase(see below). We surmise that incipient formation of themetastable intermetallic phase at the interface between theamorphous oxide layer and Pt already takes place at350 �C, since the formation of this phase is favored kineti-cally at this relatively low temperature, where diffusion ofionic species to form the perovskite phase is expected tobe low. Upon annealing at 500 �C the intermetallic phasegrows with a strong (111)-texture on (111)-Pt due to goodlattice matching (|d| = |(d111Pt � d111PtxPb)/d111Pt| = 3.3%),and this translates into the intense peak observed in Fig. 1a.

Processing of PZT on LSMO-terminated (111)-Pt sub-strates leads to a polycrystalline microstructure withoutthe prevalence of one particular orientation. This micro-structure simply reflects the polycrystalline nature of sol–gel-processed LSMO (see Fig. 1c).

The results depicted above on PZT growth on (111)-Ptare in many aspects different from those presented by Chenet al. [25]. In their study on the effects of precursor solutionand pyrolysis schedule on texture development of PZT filmsof the same stoichiometry as used here it was shown thatintermediate pyrolysis temperatures of 500 �C resulted ina high degree, almost 90%, of (100)-texture. Chen et al. dis-cussed their results in terms of effects related to annealingatmosphere, heating rate and film thickness. In all cases

b) and on LSMO-terminated platinized silicon (c). (a) XRD patterns afterase, the weak peak to the perovskite phase and the broad peak around 30�

0 �C where the strong (111) perovskite peak can be seen; (c) XRD patternsaks marked by (�) belong to the LSMO phase.

S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338 2331

thinner sol–gel PZT films (�100 nm) were shown todevelop either a (100)-texture or a random texture. Intheir opinion the (100)-texture developed in thinner sol–gel films because oxygen supply to the interface betweenfilm and Pt was more efficient and would result in the for-mation of a PbO template layer that would promote the(100)-texture. Only when thicker sol–gel films (�300 nm)are annealed under a reducing atmosphere of Ar + 2%O2 does predominance of the (111)-texture occur, and thiswas attributed to the formation of an intermetallic interfa-cial layer. Brooks et al. [2,26] and Reaney et al. [27] alsoinvestigated the effects of pyrolysis temperature and heat-ing rate on texture formation in sol–gel PZT films. Theirresults show that an intermediate pyrolysis temperatureof 525 �C and subsequent annealing at 700 �C, similar tothe temperatures used in the present work, resulted in amixed (100)/(1 11)-texture. Discussion of their resultswas mainly based on nucleation energetics related to theformation of preferred crystal orientations, on bottomelectrode effects, i.e. essentially modification of Pt-electrodesurface chemistry by Ti-diffusion from the Ti-adhesion

Fig. 2. AFM amplitude micrographs of the PZT thin films on Pt-terminated subthe homogeneous thinning of the film. Notice that sputtering alters only marginthe highest and lowest feature in the image) which remains almost unchanged bcorresponding phase micrograph after sputtering of the PZT film. The phase iintermetallic phase.

layer, and pyrochlore stability. Template layer formation,as discussed by Chen et al. [25], were not taken intoaccount, though the microstructures obtained by bothgroups were very similar. The present work shows differentresults to those presented above. In our case the intermedi-ate annealing temperature of 500 �C resulted in the strong(111)-texture shown in Fig. 1, despite the fairly low thick-ness of the 1L films, and this texture remains unaffected bythe thickness of the films. The discrepancy between ourresults and those of Chen and Chen and Brooks et al. can-not be explained at this stage, though we should bear inmind that small changes in the processing parametersmay change the microstructure of PZT films, particularlythose with high Zr/Zr + Ti ratios. Among these variationsare the way precursor solutions are prepared, thermal pre-treatment of substrate and rate of temperature increaseduring thermal annealing. All these process parameterswere different in the work of Brooks et al. and Reaneyet al. and could explain why our results differ from theirs.It should, however, be pointed out that similar results toours were obtained by Huang et al. [7].

strate in (a) the as-deposited state, and (b) after sputtering 17 nm, showingally the roughness of the film, expressed by the Z-range (distance betweenetween the micrographs; (c) is an AFM amplitude micrograph and (d) themage indicates some material contrast which might suggest some residual

Table 1Standard atomic sensitivity factors (ASFs) from Scofield and calibratedvalues from Fig. 3.

Element Pb(4f7/2) Zr(3d5/2) Ti(2p3/2) O(1s)

ASF Scofield 12.73 4.17 5.22 2.93ASF calibrated 6.63 6.61 5.22 2.93

2332 S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338

3.2. XPS depth profiling

In previous investigations on PbTiO3 thin films [19] wehave shown that the depth profiling conditions adapted inthis work did not lead to preferential sputtering of the con-stituent elements as in other studies where metallic Pb built-up during sputtering. Using AFM at different sputteringstages we have also shown that homogeneous thinning ofthe films occurs with negligibly small variation in the filmroughness. Similar results were obtained on PZT thin films.As shown in Fig. 2, the Z-ranges (which characterize themaximum height between the lowest and highest featurein the field under investigation) of topography amplitudeimages of the as-processed and sputtered films differ onlyslightly. The root mean square (rms) roughness values ofthe different surfaces are 1.8 nm for the native film,1.65 nm for intermediately sputtered and 1.6 nm for thefully sputtered film. For the latter, the phase image is alsoshown. The bright contrast of the worm-like features sug-gests the presence of residual particles of a material with dif-ferent properties (see below for discussion).

Calculation of the atomic compositions of the PZT thinfilms was first performed following the usual procedure.The intensity of the single elements is obtained from themeasured spectra after Gauss–Lorentz fitting and Shirley-background subtraction [28]. The relative atomic fractionsof Pb, Zr, Ti, O and Pt were calculated using the usualequation:

X A ¼IA=SAP

iI i=Siði ¼ Pb;Pt;Ti;Zr and OÞ; ð1Þ

where IA is the measured intensity and SA is the ASF of ele-ment A. Standard ASF values were taken from Scofield [29]for the sub-bands Pb(4f7/2), Ti(2p3/2), Zr(3d5/2), O(1s) andPt(4f7/2). While this method was successful in the case ofPbTiO3 [18] where the stoichiometry found using XPSwas very similar to the nominal one, the analysis of PZTshowed an overestimation of Zr composition and an under-

Fig. 3. Calculated elemental concentration ratios (Pb exp, Zr exp, Ti exp)from the XPS peaks using the standard ASF values from Scofield andcomparison with the nominal ratios (Pb theo, Zr theo, Ti theo).

estimation of Pb and, to a lesser extent, Ti. Therefore, wehave undertaken an empirical correction of the ASFs usingthin films with increasing Zr concentrations. Fig. 3 shows acomparison between the nominal and experimentally deter-mined elemental concentrations (using a standard ASFfrom Scofield) where the discrepancy mentioned abovecan be clearly seen. The standard ASFs were correctednumerically using Fig. 3 to yield the calibrated values givenin Table 1. With these values the calculated elemental com-positions largely agree with the nominal ones (see Fig. 5). Itshould be pointed out that correction of the ASFs usingcalibration with different Zr/Ti + Zr ratios was performedassuming no change in Pb stoichiometry and not takinginto account the Pb excess used in the precursor solution.However, as outlined above and discussed in detail below,there is a gradient of Pb stoichiometry in the films, e.g. viaPbO segregation and loss of Pb to the electrode during theformation of the intermetallic layer. We therefore assumedthat these Pb losses, including losses via evaporation ofPbO during annealing, were compensated by the PbO ex-cess of 10 mol% used. While this assumption is quite diffi-cult to verify, previous work by Reaney et al. [27] hasshown using X-ray fluorescence analysis of PZT films withdifferent PbO excess that a 10% PbO excess was sufficientto compensate for PbO losses, so that the resulting PZTfilms had about 98% Pb site occupancy.

XPS analysis and depth profiling of the 1L film annealedat 500 �C indicate a Pt(4f7/2) peak shift of approximately�1 eV at the interface between the oxide layer and the Ptelectrode, i.e. at the intermetallic PtxPb layer (see Fig. 4a).Furthermore the full width at half maximum (FWHM) ofthe Pb(4f7/2) peak increases at the interface with the appear-ance of a shoulder at low binding energy that can be attrib-uted to metallic Pb (see Fig. 4b and Table 2). We do notbelieve that this is an artifact from preferential sputteringof oxygen since the Pb peak remains unchanged after sput-tering up to 90% of the film thickness. This suggests that theelectronic configurations of Pt and Pb are changed due tobinding between these elements. Since the electronegativityof Pb(1.8) is smaller than that of Pt(2.2), electron transferfrom Pb to Pt takes place, which translates into a shift inthe Pt binding energy to lower values [30]. The depth profiledisplayed in Fig. 4c shows a diffuse interface between Pt andthe oxide layer, the thickness of which can be estimated tobe of the order of 10 nm. If we consider the Pt(4f7/2) peakshift as the signature of the intermetallic phase and plotthe peak position vs. sputter depth in Fig. 4d, it is possibleto extract the intermetallic layer thickness, which should bein the range between 6 and 9 nm. This is a reasonable value

Fig. 4. XPS analysis of PZT on Pt-terminated silicon substrate after annealing at 500 �C in air (predominance of the interfacial intermetallic phase, seeFig. 1a). (a) Comparison between Pt 4f peak positions at the interface and in bulk Pt film; (b) Pb(4f7/2) peak shape at different sputter depths, indicatingpeak shape change at the interface; (c) calculated elemental concentration as function of the sputter depth using calibrated ASFs. Errors on Ti and Zrconcentrations of ±4% are too small and cannot be visualized as error bars. The interface position is indicated by a hatched area; (d) peak positions asfunction of sputter depth (E0 denotes nominal peak positions corresponding to PZT or in the case of Pt to its bulk metallic state).

Table 2Binding energies of Pb(4f7/2), O(1s) and Pt(4f7/2) at the surface after argonion cleaning, in the film bulk, at the oxide–Pt interface and in the Pt bulk.

XPS signal E (eV) onsurface

E (eV)in film

E (eV) atinterface

E (eV) beneathinterface

Pb(4f7/2) 137.4 137.6 137.5 137.5135.8

O(1s) 529.2 529.9 529.7 –Pt(4f7/2) 70.2 71.1

S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338 2333

taking into account the resolution of the system of approx-imately 3 nm.

However, it should be pointed out that due to the fairlylow film thickness used of roughly 30 nm, the thickness ofthe interfacial layer should make up to 10% of the totalthickness of the film. Since the formation of the interfaciallayer is limited by diffusion, the oxide film can no longerbe considered as an infinite source for Pb supply. In otherwords we should expect thicker interfacial layers for thickeroxide films. Indeed, from the work of Huang et al. [31], we

can infer a thicker intermetallic layer from their TEM cross-section of a three-layer PZT film annealed at 440 �C.

Fig. 5a shows the composition profile of the film afterannealing at 700 �C, where according to Fig. 1b only theperovskite phase is present. The composition profile is prin-cipally similar to that shown above for the film treated at500 �C and the diffuse interface has about the same thickness.If we consider again the binding energy of Pt(4f7/2) vs. sputterdepth shown in Fig. 5b, it becomes clear that the annealingtemperature of 700 �C did not lead to full decompositionof the intermetallic layer as claimed by Huang et al. [31]and Chen [8]. The highest shift of the Pt(4f) peak energyoccurs between 20 and 28 nm, so that the thickness of theresidual intermetallic layer should be of the order of 8 nm.The distribution of binding energy shifts of Pt(4f) might beexplained in terms of Pb diffusion into Pt without the forma-tion of an ordered intermetallic phase.

Furthermore, the analysis of the film shows evidence ofthe existence of a segregation layer of lead oxide, probablyPbO. According to the general equation of segregation atinterfaces for a binary system [32]:

Fig. 5. (a) XPS depth profiling of PZT on Pt after annealing at 700 �C (perovskite phase); (b) indicates Pt 4f peak position as function of sputter depth; (c)XPS depth profiling of PZT on LSMO-terminated substrate.

2334 S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338

CSB ¼ �1

kT@c

@LnX B

� �; ð2Þ

where CSB is the surface coverage of solute B, in our casePbO, XB is its concentration in the bulk, c is the surfacetension, k is the Boltzmann constant and T is the absolutetemperature.

Segregation occurs if the solute lowers the surface ten-sion of the system. The high vapor pressure of PbO andits low heat of sublimation [33] indicate a low surface ten-sion of the order of 0.2 J m�2. In comparison, TiO2 isreported to have a surface tension of 0.4 J m�2 [34]. Thussegregation of PbO at the surface of PZT is expected, inagreement with XPS analysis.

Depth profiling of PZT/LSMO is shown in Fig. 5c. Inthis case too there is a slight enrichment of PbO on the sur-face for the reasons discussed above. The formation of a dif-fuse interfacial layer is also evident. Interdiffusion betweenLSMO and PZT takes place over approximately 30–40 nmwith diffusion of La and Mn into the PZT layer and Pb, Tiand Zr into the LSMO layer. The Sr was not detectable inthe diffusion layer within the sensitivity limit of XPS analy-sis, generally in the range of 0.1–1 at.% depending on theelements analyzed. These results are in agreement with pre-vious work [12]. We surmise the formation of a gradient

Fig. 6. Dielectric constant vs. DC bias field of PZT on Pt for different thicknesplot of the reciprocal of the measured capacitance vs. PZT layer thickness.

interfacial layer with a stoichiometry corresponding to(Pb1�x,La3x/2)((Ti,Mn)y, Zr1�y)O3 which could be ferro-electric on the PZT side and probably non-ferroelectric onthe LSMO side (in the La2O3–PZ–PT system a cubic phaseis expected for high La content [35]).

3.3. Thickness dependence of the dielectric and ferroelectricproperties

C–V curves with the butterfly shape characteristic for fer-roelectric switching are shown in Fig. 6a for different filmthicknesses on Pt. The dielectric constant varies substan-tially with thickness. This thickness effect on the dielectricconstant may be related to the presence of interfacial layerswith lower dielectric constants. If we suppose two platecapacitors in series consisting of a bulk ferroelectric filmwith a capacitance Cb and an interfacial layer with Ci, wecan write:

1

Cm¼ 1

Cbþ 1

Ci¼ db

e0e0bAþ 1

Ci; ð3Þ

where Cm is the measured capacitance, db is the thicknessof the ferroelectric bulk layer, e0 is the permittivity of freespace, e0b is the dielectric constant of the ferroelectric bulkfilm and A is the area of the capacitor. If we consider the

ses showing the butterfly shape characteristic of ferroelectric switching; (b)

Fig. 7. Dielectric constant vs. DC bias field of PZT on LSMO for differentthicknesses also showing the butterfly shape characteristic for ferroelectricswitching.

S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338 2335

thickness of the interfacial layer, di, being negligible incomparison to db—which is to some extent legitimate basedon the XPS results above, particularly for the thicker films,where an interfacial layer thickness of approximately10 nm was found—we can write dm � db, where dm is themeasured film thickness. The plot of the reciprocal of Cmvs. film thickness yields a straight line as shown inFig. 6b. The slope allows a thickness-independent dielectricconstant of the ferroelectric film, e0b, to be calculated andthe intercept with the ordinate axis yields the reciprocalof Ci from which the dielectric constant of the interfaciallayer, e0i, can be obtained. From the Pt peak position shownin Fig. 4d, the interfacial layer thickness was estimated tobe of the order of 10 nm which yields approximately 90for e0i. This value indicates that the interfacial layer is notferroelectric in nature. The existence of the interfacial layercauses a negative bias of the dielectric constant and coer-cive field (imprint) particularly for thinner films: the obser-vation of imprint suggests the onset of a bias field at thefilm–Pt interface which could build-up due to the presenceof charged defects, e.g. oxygen vacancies [36]. It is likelythat Pb depletion of the PZT films at the interface, whichis necessary for the formation of the intermetallic layer atthe bottom electrode and the segregated PbO at the sur-face, is balanced by the creation of oxygen vacancies thatcould lead to the observed imprint. These findings are inagreement with the results presented in previous work[20,36,37] and have been rationalized in terms of the exis-tence of a space charge layer at the electrode interface, fol-lowing the model developed by Miller et al. [38], andmodified by Le Rhun et al. [36], where in the expressionof the dielectric displacement the potential barriers andthe thicknesses of the film–electrode interfaces are takeninto account (for details, see Bouregba et al. [37]). The lar-ger imprint observed as the PZT film becomes thinner, e.g.100 nm, can be explained in terms of the relatively higherthickness of the disturbed interfacial layer.

The C–V curves of PZT on LSMO are shown in Fig. 7for different thicknesses. The maximum dielectric constantobtained for the thicker film is very similar to that of PZTon Pt. Furthermore fairly symmetric butterfly shapes areobtained for the C–V curves, regardless of film thickness;coercive fields are slightly higher for the thinner films.

Applying the same analysis as above, a small signal,thickness-independent dielectric constant of 893 isobtained. The capacitance of the interfacial layer of 14 nFleads to a dielectric constant of 157 assuming an interfaciallayer thickness of 30 nm. This value can be explained if weassume the formation of a non-ferroelectric, high La-con-taining cubic phase as reported in the La2O3–PZ–PT system[35].

The model discussed above describing thickness effectsand the derived values of the dielectric constants of theinterfacial layers relies on coarse assumptions about dis-crete layers in series. However, despite the fact that theinterfacial layer thickness could be derived from the Pt peakposition outlined above, it is hardly believable that the

properties and chemical composition of this layer are con-tinuous over its entire thickness. Rather, a diffuse layer witha gradient of composition and properties is expected.Indeed the Pt peak position takes a range of values thatshould reflect the nature of the diffusiveness of the interfa-cial layer. The dielectric constant of the interfacial layerthus derived should be taken with caution. The same couldbe said about the interfacial layer with LSMO. Addition-ally, for the latter the assumption of negligible thicknessof the interfacial layer is probably not valid as this layercomprises more than 15% of the entire thickness in the caseof thinner films.

At higher electrical fields, where the films are expected tobe in a state of saturated polarization, the dielectric con-stant converges to a value of approximately 150, regardlessof film thickness and substrate heterostructure. This valueprobably reflects the intrinsic dielectric constant of thePZT film stoichiometry under investigation.

The ferroelectric hysteresis loops are shown in Fig. 8a forfilms on Pt and in Fig. 8b for films on LSMO. Principallythe films show different behavior depending on substrateheterostructure and thickness. Thinner films of PZT on Ptshow ‘‘tilted” hysteresis loops with higher coercive fields,Ec, and lower saturation polarization, Ps (45 lC cm�2

instead of 61 lC cm�2 for the thicker film). However, theremnant polarization (Pr, polarization at zero field) isbarely affected by film thickness (26 lC cm�2 instead of27 lC cm�2 for the thicker film). Thus it seems that thedefects created at the electrode interface, probably mainlyoxygen vacancies, could be responsible for domain bound-ary pinning; a higher electric field is then necessary for fer-roelectric switching, leaving the remnant polarization of thefilms otherwise unaffected. In the case of PZT on LSMO, Ecis almost independent of thickness, whereas Pr and Ps aresubstantially affected. The thinner film shows lower polari-zation values, a result that might be amenable to the highercontribution of the diffuse interfacial layer which as men-tioned above could include non-ferroelectric phases. Fur-

Fig. 8. (a) P–E ferroelectric hysteresis loops for PZT on Pt and different thicknesses; (b) the same for PZT on LSMO.

2336 S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338

ther, microstructural effects, such as the random nature offilm orientation, could also contribute to the lower polariza-tion observed.

4. Further discussion

We first consider the PZT thin films deposited on(111)Pt/Ti/SiO2/Si. The formation of an interfacial inter-metallic layer between Pt and Pb could be unambiguouslyshown after an annealing step at 500 �C. It is likely that thislayer already nucleates during the pyrolysis step at 350 �Cand grows at 500 �C on (111)-Pt, resulting in the intensepeak shown in Fig. 1a. The remainder of the oxide filmtransforms to nanocrystalline quasi-amorphous pyroch-lore, as evidenced by the broad peak around 2H = 32�. Inprevious work [31] the existence of a transient intermetallicPt3Pb between PZT and Pt has been shown using XRD andtransmission electron microscopy. Others [18] haveattempted to characterize lead-titanate/Pt film stoichiome-try using XPS and Rutherford backscattering spectroscopy(RBS). However, besides the preferential sputtering prob-lems mentioned above no details of the peak energies ofPt and Pb were presented. In the present work, we tookadvantage of the formation of the intermetallic layer at500 �C to characterize it in detail using XPS. This hasallowed us to establish the interfacial composition andinterfacial layer thickness in films with crystallized perov-skite structure and relate them to the functional propertiesof the films and their thickness dependence. The crystalliza-tion of a highly textured thin interfacial layer of PtxPb wasevidenced by XRD measurements. Using XPS depth profil-ing and a narrow pass energy for collecting the relevantpeaks with high resolution we could determine the XPS sig-nature of the intermetallic phase; its attributes are a nega-tive shift of Pt(4f7/2) binding energy by about 1 eV and alarge FWHM of Pb(4f7/2) with eventually the appearanceof metallic Pb (as evidenced by the appearance of a shoulderon Pb(4f7/2) at lower binding energy).

In specific studies [30,39] on Pt/Ni and Pt/Ru systems forcatalytic applications a shift in the binding energy of Pt

towards lower values was shown. It was concluded thatthe electronic structure of Pt was changed upon alloy for-mation with electron transfer from Ni or Ru to Pt (follow-ing the electronegativity series for these elements). In arecent theoretical study of Pt(111) and Pt3Pb(111) Ranjanet al. [40] showed that the substitution of one Pt by one Pbatom in the Pt lattice leads to the lowering of the Pt d bandbelow the Fermi level, whereas the Pb s band is ‘‘pusheddown in energy due to Pt–Pb interaction”. They also notesome electron transfer from Pb to Pt. These considerationsconstitute further support for our results.

The existence of an interfacial intermetallic layer is inti-mately related to the strong (111)-texture obtained forPZT. The lattice matching between the two phases is verygood as can be judged from the low value of d = (d111(PtxPb) � d111(PZT))/d111(PtxPb) = 0.34% (d-values arecalculated from PtxPb-JCPDS#06-0574) and PZT-JCPDS-33-0784).

Because the formation of the perovskite phase at highertemperatures is thermodynamically favored we shouldexpect a complete transformation of the intermetallic phaseat the interface with the oxide bulk according to the follow-ing reaction:

Pt3Pbþ1

2O2 þ ðTi;ZrÞO2 ! PbðTi;ZrÞO3 þ 3Pt ð4Þ

However, the kinetics of nucleation and growth couldlead to residual intermetallic phase that is not detectablewith XRD (due to the low volume fraction), but can bedetermined with XPS depth profiling. A nucleation andgrowth scenario is expected to start with the formation ofcoherent perovskite nuclei with a definite orientation rela-tionship with the intermetallic phase. Growth proceeds intothe oxide matrix with the preferred (111) orientation whilethe intermetallic phase decomposes due to Pb incorporationinto the perovskite phase. If nuclei cusps form at the inter-face, then it is legitimate to assume that lateral growth of thenuclei is far less favored due to incoherency of the interfaceand higher interfacial barriers for nucleation and growth.That is, some residual intermetallic particles are left at the

S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338 2337

interface with the substrate. The AFM phase image pre-sented above (Fig. 2d) should provide some support to thisassumption. Recording the phase shift between excitationand oscillation of the cantilever generates phase images witha complex contrast formation, mainly discussed in terms ofdissipation [41]. While a quantitative analysis is difficult dueto the many sources of the phase shift and their nonlinearinteraction, a careful qualitative interpretation allows mate-rial inhomogeneities to be resolved in detail [42–44]. If wecombine the observations made on AFM phase images,which provide a high lateral resolution, with the results ofthe XPS analysis, with the high surface and chemical sensi-tivities peculiar to this method, it is possible to infer that theworm-like features observed in the AFM phase image areresidual particles of the interfacial layer that probably delin-eate the former PZT grains and constitute a continuous net-work between the PZT grains and the Pt electrode.

Based on the experimental observations depicted aboveit should appear obvious that nucleation and growth of(111)-textured PZT is mediated by the (111)-texturedPtxPb layer that forms during pyrolysis and intermediateannealing at 500 �C. In previous work, Brooks et al. [2]claimed that (11 1)-PZT directly forms on (111)-Pt,although they showed that an intermetallic phase wasformed under a reducing gas atmosphere. We disagree withthem as since it is clearly shown that residual intermetallicphase was still present even at the final annealing tempera-ture of 700 �C, this phase should control the formation of(111)-PZT. In a recent work Muralt [45] analyzed texturecontrol in seeded, reactive sputter-deposited PZT thin filmsin terms of nucleation kinetics using the classical formalismof nucleation theory as amply developed by Christian [46],and adapting it to reactive sputtering. The general outcomeis that the rate of chemisorption/desorption of PbO on Ptand how this is affected by seeds, e.g. TiO2, PbTiO3, shouldcontrol texture formation. In particular, it was shown thatdirect nucleation of (11 1)-PZT on Pt was not favored andthat nucleation on TiO2 seeds was faster. This wasexplained by the affinity of TiO2 for PbO, thus favoringits chemisorption. In principle this model describes mecha-nisms similar to those discussed by Chen et al. [25] in termsof PbO effects on texture formation in sol–gel PZT films,though metastable reactions between Pt and PbO, e.g. for-mation of intermetallic phases, were not considered in thework of Muralt. While the model developed by Muraltshould have only limited applicability to our results, the for-mation of (111)-texture on the TiO2-rich stoichiometrydescribed in that work could help understanding the mech-anisms of (111)-PZT growth on PtxPb. It is conceivablethat the higher affinities of PbO and TiO2 and the effectsof Zr on the activity of PbO in the PbO–TiO2–ZrO2 systemshould first lead to the formation of tetragonal (111)-PZT(at the PTO-rich side of the PTO–PZ system) via the wettingof the PtxPb layer with a TiO2-rich film and subsequentdecomposition of PtxPb as outlined above. Incorporationof Zr should then proceed after the formation of the tetrag-onal phase. The formation of PZT powder from individual

oxides is well known to proceed in this manner [47] andshould translate to PZT films in their early crystallizationstages, though on a much smaller microstructural scale.Careful experiments are, however, necessary to encompasscrystallization steps in sol–gel PZT films.

The thickness dependence of the dielectric properties ofthe PZT thin films on Pt can now be linked to the existenceof the residual intermetallic phase discussed above. Thedielectric constant of the interfacial layer of approximately90 estimated from Fig. 6b probably characterizes a layerthat encompasses all disturbed regions at the interface,including Pb-deficient perovskite and other disturbed oxidelayers and would rather be better described by more thantwo capacitances in series.

The difference in the C–V behavior of PZT on Pt andLSMO may be explained by the different nature of the inter-facial layers formed. The XPS results depicted above showthat to some extent there is interdiffusion between PZT andLSMO, resulting in a diffuse interface (see above) with var-iable stoichiometry of the type (Pbx,La1�x)((Ti,Mn)y,Zr1�y)O3, where Pb

2+ is substituted by La3+ and Ti4+ (orZr4+) by Mn3+. The incorporation of La ions into theperovskite lattice should lead to the annihilation of oxygenvacancies, whereas Mn ions are expected to induce their for-mation, according to the following reactions, using Vink–Kröger notation [48]:

Mn2O3 ���!2ðTiO2Þ 2Mn0Ti þ 3OO þ V ��O ð5ÞLa2O3 þ V ��O ���!2ðPbOÞ 2La�Pb þ 3OO ð6Þ

Although the XPS analysis suggests similar amounts ofLa and Mn in the interfacial layer, this should be taken withcaution because of the low signal-to-background ratiosobtained in this area. Nevertheless we can state that theinterface between PZT and LSMO is far less disturbed,e.g. with fewer oxygen vacancies, than that between PZTand Pt, and this translates into symmetric C–V curves andthe absence of imprint.

Regarding the ferroelectric properties, the substrate het-erostructure and film thickness all affect the magnitude ofpolarization and coercive field. The highly (111)-texturedfilms obtained on Pt show high remnant polarization valuesregardless of film thickness. However, high coercive fieldsare necessary for ferroelectric switching of the thinner films,which is an indication of ‘‘hard” ferroelectric domains.Since all films exhibit the same (111)-texture, the highercoercive field could be the result of domain pinning by dipo-lar V ��O � V 00Pb defects [47] (which are expected to be moreabundant with respect to the film thickness in thinner films)and/or tensile stress effects (due to mismatch of the thermalexpansion coefficients of PZT and Pt) that are expected tobe less relaxed for the thinner films. Saturation polarization,which should reflect the polarization of a fully switchedfilm, is found to be smaller for the thinner film, and thisagain can be rationalized in terms of domain boundary pin-ning by charged defects. The situation looks different for

2338 S. Habouti et al. / Acta Materialia 57 (2009) 2328–2338

PZT on LSMO, where the ferroelectric values are overallsmaller that those of PZT on Pt. This is first amenable tothe random microstructure obtained on LSMO that proba-bly involves a higher density of in-plane, nonswitchabledomains. In the case of the thinner film on LSMO, thehigher contribution from the diffuse interfacial layer addi-tionally contributes to the poor polarization values.

5. Conclusions

Thin PZT films of different thicknesses ranging from 100to 600 nm were deposited on two different substrate hetero-structures, namely Pt/Ti/SiO2/Si and LSMO/Pt/Ti/SiO2/Si.Film characterization included XRD, AFM, XPS, dielectricand ferroelectric properties. The following results may beinferred:

– Films deposited on Pt-terminated substrates show ahighly (111)-textured intermetallic phase, PtxPb, afterannealing at an intermediate temperature of 500 �C inair. XPS analysis of these films reveals that a negativeshift of 1 eV in the binding energy of Pt(4f7/2) and alarge FWHM of Pb(4f7/2), with eventually the appear-ance of metallic Pb, constitute the main attributes ofthis phase.

– PZT films crystallize with preferential (111)-texturewhen deposited on Pt. This is amenable to nucleationand growth on the PtxPb template layer. In contrast,crystallization on LSMO-terminated substrates leadsto randomly oriented grains.

– XPS depth profiling shows that the interfacial inter-metallic phase remains at the interface between Ptand PZT, probably in form of residual particles. A dif-fuse interface forms between LSMO and PZT as aresult of interdiffusion. Variable amounts of La andMn were found at the PZT side.

– It is shown that the interfacial layers are responsiblefor the thickness dependence of the dielectric and fer-roelectric properties. Analysis of the results suggeststhat these interfacial layers are non-ferroelectric innature and eventually contain charged defects.

Acknowledgment

The Deutsche Forschungsgemeinschaft (DFG) is grate-fully acknowledged for financial support, Grant # ES119/6-1(2).

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Substrate heterostructure effects on interface composition, microstructure development and functional properties of PZT thin filmsIntroductionMaterials and methodsResults and discussionXRDXPS depth profilingThickness dependence of the dielectric and ferroelectric properties

Further DiscussiondiscussionConclusionsAcknowledgmentsAcknowledgmentReferences