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This article appeared in a journal published by Elsevier. The attachedcopy is furnished to the author for internal non-commercial researchand education use, including for instruction at the authors institution
and sharing with colleagues.
Other uses, including reproduction and distribution, or selling orlicensing copies, or posting to personal, institutional or third party
websites are prohibited.
In most cases authors are permitted to post their version of thearticle (e.g. in Word or Tex form) to their personal website orinstitutional repository. Authors requiring further information
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Microstructure and elastic properties of atomic layer depositedTiO2 anatase thin films
L. Borgese a, E. Bontempi a,⇑, M. Gelfi a, L.E. Depero a, P. Goudeau b, G. Geandier c,D. Thiaudiere d
a INSTM and Dipartimento di Ingegneria Meccanica ed Industriale, Universita di Brescia, via Branze 38, 25123 Brescia, Italyb Institut Prime, UPR 3346, CNRS, Universite de Poitiers, ENSMA, SP2MI, Bd Marie et Pierre Curie, BP 30179, 86962 Chasseneuil,
Futuroscope Cedex, Francec LETAM, CNRS FRE 3143, Universite de Metz, ISGMP Bat. B, Ile du Saulcy, F-57012 Metz Cedex, France
d Synchroton Soleil, Saint-Aubin, BP48, 91192 Gif-sur-Yvette Cedex, France
Received 13 July 2010; received in revised form 19 January 2011; accepted 19 January 2011
Abstract
Amorphous TiO2 thin films were deposited by means of atomic layer deposition on Kapton substrates and then crystallized ex situ byannealing at 300 �C to obtain the anatase phase. The morphology, structure and microstructure of films treated for 12, 24, 72 and 90 hwere investigated. The local Ti coordination changes were studied by X-ray near-edge structure (XANES).
On the basis of X-ray diffraction residual stress calculations, the elastic anisotropy of the films is experimentally determined for thefirst time (A�comp ¼ 0:07, A�shear ¼ 0:03). The film macro-strains increased with the time of treatment, while the micro-strains decreased.This effect may be correlated with the incipient anatase-to-rutile transformation as suggested by the changes observed in the XANESpattern of the film treated for 90 h. However, the contribution of the substrate cannot be excluded.� 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Atomic layer deposition (ALD) facilitates the depositionof ultrathin films at the nanometre scale with extremelywell-defined thicknesses and compositions as well as excel-lent conformity [1]. For chemical vapour deposition, theprecursors and the deposition temperature significantlyaffect the microstructure and surface properties of thefilms. Usually, a high adhesion of the deposited thin filmis obtained owing to the formation of strong chemicalbonds with the substrate [2]. Moreover, since films aregrown in the equilibrium regime, residual stress should beabsent or very low. ALD functional oxides (such as ZnO,ZrO, TiO2, Al2O3) are widely used in microelectronics;
recently, they have been studied extensively as biocompat-ible coatings. However, the mechanical behaviour of thefilms is critical for various commercial applications [3,4].The accurate evaluation of the mechanical properties andresidual stress of the film are then mandatory, even if goodmechanical performance of ALD deposited thin films isexpected.
Today, titanium dioxide is one of the most importantphotocatalyst materials, having a wide range of applica-tions (for an extensive review, see Ref. [5]) because it exhib-its high durability, good corrosion resistance and theability to activate different reactions. It is also biocompat-ible, and therefore it is widely employed in biomedicalimplants [6]. Furthermore, polymer–inorganic hybridmaterials containing TiO2 are major issues owing to theirdiphasic structure, which promotes new multifunctionalmaterials [7].
1359-6454/$36.00 � 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
doi:10.1016/j.actamat.2011.01.032
⇑ Corresponding author.E-mail address: [email protected] (E. Bontempi).
www.elsevier.com/locate/actamat
Available online at www.sciencedirect.com
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The functional properties of titanium oxide are stronglydependent on the phase and, in particular, on the micro-structure, which plays a fundamental role. Indeed, the crys-tallite size is critical for phase stability [8], and the presenceof micro- and/or macro-strains may affect the photoin-duced hydrophilicity [9].
There are several natural polymorphs of titania: rutile isthe most stable form of TiO2, whereas anatase and brook-ite are metastable and transform to the rutile phase onheating (see for example Ref. [10] and references therein).The temperature of anatase-to-rutile phase transition isusually <600 �C, depending on different factors such asdefects, impurities and grain sizes. The relationshipbetween the microstructure and elastic properties for therutile phase has been widely discussed [11–13], whereasfew studies have been conducted for the anatase phase;i.e., only the bulk modulus B of anatase at high pressurehas been determined experimentally [8], while the Cij elasticstiffnesses have only been determined by numerical simula-tions [14–16].
X-ray diffraction (XRD) is generally used for the struc-tural analysis of films and surface layers. In addition to thephase determination, XRD is a crucial tool for microstruc-tural analysis. In particular, the “sin2 w method” is popularfor the determination of the stress field in bulk [17] and thinfilms [18]. This analysis is phase selective rather thandestructive, and it does not influence the microstructureand stress yield. To minimize experimental errors, high-2h peaks are usually chosen. However, for thin film, thescattering signal is too low to allow an accurate determina-tion of peak positions by means of laboratory equipment.The synchrotron X-ray source reduces or eliminates manyproblems resulting from low intensity and instrumentalbroadening, and thus it is more suitable for studying verythin films.
This paper reports and discusses the structural andmicrostructural properties of TiO2 anatase thin filmsdeposited onto organic substrates by means of a low-tem-perature ALD technique. Preliminary XRD analyses withlaboratory equipment [19] showed that, on annealing at300 �C from the amorphous as-deposited phase, a non-ori-ented anatase phase crystallized. The film characterizationand the analysis of changes in the morphology and micro-structure were performed by scanning electron microscopy(SEM), atomic force microscopy (AFM), XRD and X-raynear-edge structure (XANES). The experiments performedat Soleil Synchrotron facilitate the evaluation of changes inthe Ti coordination and elastic anisotropic behaviour ofthe film.
2. Experimental
ALD was performed using the Cambridge NanotechInc., Savannah 100 system. The reactor is a stainless steelcylinder (18.1 cm diameter and 3.6 cm high) with a bot-tom-heated plate 179 cm2. The precursors are tetra-kis(dimethylamido)titanium(IV) (TDMAT) and milliQ
water (H2O). TDMAT (99.999%) was purchased fromSigma Aldrich Chemical Co. (Germany) and used withoutany further purification. The Millipore DirectQ-5 purifica-tion system was used to obtain milliQ water. The precur-sors were injected into the reactor. The TDMAT andwater stainless steel reservoirs were kept at 80 �C and roomtemperature, respectively. Nitrogen (99.999% purity) wasused as a carrier gas. The deposition temperature was keptat 90 �C and the base pressure of the reactor at 0.5 torr. Toavoid precursor condensation, valves and delivery lineswere kept at 120 �C. The carrier gas and reaction productswere evacuated by a rotary vane vacuum pump to keep thepressure in the system. The processing cycle consisted of a0.1 s TDMAT pulse, 10 s for purging, 0.1 s H2O pulse and10 s for purging. The deposition rate of 0.0667 nm per cyclewas evaluated by means of X-ray reflectivity (data arereported elsewhere [19]). TiO2 thin films 200 nm thick weredeposited onto Kapton HN substrates 125 lm thick. Thesamples were annealed at 300 �C for 12 (sample A), 24(sample B), 72 (sample C) and 90 h (sample D) (see Table1). This temperature is below the Kapton Tg; it wasreported that these treatments do not change the substratestructure [19]. At least three specimens were deposited foreach treatment, which confirmed the reported results.
The surface morphology of the samples was examinedby a conventional optical microscope, a LEO EVO40XVP scanning electron microscope, equipped with aLink Analytical probe for energy dispersive X-ray spectros-copy and a Jeol 4210/TM-4210BU atomic force micro-scope. The root mean square roughness was calculatedon an area 1 � 1 lm2.
For structural and microstructural characterization,XRD and XANES measurements were performed at SoleilSynchrotron on the Diffabs beamline [20,21]. Diffractionpatterns were acquired in reflection geometry. The selectedenergy was 8 keV (k = 1.54 A) with beam size 0.3 � 1 mm(vertical � horizontal). A scintillation detector was usedto scan the diffraction angle range between 24� and 90�.The diagrams acquired were processed by commercial soft-ware to accurately determine diffraction peak position,intensity, area and full width at half maximum (FWHM).
Two-dimensional X-ray diffraction (XRD2) images werecollected in the laboratory on a D/max-RAPID Rigakumicro-diffractometer with Cu Ka radiation. This systemwas equipped with a cylindrical imaging plate detectorwhich can record XRD2 data in angular diffraction rangesfrom �45� to 160� (2h) horizontally and ±45� (2h) verti-cally with respect to the direct beam position (0�). The irra-diated area (chosen using a collimator) was 300 lm.
Table 1Samples description.
Sample name Annhealing time (h) at 300 �C
A 12B 24C 72D 90
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3. Results and discussion
3.1. Morphology, structure and microstructure
Optical microscope images of samples A and D areshown in Fig. 1. The morphology of samples B and C issimilar to the morphology of sample A. Only sample D(annealed for 90 h) shows a cracked network surface.
To evaluate the film adhesion, the scotch test was per-formed. In the SEM image of sample D (Fig. 2a), the pres-ence of cracks is evident. However, after the test (Fig. 2b),the delamination degree of the film is low and demonstratesgood adhesion to the substrate.
AFM images of samples A and D (20 � 20 lm2 area)are reported in Fig. 3. A non-cracked area was chosen toevaluate the roughness. In addition to the cracks, the mor-phology of samples A and D is similar and independentfrom the thermal treatment. In particular, the averageroughness values (calculated on a 10 � 10 lm2 area) are10 nm for sample A and 13 nm for sample D.
On the basis of these results, TiO2 thin films have goodadhesion, which can be attributed to the formation ofstrong chemical bonds with the substrate.
XRD2 analyses of the four samples were performed, andthe corresponding images are reported in Fig. 4. All Debyerings belong to the anatase phase [22], and the films arepolycrystalline. The large diffused halos near the centre ofthe image are due to the Kapton substrate. Fig. 5 showsthe XRD patterns of samples A and C collected at the syn-chrotron. Large halos at 27�, 37� and 42� were produced bythe Kapton substrate [23]. Reflections of the anatase phaseare reported (see Refs. [22,23]). The inset highlights thehigh angle diffraction peaks, which can be clearly identifiedin spite of the low thickness of the film caused by the highintensity of the synchrotron source. The refined cell param-eters for sample A do not differ from the tabulated valuesof polycrystalline anatase (symmetry: I41/amd and unit celllattice parameters a = 0.3791 (1) and c = 0.9515 (1) nm)[22,8]. The positions of the peaks are shifted in samplethe C pattern, because of the stress presented. From datashown in Fig. 4c and d, an accurate evaluation of cellparameters is not meaningful, because the integration of
the pattern over the Debye rings averages the peak positionin different directions, and the presence of residual stresswill give large errors, different for each reflection.
Peak broadening comes from several sources: namely,instrumental effects, coherently diffracted domain sizes(<100–500 nm) and micro-strain.
Strain is defined as the deformation of an object dividedby the ideal length, Dd/d. Two types of strain can distin-guished: (i) uniform strain (macro-strains), which causes
Fig. 1. Optical microscopy images of samples: (a) A and (b) D.
Fig. 2. SEM image of sample D: (a) before and (b) after the scotch test,showing the presence of cracks in the film and its low delamination.
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Fig. 3. AFM images of samples: (a) A and (b) D performed in a 20 � 20 lm2 area.
Fig. 4. XRD2 measurements of sample A, B, C and D.
Fig. 5. XRD pattern of samples: (a) A and (b) C. The inset highlights the high angle diffraction peaks of the TiO2 crystalline phase.
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the unit cell to expand/contract, thus leading to a change inunit cell parameters and a shift in the peaks; there is nobroadening associated with this type of strain; (ii) non-uniform strain (micro-strains), which leads to systematicshifts of atoms, i.e., distortions of the crystal structure.Micro-strains arise from different sources: point defects(vacancies, site-disorder), plastic deformation (cold workedmetals, thin films) or poor crystallinity.
Size and micro-strain information can be extracted byanalysing peaks using different methods and often subtract-ing instrumental effects that are negligible (in the case ofthe synchrotron source).
To evaluate the crystalline size change and micro-strains, the diffraction patterns of samples A and C werecollected from 24� to 90� 2h. In Fig. 6a, the FWHM ofthe anatase reflection as a function of 2h are compared.
When micro-strains can be neglected, the coherently dif-fracted domain size can simply be evaluated by the Sherrerequation [24]. In Fig. 6a, the Sherrer equation for sphericalcoherently diffracted domains of 60 nm is drawn as adashed line. Note that the FWHM of sample C reflections
roughly agree with this equation, while data for sample Aexhibit different behaviour, which suggests the presenceof significant micro-strains and/or structural defects.
Williamson and Hall proposed a method for deconvo-luting size and strain broadening by looking at the peakwidth as a function of 2h [24]. If a linear fit is obtained fromthe equation
b ¼ 1
Dþ 4e � sin h
k
it is possible to derive the crystallite size D from the y inter-cept, and the micro-strain (e) from the slope. The dimensionof coherently diffracted domains of sample D obtained bythe Williamson–Hall equation is �70 nm, which is quiteconsistent with the value obtained by the Sherrer equationin the hypothesis of spherical coherently diffracted domains.
Williamson–Hall plots of samples A and C are shown inFig. 6b. The linear fit slopes of the two samples are signif-icantly different, but the intercept is similar, which suggeststhat the average crystallite dimension for the two samples isabout the same, while the micro-strain distribution differs.
0
0.05
0.1
0.15
0.2
0.25
0.3
0.35
0.4
2θ
FWH
M
Sherrersample Asample C
0
0.0005
0.001
0.0015
0.002
0.0025
0.003
0.0035
0.004
0.0045
0.005
10 20 30 40 50 60 70 80 90
0 0.1 0.2 0.3 0.4 0.5 0.6 0.7
sample C
sample A
sin(θ)
βco
s(θ)
(a)
(b)
Fig. 6. (a) Evolution of the FWHM for samples A and C, as a function of 2h. The particles size contribution (bD) to the broadening for a peak (using theSherrer equation [22]) is also reported (as a dashed line). (b) b*cos(h) as a function of the sin h plot for samples A and C. From the intercept of the graph, itis possible to approximate the crystallite size dimension. From the slope, it is possible to approximate the micro-strains (often referred to as a Williamson–Hall plot).
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Considering all the reflections of the anatase phase for aqualitative calculation, the average micro-strain value forsample A is �0.13% and �0.07% for sample C, which sug-gests that micro-strains decrease when the thermal anneal-ing time is increased.
3.2. Local atomic structure
XANES spectroscopy is a fruitful tool for providinginformation about the local atomic structure. Fig. 7a showsXANES spectra collected at the Ti–K absorption edge onsamples A, C and D. Spectra of sample B are similar tospectra of sample A, which correspond to pure anatasespectra [25]. Spectra of sample C show an increase in thetypical XANES structure of anatase (marked B) [25], whichsuggests very good crystallization.
In XANES spectra of sample D (Fig. 7b), distinguishingfeatures characteristic of the rutile phase appear (markedwith B2 and D) [25]. Thus, XANES spectra reveal localchanges in the structure which might be related to anembryonic anatase–rutile phase transition. As discussedbelow, this change can also be related to residual compres-sive stress present in samples C and D and to the formationof cracks in sample D [26].
3.3. Elastic anisotropy of the anatase phase and residual
stresses
Macro-strains (e) were studied by the “sin2 w method”
[17,27,28]. Fig. 8 shows e as a function of sin2 w for two dif-fracting planes: namely, (0 0 4) and (2 0 0). The measure-ments reveal the presence of significant residual stresses
Fig. 7. (a) Ti K-edge XANES spectra of samples A, C and D; (b) the differences between spectra C and A (C–A) as well as D and A (D–A) are reported.
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(macro-stresses) only for samples C and D (after 72 and90 h of thermal treatment). In particular, the slope value(and thus the stress value) is lower for sample D, whichwas annealed for the longest period.
The residual stress values can be calculated by assuminglinear elastic behaviour (Hooke’s law). From the slope ofthe linear fitting, the residual strain in all samples can beextracted. Apparently, samples A and B do not show resid-ual stresses, while samples C and D have compressive resid-ual stresses; samples A and B have the highest value, whichcan be related to surface relaxation effects in sample Dowing to crack formation, as revealed by optical micros-copy, SEM and AFM.
To quantify the residual stress, elastic constants of thephase are needed. However, for the anatase phase, onlytheoretical elastic constant values are available in the liter-ature [14–16]. The values reported in Refs. [14] and [16]were used to calculate the residual stress using differentcrystallographic planes.
The elastic constants (stiffness and compliances) of ana-tase are reported in Table 2 (in GPa). The last two columnsshow the calculated bulk and shear modulus, which are theaverage of Reuss and Voigt values [29]. A combination ofReuss and Voigt simple models gives a good description ofX-ray elastic constants. Anisotropy constants are alsoreported in the last two columns of Table 2. It is possibleto experimentally verify the anisotropy of anatase bycalculating the elastic modulus along different crystallo-graphic planes. The residual stress can be calculated usingdiffraction average elastic compliances, as calculated byelastic constants (see Appendix) and experimental strainvalues.
Within the Reuss approximation, diffraction averagecompliances in the tetragonal system are:
Fig. 8. Evolution of the intra-granular macro-strains e for (a) (0 0 4) and(b) (2 0 0) diffracting planes as a function of sin2 w for all the samples. Thedashed lines represent the linear regression plotted for samples C and D. A(}), B (h), C (M) and D (�).
Table 2The left first line of the table gives the elastic stiffnesses and the left second line the elastic compliances. The last two columns on the right show thecalculated bulk and shear modulus as the arithmetic average of Reuss and Voigt.
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The diffraction average compliances hs011i and hs013i are iden-tified by the “X-ray elastic constants” through the following
relations: ð1þmÞE
h i¼ hs011ihkl � hs013ihkl and �m
E
� �¼ hs013ihkl.
Calculated X-ray elastic constants for different crystallo-graphic directions of anatase are reported in Table 3.Young’s modulus along the [1 0 0] direction is E100 =(S11) � 1 = 195 GPa while along the [0 0 1] direction it isE001 = (S33) � 1 = 93 GPa, which shows strong anisotropy.
Using these calculated X-ray elastic constants, residualstress values were evaluated for different (h k l) planes.The residual stress is expected to be independent fromthe particular family of planes. Using data from Ref. [14]for sample C (where the residual stress is the highest), avalue of �0.7 GPa in compression was calculated for any(h k l) plane considered. For (001), a significantly lowervalue is obtained (0.59 GPa) (see Table 3). Thus, theoreti-cal elastic constants seem to be slightly different from theexperimental constants. In particular, setting C33 =208 GPa as the stress values in all directions produces verysimilar results (see Table 3). The bulk modulus B evaluatedby C33 = 190 GPa is equal to 173 GPa while, if C33 = 208,the GPa is 178. Since the film texture is isotropic as well asthe grain morphology, large differences are not expectedusing approaches which take into account the inter-granular elastic interaction and thin film microstructure[30]. Indeed, the experimental bulk modulus B of anatase(178 GPa [20]) confirms the calculation, and thus the exper-imental C33 seems to have a higher value with respect to thevalue found in the theoretical calculation [14].
For sample D, the residual stress value is ��0.5 GPa.For sample C, the same value for all crystallographic direc-tions is obtained by considering C33 = 208 GPa. The lowervalue of the residual stress can be correlated to the forma-tion of the cracking network and the surface relaxationafter 90 h of treatment.
The anisotropy constants evaluated on the basis of thisdiscussion are A�comp = 0.07 and A�shear = 0.03, which areslightly smaller than the theoretical anisotropy constants[14].
3.4. Origins of residual stresses and micro-strains
The compressive residual stress appears in the films after72 h of thermal treatment, and thus is not due to thedeposition process [31,32]. The micro-strain as calculatedby the Williamson–Hall plot decreases with thermal treat-ment (see above). Moreover, after 72 h (sample C), the filmhas a compressive residual stress, while the micro-strainsrelax.
It is suggested that the macro- and micro-stress in thesamples are related to the transformation of anatase-to-rutile. Bulk anatase transforms to rutile at temperaturesbetween 400 and 1200 �C [33], depending on different fac-tors such as grain size, impurities and synthesis technique[34–36]. The strain energy of this transformation is quitelarge, since the molar volume of rutile is �10% lower thanthat the molar volume of anatase [37,8]. One of the factorsinducing the anatase-to-rutile transformation is the pres-ence of large micro-strains [38] as in sample A, in whichthe anatase phase is formed from the amorphous as-deposited film. The origin of these micro-strains can beoxygen deficiency during deposition, which also is claimedto favour the formation of the anatase phase [39].
These considerations suggest that compressive residualstress and microstructure are related to local changes thatwill drive the transformation to rutile. This hypothesis isalso consistent with XANES data, as discussed above.
Residual stress may also be related to different thermalexpansion coefficients of the anatase film and Kapton sub-strate. However, at a fixed temperature (300 �C in the pres-ent case), one should expect the macro-strain due to thisdifference to be independent of the annealing time. In thepresent experiment, the macro-strain value depends onthe annealing time: after 72 h (sample C), the maximummacro-stress that relaxes after 90 h of treatment (sampleD) because of surface cracking was evaluated.
It has been reported that, when the temperature is closeto the glass transition temperature (second-order transitionTg between 360 �C and 410 �C [40]), the elastic propertiesof the Kapton increase during thermal annealing [40] andthus give rise to macro-stress in the film. Thus, one cannotrule out that the substrate changes during treatment mayalso influence the residual stress of the film.
4. Concluding remarks
In this paper, the structural and microstructural proper-ties of TiO2 anatase thin films deposited by the ALD tech-nique and annealed at 300 �C are reported and discussed.
The high adhesion of titania thin films is excellent. Inspite of the presence of compressive residual stress, the filmdoes not show any delamination. Crack networks appearafter 90 h of annealing time with the relaxation of residualstresses.
Residual stresses for different crystallographic directionswere calculated on the basis of theoretical elastic constants[14]. The anisotropy experimentally determined is A�comp ¼0:07 and A�shear ¼ 0:03, which is slightly smaller than the
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anisotropy exhibited by the ab initio calculation (A�comp ¼0:09, A�shear ¼ 0:04).
Micro-strains are present in the sample after 12 h ofannealing when the anatase phase is formed and arerelaxed after 72 h of thermal treatment at 300 �C. Con-versely, the residual stresses increased with annealing timeand relaxed after 90 h when cracks appeared.
The origin of changes in micro- and macro-stresses maybe related to the phase transition from anatase-to-rutile.This conclusion is supported by local changes in the Ticoordination, as revealed by XANES spectra.
Acknowledgements
The authors are grateful to Dr. Paolo Colombi for use-ful discussions about thin film morphology and to StefaniaFederici for AFM measurements. This project was partiallyfinanced by Galileo project 2007/2008 No. 29. The authorsmust also thank Soleil, the French synchrotron radiationfacility (Saint Aubain, France) for providing beam timethrough the program committee: Proposal No. 20080450“Mechanical behaviour of thin ceramic layers depositedby ALD”.
Appendix A
In tetragonal crystal, elastic anisotropy arises fromanisotropy of the linear compressibility in addition to theusual shear anisotropy [41].
A�compression ¼ðKV � KRÞðKV þ KRÞ
A�snear ¼ðGV � GRÞðGV þ GRÞ
(V refers to Voigt, and R refers to Reuss averaging sche-mes)where [42]
Cij are the elastic stiffness constants, and Sij are the elasticcompliances. Elastic compliances can be calculated by theinversion of the matrix of elastic stiffness. The equationsto transform Cij to Sij constants are given by Nye [43]and, for the tetragonal symmetry, they become [44]
S11 ¼ S22 ¼1
2
C33
C0þ 1
C11 � C12
� �
S12 ¼1
2
C33
C0� 1
C11 � C12
� �
S13 ¼ S23 ¼ �C13
C0
S33 ¼C11 þ C12
C0
S44 ¼ S55 ¼1
C44
S66 ¼1
C66
C0 ¼ C33ðC11 þ C12Þ � 2C212
A* is zero for an isotropic crystal. A* is always a positivequantity for an anisotropic crystal, and it is a measure ofthe relative magnitude of the elastic anisotropy.
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