structural evolution of lithium niobate deposited on sapphire (0 0 0 1): from early islands to...

Journal of Crystal Growth 196 (1999) 141—150

Structural evolution of lithium niobate deposited on sapphire(0 0 0 1): from early islands to continuous films

Franck Veignant!,*, Madeleine Gandais!, Pascal Aubert",1, Guy Garry"

! Laboratoire de Mine& ralogie-Cristallographie, CNRS URA 09, Universite& s Paris VI et Paris VII, 4 Place Jussieu, 75252 Paris Cedex 05,France

" Thomson CSF, Laboratoire Central de Recheche, Domaine de Corbeville, 91404 Orsay Cedex, France

Received 3 February 1998; accepted 16 September 1998

Abstract

Lithium niobate films were deposited on sapphire substrates at temperature ¹$"750°C by pulsed laser deposition.

Their structure was investigated at different stages of growth by transmission electron microscopy. The epitaxialrelationship already found by X-ray diffraction has been confirmed to be (0 0 0 1)

A-2O3E(0 0 0 1)

L*N"O3and

[1 1 21 0]A-2O3

E[1 1 21 0]L*N"O3

. A second cristallographic variant, rotated by 60° about c-axis, was also present in somesamples. At the early stages of growth, lithium niobate constitutes pyramidal islands with M0 1 11 2N faces that have meanheight of 40 nm and are misoriented from each other by 0.5° on average. At the coalescence stage, surface migration ofalready deposited and incoming atoms from the source leads to the disappearing of pyramidal faces and the formation of20 nm height plateaux; the misorientation between initial nuclei is accommodated by formation of subgrain boundaries.Continuous films have a mosaic texture also inherited from the initial misorientation of islands. They display mechanicaltwins and cracks induced by tensile stresses occurring at the cooling stage from deposition to room temperature. Crackshave variable width according to the film thickness and they are oriented preferentially along S1 0 11 0T directions. Twinsare located at the border of cracks and seem to be preferential sites for their initiation. ( 1999 Elsevier Science B.V.All rights reserved.

PACS: 68.55.a; 81.15.Fg; 61.16.Bg

Keywords: Lithium niobate; Thin film; Sapphire; Crystal growth; Epitaxy; Mechanical strains

*Corresponding author: Fax: #33 01 44 27 37 85; e-mail:[email protected].

1Present address: Laboratoire Multicouches Nanometriques,Institut des Materiaux de l’Universite d’Evry Val d’Essone, Bd.Franiois Mitterand, 91025 Evry Cedex, France.

1. Introduction

Lithium niobate (LiNbO3) has been widely

studied during the last decades because of thegreat number of applications associated with itsseveral specific physical properties: ferroelectricity,piezoelectricity, high electro-optics and second

0022-0248/99/$ — see front matter ( 1999 Elsevier Science B.V. All rights reserved.PII: S 0 0 2 2 - 0 2 4 8 ( 9 8 ) 0 0 9 0 5 - 1

harmonic generation coefficients. Bulk single crys-tals are already commonly used in acoustic andoptical devices. The growth of thin films has beenrecently envisaged for the fabrication of highly in-tegrated optical devices. As a starting point tostudy the properties of such films, sapphire(aAl

2O

3) is a good candidate as substrate because

its refractive indices are smaller than the refractiveindices of lithium niobate, which is a necessarycondition for the optical confinement in waveguidedevices. Furthermore, sapphire and lithium niobatehave a close crystallographic structure (both crys-tallise in the trigonal system). The lattice para-meters of the two materials are: a"4.758 A_ andc"12.991 A_ for aAl

2O

3[1]; a"5.147 A_ and

c"13.862 A_ for LiNbO3

[2].The lattice mismatch in the (0 0 0 1) plane equals

approximately 8%, which seems high for good epi-taxy. Nevertheless, epitaxial growth of lithium nio-bate on sapphire has already been successfullyachieved by several methods: MOCVD [3], sput-tering [4], sol—gel method [5—7] and pulse laserdeposition (PLD) [8—10]. However, films havea mosaic structure and exhibit cracks of severaltens of nanometers as soon as their thickness ex-ceeds 180 nm. This paper reports on the structuralaspect of thin films of lithium niobate studied byTEM. The films were prepared using the PLDmethod and were deposited on sapphire(0 0 0 1).Previous works [9,10] have specified the firststructural and optical properties of these films,in particular the epitaxial relationship betweensapphire and lithium niobate: (0 0 0 1)

L*N"O3E

(0 0 0 1)A-2O3

and [1 1 21 0]L*N"O3

E[1 1 21 0]A-2O3

. Thepresent work concerns the microstructural evolu-tion at different stages of growth in order to under-stand the origin of defects and to prevent theirappearance in future experiments.

2. Experimental procedure

The films were deposited by using a KrF excimerlaser operating at 248 nm and 5 Hz pulse fre-quency. High density and stoichiometric ceramicsof lithium niobate were used as targets. The sub-strates are sapphire(0 0 0 1) platelets: they weremechanically polished by the manufacturer (Joh-

nson Matthey) and cleaned in organic solvent. Theaverage roughness of the substrate surface has beenmeasured by AFM to be 0.5 nm. The influence ofdeposition conditions (deposition temperature (¹

$),

oxygen pressure (PO2

)) on the composition andtexture of the thin films was investigated in otherworks [8,11], especially the formation of the unde-sirable LiNb

3O

8and Li

3NbO

4compounds. All the

films presented here were grown under the follow-ing conditions: P

O2"1.33 mbar, ¹

$"750°C, the

distance target—substrate was 45 mm. The coolingrate is approximately 10°C/min. In these condi-tions, X-ray diffraction experiments have shownthat the films are only composed of lithium niobate,and that the film grows epitaxially [9]. The Li/Nb,ratio in the films is not accurately known. Thenonstoichiometry is not expected to have signifi-cant influence in the phenomena observed in thisstudy. In particular, the lattice parameter vari-ations over the range of the nonstoichiometry donot exceed 10~2 A_ [12] This is smaller than thec parameter variation considered in our films andwhich will be described later. The parameters wewill use as reference are those of stoichiometriccompound.

Films obtained at different deposition times wereexamined: 20 s, 1 min, 3 min, 10 min, 20 min. Thefilms deposited during the shortest times (t"20 s—3 min) are discontinuous whereas those ob-tained at the latter two times (t"10—20 min) arecontinuous and their thicknesses were found byRBS to be 110 and 220 nm, respectively.

Conventional transmission electron microscopy(CTEM) and high resolution transmission electronmicroscopy (HRTEM) were performed usinga JEOL 2000 EX microscope and a TOPCON002B microscope, respectively. The latter givingpoint to point resolution of 0.18 nm. Both micro-scopes operating at 200 kV. Specimens parallel tothe film were prepared for plane view observation.They were obtained by grinding, dimpling and Ar`ion milling the substrate. Cross-sectional specimenswere also obtained by fabricating a sandwich oftwo film—substrate pieces glued together with thefilms facing each other, then cutting slices normalto the corresponding directions and finally grindingand ion milling the slices. Ion milling was per-formed on a GATAN PIPS operating at 5 kV.

142 F. Veignant et al. / Journal of Crystal Growth 196 (1999) 141–150

3. Results and discussion

3.1. Morphology and texture

Much knowledge of the film formation can beobtained by following the morphology and texturefrom the early to the later stages of growth. Planeviews of discontinuous films (t"20 s, t"1 min,t"3 min) in the [0 0 0 1] orientation and theirdiffraction patterns are shown in Fig. 1. Sample I,obtained after the shortest deposition time (20 s)shows triangular islands whose sides are orientedalong the S1 1 21 0T axis (Fig. 1a). Samples II andIII, deposited during 1 and 3 min, respectively, rep-resent stages of coalescence before the formation ofa continuous thin film. Both are composed of co-alesced grains, but in sample II sides with S1 1 21 0Torientation remain present (Fig. 1b), whereas theyhave completely disappeared in sample III(Fig. 1c). The determination of the three-dimen-sional morphology of the overgrown nuclei hasbeen made using dark-field images in weak beamconditions. This method consists of performingdark-field images with a strong lithium niobatediffracted beam far from the Bragg orientation.

Fig. 1. Bright-field micrograph of (a) sample I, (b) sample II, (c)sample III, in plane view along the [0 0 0 1] zone axis, and (d)selected area diffraction pattern of the sample II.

Fig. 2 shows the corresponding micrographs ofsamples I, II, III performed with the 0 1 11 2 diffrac-ted beam. The periodic thickness fringes appearingalong the island edges of sample I (Fig. 2a) andextending into the whole island indicate that theislands have pyramidal-like shapes. The pyramidfaces have been identified as being M0 1 11 2N bymeasurement of the fringes periodicity which isa known function of the angle between the islandface and substrate plane, the structure factor of thereflection, and the deviation parameter to Braggorientation (see the appendix for details). The meanheight of the islands has been deduced from thenumber of thickness fringes to be 40 nm on average.On the other hand, groups of coalesced grains insamples II and III (Fig. 2b and Fig. 2c) appear asplateaux limited by inclined surfaces that are notcristallographically well defined. The height of theplateaux is 22$5 nm for sample II and 25$5 nmfor sample III. The effect of coalescence is thus theerosion of the tops of the initial pyramids, thedisappearance of the pyramidal faces and the cre-ation of faces parallel to the substrate. These phe-nomena can be explained by surface migration ofatoms towards island junctions which constitutesites of high adsorption energy. Then the overallsurface energy is reduced.

The plane view of a continuous film is presentedin Fig. 3a. The film has a mosaic structure withgrain size ranging between 50 and 200 nm in lateralsize. The grains are slightly misoriented with re-spect to each other in the (0 0 0 1) plane, and theirmisorientations are accommodated by subgrainboundaries. Moire fringes (MF), whose orientationchanges strongly as a function of small misorienta-tions of the island-substrate system, give an accu-rate measurement of these misorientations. In thepresent case, MF misorientation associated withdifferent grains is found to be 10° on average, whichcorresponds to a mean value of misorientation be-tween island lattices of 0.5°. The mosaic texture ofthe continuous films is inherited from the earlystage of growth as shown from the features found insample I. The MF indicate that the initial islandsare misoriented with respect to each other by a sim-ilar value as in the continuous films (Fig. 3b). Theformation of continuous films may thus be ex-plained by the growth of individual islands initially

F. Veignant et al. / Journal of Crystal Growth 196 (1999) 141–150 143

misoriented from each other and conserving theirown orientation until the coalescence stage. At co-alescence, the initial misorientation is accommod-ated by the formation of subgrain boundaries at theisland junctions. Furthermore, although a largepart of initial islands are free from defects, some ofthem already contain subgrain boundaries thatprobably remain during the growth and contributeto the texture of the continuous films (Fig. 3c).

The mosaic structure existing in lithium niobatefilms on sapphire has been related to the largelattice mismatch between the two materials by De-rouin et al. [7]. Furthermore, some defects haveprobably been induced by the distortions at thesubstrate surface: indeed, sapphire substrates areperfectly crystalline in the bulk, as shown by X-raydouble diffraction experiments [11] but they maybe superficially damaged by mechanical polishingup to depth of about 1 lm from the surface.

3.2. Crystallographic variants

X-ray /-scans previously performed have shownthat two crystallographic variants rotated fromeach other by 60° around the c-axis coexist inthe film [10]. The epitaxial relationships are(0 0 0 1)

&E(0 0 0 1)

4and [1 1 21 0]

&E[1 1 21 0]

4for the

first variant (now labelled A) and (0 0 0 1)&E(0 0 0 1)

4and [11 11 2 0]

&E[1 1 21 0]

4for the second one (now

labelled B), where s and f subscripts designate,respectively, the substrate and the film. The exist-ence of these two variants has been observed anddiscussed by several authors [7,13]. In our case, itshould be noted that the B-variant is not observedin all of the specimens, although they were all grownin the same conditions. We do not know at thismoment the reason of this behaviour. Our purposein this section is to confirm that the B-variantappears at the nucleation stage and to specify themorphology of B-grains in continuous films.

The variants cannot be discriminated in the[0 0 0 1] orientation since the diffraction patternhas sixfold symmetry. Therefore, the specimen hasto be tilted into an orientation giving rise to distinct

§——————————————————————Fig. 2. Dark-field micrograph in weak beam conditions usingthe 0 1 11 2 reflection showing thickness fringes of (a) sample I, (b)sample II, (c) sample III, in plane view.


Fig. 3. Bright-field micrograph of (a) continuous film, (b) twoisolated islands misoriented with respect to each other, in planeview orientation along the [0 0 0 1] zone axis, (c) isolated islandpresenting a subgrain boundary.

Fig. 4. (a) Selected area diffraction pattern along the [2 41 2 3]zone axis showing the diffraction spots of the two variants.(b)—(d) Dark-field micrograph performed with 0 1 11 2 B reflec-tion showing the B variant (in bright contrast) in: (b) isolatedisland sample in plane view; note the B island is rotated by 60°from the A island, (c) continuous 110 nm thick film in planeview, (d) continuous 110 nm thick film in cross section along the[2 11 11 0] zone axis.

diffraction patterns for each variant. This is realisedin the S2 41 2 3T zone-axis orientation, after tiltingby 33° around the a-axis (Fig. 4a). In this orienta-tion, A-diffraction spots are easily recognisable be-cause they are coupled with sapphire spots of thesame indices, whereas the B-spots are isolated.Dark-field images have been performed with the0 1 11 2

Bdiffracted beam of the B-variant in the

isolated islands sample (Fig. 4b) and in the continu-ous 110 nm thick film (Fig. 4c). It can be seen inFig. 4b that the B-variant already exists as isolatedislands, and that they have been formed during thenucleation stage. In the continuous 110 nm thickfilm (Fig. 4c), the B-grains do not have any particu-lar shape and have sizes ranging between 50 and200 nm. Experimental images being a projection ofthe three-dimensional structure, the volume frac-tion of B-variant in the film can be evaluated bymeasuring the surface ratio of B-grains in the planeview micrographs, provided that the following con-ditions are fulfilled: (1) the film thickness is constant


and well known and (2) all B-grains are constitutedof single crystals extending from the bottom to thefilm surface and limited by boundaries normal tothe surface. Assuming these conditions fulfilled andintroducing a correcting term for taking into ac-count the tilt angle of the specimen, a ratio of 15%of the B-variant was found. This value is somewhathigher than the value found by measuring the rela-tive intensity of diffracted beams in the X-ray /-scan (10%). Cross-sections were therefore investi-gated for determining the exact morphology of theB-variant. They were observed in dark field condi-tions with the 0 1 11 2

Bdiffracted beam (Fig. 4d).

The micrographs show that part of the boundariesbetween A- and B-grains are straight and extendfrom the surface to the interface. However, someboundaries are inclined and some B-grains do notreach the surface. These features may explain theoverestimation of the B-volume ratio by CTEMmeasurements.

3.3. Stress-induced defects

The Al2O

3/LiNbO

3system is expected to store

considerable stress because there is a large latticemismatch and a large difference between the ther-mal expansion coefficients of the two materials.The lattice parameter is a

&"5.22 A_ [14] and

a4"4.79 A_ [15], respectively, for lithium niobate

and sapphire at the deposition temperature. Themisfit parameter, defined as d"(2(a

&!a

4))/

(a4#a

&), is then 8.6% at ¹

$instead of 7.7% at

room temperature. HRTEM images of a cross-sectional specimen viewed along the S1 0 11 0T di-rection show misfit dislocations located at the in-terface (Fig. 5). Their Burgers vector is not welldefined. The dislocation core extension is aboutthree or four planes, whereas the continuity be-tween the substrate and the film planes is perfectelsewhere. This result agrees with the model ofa “pseudo semi coherent” interface, describing thecase of a misfit parameter greater than 5% [16] andused for lithium niobate on sapphire by Terabe etal. [17]. The misfit dislocations have in this case thesame spacing D as the Moire fringes, i.e. D"

(d4d&)/(d

&!d

4), where d

&and d

4are the spacing of

lattice planes for the film and the substrate, respec-tively. In our case, the lattice planes imaged

Fig. 5. HRTEM micrograh of a cross sectional specimen in the[1 11 0 0] zone axis orientation. Misfit dislocations are indicatedwith black arrows.

in the micrograph are the M1 1 21 0N planes withd4"1

2a4and d

&"1

2a&. The spacing expected is then

31.5 A_ for a fully relaxed film at room temperature.The spacing measured is 29$1 A_ . It is close to28.6 A_ , which is the theoretical spacing expected forthe accommodation of the lattice mismatch at thegrowth temperature (¹

$). This result confirms that

the film was relaxed at ¹$, and has acquired tensile

stresses during the cooling stage from ¹$

to roomtemperature. The elastic energy stored in the film isproportional to its thickness and the stress effectsare expected to be thickness dependent. A thicknessdependent stress effect-has been shown by X-raymeasurements at room temperature. The c latticeparameter of films thinner than 180 nm is a littlesmaller than in bulk crystals, i.e. c"13.81 A_ in-stead of 13.86 A_ . This means that these films are incompression along the c-axis and consequently inextension in the (0 0 0 1) plane [3,9,11]. Thickerfilms are fully relaxed by the formation of cracks ofseveral tens of nanometers seen by optical micros-copy [11]. TEM studies show different kinds ofthickness dependent defects. Very thin and discon-tinuous films do not show any defects other thanelastic deformation. Continuous 110 and 220 nmthick films show two types of stress-induced defectsin addition to elastic deformation: cracks and twin-ned lamellae. Cracks observed in the 110 nm thickfilm in plane view have a mean width of 10 nm,which explains why they have not been seen byoptical microscopy (Fig. 6a). They appear to liepreferentially along grain boundaries, which ex-plains their winding. However, when they arecrossing a single-crystalline area, they are orientedalong the S1 0 11 0T direction. In these films,


Fig. 6. Bright-field micrograph of (a) 110 nm thick film, (b)220 nm thick film, in plane view along the [0 0 0 1] zone axis,showing cracks.

nanocrack formation is not sufficient to relaxthe elastic deformation. As mentioned before, thelattice parameters remain different from those ina bulk crystal. The elastic deformation is fully re-laxed in the 220 nm thick films (Fig. 6b). In thethicker films, the cracks have a mean width of60 nm and are clearly oriented along the S1 0 11 0T

Fig. 7. HRTEM micrograph of a cross-sectional specimen inthe [2 11 11 0] zone axis orientation showing the twin matrixinterface.

direction. The abrupt stopping of the moire fringesat the cracks-film limit indicates that the crackwalls are perpendicular to the substrate.

Concerning twin lamellae, their orientation rela-tionship relative to the matrix film has been deter-mined by examining cross-sections along the[2 11 11 0] direction. HRTEM images show the twin-matrix interface with (0 1 11 2) as the compositionplane (Fig. 7). Plane view observations give furtherinformation on twin morphologies. The imageswere performed after a small tilt (3°) around theS2 11 11 0T direction. In this condition the M0 11 1 4Nplanes come into diffraction position for two twinsystems and permit us to observe these two systemssimultaneously (see stereogram in Fig. 8a andFig. 8b). Twins in the 110 nm thick film appear aslamellae having average length of 300 nm, andaverage width of 60 nm (Fig. 8c). It is important tonote that twins are always located at the border ofthe cracks.

The M0 1 11 2N twin system has already been ob-served in bulk single crystals [18]. It is a conse-quence of stress in the crystal, i.e. stress during thegrowth process in the Czochralsky method [19] orstress applied in deformation experiments per-formed at temperatures (1000°C [20]. Our obser-vations indicate that the film is relaxed during thegrowth. So twinning most likely arises during thecooling stage. Concerning the relation betweencracks and twins, the junction between twin andmatrix as well as the junctions between two twinsare expected to store large stress and therefore to be


Fig. 8. (a) Stereographic projection along the [0 0 0 1] matrix direction showing the M0 1 11 2N and M0 11 1 4N poles of the matrix and thethree equivalent twin systems. (b)—(d): 110 nm thick film in plane view: (b) selected area diffraction pattern showing the 0 11 1 4 reflectionsof the two equivalent twin systems, I and II; (c) dark-field micrograph performed with the 0 11 1 4 I reflection (lamellae in bright contrastare type I twins). (Note that twins are in border of cracks.) (d) bright-field micrograph showing the intersection of two twinned lamellaewhich give rise to a crack.

preferential sites for cracks initiation. It has thusbeen demonstrated that the intersection of twoM0 1 11 2N twins leads to the formation of a crack[20,21]. To a lesser degree, grain boundaries arealso stressed areas, but it remains unclear if they aresites for crack initiation or favourable sites forcrack propagation.

4. Conclusions

Transmission electron microscopy observationsof LiNbO

3films on aAl

2O

3(0 0 0 1) single crystals

corroborate the existence of close epitaxial rela-tionships between film and substrate, previouslyfound by X-ray diffraction. At the electron micro-scopy scale, the lithium niobate films show struc-tural imperfections that may be classified into twocategories, according to their origin. (1) The filmspresent a mosaic-type texture, defined by grainswhose size ranges between 50 and 200 nm, mis-oriented from each other by 0.5° on average. Thesubgrain misorientation is accommodated by a net-work of subgrain boundaries. It has been shownthat the film texture has been formed at the veryearly stages of the film deposition. (2) There are also


mechanical defects induced by thermal stress oc-curring during the cooling from deposition to roomtemperature. The lattice mismatch betweenLiNbO

3and Al

2O

3is accommodated by misfit

dislocation at the deposition temperature,¹

$"750°C, but tensile stresses appear during the

cooling process because the two materials havedifferent thermal coefficients. These stresses are re-sponsible for the formation of twin lamellae insmall amounts (1% in volume) and nanocrackswhose width, w, depends on the film thickness, t (w:10 and 60 nm, respectively, for t"110 and220 nm). Although directly caused by different phe-nomena, the two classes of defects may interact.Subgrain boundaries present at the depositionstage constitute zones of stress concentration inaddition to the thermal stresses appearing at thecooling stage. Thus, they may contribute to theovercoming of the film elastic limit giving rise tonanocrack formation.

The influence of each kind of defect on the op-tical properties of the lithium niobate films is notcompletely elucidated. It is only known that opticalwaveguiding has not been obtained in thicker films(220 nm) which contain large nanocracks (60 nm).On the contrary, optical waveguiding has beenobtained in thinner films (110 nm) containing grainboundaries, twins and narrow nanocracks (10 nm).However, optical losses have been found in thesefilms. Each kind of the above mentioned defectsmay contribute to these losses, together with sur-face and interface roughness. The contribution oftwin lamellae is probably insignificant since theyare present in small amounts, but grain boundariesand the narrow nanocracks uniformly distributedthrough out the films may have a larger effect.Attempts will be made in the future to reduce thecreation of defects and in this way to understandtheir influence on the optical losses.

Acknowledgements

HRTEM has been performed at the Centred’Etudes de Chimie Metallurgique (CECM/CNRS), Vitry sur Seine, France. The authors wouldlike to thank its director J.P. Chevalier and theelectron microscopy staff for their hospitality. We

are also grateful to Dr Beaunier from the Labora-toire de Physique des Liquides et Electrochimie forproviding us the possibility of using the GATANPIPS for the ion milling of our samples.

Appendix A

The determination of the island morphology hasbeen made in the frame of the dynamical theory ofthe electron diffraction. In the two beam condition(one transmitted beam and one diffracted beam g),the beam intensity oscillates with depth in the crys-tal. The period of the oscillation is given by therelation [22]

m'%&&

"

1

(s2#1/m2')1@2

,

where s is the deviation parameter to Bragg condi-tion. m

'"p»

#/jF

'is the extinction distance for the

beam used. »#, j, F

'are, respectively, the cell vol-

ume, the beam wavelength and the structure factor.Thickness fringes are produced by wedge-shaped

crystals. The islands height has been found bycounting these fringes. The angle / between theface of the pyramids and the (0 0 0 1) plane has beendeduced by the simple geometric relationshiptan /"m

'%&&/i, where i is the mean distance be-

tween the fringes in the image plane.Because of the uncertainty in the s measurement,

several experiments, using different diffracted be-ams and different orientations of the sample, havebeen performed for an accurate determination ofthe morphology. The dark field images shown inthis paper have been performed with 1 1 21 0 beam,whose extinction distance is 85 nm. The parameters has been determined with the help of the Kikushilines of the substrate, taking into account the differ-ence between the film and the substrate. The valuesof s are 0.10 nm~1 for sample I, and 0.16 nm~1 forsamples II and III. In sample I, i has been found tobe 6.5$0.5 nm, which leads to an angle of 57$3°.This value nearly corresponds to the angle of 57.2°between the M0 1 11 2N faces and the (0 0 0 1) plane.

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structural evolution of lithium niobate deposited on sapphire (0 0 0 1): from early islands to...

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