Defects and Diffusion in Ceramicspearton.mse.ufl.edu/semic_properties/data/5037.pdf65 Defects and...

Defects and Diffusionin Ceramics

An Annual Retrospective I

Editor:

D.J. Fisher

SCITEC PUBLICATIONS

Notes:Each item in this section of the volume begins with a graphical compilation of relevant diffusion data whichhave been reported during the past decade. The plotted data are also tabulated as indicated on the graph. In somecases, the tabulated data have been obtained by digitizing published graphs and the values may not correspondexactly with the author's unpublished raw data.

3N Bulk Diffusion - Quantitative DataThe migration of Ag from epitaxial layers and into (111) samples of Si,during annealing at temperatures of between 450 and 500C, was studiedby means of secondary ion mass spectrometric depth profiling. It wasfound that the diffusivities lay between 8 x 10 -16 and 1.6 x 10-15cm2/s(table N). These values were lower than were expected on the basis ofprevious data.T.C.Nason, G.R.Yang, K.H.Park, T.M.Lu: Journal of Applied Physics,1991, 70[3], 1392-6

[446-91/92-027]

Refers to table N

Indicates volume and page number inDDF where abstract first appeared

65

Defects and Diffusion in CeramicsAn Annual Retrospective

Carbon and Carbides

C: Electron Irradiation, Point Defects, and Surface ReconstructionThe response of nanotubes to uniform atom loss was observed. The response was foundto involve surface reconstruction and marked dimensional changes. Experimentsperformed using electron irradiation led to nanotube diameters shrinking from some 1.4 to0.4nm. Molecular dynamics simulations showed that such surface reconstruction and sizereduction occurred via dangling-bond saturation, thus forming non-hexagonal rings and 5-7 defects in the lattice. Non-uniform atom removal resulted in inhomogeneous tubedeformation and local necking, and the formation of linear atomic C chains in thenanotube body.P.M.Ajayan, V.Ravikumar, J.C.Charlier: Physical Review Letters, 1998, 81[7], 1437-40

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C: Point DefectsTransmission electron microscopy was used to investigate in situ the structure of laser-ablated single-wall C nanotubes which had been deposited onto Ni grids. It was foundthat most of the nanotubes formed highly curved bundles, which split or joined to formvarious structures. Individually grown nanotubes were found to contain more defects.Y.Zhang, S.Iijima: Philosophical Magazine Letters, 1998, 78[2], 139-44

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C: Point DefectsThe electronic structure of amorphous diamond-like C was considered. The sp2 sitesformed small, mainly chain-like clusters, which controlled the band-gap. Detailedanalysis showed that all p states within the s -s * gap were localized, so that the mobilitygap greatly exceeded the optical gap; as reflected by the photoluminescence excitationspectrum. Dangling-bond states were predicted to possess some s-orbital character. Thislowered their energy level in the gap, and accounted for the p-type conduction of

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C Carbon and Carbides C60

tetrahedral amorphous C. The defect density was high in hydrogenated amorphous C andwas found to be quite well described by the weak-bond to dangling-bond conversionmodel. It was concluded that paramagnetic defects were the predominant recombinationcenter.J.Robertson: Philosophical Magazine B, 1997, 76[3], 335-50

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C60: Electron Irradiation and Point DefectsRate equations were proposed for the generation and recombination of vacancies and self-interstitial atoms in a spherical bucky onion under electron irradiation. These were usedto study the temporal development of the point defect concentrations in various onionsshells. The self-interstitial concentration was found to remain low, and to be only slightlyenriched near to the onion's center. The loss of atoms from the outer surface of the buckyonions, via sputtering, acted as a vacancy source for the whole onion. Irradiation-induceddiffusion led to a net migration of vacancies towards the onion core. The coupling,between neighboring shells, which was necessary for such radial diffusion could bedescribed by invoking an intermediate-shell concept.W.Sigle, P.Redlich: Philosophical Magazine Letters, 1997, 76[3], 125-32

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C60: DefectsThe defect structures of fullerenes and fullerene-like clusters, produced by two C60molecules colliding at various rates and directions (with energies ranging from 50 to600eV) were studied by using molecular dynamics simulations. The electronic structuresaround the defects were deduced by using a self-consistent field Hartree-Fock scheme.The results showed that the coordination-number defects obeyed a magic-number rule,and that the numbers of rings which formed the closed-cage structure of the fullerene orfullerene-like products could be closely described by a modified form of Euler's theorem.The electronic structures around the defects were substantially changed when comparedwith those of C atoms on normal fullerene cages.Y.Xia, Y.Mu, Y.Xing, C.Tan, L.Mei: Physical Review B, 1997, 56[8], 4979-86

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C60: DislocationsSingle and double spiral growth of C60 epitaxial films on a KBr (001) substrate wereobserved via atomic force microscopy. Single and double spirals were found on filmsgrown at different rates. All of the spiral islands exhibited the 3-fold symmetry of theface-centered cubic structure. In the case of single spirals, many fringes were observed in3 equivalent [11̄0] directions on the face-centered cubic (111) surface. On the basis of theatomic force microscopy images, it was concluded that the growth of these spirals couldbe explained in terms of the Burton-Cabrera-Frank crystal growth theory, and that screwdislocations formed without any direct influence of the substrate.

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C60 Carbon and Carbides C (Diamond)

Y.Kim, L.Jiang, T.Iyoda, K.Hashimoto, A.Fujishima: Applied Physics Letters, 1997,71[24], 3489-91

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C60: DislocationsChanges in dislocation mobility and micro-hardness, on the (111) faces of crystals whichwere exposed to the atmosphere, light irradiation and heat, were investigated. A markedincrease (up to 300%) in hardness was observed after irradiation-assisted aging in air. Itscause was suggested to be photo-oxidation of the near-surface (1 to 5µ) layer. Theintercalation of absorbed O during aging in air (in the dark) resulted in a slight (about11%) increase in the hardness, while the dislocation mobility in the oxygenated crystalswas reduced by a factor of 3. Abrupt softening of the photo-oxidized layer upon heatingto 310 or 480K was observed; thus indicating a 2-stage structural transformation. Ananomaly was found in the temperature-dependence of the hardness of pristine samples attemperatures above 470K.I.Manika, J.Maniks, J.Kalnacs: Philosophical Magazine Letters, 1998, 77[6], 321-6

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C60: DislocationsThe equilibrium molecular configuration around a ½[1̄01] screw dislocation in a face-centered cubic phase was calculated by using molecular dynamics techniques. Themolecules were assumed to be rigid, and the Girifalco spherical intermolecular potentialwas used. The dislocation was found to dissociate into 2 Shockley partials. The partialdislocations had spread-out cores, and the width of the Burgers vector density at the half-peak height attained 13 times the magnitude of the Burgers vector.S.Tamaki, N.Ide, I.Okada, K.Kojima: Japanese Journal of Applied Physics - 1, 1998,37[5A], 2608-9

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C (Diamond): B DiffusionThree natural type-IIa crystals were studied with regard to the forced diffusion of B in anelectric field. The B diffusivity was found to be equal to 8.4 x 10-15 or 4 x 10-14cm2/s at1000C, depending upon the direction of the electric field. The drift velocity of B in a fieldof 850V at 1000C was about 1.2 x 10-8cm/s.T.Sung, G.Popovici, M.A.Prelas, R.G.Wilson, S.K.Loyalka: Journal of MaterialsResearch, 1997, 12[5], 1169-71

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y C (Diamond): D DiffusionThe diffusion of D in diamond-like C films was studied. The D concentration profiles inD+-ion implanted films were measured by means of secondary-ion-mass spectrometry. Amodel was proposed, for describing the experimental depth profiles, in which it wasassumed that atomic D was the diffusing species, whereas D in clusters was immobile.

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C (Diamond) Carbon and Carbides C (Diamond)

The results showed that the concentration of D clusters, relative to the total Dconcentration, increased when the total D concentration decreased; thus leading toconcentration-dependent diffusion. The diffusion coefficients which were deduced foratomic D (figure 1) indicated an associated activation energy of 2.9eV. The solidsolubility of D decreased with increasing temperature.T.Ahlgren, E.Vainonen, J.Likonen, J.Keinonen: Physical Review B, 1998, 57[16], 9723-6

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Figure 1Diffusivity of D in Diamond

C (Diamond): Dislocations and Grain BoundariesTransmission electron microscopy of polycrystalline natural diamond (carbonado)revealed the existence of voidites. These were different to the voidites in type-Ia naturaldiamonds, not only with regard to shape but also content and origin. Other defects werealso observed, such as dislocation networks, micro-twins and large-angle grainboundaries. It was concluded that voidites could be formed as a sink for impurities; notonly within diamond grains but also between grains with a closely-related orientation. Inthe former case, the nucleation of voidites along dislocation lines was preferred while, in

1.0E-02

1.0E-01

1.0E+00

1.0E+01

1.0E+02

7 8 9 10

E = 2.9eV

104/T(K)

D (n

m2 /s

)

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the latter case, planar-like voidites formed so as to permit direct bonding betweendiamond grains.J.H.Chen, D.Bernaerts, J.W.Seo, G.Van Tendeloo, H.Kagi: Philosophical MagazineLetters, 1998, 77[3], 135-40

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C (Diamond): Point DefectsFirst-principles plane-wave pseudopotential local-density theory was used to model Insites. The most stable site involved In which was located at an off-center substitutionalposition, between 2 adjacent C vacancies. This was thought to be the most stablefundamental defect which formed following implantation. Initial formation of this centerwas associated with two C interstitials. However, the total energy of such a defect wassignificantly higher than that without C interstitials. Thus, only after these C interstitialatoms were removed by annealing would it be possible to observe the most stablepredicted form of the In defect. This had a VC-In-VC [111] structure, and was expected tobe associated with an observed 117MHz quadrupole frequency.B.P.Doyle, J.K.Dewhurst, J.E.Lowther, K.Bharuth-Ram: Physical Review B, 1998, 57[9],4965-7

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C (Diamond): Point DefectsThe characteristic features of photo-current and electron paramagnetic resonance spectrawere attributed to the principal defects within the gap of optical-quality chemical vapordeposited samples. A shoulder in the photo-current spectra, with an onset at about 2.2eV,was attributed to the single substitutional N defect (g = 2.0024). A second feature in thephoto-current spectra, with an onset of about 1.3eV, was observed in as-grown sampleswith a H-terminated surface. The defect level which was associated with this feature wasH-related. This defect disappeared after oxidation of the sample surface. An electronparamagnetic resonance g-value of 2.0028 was also suggested to be H-related.M.Nesládek, L.M.Stals, A.Stesmans, K.Iakoubovskij, G.J.Adriaenssens, J.Rosa,M.Vanecek: Applied Physics Letters, 1998, 72[25], 3306-8

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C (Diamond): Point DefectsOn the basis of 11B and 14N hyperfine interactions, as determined by using the electronparamagnetic resonance technique, the NIRIM-4 center in electron-irradiated B-dopedsynthetic diamond crystals was identified as being a <100>-split [B-N]+ interstitialcy. TheC2v symmetry of the spectrum required that the N and B atoms lay on the same [100]axis. The contributions to the wave function of unpaired electrons from, the atomicorbitals of N and B, were almost entirely p-like, with the directions of p-orbitals along[011] and [01̄1], respectively.J.Isoya, H.Kanda, Y.Morita: Physical Review B, 1997, 56[11], 6392-5

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C (Diamond): Point DefectsIt was recalled that the vacancy model for impurity vacancy defects in semiconductorsassumed that the ground and low-energy excited states could be derived from the four sp3

hybrid orbitals on atoms which bordered the vacancy. It was pointed out that there weremany cases where this model worked, but a counter-example was described here whichconcerned the lowest excited state of the [V-N3] defect in diamond. It was shown that ashallow electron trap, localized outside the vacancy, was involved in the first excited stateand was responsible for the N2 and N4 optical bands which were associated with thedefect.R.Jones, J.P.Goss, P.R.Briddon, S.Oberg: Physical Review B, 1997, 56[4], R1654-6

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C (Diamond): Point DefectsInfra-red absorption data on irradiated and annealed synthetic diamonds were presentedwhich confirmed the suggestion that a component which was found in the defect-induced1-phonon region of some diamonds arose from single-substitutional N+. Theconcentration ratio of N+ to neutral substitutional N0 centers could be changed by shininglight with various energies onto the samples. Changes in absorption, of the infra-redcomponent which was associated with N0 centers, were correlated with changes in the N+

component. By using a previously determined relationship, between the concentration ofN0 centers and the peak absorption coefficient at 1130/cm, a relationship was derivedbetween the peak absorption at 1332/cm and the concentration of N+ centers. Thisrelationship was that 1/cm of absorption was produced by 5.5ppm of N+ centers. Otherdefects could also give rise to absorption at 1332/cm, but the N+ component was uniquelyidentified by further peaks at 1046 and 950/cm. The significance of this component wasillustrated by the fact that some samples could contain more than 80ppm of N+ centers,and this therefore had to be allowed for when estimating the total N concentration. Byusing the above relationship, useful parameters were derived which related theconcentration of neutral vacancies, negative vacancies and negatively-charged N-vacancycenters to their respective zero-phonon line integrated absorptions.S.C.Lawson, D.Fisher, D.C.Hunt, M.E.Newton: Journal of Physics - Condensed Matter,1998, 10[27], 6171-80

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C (Diamond): Surface ReconstructionIt was recalled that a plausible (2 x 2) p-bonded-trimer structure had recently been foundfor the clean diamond (111) surface. By using the local-orbital density-functionalmolecular-dynamics method, this (2 x 2) structure was found to undergo phase transition,into the Pandey (2 x 1) p -bonded chain structure upon adsorbing H or Li. However, thissurface did not undergo phase transition upon the adsorption of F or O at the sameadsorption site. The results indicated that s-electrons played an important role indisrupting the sp2 bonding of the (2 x 2) structure.

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C (Diamond) Carbon and Carbides C (Graphite)

M.H.Tsai, J.C.Jiang, S.H.Lin: Physical Review B, 1997, 56[19], 12127-30[446-164-071]

C (Graphite): Electron Irradiation and Point DefectsElectron irradiation damage in highly-graphitized C fibres was examined by means oftransmission electron microscopy, in multi-beam lattice imaging mode, together withselected-area electron diffraction. It was found that the (00•2) lattice spacing wasincreased homogeneously up to about 11%, over the whole of the irradiated area, withincreasing irradiation time. The structural order, parallel to the basal plane, graduallydeteriorated although the order perpendicular to the basal planes was relatively wellretained. Computer image simulations based upon previously proposed model structures,in which interstitial atom clusters were incorporated, could not explain the homogeneousdilation of the (00•2) lattice fringes.S.Muto, T.Tanabe: Philosophical Magazine A, 1997, 76[3], 679-90

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C (Graphite): Ion Bombardment and Point DefectsThe surface defects which were produced by single-ion impacts of highly-charged Ar ions(charge state of up to 8) were investigated by using scanning tunnelling microscopy andatomic force microscopy. The defect appeared to be a protrusion in scanning tunnellingmicroscopic images, but appeared flat in atomic force microscopy images. On the basis ofthese contrasting images, the defects were concluded to be due to an increase, in the localcharge density of states at the surface, which was caused by C-atom sputtering. Theaverage defect size increased markedly with the charge state of the incident Ar ions. Thiswas attributed to an enhancement of potential sputtering, due to the Coulomb repulsionbetween surface holes which were generated by the neutralization of highly charged Arions.K.Mochiji, S.Yamamoto, H.Shimizu, S.Ohtani, T.Seguchi, N.Kobayashi: Journal ofApplied Physics, 1997, 82[12], 6037-40

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C (Graphite): DislocationsThe remarkable features of the local density of states at a single n-membered ring defect(n = 4, 5, 7, or 8) in mono-layered graphite were described. On the basis of the n-foldrotational symmetry at the center of the n-membered ring, the wave functions and spectraof these systems were classified into n types of component; according to the nature of therotational symmetry group. By taking account of this symmetry, the local density of stateswas concluded to be broadened energy levels of the corresponding isolated n-memberedring molecule. The n-membered ring defect, in the large-n limit, could be considered to bea screw dislocation. The value of its local density of states approached zero as the inverseof the logarithm of the energy at the band edges and at the band center.R.Tamura, K.Akagi, M.Tsukada, S.Itoh, S.Ihara: Physical Review B, 1997, 56[3], 1404-11

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K2RbC60 Carbon and Carbides SiC

K2RbC60, Rb6C60: Point DefectsThe temperature dependence of atomic displacements in superconducting K2RbC60 andnon-superconducting Rb6C60 was studied by using X-ray powder diffraction methods.Rietveld refinement of the powder diffraction data for K2RbC60 at 10.2K revealed that theoccupancy of Rb+ ions at an octahedral site was 71, rather than 100%. It was found thatthe atomic displacements of C atoms and metal ions in K2RbC60 exhibited an increasearound the superconducting critical temperature. No such increase was observed inRb6C60.Y.Yoshida, Y.Kubozono, T.Urakawa, H.Maeda, S.Kashino, Y.Murakami, T.Ohta,F.Izumi, K.Yamada, Y.Furukawa: Solid State Communications, 1998, 105[9], 557-60

[446-164-072]

SiC: Electron Irradiation and Point DefectsAn investigation was made of 2.5MeV irradiation-induced deep levels in 6H-type p+ndiodes which had been prepared by means of chemical vapor deposition. Deep-leveltransient spectroscopy revealed several overlapping peaks at temperatures ranging from140 to 650K. It was shown that the amplitudes of the deep-level transient spectroscopypeaks, E1 and E2, were significantly reduced when both holes and electrons were injectedinto the space-charge region. This reduction in peak amplitude occurred if the E1 and E2centers acted as efficient hole-electron recombination centers. Due to this behavior, it wassuggested that the E1 and E2 centers might be important during recombination inirradiated material. However, these two centers did not govern the minority carrierlifetime. The temperature dependence of the electron capture cross-section was measuredfor the E1, E2 and Ei levels. Two of these, E1 and E2, exhibited a weak temperaturedependence. On the basis of their small electron-capture cross-sections, and the absenceof a Poole-Frenkel effect, it was suggested that E1 and E2 were associated with neutralacceptor-like centers. The electron-capture cross-section of the Ei level was shown to betemperature-independent at temperatures ranging from 248 to 275K.C.Hemmingsson, N.T.Son, O.Kordina, E.Janzén, J.L.Lindström: Journal of AppliedPhysics, 1998, 84[2], 704-8

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SiC: Electron Irradiation and Point DefectsThe production of vacancies, via the 3MeV electron irradiation of 6H-type material atroom temperature, was studied by using positron lifetime spectroscopy; combined withannealing experiments. It was found that the trapping rates of positrons in the vacanciesincreased linearly with the fluence in the initial stages of irradiation. Following this linearincrease, the trapping rates were found to be proportional to the square root of thefluence. The linear and non-linear fluence dependences of the trapping rate wereexplained in terms of a reduction in the number of vacancies, due to recombination withinterstitials during irradiation. The positron trapping rate for an admixture of Si vacanciesand divacancies exhibited a tendency to saturate in the higher fluence range. The trappingrate

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SiC Carbon and Carbides SiC

of C vacancies decreased after reaching a maximum. These results were explained interms of a shift in the Fermi level, due to irradiation. It was found that, in lightlyirradiated specimens, an annealing stage was observed which was caused byrecombination between close vacancies and interstitials. Such an annealing stage was notobserved in heavily irradiated specimens. The differing results were explained in terms ofa reduction in the number of interstitials, due to recombination with vacancies and thelong-range migration of interstitials to sinks during irradiation.A.Kawasuso, H.Itoh, T.Ohshima, K.Abe, S.Okada: Journal of Applied Physics, 1997,82[7], 3232-8

[446-164-073]

SiC: Ion Implantation and Point DefectsThe annealing of damage in 6H-type material, which was caused by ion implantation to 2different fluences, was studied by using mono-energetic positron Doppler broadening andlifetime techniques. The measurements were supported by calculations of the positronlifetimes in vacancy clusters. At both fluences, 2 defective layers were identified andcharacterized, by depth and defect-type, as a function of the annealing temperature. Theresults indicated that it was impossible to remove the radiation damage by annealing attemperatures of up to 1500C.G.Brauer, W.Anwand, P.G.Coleman, J.Störmer, F.Plazaola, J.M.Campillo, Y.Pacaud,W.Skorupa: Journal of Physics - Condensed Matter, 1998, 10[5], 1147-56

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SiC: Ion Implantation and Point DefectsIn a polemical exchange, the first-named author demonstrated that the other authors'simplified analysis of Rutherford back-scattering channelling data on damage productionin SiC could not be used to calculate the atomic displacement energy. The value of 12eVwhich these authors had given was much too small. Also, their assumption of similardisplacement energies in Si and SiC was essentially wrong. The latter 2 authors repliedthat their method of estimating the atomic displacement energy was correct since theanalysis was limited to a case in which the average energy density within the collisioncascade was of the order of only 0.0001eV/atom, while the areal density of displacedatoms approached the value which was expected on the basis of the Kinchin-Pease model.V.Heera, M.G.Grimaldi, L.Calcagno: Journal of Applied Physics, 1998, 83[7], 3935-7

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SiC: Ion Implantation, Point Defects, and Defect AnnealingThe annealing of defects in N2

+- or Al+-implanted 3C-material was studied by usingmono-energetic positron beams. In the case of as-implanted specimens, the mean size ofthe open volume of defects was estimated to be close to that of divacancies. On the basisof the annealing behavior of the characteristic value of the S-parameter whichcorresponded to the annihilation of positrons which were trapped by vacancy-typedefects, the temperature range for the annealing of defects could be divided into 5 stages.The annealing behavior in stages I (20 to 500C), II (500 to 800C) and III (800 to 1006C)

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was identified with the agglomeration of vacancy-type defects due to the migration of Cvacancies, Si vacancies and vacancy complexes (such as divacancies), respectively.Stages IV (1000 to 1200C) and V (1200 to 1400C) were attributed to the formation ofextended defects and their recovery, respectively.A.Uedono, H.Itoh, T.Ohshima, R.Suzuki, T.Ohdaira, S.Tanigawa, Y.Aoki, M.Yoshikawa,I.Nashiyama, T.Mikado, H.Okumura, S.Yoshida: Japanese Journal of Applied Physics -1, 1997, 36[11], 6650-60

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SiC: Ion Implantation, Point Defects, and Defect AnnealingVacancy-type defects and their annealing behavior were studied, using mono-energeticpositron beams, in 6H-type samples which had been implanted with 200keV P+. Inspecimens which had been implanted to a dose of 1013/cm2, the mean size of the openvolume of defects was estimated to be close to that of divacancies. On the basis of theannealing behavior of the S-parameter for the annihilation of positrons that were trappedby vacancy-type defects, the temperature range for the annealing of vacancy-type defectswas divided into 3 parts. The annealing behaviors in stages I (200 to 700C) and II (700 to1000C) were attributed to the agglomeration of defects via the migration ofmonovacancies, and vacancy complexes such as divacancies, respectively. In stage II,near to the defect-free region, the agglomeration of defects was suppressed by therecombination of vacancy-type defects and interstitials. Stage III (1000 to 1300C) wasattributed to the formation of extended defects, and their recovery.A.Uedono, T.Ohshima, H.Itoh, R.Suzuki, T.Ohdaira, S.Tanigawa, Y.Aoki, M.Yoshikawa,I.Nashiyama, T.Mikado: Japanese Journal of Applied Physics - 1, 1998, 37[5A], 2422-9

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SiC: Ion Implantation, Point Defects, and Defect AnnealingThe defects which were introduced into epitaxially grown 3C-type material, byimplantation of 200keV N2

+ or Al+ to doses of between 1013 and 1016/cm2 at temperaturesranging from ambient to 1200C, were studied by using electron spin resonance,photoluminescence and positron annihilation spectroscopy. It was found that, althoughhot-implantation reduced the number of paramagnetic defects and improved thecrystallinity of implanted layers, it led to the simultaneous formation of vacancy clusters.Small vacancy clusters were produced mainly by low-dose (1013/cm2) implantation, andlarger vacancy clusters were formed by high-dose (1015/cm2) implantation. The averagesize of such clusters increased with implantation temperature. Formation of the vacancyclusters was independent of the nature of the implanted ion species. All of the resultswere explained in terms of the migration and combination of point defects such asvacancies and interstitials during hot-implantation. In the case of high-dose Al+-implantation, additional paramagnetic defects with g = 2.0035 were formed byimplantation above about 800C; thus suggesting that this defect was related to theprecipitation of Al atoms. It was suggested that the g = 2.0035 defect acted as a non-radiative recombination center. This defect was thought to be unrelated to vacancy-typedefects. Very large vacancy clusters were created by annealing samples which had beenamorphized by high-dose

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(1015/cm2) implantation at room temperature, whereas such further vacancy clustering didnot occur in hot-implanted samples.H.Itoh, T.Ohshima, Y.Aoki, K.Abe, M.Yoshikawa, I.Nashiyama, H.Okumura, S.Yoshida,A.Uedono, S.Tanigawa: Journal of Applied Physics, 1997, 82[11], 5339-47

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SiC: Neutron Irradiation and Point DefectsIt was recalled that deep-level transient spectroscopy signals, attributed to centers labelledH1, H2, H3 and E2, were long ago detected in neutron-irradiated 3C-type material. The Hcenters were believed to be the primary point defects, with the E2 center being asecondary defect which formed after the H centers began to migrate. Computationalevidence was presented here which suggested that the H centers were due to Si antisitedefects (SiC). In both the cubic (3C) and hexagonal (2H) polytypes, the Si antisite hadseveral ionization levels in the band-gap. The positions of these ionization levels in 3C-type material were calculated accurately by using the plane-wave pseudopotential methodand a 128-atom super-cell. Very good agreement with experimental data was found, andindicated that the H centers were due to the formation of SiC during neutron irradiation.The formation energies and local geometries of antisite defects were also predicted.L.Torpo, S.Pöykkö, R.M.Nieminen: Physical Review B, 1998, 57[11], 6243-6

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SiC: Neutron Irradiation and Point DefectsElectron paramagnetic resonance and electron-nuclear double resonance methods wereused to identify the negatively charged Si vacancy in neutron-irradiated 4H-type material.The identification was based upon resolved ligand hyperfine interactions with C and Sinearest-neighbors and next-nearest neighbors, and upon the determination of the spinstate (S = 3/2). The magnetic resonance parameters of VSi

- were found to be almostidentical for 3C, 4H and 6H samples. The experimental data were supported bytheoretical ligand hyperfine interaction data that were based upon a total-energycalculation which involved the standard local-density approximation of density-functionaltheory.T.Wimbauer, B.K.Meyer, A.Hofstaetter, A.Scharmann, H.Overhof: Physical Review B,1997, 56[12], 7384-8

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SiC: DislocationsThe hardnesses of the opposite basal faces of 4H-type single crystals were measured attemperatures ranging from 25 to 200C. A marked hardness discrepancy was foundbetween the Si-terminated (00•1) and C-terminated (00•̄1) faces of the polar crystal.Transmission electron microscopic investigation of dislocations in the plastic zone of the1200C indentations showed that they lay predominantly on basal planes, parallel to theindented face, and that the extra half-planes of the non-screw dislocations originated fromthe indented face. It was also found that, when the (00•1) Si-terminated face wasindented, the dislocations were either widely dissociated (with the width of the stacking-fault ribbon being much greater than the equilibrium value) or were single leading

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partials, with the corresponding trailing partials absent. In this case, all of the leadingpartials were found to have a Si core. On the other hand, the dislocations in the plasticzone of the C-terminated face were in the form of dissociated dislocations, with the widthof the associated stacking-fault ribbons being appreciably less than the equilibrium value.The leading partials of these dissociated dislocations had a C core. The results indicatedthat the hardness of the polar basal faces of 4H-crystals at high temperatures wasdetermined partly by the nature of the dislocation cores which were nucleated by theindentation process. It was suggested that this was due to the effect of the core upon thegeneration and glide of the leading partial dislocations.X.J.Ning, N.Huvey, P.Pirouz: Journal of the American Ceramic Society, 1997, 80[7],1645-52

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SiC: DislocationsThe irreversible formation of a network of linear defects was observed in images whichexhibited recombination luminescence from injection diodes in hexagonal carbidesamples. The defects were related to dislocations that had initially formed as a result ofthermal stresses near to the tip of the contact probe and which then propagated throughthe diode area. The dislocation network appeared, in electroluminescence images, asbright-line defects and contrasted with the well-known dark-line defects that were due tothe degradation of GaAs-based light-emitting devices. Higher forward currents werefound to promote dislocation growth.A.O.Konstantinov, H.Bleichner: Applied Physics Letters, 1997, 71[25], 3700-2

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SiC: DislocationsStep structures on the {00•1} facets of single crystals which had been grown using themodified Lely method were examined by means of optical and atomic force microscopy.The results were compared with those obtained using the chemical vapor depositionmethod. The step structures around micro-pipes were quite different for differingpolarities of the growing surface. Step bunching was more favored on Si faces than on Cfaces. A single large step with a height of 8c, and 18 steps with a height of c/2, weretypically observed to propagate from micro-pipes at the center of each spiral on the Si-faces and C-faces, respectively, of 6H-type crystals. No micro-pipes were detected at thecenter of spirals with a step height of 1.5nm in 6H-type material. The motion ofdislocations was monitored by successively etching and polishing a grown crystal fromthe surface to the inside. It was noted that the screw dislocations shifted outwards, by 10to 40µ, from the center of a giant spiral during 200µ of vertical growth. On the otherhand, edge dislocations could glide in any direction.N.Sugiyama, A.Okamoto, K.Okumura, T.Tani, N.Kamiya: Journal of Crystal Growth,1998, 191, 84-91

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SiC: DislocationsIt was recalled that this material suffered from microscopic hollow defects called micro-pipes. It was noted that large-angle convergent-beam electron diffraction, which was awell-established method for determining the total Burgers vector of a dislocation, hadbeen applied to nanopipes in GaN. However, this method could not be successfully usedhere because the micro-pipes in SiC usually had diameters which were some orders ofmagnitude larger than those of the nanopipes in GaN. The present results showed that themicro-pipes were hollow-core dislocations, according to Frank's model, but containeddislocations of mixed type.J.Heindl, W.Dorsch, H.P.Strunk, S.G.Müller, R.Eckstein, D.Hofmann, A.Winnacker:Physical Review Letters, 1998, 80[4], 740-1

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SiC: Dislocations and MicropipesAtomic force microscopy was used to study the (00•1) growth surface of a 6H-type singlecrystal at points where micro-pipes emerged at the growth surface. All of the micro-pipesexamined were at the origin of spiral steps; thus indicating that dislocations intersectedthe surface at these points. The dislocations which were observed at surface/micro-pipeintersections had Burgers vectors of at least 4b, where b was the Burgers vector of a unitscrew dislocation which was aligned along the c-axis (b = 1.519nm). Single and doubleunit dislocations were also observed, but they were not associated with micro-pipes.Micron-sized deposits of heterogeneous phase were observed in the vicinity of the micro-pipes. The curvature of growth steps around these heterogeneities indicated that theyimpeded step motion when the crystal was growing. On the basis of the observations, amodel was proposed for the formation of super-dislocation/micro-pipe complexes thatinvolved the coalescence of unit screw dislocations which were forced towards oneanother as large steps grew around heterogeneous surface material.J.Giocondi, G.S.Rohrer, M.Skowronski, V.Balakrishna, G.Augustine, H.M.Hobgood,R.H.Hopkins: Journal of Crystal Growth, 1997, 181, 351-62

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SiC: MicropipesA process was developed for growing micropipe-free single crystals by using a modifiedLely method. The process parameters were kept near to thermal equilibrium. Themaximum average thermal gradient, inside the furnace, which led to micropipe-freegrowth was 5K/cm. A gradient of 7.5K/cm resulted in marked defect formation andproduced a high density of micro-pipes (greater than 200/cm2). The highest growth ratehere which provided micro-pipe-free growth was 0.27mm/h. Single boule crystals of 6H-material were grown on both the C face and the Si face of 6H-SiC Lely platelets.N.Schulze, D.L.Barrett, G.Pensl: Applied Physics Letters, 1998, 72[13], 1632-4

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78


SiC, GaN: Micropipes and NanopipesIt was recalled that micro-pipes in 6H- and 4H-samples, grown by using the modifiedLely technique, and nanopipes in GaN, grown via metal-organic vapor-phase epitaxy ontosapphire, had been attributed to Frank growth dislocations which had an empty core dueto their large Burgers vectors. Such so-called killer defects had a deleterious influenceupon device performance. A formation mechanism was proposed here for nanopipes andmicro-pipes in hexagonal semiconductors.P.Pirouz: Philosophical Magazine A, 1998, 78[3], 727-36

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SiC: Point DefectsThe effects of annealing upon the structure of radio-frequency sputtered amorphous filmswhich had been prepared under H partial pressures were investigated. The infra-redresults for as-prepared films suggested that, as the H partial pressure was increased, moreH was incorporated into the film to form Si-H and C-H bonds, while fewer Si and Catoms were available to form Si-C bonds. The X-ray photo-electron spectroscopic resultsfor as-prepared films agreed with the infra-red results, in that the fraction of Si-Cdecreased and the fraction of Si-H and C-H increased with increasing partial pressure.The infra-red and X-ray photo-electron spectroscopic results for annealed films suggestedthat, as the annealing temperature was increased, dangling Si and C bonds combined toform Si-C bonds in non-hydrogenated samples. The increase in Si-C bonds inhydrogenated samples was concluded to be more probably due to the formation of Si-Cbonds from a break-up of Si-H and C-H bonds.W.K.Choi, T.Y.Ong, L.S.Tan, F.C.Loh, K.L.Tan: Journal of Applied Physics, 1998,83[9], 4968-73

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SiC: Point DefectsHigh-frequency (95GHz) and conventional (9.3GHz) pulsed electron paramagneticresonance and electron-nuclear double resonance studies were made of the deep Bacceptor in 6H-type material. The results supported a model in which the deep B acceptorconsisted of a B atom on a Si site, with an adjacent C vacancy. The latter vacancycombined with a B atom along the hexagonal c-axis. It was concluded that 70 to 90% ofthe spin density resided in the Si dangling bonds which surrounded the vacancy, andanother 9% on the neighboring C atoms. The spin-density distribution was more localizedthan in the case of the shallow B acceptor, as deduced from electron nuclear doubleresonance data.A.V.Duijn-Arnold, T.Ikoma, O.G.Poluektov, P.G.Baranov, E.N.Mokhov, J.Schmidt:Physical Review B, 1998, 57[3], 1607-19

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SiC: Point DefectsA study was made, of the effect of excimer laser annealing upon a high-qualitySi0.982C0.018 pseudomorphic layer, by using Fourier-transform infra-red spectroscopy,Raman spectrometry and X-ray diffraction techniques. The substitutional C concentrationwas found to decrease, as a function of fluence, when 50 laser pulses were used at roomtemperature or in vacuum. The evolution of the strain profile was studied by means of X-ray diffraction, and dynamic diffraction simulations. Most of the strain was released aftermelting, and new Fourier-transform infra-red and Raman peaks appeared around 830/cm.This was attributed to the formation of SiC micro-precipitates and V-O asymmetricalcenters. The proposed mechanism of substitutional C removal involved SiC precipitation,and a reaction (between V-O and substitutional C) which formed volatile CO. In the caseof pulsed laser-induced epitaxy, it was predicted that the highest substitutional C contentwould be obtained by 1 laser pulse in an O-free ambient. It was concluded that pulsed-laser induced epitaxy was suitable for the localized patterning of ultra-shallow buried SiCjunctions.C.Guedj, G.Calvarin, B.Piriou: Journal of Applied Physics, 1998, 83[8], 4064-8

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SiC: Point DefectsAn electrically active defect was detected in epitaxial layers of 4H-type material whichhad been grown by means of vapor-phase epitaxy. This defect was characterized by alevel at about 0.70eV below the conduction-band edge, an extrapolated capture cross-section of about 5 x 10-14cm2 and a concentration of approximately 1013/cm3. Secondary-ion mass spectrometry revealed no sign of Ti, V or Cr. After 2MeV electron irradiation,the 0.70eV level did not increase in concentration, but 3 new levels were observed atapproximately 0.32, 0.62 and 0.68eV below Ec; with extrapolated capture cross-sectionsof 4 x 10-14, 4 x 10-14 and 5 x 10-15cm2, respectively. However, the defects which causedthese levels were unstable, and decayed after some time at room temperature. Thisresulted in the formation of the 0.70eV level. The results strongly suggested that the0.70eV level originated from a defect of intrinsic type. Such an unstable behavior byelectron irradiation-induced defects had not been observed in the 6H polytype at roomtemperature.J.P.Doyle, M.K.Linnarsson, P.Pellegrino, N.Keskitalo, B.G.Svensson, A.Schöner,N.Nordell, J.L.Lindström: Journal of Applied Physics, 1998, 84[3], 1354-7

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SiC: Point DefectsA deep-level transient spectroscopic study was made of deep-level defect centers on then-side of p+n junction diodes which had been prepared by low- and high-temperature Al-ion implantation of n-type 6H-material. Two shallow Al-acceptor levels were detected inthe n-type region, just beyond the implantation depth, via their minority-carrier emissionsignatures. The predominant level was situated at 0.26eV above the valence band, andwas associated with a shallower level of low intensity. A comparison with

80


photoluminescence results suggested that the predominant level (labelled Ak) and theshallower level (labelled Ah) were associated with the cubic and hexagonal lattice sites,respectively. Contrary to previously reports, which had noted the presence of manydifferent implantation-induced donors within the implantation region, only a single deepdonor level, at Ec - 0.44eV, was found to occur in the post-implantation region. Thisindicated that the various crystal damage sites had differing spatial distributions.S.Fung, M.Gong, C.D.Beling, G.Brauer, H.Wirth, W.Skorupa: Journal of AppliedPhysics, 1998, 84[2], 1152-4

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SiC: Point DefectsA comprehensive study was made of C-rich films of amorphous hydrogenated material byusing optical absorption, Fourier-transform infra-red spectroscopy, thermal desorption,atomic force microscopic, positron lifetime and Doppler-broadening techniques. Theresults suggested that open volumes formed in the films due to incomplete breaking of thesource molecule during film deposition. These open volumes were interconnected andcould trap ambient gases during film growth, or afterwards. With increasing temperature,the gases were desorbed from the internal surfaces of these open volumes and werereleased from the sample. This then increased the areal density of the defects, and wasdetected via positronium formation and the annihilation of positrons with surfaceelectrons. At sufficiently high temperatures, thermal rupture of Si-H and C-H bondsoccurred and resulted in irreversible structural changes, and film densification, due to newC-C bond formation.T.Friessnegg, M.Boudreau, P.Mascher, A.Knights, P.J.Simpson, W.Puff: Journal ofApplied Physics, 1998, 84[2], 786-95

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SiC: Point DefectsThe (111)-(3 x 3) phase was analyzed by using scanning tunnelling microscopy, low-energy electron diffraction holography, density functional theory, and conventional low-energy electron diffraction. A single adatom per unit cell, which was found in scanningtunnelling microscopy, acted as a beam-splitter for the holographic inversion of discretelow-energy electron diffraction spot intensities. The resultant 3-dimensional image guidedfurther detailed low-energy electron diffraction and density functional theory analyseswhich identified a Si tetramer on a twisted Si adlayer, with clover-like rings. This twistmodel, with one dangling bond left per unit cell, represented a novel (n x n)reconstruction mechanism for group-IV (111) surfaces.U.Starke, J.Schardt, J.Bernhardt, M.Franke, K.Reuter, H.Wedler, K.Heinz, J.Furthmüller,P.Käckell, F.Bechstedt: Physical Review Letters, 1998, 80[4], 758-61

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SiC: Stacking FaultsDeformation tests were carried out on 6H-type samples with an orientation that favoredactivation of the <21̄•0>(00•1) slip system. The tests were performed at temperatures of

81

SiC Carbon and Carbides (Si,Ge)C

between 550 and 1400C, using a strain rate of 3.1 x 10-5/s. Optical and transmissionelectron microscopy were then used to study deformation-induced defects such asstacking faults, deformation kinks, and cracks. On the basis of the observations, amechanism was proposed for the formation of deformation kinks, nucleation andpropagation of cracks, and temperature dependence.A.V.Samant, X.L.Wei, P.Pirouz: Philosophical Magazine A, 1998, 78[3], 737-46

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SiC: Stacking FaultsIt was recalled that the X-ray diffraction patterns from nominally ß-phase specimens oftendiffered from those which were expected for a cubic crystal structure. These differencesincluded the presence of additional peaks, enhanced background intensities, peak-broadening, changes in relative peak heights, and shifts in peak positions. Thediscrepancy had long been recognized as being due to the presence of stacking faults.Computer simulations were used here to showed that the variations were closely relatedto differences in the types and spatial distributions of stacking faults. In these simulations,stacking sequences were generated by using a selectively activated 1-dimensional Isingmodel which permitted a wide variety of fault configurations to be generated. Thesimulation results showed that it was necessary to suppress 2-layer twins, but to promotethe formation of 3-layer twins, so as to reproduce the gradual increase in backgroundintensity which was observed experimentally. The greatest puzzle which was associatedwith the present strategy for generating model stacking sequences was the need for a so-called switch which would produce a bimodal distribution of stacking errors.V.V.Pujar, J.D.Cawley: Journal of the American Ceramic Society, 1997, 80[7], 1653-62

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SiC: Surface ReconstructionThe Si-rich 3C-type (001) 3 x 2 surface was studied by using high-resolution core-levelphoto-emission methods. Well-resolved Si 2p and C 1s core-level spectra were measuredat a temperature of about 120K. Three different Si 2p surface components were clearlyidentified, with binding-energy shifts of -0.58, -0.92 and -1.27eV. The presence of thesecomponents, and their intensity ratios, were consistent with a structural model whichinvolved 2/3 of a monolayer of additional Si dimers, but were incompatible with anothermodel which assumed only 1/3 of a monolayer of Si dimers.H.W.Yeom, Y.C.Chao, S.Terada, S.Hara, S.Yoshida, R.I.G.Uhrberg: Physical Review B,1997, 56[24], R15525-8

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(Si,Ge)C/Si: Point DefectsDeep-level transient spectroscopy was used to measure the activation energies of deeplevels in n-type heterostructures which had been grown by means of solid-sourcemolecular beam epitaxy. Four deep levels were observed at energies ranging from 0.231to 0.405eV below the conduction band. The largest deep-level concentration was foundfor the deepest level, and was equal to about 2 x 1015/cm3. Although a large amount (1 to

82

(Si,Ge)C Carbon and Carbides TiC

2at%) of non-substitutional C was present in the alloy layers, no deep levels wereobserved at any energy levels that had apparently been previously attributed to interstitialC.B.L.Stein, E.T.Yu, E.T.Croke, A.T.Hunter, T.Laursen, J.W.Mayer, C.C.Ahn: AppliedPhysics Letters, 1998, 73[5], 647-9

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Figure 2Diffusivity of N in Various Nitrides

y TaC, Ta2C, TaN, Ta2N: C, N DiffusionThe diffusivity of C in the carbides was investigated at temperatures ranging from 1700 to2200C (figure 3), and the diffusivity of N in the nitrides was investigated at temperaturesranging from 1700 to 1950C (figure 2). The concentration-independent diffusioncoefficients were obtained in each case by investigating enhanced layer growth in wedge-shaped specimens. In the case of the non-metal rich phases with a broad homogeneityrange (d-TaC, ß-Ta2C, ß-Ta2N), the concentration-dependent diffusion coefficients werecalculated and were compared with the concentration-independent diffusion coefficients.The calculation of the concentration-dependent diffusion coefficients was carried out byfitting a modified error function to the measured concentration profiles, while assumingthat the non-metal diffusivity was an exponential function of the non-metal concentration.A marked dependence of the diffusion coefficient upon the non-metal concentration wasfound for d-TaC, whereas the non-metal diffusion coefficients were almost independentof concentration in ß-Ta2C and ß-Ta2N.D.Rafaja, W.Lengauer, H.Wiesenberger: Acta Materialia, 1998, 46[10], 3477-83

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TiC/Cu: DislocationsSamples of TiC dispersion-strengthened Cu alloy were prepared by mechanical alloyingand hot extrusion, and the evolution of the microstructure was monitored by using

1.0E-09

1.0E-08

1.0E-07

4.4 4.5 4.6 4.7 4.8 4.9 5 5.1

beta-Ta2Ndelta-TaNalpha-Ta(N)epsilon-TaN

104/T(K)

D (c

m2 /s

)

83

TiC Carbon and Carbides Ti3SiC2

transmission electron microscopy. The TiC dispersoids were formed in situ by reactionbetween Ti and graphite. The Ti diffused from the pre-alloyed CuTi matrix, and into Cinclusions which were embedded in the matrix after high-energy milling. Heat treatmentof the powder mixtures at 400C led to the heterogeneous nucleation of TiC at the C/Cuinterface. A well-defined cube-on-cube orientation relationship was established betweenthe TiC and the Cu matrix. A study of the morphology of the TiC dispersoids showed thatthey were faceted on the {111}TiC, {110}TiC and {100}TiC planes, and exhibited ledges atthe atomic scale. The TiC/Cu interfaces were atomically sharp, and were free of interfacephases. A {100}TiC||{100}Cu and <110>TiC||<110>Cu topotaxy led to a misfit of 17.6%between the lattices. This misfit was accommodated by a dislocation network along the<100>Cu directions.G.Dehm, J.Thomas, J.Mayer, T.Weissgärber, W.Püsche, C.Sauer: PhilosophicalMagazine A, 1998, 77[6], 1531-54

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Figure 3Diffusivity of N in Various Nitrides

Ti3SiC2: C, Si DiffusionIt was found that the reaction of Ti3SiC2 with monocrystalline Si wafers, at temperaturesranging from 1200 to 1350C, resulted in the formation of a dense surface layer thatcomprised a 2-phase mixture of TiSi2 and SiC. This layer grew as an outer layer with fine(1 to 5µ) SiC particles and inner coarser (10 to 15µ) ones. The overall growth rates of thelayers were parabolic. A comparison with previously published results supported theconclusion that the diffusion of Si through TiSi2 was rate-limiting. At temperaturesranging from 1400 to 1600C, the reaction of Ti3SiC2 with graphite foils resulted in theformation of a 15vol% porous surface layer of TiCx, where x was greater than 0.8. It wasshown that the carburization rate was limited by the diffusion of C through TiCx.T.El-Raghy, M.W.Barsoum: Journal of Applied Physics, 1998, 83[1], 112-9

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1.0E-12

1.0E-11

1.0E-10

1.0E-09

1.0E-08

1.0E-07

4 4.1 4.2 4.3 4.4 4.5 4.6 4.7 4.8 4.9 5 5.1

alpha-Ta(C)beta-Ta2Cdelta-TaCTa4C3

104/T(K)

D (c

m2 /s

)

84

Ti3SiC2 Carbon and Carbides Yb2.75C60

Ti3SiC2: Dislocations and Stacking FaultsThe carbide was prepared by reactive hot-pressing and was investigated by means oftransmission electron microscopy. The material consisted mainly of large elongated grainswith planar boundaries, and contained a low defect density. Dislocations were observedwithin the grains and at grain boundaries. This was the first reported detailed study ofdislocations in this material. Perfect dislocations, with a Burgers vector of 1/3<11•0>,were seen to be lying in (00•1) basal planes. These dislocations were mobile, andmultiplied during room-temperature deformation. All of the observed stacking faults layin basal planes. It was noted that, on the basis of the structure of this carbide, the fact thatthe defects were confined to basal planes was not surprising. The carbide was a layeredhexagonal material in which almost close-packed planes of Ti were separated from eachother by hexagonal nets of Si, with every fourth layer being a Si layer. A basalinteratomic vector was the shortest full translation vector in the structure. Therefore,perfect dislocations could be expected to have a Burgers vector of 1/3<12̄•0> and to lie inthe basal planes. Other dislocations were much less likely to exist because their Burgersvectors would be relatively large. The (00•1)[11•0] slip system was also common to allhexagonal metals, and it was not surprising that perfect basal plane dislocations with aBurgers vector of 1/3<11•0> should exist in this carbide. It was assumed that, afterlengthy annealing at 1600C, any dislocations which were created by plastic flow duringhot-pressing would have annealed out. This was consistent with the low defect densitywhich was observed. The arrays of perfect basal dislocations were suggested to haveformed during cooling, as a result of thermal residual stresses due to an anisotropy in thethermal expansion coefficients along the c- and a-axes. These dislocations appeared to beemitted from triple junctions of grain boundaries. The stacking faults were suggested tohave formed via the dissociation of a perfect dislocation into 2 partials which thenbounded the resultant stacking fault. The dissociation reaction was suggested to be:1/3<11•0> ? 1/3<10•0> + stacking fault + 1/3<01•0>. Because the perfect dislocationand the 2 partials had Burgers vectors which lay in the same basal plane, the stackingfault bounded by these partials also lay in that plane.L.Farber, M.W.Barsoum, A.Zavaliangos, T.El-Raghy, I.Levin: Journal of the AmericanCeramic Society, 1998, 81[6], 1677-81

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Yb2.75C60: Point DefectsThe well-characterized vacancy-ordered structure of this material was interpreted in termsof a simple electrostatic-energy model. The results provided useful insights into thestructural distortions that were introduced locally by tetrahedral cation vacancies, and themodel could be applied to alkali-metal and alkaline-earth fullerides. When applied tomonovalent fullerides, the model clearly showed that the electrostatic force between an Ocation and a neighboring T-site vacancy was very large. Therefore, in materials such asRb3C60, which exhibited T-site vacancies, the surrounding O-site cations would undergooff-center displacements towards the vacancy; within the limits set by cation size. Suchdisplacements were expected to be more difficult to detect there, than in Yb2.75C60,

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Yb2.75C60 Carbon and Carbides/Oxides Al2MgO4

because of the disordered arrangement of the vacancies. The displacements were alsoexpected to inhibit strongly any vacancy hopping; a process which had been proposed inorder to explain the splitting of T-site nuclear magnetic resonance lines in Rb3C60.K.M.Rabe, P.H.Citrin: Physical Review B, 1998, 58[2], R551-4

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Oxides

Ag2O-B2O3-SiO2-AgI: Ionic ConductionIonically conducting glasses in the system, 2AgI-Ag2O-2(0.95B2O3,0.05SiO2), wereprepared by melting in a microwave oven (5GHz, 900W). The homogeneity of thequenched glasses was equal to, or better than, that of glasses obtained by conventionalmelting at 730C. The microwave-melted glasses were reddish, while glasses of the samecomposition which were melted in a conventional furnace tended towards yellow. Thesecolours could be switched by using microwave and conventional re-melting. Thisbehavior was attributed to structural changes and/or changes in Ag+ mobility. The latterwere suggested by solid-state nuclear magnetic resonance data.D.J.Duval, B.L.Phillips, M.J.E.Terjak, S.H.Risbud: Journal of Solid State Chemistry,1997, 131[1], 173-6

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Ag2VP2O8: Ionic ConductionThe structure was studied by means of X-ray powder diffraction techniques, including theRietveld method. This phase, which was isostructural with Na2VP2O8, was monoclinic(P21/c, a = 0.7739, b = 1.3611, c = 0.6294nm, ß = 99.01º, Z = 4, V = 0.6548nm3). Itconsisted of [VP2O8]∞ layers, parallel to the (010) plane, which were interleaved with Agcations. This structure contained rather large tunnels which ran along [010] direction,where Ag cations were also located. Complex impedance electrical measurements wereused to determine the ionic conductivity, and its variation with temperature andfrequency. The activation energy for conduction was 0.58eV, and the conductivity wasequal to 2.08 x 10-5 and 1.43 x 10-3S/cm at 558 and 673K, respectively.A.Daidouh, M.L.Veiga, C.Pico: Journal of Solid State Chemistry, 1997, 130[1], 28-34

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Al2MgO4: Electron Irradiation, Ion Bombardment, and DislocationsThe growth of defect clusters was observed in situ during irradiation. Various 30 or300keV ions (He+, O+, Mg+, Ar+, Xe+) and 200keV or 1MeV electrons were used toproduce a range of displacement cascade effects. Dislocation loops were created bothinside and outside of the bombarded region, at 870K, under concurrent irradiation with30keV ions and 1MeV electrons. Various phenomena occurred during concurrent

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Al2MgO4 Oxides Al2O3

irradiation with 300keV ions and 200keV electrons. These included the formation ofcavities, the suppression of dislocation loop formation and the preferential formation ofloops at the periphery of the focussed electron beam. The irradiation spectrum effectswere explained in terms of ionizing, displacive and sub-threshold nuclear stoppingpowers, localized energy densities within displacement cascades, and flux distributionswithin the focussed electron beams. It was noted that cation diffusion from the inside tothe outside of the electron beam, which was probably caused by the sub-threshold nuclearand ionizing stopping powers, played an important role in microstructure evolution duringconcurrent irradiation.K.Yasuda, C.Kinoshita, R.Morisaki, H.Abe: Philosophical Magazine A, 1998, 78[3], 583-98

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Al2O3: Ni Grain Boundary DiffusionComposites, which consisted of a Ni3Al matrix that was reinforced with continuous Al2O3fibres, were investigated by means of transmission electron microscopy, scanning electronmicroscopy, X-ray energy-dispersive spectroscopy and electron-probe micro-analysis.The specimens were hot-pressed (1280C, 20MPa, 2h, vacuum and air), and theexperimental results revealed a noticeable diffusion of Ni into the fibres during pressing.The diffusion coefficient of Ni in the fibres was estimated from electron-probe micro-analyses. The estimated coefficients at 1280C were equal to 5 x 10-14cm2/s for volumediffusion and to 6 x 10-9cm2/s for grain-boundary diffusion. On the basis of the results, itwas proposed that Ni diffusion in the fibres involved both volume and grain boundarydiffusion. The diffusion of O in the fibres during hot pressing in air was suggested tooccur via diffusion along grain boundaries as well as through the bulk.W.Hu, P.Karduck, G.Gottstein: Acta Materialia, 1997, 45[11], 4535-45

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Al2O3: Y DiffusionThe diffusion of Y in a-phase alumina was measured, at temperatures ranging from 1150to 1500C, by means of secondary ion mass spectrometry. It was found that the diffusiondata could be described by:

D (m2/s) = 1.2 x 10-10 exp[-295(kJ/mol)/RT]These results were similar to those for Cr diffusion in this material, although the Y3+ ionwas much larger than the Cr3+ ion.E.G.Moya, F.Moya, B.Lesage, M.K.Loudjani, C.Grattepain: Journal of the EuropeanCeramic Society, 1998, 18[6], 591-4

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Al2O3: Grain Boundary DiffusionThermally grown a-phase external scales which were formed on alloys by oxidation inpure O, at temperatures of between 1000 and 1500C, were analyzed. Alloy dopants, suchas Y, Zr, La, Hf and Ti, were found to segregate to the a-phase grain boundaries and to

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Al2O3 Oxides Al2O3

the alloy/scale interface. With increasing oxidation time and temperature, the amount ofsegregant at oxide grain boundaries, near to the gas interface, increased until a criticallevel was reached and precipitates began to grow. This was a result of the outwardtransport of dopants from the alloy, and through the external alumina scale, to the gasinterface. The apparent driving force for dopant diffusion was the O potential gradient inthe growing oxide scale.B.A.Pint, A.J.Garratt-Reed, L.W.Hobbs: Journal of the American Ceramic Society, 1998,81[2], 305-14

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Al2O3: Dislocations, Stacking Faults, and TwinsBy combining diffraction contrast and high-resolution transmission electron microscopicimaging, a study was made of the relationship between the 2 most commonly observedfeatures of shock-deformed sapphire. These were: 1/3<10•0> dislocations, and basaltwins. It was noted that the triple partial dislocation could be identified by the fact thatstacking-fault fringes appeared at the first and second partials, but disappeared after thethird. High-resolution transmission electron microscopic imaging of this dislocationshowed that the partial dislocation was not a co-planar partial, but was separated by onethird of a unit cell along the [00•1] direction. This was the same as the height of the twinstep. The twin-boundary dislocation which was associated with the twin step also had aBurgers vector of 1/3<10•0>. High-resolution transmission electron microscopy anddiffraction contrast imaging suggested that the basal twin might be a mirror twin, with aglide of 1/3<10•0>. The overall conclusion drawn was that the partials were twinningdislocations for the basal twins.S.J.Chen, D.G.Howitt: Philosophical Magazine A, 1998, 78[3], 765-76

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Al2O3: Dislocations, Stacking Faults, and TwinsIt was recalled that 1/3<10•0> partial dislocations played a crucial role in the plasticdeformation of sapphire. During deformation at high temperatures, basal slip(1/3<11•0>{00•1}) required the lowest critical resolved shear stress. The 1/3<11•0>perfect dislocations underwent dissociation (probably restricted to the dislocation core) togive 1/3<10•0> and 1/3<01•0> half-partial dislocations. These partials glided on anelectrically neutral so-called motion plane within a puckered cation array. The 1/3<10•0>partial also acted as the twinning partial when basal twinning occurred at 600 to 1000C.Twinning occurred when a pinned screw partial, sweeping over the same motion plane,made a complete loop of a micro-twin and then cross-slipped onto the next availablemotion plane to start twin-thickening. New transmission electron microscopic evidenceconfirmed several predictions of a new model for basal twinning. Prism-plane slip(1/3<10•0>{12̄•0}) was the preferred slip system at temperatures below about 600C, inspite of the very large Burgers vector (0.822nm) of the <10•0> dislocation. This occurredbecause the latter dislocation dissociated into 3 co-linear 1/3<10•0> partials which wereseparated by 2 relatively low-energy stacking faults. The stacking-fault energy in sapphirewas much lower on prism planes than on basal planes. The motion plane for prism-plane

88

Al2O3 Oxides Al2O3

slip lay between 2 puckered O layers, but also permitted dislocation motion with no netchange.A.H.Heuer, K.P.D.Lagerlöf, J.Castaing: Philosophical Magazine A, 1998, 78[3], 747-63

[446-164-088]

Al2O3: Grain BoundariesThe effect of Zr doping (1000ppm) upon fine-grained alumina was studied by means ofhigh-resolution transmission electron microscopy and field emission scanningtransmission electron microscopy. The high-resolution transmission electron microscopicobservations revealed that the atomic structure of grain-boundary regions was not stronglydistorted in comparison with the surrounding bulk. It was found that there was noamorphous phase at any of the grain boundaries. It was shown, using energy-dispersiveX-ray spectroscopy, that Zr segregated to the grain-boundary regions; with a Zr/Al atomicratio of 1.6%. Normalization and subtraction of matrix spectra from the interface spectrarevealed features which were attributed to heavily misshapen defects at the boundary.K.Kaneko, T.Gemming, I.Tanaka, H.Müllejans: Philosophical Magazine A, 1998, 77[5],1255-72

[446-164-088]

Al2O3: Grain BoundariesHigh-resolution transmission electron microscopy and analytical electron microscopywere carried out on Si-doped sintered a-phase material. High-resolution transmissionelectron microscopy showed that there was no amorphous phase at the grain boundaries.The Si-segregated boundaries were found to be much more sensitive to irradiation damagethan were undoped alumina grain boundaries. Analytical electron microscopy, withenergy dispersive X-ray spectroscopy, revealed significant Si segregation at grainboundaries, and analytical electron microscopy, with electron energy-loss spectroscopy,revealed the existence of 6-fold coordinated Si at the grain boundaries. Theoretical resultswhich were obtained by using the molecular orbital method supported the data whichwere obtained by using electron energy-loss spectroscopy.K.Kaneko, I.Tanaka, M.Yoshiya: Applied Physics Letters, 1998, 72[2], 191-3

[446-164-088]

Al2O3: Grain BoundariesThe microstructure of sol-gel-derived a-alumina, which was doped with 0.6wt%TiO2 andsintered (1450C, 1h), consisted of thin platelets, with (00•1) faces, in a matrix ofequiaxed grains. Short facets at the edges of the platelets developed mainly parallel to the{10•2} planes, while some were parallel to the {11•3} planes. Other edges exhibitedirregular curved boundaries. The basal surfaces of the platelets were coated with thinlayers (0.5 to 6nm) of an amorphous Ti-containing aluminosilicate phase. This was alsopresent at triple-points. No amorphous phase was found on the short faceted boundaries,on curved boundaries at platelet edges, or at the grain boundaries of equiaxed matrixgrains. However, Ti enrichment was observed at all of the boundaries; thus suggestingthat Ti segregation alone did not account for the development of the anisotropicmicrostructure.

89

Al2O3 Oxides Al2O3

A.Kebbede, G.L.Messing, A.H.Carim: Journal of the American Ceramic Society, 1997,80[11], 2814-20

[446-164-089]

Al2O3: Point DefectsA study was made of micro-defects and interfaces, in oxide films on FeAl or NiAlsubstrates, by using variable-energy positron lifetime spectroscopy. Di-vacancies,vacancy clusters and micro-voids were observed in the oxide scales. Their sizes anddistributions were governed by the process which was used to synthesize the oxide film,and were influenced by the composition of the substrate. In the case of oxide/FeAlinterfaces, the positron lifetimes were longer than those for the alumina layer itself; thussuggesting the existence of a greater defect concentration at such sites.J.Xu, B.Somieski, L.D.Hulett, B.A.Pint, P.F.Tortorelli, R.Suzuki, T.Ohdaira: AppliedPhysics Letters, 1997, 71[21], 3165-7

[446-164-089]

Al2O3: Point DefectsThe microstructures of transition aluminas which were prepared by the dehydration ofboehmite were characterized by using transmission electron microscopy. The presence of?-, d- and ?-aluminas was identified by means of selected-area electron diffraction.Modifications that resulted from the re-ordering of Al vacancies on octahedral sites in acubic close-packed O network were detected and were analyzed by using high-resolutiontransmission electron microscopy, combined with image simulation. A closecorrespondence between observed and calculated images confirmed the ordering ofvacant octahedral sites, located on {011} and {011̄) planes, that formed a zig-zagconfiguration along the <010> direction. Two more arrangements of empty octahedralsites, but concentrated on {001} planes, were detected in sintered powder-gelagglomerates. Structural analysis suggested that the modifications were all associatedwith the rearrangement of vacant sites during the phase transformation from ?-alumina tod-alumina to ?-alumina, and were probably driven by configurational entropyminimization.Y.G.Wang, P.M.Bronsveld, J.T.M.De Hosson, B.Djuricic, D.McGarry, S.Pickering:Journal of the American Ceramic Society, 1998, 81[6], 1655-60

[446-164-089]

Al2O3: Point DefectsThe behavior of a single O vacancy in a-phase material was studied by means of super-cell total-energy calculations, using a first-principles method that was based upon density-functional theory. The super-cell model, with 120 atoms in an hexagonal lattice, wassufficiently large to give realistic results for an isolated single vacancy, o . Self-consistentcalculations were performed for each assumed lattice relaxation configuration whichinvolved nearest-neighbor Al atoms and next-nearest neighbor O atoms of the vacancysite. The total-energy data which were obtained were used to construct an energyhypersurface. A theoretical zero-temperature vacancy formation energy of 5.83eV was

90

Al2O3 Oxides Al2O3

deduced. The results revealed a large relaxation of Al atoms, away from the vacancy site,by about 16% of the original Al-o distance. There was a similarly large relaxation of Oatoms away from the vacancy site by about 8% of the original O-o distance. Therelaxation of the neighboring Al atoms exhibited a much weaker energy dependence thandid the O atoms. The O vacancy introduced a deep and doubly-occupied defect level, oran F-center in the gap, and 3 unoccupied defect levels near to the conduction band edge.The positions of the latter were sensitive to the degree of relaxation. The defect-statewave-functions were not so localized, but extended up to the boundary of the super-cell.Defect-induced levels were also found in the valence-band region below the O 2s and theO 2p bands. The case of a singly occupied defect level (an F+ center) was alsoinvestigated. This was done by reducing the total number of electrons in the super-cell,and the background positive charge by one electron, in self-consistent electronic structurecalculations. Optical transitions between occupied and excited states of the F and F+

centers were also investigated and were found to be anisotropic; in agreement with opticaldata.Y.N.Xu, Z.Q.Gu, X.F.Zhong, W.Y.Ching: Physical Review B, 1997, 56[12], 7277-84

[446-164-090]

Al2O3: Point DefectsPolymorphic phase transitions, between alumina structures that were based upon face-centered cubic packings of O, were studied by means of electron diffraction and high-resolution electron microscopy. A new metastable alumina polymorph, with monoclinicsymmetry, was identified in alumina samples which were obtained by plasma sprayingand thermal oxidation of Al. The structure of this phase belonged to the P21/c spacegroup, with a unit cell that contained 32 formula units and had the lattice parameters: a =0.845, b = 1.6, c = 1.264nm, ß = 115º. It was suggested that the new phase evolved fromcubic ?-Al2O3 via the ordering of Al cations on the interstitial sites of the face-centeredcubic O sub-lattice. The transformation occurred via an orthorhombic distortion of thecubic anion structure. Crystallographic domain and inter-domain boundaries in ?- Al2O3were identified and were related to symmetry changes which accompanied the phasetransition.I.Levin, D.G.Brandon: Philosophical Magazine Letters, 1998, 77[2], 117-24

[446-164-090]

Al2O3: Point DefectsThe thermoluminescence of a-phase samples was studied at temperatures ranging from100 to 480K. The crystals had been subjected to X-ray and ultra-violet irradiation at anaverage temperature of about 170K. The thermoluminescence signal was characterized by4 bands which were detected at wavelengths of 300, 418, 696 and 750nm; thus reflectingthe emission of F+ centers, F centers, Cr ionic impurities and Ti ionic impurities,respectively. A profile change in the trap levels was proposed in order to explain thedifference in thermoluminescence response after the exposure of the sapphire crystals toX-ray and ultra-violet radiation. On the basis of the thermoluminescence results for a-Al2O3 which had been subjected to ultra-violet illumination (4.8eV), it was concludedthat

91

Al2O3 Oxides Al2O3

the energy level of the ground state of the F center lay in the band-gap at about 4.5eVabove the top of the valence band, while the F+ center was associated with a fundamentallevel that was lower than that of the F center by about 2eV.M.Ghamnia, C.Jardin: Philosophical Magazine B, 1997, 76[6], 875-85

[446-164-091]

Al2O3/Cu: O InterdiffusionInterface reactions were studied in situ by using a high-temperature X-ray diffractometerwhich was capable of furnishing a diffraction pattern every 1 to 2s. It was found thatCuAlO2 formed at the interface, at temperatures of between 1411 and 1467K, in air. Itsformation obeyed the parabolic rate law. The value of the associated activation energy(185kJ/mol) suggested that diffusion of O through the CuAlO2 controlled its rate offormation.T.Fujimura, S.I.Tanaka: Acta Materialia, 1998, 46[9], 3057-61

[446-164-091]

Al2O3/Ti: InterdiffusionA new scheme was suggested for the prediction of interface reaction products atmetal/ceramic interfaces. This was based upon thermodynamic calculations and diffusionsimulations. Diffusion-controlled reaction and local equilibrium were assumed to exist atthe interface. The thermodynamic state of the interface, before the formation of reactionproducts, was assumed to correspond to metastable equilibrium between the 2 initialphases. In order to determine the boundary compositions, multi-component diffusionsimulation was performed. The driving forces for the formation of all of the other phasescould be calculated under the assumed metastable equilibrium conditions. By selectingthe phase with the highest driving force for formation to be the first-formed interfacereaction product, the order of formation as well as the interface layer sequence could bepredicted. The present scheme was applied to interface reaction between pure Ti andalumina at 1100C. It was predicted that TiAl would always form first at the onset of theinterface reaction, but the stability of the TiAl depended upon the O potential in the Timatrix.B.J.Lee: Acta Materialia, 1997, 45[10], 3993-9

[446-164-091]

Al2O3-MgO: Dislocations and Point DefectsIt was noted that deviations from stoichiometry, in Mg-Al spinels (MgO•nAl2O3), causeda marked decrease in the high-temperature critical resolved shear stress and in the steady-state flow stresses for both {111}<1̄01> and {101}<1̄01> slip. However, Arrhenius plotsgave activation energies and stress exponents which were essentially the same for bothstoichiometric and non-stoichiometric crystals. The dislocations which were observedafter the deformation of non-stoichiometric specimens via {101} slip were ofpredominately edge type, while 60º climb-dissociated dislocations were found inspecimens which had undergone {111} slip. On the other hand, edge and 30º dislocations

92

Al2O3 Oxides B2O3

were found in stoichiometric (n = 1) spinel which had undergone {111} slip. Most of thedislocations in non-stoichiometric crystals were dissociated by climb, although somepartial dislocations were seen to be bounding widely separated glide faults. Somedislocations had segments which were alternately dissociated by a combination of glideand climb. Further analysis of the critical resolved shear stress revealed a linearrelationship, between its logarithm and the test temperature, which was at least as good asthe usual Arrhenius plot. Moreover, the critical resolved shear stress decreased with [Vc]-

2. The concentration, [Vc], of cation vacancies was here given by (n - 1)/[3(3n + 1)].These relationships implied that the critical resolved shear stress was controlled by aPeierls stress which was reduced by kink nucleation at cation vacancies. The activationenergy was slightly lower for {101}<1̄01> slip, so that this system was favored in non-stoichiometric samples at lower temperatures.W.T.Donlon, A.H.Heuer, T.E.Mitchell: Philosophical Magazine A, 1998, 78[3], 615-41

[446-164-092]

AlPO4: Grain BoundariesEvidence was presented for finite and/or imperfect low-angle grain boundaries, and theirassociated long-range stress fields, which formed during the hydrothermal growth ofberlinite crystals on multiple-seed arrays. These defective boundaries originated atangularly misaligned junctions between adjacent seeds, as a result of the failure of somecomponent sets of dislocations to propagate epitaxially into the new growth. Since theylacked a full complement of dislocations, the newly-grown boundaries could be imperfectand could generate long-range stresses; thus leading to fracture and/or plastic flow.R.C.Morris, B.H.T.Chai: Journal of Crystal Growth, 1998, 191, 108-12

[446-164-092]

B2O3-Li2O-NaBr, B2O3-Li2O-NaCl, B2O3-Na2O-NaBr, B2O3-Na2O-NaCl: IonicConductionThe structures of these fast-ion conducting glasses were studied by using neutrondiffraction techniques and Monte Carlo simulations. It was found that the short-rangestructure of the B-O network was almost unchanged upon increasing the dopant saltconcentration, and was independent of the nature of the dopant salt. On the other hand,the intermediate-range order of the B-O network decreased significantly with increasingdopant salt concentration. The Na borate glasses tended to be slightly more ordered thanthe corresponding Li borate glasses. The differences were attributed to the fact that theLi-borate glasses consisted of a disordered random mixture of many different types ofborate configuration, while the Na borate glasses were built up of randomly distributeddiborate groups, as suggested by previous nuclear magnetic resonance results. However,the Monte Carlo simulations of the most highly LiCl- and NaCl-doped glasses showedthat large density fluctuations occurred within the B-O network. The voids were of widelydiffering sizes and geometries. The present results demonstrated that the intermediate-range order of the LiCl- and NaCl-doped glasses was significantly different to thatreported for the analogous AgI-doped glasses. In the latter glasses, the B-O networkformed a more ordered chain-like structure; with the salt ions cross-linking chains.

93

B2O3 Oxides (Ba,La)(Mg,Nb)O3

J.Swenson, L.Börjesson, W.S.Howells: Physical Review B, 1998, 57[21], 13514-26[446-164-093]

Ba(Bi,La)O3: O PermeationThe O permeability and electrical conductivity of BaBi1-xLaxO3 perovskite-like solidsolutions, where x was equal to 0, 0.2 or 0.4, were found to decrease with the La content.The O transport through dense ceramic membranes, with thicknesses ranging from 0.6 to1.2mm, was shown to be limited both by the bulk ionic conductivity and by the surfaceexchange rate.A.A.Yaremchenko, V.V.Kharton, A.P.Viskup, E.N.Naumovich, V.V.Samokhval:Materials Research Bulletin, 1998, 33[7], 1027-33

[446-164-093]

Ba2In2O5: Point DefectsAtomistic simulation techniques were used to investigate crystal structures and defectenergetics. An interatomic potential model was developed which reproduced theorthorhombic brownmillerite structure, with alternating layers of O octahedra andtetrahedra. Defect calculations showed that the most energetically favorable intrinsicdefects were of Frenkel type, with an O vacancy on the O(1) site and an O ion in aninterstitial position in the tetrahedral layer. The predicted formation energies of electronicdefects suggested that this material would oxidize with the formation of positive holeswhich then contributed to the observed electronic conductivity. An analysis of thepossible O-ion migration pathways suggested that the energy barriers to migration werelowest between the equatorial sites of O octahedra in the [001] direction.C.A.J.Fisher, M.S.Islam, R.J.Brook: Journal of Solid State Chemistry, 1997, 128[1], 137-41

[446-164-093]

(Ba,La)(Mg,Nb)O3: Domain BoundariesA microstructural study was made of the interfacial boundaries of 1:1 and 1:2 ordereddomains in (Ba0.9La0.1)(Mg0.37Nb0.63)O3 by using high-resolution transmission electronmicroscopy and X-ray diffractometry. Both 1:2 and 1:1 ordered domains were found tocoexist in a fully-ordered single grain. Each ordered domain occupied its own region, andthe interfaces were atomically sharp and coherent. The wavelength of the superlatticemodulation was about 0.47nm in the 1:1 ordered domains, and about 0.71nm in the 1:2ordered domains. The transition from the 1:2 ordered region to the 1:1 ordered region wasclearly seen at the interfaces. These observations supported previously proposedstructural models.H.J.Lee, H.Ryu, H.M.Park, J.H.Paik, S.Nahm, J.D.Byun: Journal of the AmericanCeramic Society, 1998, 81[6], 1685-8

[446-164-093]

94

Ba0.5Sr0.5TiO3 Oxides BaTiO3

Ba0.5Sr0.5TiO3: Point DefectsThe current density was measured as a function of time in thin (less than 100nm)capacitors with Pt electrodes. The current density curve exhibited a peak ahead of theonset of resistance degradation. The peak position on the time axis varied with the appliedvoltage and temperature. The data were explained in terms of space-charge limitedcurrent transients, and the measured current was identified as being an ionic currentwhich was associated with O vacancies. By means of space-charge limitation analysis, themobility of the O vacancies was measured as a function of temperature. The mobility, asdeduced from current measurements, was shown to be compatible with the Einsteinrelationship for mobility and diffusivity. The ionic current that was associated with Ovacancies was shown to be an important component of the measured current in thesesamples.S.Zafar, R.E.Jones, B.Jiang, B.White, P.Chu, D.Taylor, S.Gillespie: Applied PhysicsLetters, 1998, 73[2], 175-7

[446-164-094]

BaTiO3: Dislocations and TwinsThe structure and configuration of {111} twins, in polycrystalline films which had beenprepared by laser ablation deposition, were investigated by means of high-resolutiontransmission electron microscopy. The number density of twins in the films was high, andthe thickness of the twin lamellae was of the order of nm. In some grains, ordering of thenano-twins gave rise to a local structure which resembled that of the hexagonal phase.These twin lamellae usually extended so as to penetrate the whole grain. Twin lamellaewhich terminated within the grain were also found. Some of them were very small (lessthan 1nm in thickness) and were accompanied by partial dislocations with a Burgersvector of (a/3)<111>. These small twins were considered to represent an early stage oftwin formation. It was deduced that the nucleation of a {111} twin via the dissociation ofan a<111> edge dislocation was one of the mechanisms of twin formation in thesepolycrystalline films.C.L.Jia, K.Urban, M.Mertin, S.Hoffmann, R.Waser: Philosophical Magazine A, 1998,77[4], 923-39

[446-164-094]

BaTiO3: TwinsThe formation of (111) twin lamellae in thin films was studied by using transmissionelectron microscopy. Voids in the thin films were related to the (111) twin lamellae. Thevoids were formed due to the incomplete coalescence of islands, and it was found that(111) twin lamellae appeared in order to accommodate the small misfit between growingislands during contact. In the case of planar defects, the formation of (111) twin lamellaewas attributed to the maintenance of a TiO6 octahedron and its relationship to anhexagonal BaTiO3 structure. The very narrow width of the twin lamellae was attributed tothe contact of 2 growing islands, and to the very small grain size of the thin films.

95

BaTiO3 Oxides Bi12GeO20

J.W.Jang, W.J.Cho, J.H.Lee, S.S.Choi, T.S.Hahn: Japanese Journal of Applied Physics -1, 1997, 36[11], 6942-5

[446-164-095]

BaTiO3: TwinsNanocrystals were investigated by means of high-resolution transmission electronmicroscopy. The nanocrystals were found to exhibit a cubic-tetragonal structure, with(111) planar defects. The resultant images, apart from cubic-tetragonal material, revealedcoherent (111) twins which often formed lamellae that consisted of 3 to 5 atomic layers,periodic images of a 0.70nm x 0.49nm rectangle, and incoherent (111) planar images. Therectangular lattice and the incoherent (111) planar images originated from superpositionof the matrix and the (111) twin. Peaks which were attributed to the hexagonal phase, inthe electron diffraction pattern, could be explained in terms of the existence of nano-sizedtwins in the cubic tetragonal matrix.E.Hamada, W.S.Cho, K.Takayanagi: Philosophical Magazine A, 1998, 77[5], 1301-8

[446-164-095]

(Bi,Er)2O3: O PermeationPermeation experiments were carried out on dense mixed-conducting membranes (0.2 to2mm) of Bi1.5Er0.5O3, which contained 10, 27.8 or 40vol%Ag, at temperatures rangingfrom 873 to 993K, under O partial pressures ranging from 10-3.5 to 1bar. It was found thatthe O flux increased with increasing Ag content. In the case of samples with a non-percolative Ag phase (10 or 27.8vol%Ag), an increased O flux relative to that of the pureoxide was attributed to the more rapid kinetics of surface O exchange in the presence ofAg. The percolative nature of the Ag phase in the 40%Ag material enhanced theambipolar diffusion of O and electrons. High O fluxes, of the order of 0.25mol/m2s at873K, were detected for this composition. The activation energy for O permeation attemperatures ranging from 848 to 1003K was between 85 and 95kJ/mol for specimenswithout percolation, and was equal to 115kJ/mol for the sample with 40vol%Ag. Thisdifference was attributed to a change in the rate-limiting process.J.E.ten Elshof, N.Q.Nguyen, M.W.den Otter, H.J.M.Bouwmeester: Journal of theElectrochemical Society, 1997, 144[12], 4361-6

[446-164-095]

Bi12GeO20, Bi12SiO20: Point DefectsChanges in the defect structure, which were caused by heat treatment and laserirradiation, were studied by means of cathodoluminescence scanning electronmicroscopy. The results were compared with those for untreated and electron-irradiatedsamples. It was found that annealing of the Bi12GeO20 samples led to the appearance of anew luminescence band at about 390nm. The centers which were responsible for thisband decorated the deformation slip-bands in quenched samples, as observed incathodoluminescence images. An emission which was observed in Bi12SiO20 over thesame spectral range was quenched during annealing. An annealing-induced reduction of

96

Bi12GeO20 Oxides Bi2O3

Bi ions to metallic Bi was suggested to be related to the quenching of a band at 640nm inuntreated samples.A.Cremades, J.Piqueras, A.Remón, J.A.García, M.T.Santos, E.Diéguez: Journal ofApplied Physics, 1998, 83[12], 7948-52

[446-164-096]

Table 1Ionic Conductivity of Bi2V0.9Cu0.1O5.35

Theoretical Density (%) Temperature (C) s (S/cm) E (eV)85.9 200 0.20 x 10-3 0.6385.9 400 1.46 x 10-2 0.6385.9 600 1.05 x 10-1 0.4893.7 200 0.30 x 10-3 0.6293.7 400 1.68 x 10-2 0.6293.7 600 9.97 x 10-2 0.4692.1 200 0.70 x 10-3 0.5392.1 400 1.96 x 10-2 0.5392.1 600 1.13 x 10-1 0.4787.9 200 0.30 x 10-3 0.5987.9 400 1.46 x 10-2 0.5987.9 600 9.25 x 10-2 0.4792.9 200 0.20 x 10-3 0.7092.9 400 1.70 x 10-2 0.7092.9 600 9.74 x 10-2 0.5695.6 200 0.40 x 10-3 0.5895.6 400 1.93 x 10-2 0.5895.6 600 1.11 x 10-1 0.4697.3 200 0.30 x 10-3 0.5197.3 400 1.61 x 10-2 0.5197.3 600 8.62 x 10-2 0.5196.8 200 0.30 x 10-3 0.6096.8 400 1.87 x 10-2 0.6096.8 600 1.13 x 10-1 0.52

Bi2O3-Gd2O3, Bi2O3-Y2O3: Ionic ConductionIn order to investigate the conductivity, below 400C, in high-temperature oxide-ionconductors, impedance spectra were determined at frequencies ranging from 20Hz to

97

Bi2O3 Oxides CeO2

1MHz. The conductivity relaxation phenomena which were observed were explained interms of the electric modulus formula. The relaxation frequencies were found to exhibitArrhenius-type behavior, and the activation energies were in good agreement with thoseobtained by means of high-temperature conductivity measurements.S.Takai, N.Kohno, T.Esaka: Materials Research Bulletin, 1998, 33[6], 945-53

[446-164-097]

3 Bi2(V,Cu)O5.35: Ionic ConductionIt was recalled that the Cu-substituted vanadate, Bi2V0.9Cu0.1O5.35 (BICUVOX), exhibitedhigh O-ion conductivities at low temperatures. The ionic conductivity (table 1) at 300C(about 0.001S/cm) was 50 to 100 times greater than that of any other solid electrolytewithin this temperature range. The preparation of sinterable powder by means of solid-state reaction and conventional ceramic powder processing techniques was reported here.The results indicated that densities which were greater than 95% of the theoretical valuecould be obtained by sintering (800C, 10 to 20h).S.P.Simner, D.Suarez-Sandoval, J.D.Mackenzie, B.Dunn: Journal of the AmericanCeramic Society, 1997, 80[10], 2563-8

[446-164-097]

Calcium Hexaluminate/Al2O3: DislocationsThe epitactic growth of calcium hexaluminate (CA6) on the basal plane of alumina wasstudied. Two distinctly different orientational relationships were found. The first could bedescribed as being perfect hexagon-on-hexagon epitaxy with (11•0)CA6||(10•0)alumina and(00•1)alumina||(00•1)CA6. This relationship was associated with a 1.1% lattice misfit in thesubstrate plane. The misfit was accommodated at the interface by an hexagonal networkof misfit dislocations. In the case of the second relationship, the CA6 grains were rotatedaround the substrate normal, and could be explained by the coincident-site lattice modelfor phase boundaries.M.P.Mallamaci, K.B.Sartain, C.B.Carter: Philosophical Magazine A, 1998, 77[3], 561-75

[446-164-097]

CeO2: Surface DefectsThe (111) surface structure of non-stoichiometric material was studied by means ofscanning tunnelling microscopy. By using extremely low tunnelling currents, it waspossible to obtain the first reported atomically resolved images of this oxide. Scanningtunnelling microscopic imaging was possible at a sample bias voltage of between -2 and -3.5V. Upon comparing this with the band-structure of ceria, it was claimed that the maincontribution to image contrast resulted from O in the uppermost layer. The predominantdefect type on the surface at room temperature was triangular, and comprised three Ovacancies. Scanning tunnelling microscopy, at a substrate temperature of 500C, revealedan alignment of O vacancies on the surface. The defect shapes were in qualitativeagreement with previous energy calculations.

98

CeO2 Oxides Cr2O3

H.Nörenberg, G.A.D.Briggs: Physical Review Letters, 1997, 79[21], 4222-5[446-164-098]

Table 2Bulk Diffusivity of Y in Cr2O3

Temperature (C) D (cm2/s)800 2.0 x 10-18

850 1.7 x 10-18

900 3.2 x 10-18

950 7.9 x 10-18

1000 2.4 x 10-17

CoLixO2: Point DefectsA first-principles technique was presented for predicting ordered vacancy ground states,intercalation voltage profiles, and voltage-temperature phase diagrams. Its application tothe present system correctly predicted the observed ordered-vacancy phases. Theexistence of additional ordered phases was also predicted.C.Wolverton, A.Zunger: Physical Review Letters, 1998, 81[3], 606-9

[446-164-098]

CoO: Cr DiffusionSurface segregation-induced Cr depth-profiles were determined for Cr-doped singlecrystals at various temperatures (1373 to 1673K) and O partial pressures (102 to 105Pa). Itwas shown that the shape of the depth profiles depended upon both the annealingconditions and upon the subsequent cooling procedure. It was observed that during theslow cooling of specimens, from the high equilibrium segregation temperature to roomtemperature, there was a substantial change in the concentration profile. This involved Crdiffusion from the surface to the bulk, and resulted in its depletion. It was also shown thatthe rate of the de-segregation which resulted in a decrease in surface concentration wasslower than the lattice transport of Cr. It was concluded that the de-segregation was rate-controlled by the decomposition of a spinel-type bi-dimensional surface structure whichformed at high temperatures due to Cr segregation.A.Bernasik, J.Nowotny, S.Scherrer, S.Weber: Journal of the American Ceramic Society,1997, 80[2], 349-56

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3 Cr2O3: Y DiffusionThe diffusivity of Y in chromia scales which had grown on a Ni-30wt%Cr alloy duringhigh-temperature oxidation was determined in air at temperatures ranging from 800 to1000C. A thin Y-containing film was produced at the oxide surface by depositing a liquidsolution of YCl3. Following the diffusion treatment, penetration profiles were measured

99

Cr2O3 Oxides Cu3Ba2NdO7

by means of secondary-ion mass spectrometry. The apparent diffusion coefficients werecalculated by using thin-film solutions to Fick's equation. The bulk and grain-boundarydiffusion coefficients were obtained by using a new model which took account of oxidesurface roughness and of the relationship between bulk, grain-boundary and apparentdiffusion coefficients. The results (tables 2 and 3) showed that Y diffusion in Cr2O3scales obeyed the relationships:

bulk: D (cm2/s) = 1.2 x 10-11exp[-144(kJ/mol)/RT]grain boundaries: D (cm2/s) = 2.56 x 10-4exp[-190(kJ/mol)/RT]

The difference between the bulk and grain-boundary diffusion coefficients was of theorder of 5 orders of magnitude within the temperature range studied. The bulk and grain-boundary diffusion coefficients for Y were both lower than the equivalent values for Oand Cr diffusion in the same oxide scale.J.Li, M.K.Loudjani, B.Lesage, A.M.Huntz: Philosophical Magazine A, 1997, 76[4], 857-69

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Table 3Grain-Boundary Diffusivity of Y in Cr2O3


850 5.4 x 10-13

900 6.7 x 10-13

950 1.8 x 10-12

1000 4.6 x 10-12

Cu3Ba2GdO6: O Diffusion and Point DefectsThe diffusion of O in this perovskite-related non-stoichiometric oxide was investigated insitu by using thin films and high-temperature resistometry. Apart from a step change inthe vicinity of the orthorhombic-tetragonal transition, the diffusivity exhibited aconcentration-dependence for O stoichiometry deviations ranging from 0.3 to 0.8. Theactivation energy was constant (1.2eV) over the whole interval. It was proposed that Odiffusion occurred via a vacancy mechanism which involved the O atoms at the extremesof the -O-Cu-O- chains in the basal plane.V.Dediu, F.C.Matacotta: Physical Review B, 1998, 57[13], 7514-7

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Cu3Ba2NdO7: TwinsAn in situ study was made, by means of high-temperature polarized optical microscopy,of twin formation during the annealing of single crystals. The latter were grown under a21%O partial pressure atmosphere by using a top-seeded solution-growth method. Thecut samples were heated to 700C in a pure Ar gas flow, followed by cooling (10C-steps)from

100

Cu3Ba2NdO7 Oxides Cu3Ba2YO6.5

700C (in an O gas flow) until a twin formed. Its formation, which involved a tetragonal-orthorhombic phase transition, was observed at 660 and 620C in a 100 and a 21%Opartial-pressure gas flow, respectively.A.Oka, S.Koyama, H.Kutami, Y.Shiohara: Applied Superconductivity, 1997, 5[1-6], 79-86

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Cu3Ba2YOx: TwinsThe characterization of a twin-free orthorhombic crystal was described. A tetragonalsingle crystal was grown by crystal pulling, and was annealed in an O atmosphere underuniaxial compression in order to prevent twin formation. The resultant crystal had a twin-free structure. The critical temperature of the crystal was about 91.8K and the criticalcurrent was only some 20% of that of the twinned crystal at 1T.H.Kutami, Y.Yamada, S.Koyama, J.G.Wen, T.Egi, Z.Nakagawa, Y.Shiohara: JapaneseJournal of Applied Physics - 1, 1997, 36[2], 674-5

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Cu3Ba2YO6: Point DefectsBy using the Green's function method, the localized state energy was calculated for anelectron which was trapped in the chain-O vacancy. For reasonable parameter values, itlay above the chain-O and plane-O band top. The formation of a stable chain F+ centerwas proposed, as in the mechanism of photo-induced conductivity.P.Rubin, N.Kristoffel: Journal of Physics - Condensed Matter, 1998, 10[8], L127-9

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Cu3Ba2YO6: TwinsTwin boundaries in polycrystalline samples were studied by means of atom-probe field-ion microscopy. Depletion of O was observed at a twin boundary in Cu3Ba2YO6.6. Thewidth of the depleted region was 6 to 7nm. This agreed well with the O depletion widthwhich was calculated by using the Jou-Washburn O-depleted twin boundary model.However, a twin boundary in Cu3Ba2YO6.9 did not exhibit any O depletion. This workprovided direct evidence for O depletion at twin boundaries in polycrystallineCu3Ba2YO7 ceramics.Q.H.Hu, K.Stiller, E.Olsson, H.O.Andrén, P.Berastegui, L.G.Johansson: Physical ReviewB, 1997, 56[18], 11997-2003

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Cu3Ba2YO6.5: TwinsThe remanent and in-field magnetic field distributions of 2 crystals, one with a singletwin boundary and one completely untwinned, were observed by using a non-destructivemagneto-optical technique. Geometrical barrier effects upon the flux penetration of bothsamples were found. Two distinct penetrating field patterns were observed, whichresulted from the differing pinning characteristics of the samples. A concentration of

101

Cu3Ba2YO6.5 Oxides Cu3Ba2YO7

penetrating vortices in the center of one sample indicated an extremely low degree ofpinning. A comparison was made between the observed flux penetration profiles of bothsamples, and those calculated by using a theoretical model for a thin strip superconductor.A single twin boundary in the extremely low-pinning crystal was observed to be a barrierto transverse vortex motion.M.W.Gardner, S.A.Govorkov, R.Liang, D.A.Bonn, J.F.Carolan, W.N.Hardy: Journal ofApplied Physics, 1998, 83[7], 3714-9

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Figure 4Diffusivity of O in Cu3Ba2YO7 Thin Films

Cu3Ba2YO7: O DiffusionThe intensities of the 63Cu nuclear quadrupole resonance and quadrupolar-satellitenuclear magnetic resonance at high temperatures were determined for the planar Cu(2)sites. It was observed that the spin-echo intensities decreased markedly between room

1.0E-15

1.0E-14

1.0E-13

1.0E-12

18 19 20 21

E = 1.10eV

1/kT(eV)

D (c

m2 /s

)

102

Cu3Ba2YO7 Oxides Cu3Ba2YO7

temperature and 500K, and that the onset of intensity-loss occurred at lower temperaturesin more defective samples. Measurements of the spin-echo decay rates revealed that theintensity loss was associated with a loss of phase coherence which exceeded that whichwas expected for spin-spin and spin-lattice relaxation. By using simple models, theseeffects could be attributed to changes, in the local atomic environment, which werepresumed to be caused by the motion of O atoms in the Cu(1) layer. These motionsoccurred over a time-scale which was of the order of 10µs at 500K. Their probability wasenhanced, at lower temperatures, by the presence of structural defects.S.P.Klein, R.P.Wang, A.W.Sleight, W.W.Warren: Physical Review B, 1997, 56[10],6335-42

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y Cu3Ba2YO7: O DiffusionBy using isothermal 4-point electrical resistivity measurements, chemical diffusioncoefficients were determined for the in-diffusion and out-diffusion of O to and fromepitaxial c-axis oriented films, at temperatures ranging from 500 to 650K (figure 4).Temporal changes in resistivity, at constant O partial pressure, were monitored followingsmall upward or downward steps in temperature. The same coefficient, 2 x 10-14cm2/s,was found at 600K for both the in-diffusion and out-diffusion of O. The temperaturedependence of the diffusivity could be closely fitted by assuming an Arrhenius law. Thisyielded the same activation energy (1.1eV) for in-diffusion and out-diffusion.S.Kittelberger, O.M.Stoll, R.P.Huebener: Superconductor Science and Technology, 1998,11[8], 744-50

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Cu3Ba2YO7: O DiffusionA study was made of c-axis oriented films with superconducting transition temperaturesof up to 89K. These were prepared by ion beam sputter-deposition onto monocrystalline(100) MgO or (100) SrTiO3 substrates. The variation in the O concentration duringannealing in vacuum or O was measured, in situ and in real time, by means ofspectroscopic ellipsometry. Changes in O concentration of less than 1% could beresolved. Measurements on films with thicknesses ranging from 15 to 100nm were usedto elucidate the mechanisms of O out-diffusion and in-diffusion. It was found that the Oout-diffusion rate depended upon the layer thickness; thus confirming the operation of abulk diffusion mechanism. The O in-diffusion was relatively independent of the layerthickness; thus suggesting that a surface cum grain-boundary diffusion mechanismpredominated in this case. It was noted that H2O enhanced O out-diffusion, while the in-diffusion remained unchanged; or even decreased.A.Michaelis, E.A.Irene, O.Auciello, A.R.Krauss: Journal of Applied Physics, 1998,83[12], 7736-43

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103


Cu3Ba2YO7: O DiffusionIt was recalled that diffusion-based processes such as sintering were substantiallyenhanced in the presence of a microwave field, and it was suggested here that thediffusion of O into the present material might also be affected in this way. Theoxygenation of melt-processed material, using conventional and microwave-assistedheating, was therefore compared. The diffusivity at 400 and 450C was shown to beenhanced by about 30% by applying a high-frequency microwave field. The equilibriumvalue of the O content was not affected by using the field.A.T.Rowley, R.Wroe, D.Vázquez-Navarro, W.Lo, D.A.Cardwell: Journal of MaterialsScience, 1997, 32[5], 4541-7

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3 Cu3Ba2YO7: O DiffusionThe diffusion of O in undoped and Ag-doped single crystals was investigated byperforming isothermal resistivity measurements at temperatures ranging from 550 to750C. It was found that the diffusion coefficients for Ag-doped crystals were an order ofmagnitude lower than those for undoped ones (table 4). Furthermore, the activationenergy for O diffusion was found to be considerably higher for Ag-doped crystals.Further analysis of the data indicated that doping with Ag stabilized the phase at higher Ocontents.D.K.Aswal, S.K.Gupta, P.K.Mishra, V.C.Sahni: Superconductor Science andTechnology, 1998, 11[7], 631-6

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Table 4Diffusivity of O in Cu3Ba2YO7

Sample Experiment Temperature (C) D (cm2/s)undoped in-diffusion 550 2.77 x 10-6

undoped in-diffusion 600 5.20 x 10-6

undoped in-diffusion 650 9.24 x 10-6

undoped out-diffusion 550 1.66 x 10-6



Ag-doped in-diffusion 650 2.89 x 10-7



Ag-doped out-diffusion 650 1.01 x 10-7



104


Cu3Ba2YO7: Electron Irradiation, Point Defects, and Stacking FaultsSamples of Ca-doped material were investigated by using transmission electronmicroscopy. Irradiation led to the formation of stacking faults which lay on {001} planes,and extended in <100> and <010> directions. It was proposed that vacancies formedunder the electron irradiation, and that relaxation then proceeded via the coalescence of 2portions of crystal which were separated by clustered vacancies. Slippage of the bondingbetween two BaO layers could then occur with a displacement of 1/6[301] or 1/6[031].Stacking faults could therefore form which had either of these displacements. Thestacking faults which were formed in this way clearly lay on {001} planes, and seemed tostretch along their longitudinal direction during irradiation. Formation of the stackingfaults generally began from the edge, and propagated towards the interior of the parentgrain.C.C.Lam, G.J.Shen, C.L.Fu, K.C.Hung: Superconductor Science and Technology, 1997,10[8], 807-12

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Cu3Ba2YO7: Dislocations and Grain BoundariesThe dependence of the critical current density upon the orientation of the magnetic fieldwas studied in bicrystalline films. When the field was rotated in the plane of the grainboundary, the critical current density exhibited a maximum, when the orientation of themagnetic field was close to the c-axis, in fields of up to 7T. The same measurements,when performed within the interior of the grains, revealed no maxima. The pinning forceat the maximum varied as the square root of the magnetic field strength. This wasconsistent with vortex-pinning by a dense planar distribution of line pinning-points. It wasconcluded that these results were a clear demonstration of flux-pinning by dislocations inlow-angle grain boundaries.A.Díaz, L.Mechin, P.Berghuis, J.E.Evetts: Physical Review Letters, 1998, 80[17], 3855-8

[446-164-104]

Cu3Ba2YO7: Dislocations and Stacking FaultsThe microstructures of Cu3Ba2YO7-Y2BaCuYO5 melt-textured composites which hadbeen deformed into the secondary and tertiary creep regimes were investigated by meansof transmission electron microscopy. A high density of Y2BaCuYO5 precipitates playedan important role as pinning sites for gliding dislocations. In the secondary regime,trapped dislocations dissociated to leave a stacking fault with a displacement vector of[1/2-d0(1/3)]. A second stacking fault, 1/6<301>, tended to be associated with the formerstacking fault. In this stage, deformation was dominated by diffusion between precipitatesthat were interconnected by trapped dislocations. In the tertiary stage, dislocationmultiplication was the main factor which controlled the microstructure. This wascharacterized by a marked increase in the density of perfect dislocations with Burgersvectors of <100> and <110>. It was noted that oxygenation, performed at 450C, led to amarked modification of the as-deformed microstructure.

105


N.Vilalta, F.Sandiumenge, J.Rabier, M.F.Denanot, X.Obradors: Philosophical MagazineA, 1997, 76[4], 837-55

[446-164-105]

Cu3Ba2YO7: Grain BoundariesA study was made of high-current samples which had been processed by means of melt-texturing, using the liquid-phase removal method. The misorientation (kink angle) of thea-b planes across each boundary was measured. The results indicated that high-anglegrain boundaries predominated in the samples. The high current transport capacity of thebulk polycrystalline samples indicated that the liquid-phase removal method was capableof producing good grain boundary links which permitted the passage of a high currents.S.Sathyamurthy, A.S.Parikh, K.Salama: Superconductor Science and Technology, 1997,10[9], 651-6

[446-164-105]

Cu3Ba2YO7: Grain BoundariesA method was presented for the separate study of the electrical transport properties of thegrain boundaries which were formed at the top and bottom edges of step-edge Josephsonjunctions. Step-edge junctions were prepared on (100) LaAlO3 steps by using tilted Ar ionmilling to define the electrodes and micro-bridges. Due to the shadowing effect of thestep, a continuous Cu3Ba2YO7 stripe remained along, and at the bottom of, the step onboth sides of a micro-bridge. It was found that the top grain boundary was responsible forthe weak-link behavior of the step-edge junctions. The transport properties were related tothe differing microstructural properties of the 2 grain boundaries which formed at theedges of the step.F.Lombardi, Z.G.Ivanov, G.M.Fischer, E.Olsson, T.Claeson: Applied Physics Letters,1998, 72[2], 249-51

[446-164-105]

Cu3Ba2YO7: Planar DefectsElectron-phonon interaction was considered with regard to the positions of atoms in theunit cell. It was found that the strength of the interaction depended upon the number ofCu-O layers per unit cell only when defects were present. The transition temperaturetherefore depended upon the variable number, n, of Cu-O layers per unit cell. The largerthe number of layers, the larger was the transition temperature. It was found that thelattice waves could be treated by using spherical Bessel functions, and that the electron-phonon interaction oscillated as a function of the number of Cu-O planar layers. Thetransition temperature of the superconductor then oscillated as a function of n. It had ahigh value at n = 3 and lower value at n = 4. It was predicted that the next maximumvalue would occur at n = 7. The predicted variation, in transition temperature as afunction of n, was in reasonable agreement with experimental data.

106

Cu3Ba2YO7 Oxides (Fe,Cr)2CuO4

N.M.Krishna, K.N.Shrivastava: Superconductor Science and Technology, 1997, 10[5],278-83

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Cu3Ba2YO7: Point Defects and TwinsA mosaic of non-twinned single crystals was studied by means of muon spin rotation. Theflux-line lattice in these crystals was found to be pinned. This pinning then led to disorderin the lattice. The average relative breathing mode lattice displacement was found to beabout 0.8%. This value was comparable to that found for twinned samples. It wassuggested that the origin of the pinning might be O vacancies.W.J.Kossler, A.D.Goonewardene, A.J.Greer, D.L.Williams, E.Koster, D.R.Harshman,J.Z.Liu, R.N.Shelton: Physical Review B, 1997, 56[5], 2376-8

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Cu3Ba2YO7/SrTiO3: TwinsThe twinning in thin films which had been grown onto vicinal SrTiO3 (001) substrates bymeans of pulsed laser deposition was studied, as a function of the vicinal mis-cut angle,by using X-ray diffraction; involving the 2-dimensional q-scan technique with a 4-circlediffractometer. The results revealed a strong correlation between the mis-cut of theSrTiO3 (001) substrate, and the strain which occurred at the substrate/film interface.Depending upon the proper choice of mis-cut orientation and angle, one of the 2 twinsystems could be suppressed and about 70% untwinned growth could be achieved.J.Brötz, H.Fuess, T.Haage, J.Zegenhagen: Physical Review B, 1998, 57[6], 3679-82

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Cu2Y2O5: Stacking FaultsThe crystal structure of Cu2Y2O5 was confirmed, using transmission electron microscopy,to be orthorhombic (Pna21, a = 1.08, b = 0.35, c = 1.23nm, V = 0.46nm3). Planar defectswere observed to lie parallel to the a-b plane, with a displacement vector of ½[100], andarose from the movement of O atoms within the structure. These defects arose fromdisturbances of the Y-O bonding during formation of the crystal; that is, from a stackingfault. The density of the planar defects was further studied by varying the Ca dopantcontent from 0.05 to 0.4. It was found that there was no direct relationship between theplanar defect density and the Ca concentration.C.N.Feng, D.R.Lovett: Journal of Physics - Condensed Matter, 1998, 10[16], 3497-507

[446-164-106]

(Fe,Cr)2CuO4: DiffusionThe Mössbauer spectra and X-ray diffraction patterns of CuCrxFe2-xO4, where x rangedfrom 0.1 to 0.8, were studied with especial regard to the atomic distribution and Debyetemperature. At room temperature, the crystal structure was found to be a cubic spinel forx-values of 0.2 to 0.8, and was tetragonal for an x-value of 0.1. The increasing intensityratio of tetrahedral to octahedral patterns, with increasing temperature at temperatures

107

(Fe,Cr)2CuO4 Oxides Feldspar

below 250K, was explained in terms of large differences in the Debye temperature oftetrahedral and octahedral sites. A decreasing intensity ratio at high temperatures wasattributed to the migration of Fe ions from tetrahedral to octahedral sites, which wasfound to increase sharply with increasing temperature and Cr content. Even at 150K,atomic migration began for an x-value of 0.2 and increased - with increasing x-value - tosuch an extent that about 10% of the Fe3+ ions at tetrahedral sites migrated to theoctahedral sites when x was equal to 0.6. The atomic migration increased markedly withincreasing temperature, and about 50% of the Fe3+ ions at tetrahedral sites migrated at300K when x was equal to 0.5.Y.K.Kim: Japanese Journal of Applied Physics - 1, 1997, 36[10], 6339-43

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Figure 5Diffusivity of O in Ni1-xFe2+xO4 Ferrites

Feldspar: TwinsAtomic force microscopy of a (001) cleavage surface revealed exsolution lamellae, wave-like (001) surfaces on Na-rich lamellae, and surface steps with heights of about 0.6 or0.3nm. The wave-like shape of the (001) surfaces of albite twin domains was traced to thesurface relaxation of periodic twin domains. Both twin boundary and twin composition

1.0E-19

1.0E-18

1.0E-17

1.0E-16

1.0E-15

9 10 11 12 13 14 15 16

x = 0.1x = 0.2x = 0.4

104/T(K)

D (c

m2 /s

)

108

Feldspar Oxides Fe2(Ni,Zn,Cu)O4

planes were parallel to (010). Three-dimensional atomic force microscopic imagesrevealed surface height differences between Na-rich and K-rich lamellae in some areas.This indicated that the boundaries were semi-coherent along the c-axis.H.Xu: Materials Research Bulletin, 1997, 32[9], 1221-7

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y (Fe,Ni)2O4: O Diffusion and Point DefectsThe O tracer diffusivity in polycrystalline samples of Ni1-xFe2+xO4, where x was equal to0.1, 0.2 or 0.4, was measured at temperatures ranging from 450 to 700C. The number of18O atoms which diffused to the sample was determined by exploiting the 18O(p,a)15Nnuclear reaction. It was established that O diffusion, for x = 0.2, was characterized by arelatively low activation energy of 0.89eV (figure 5). The corresponding pre-exponentialfactor was also low (10-13cm2/s). The results suggested that O diffusion in these ferritesoccurred via structural vacancies with a concentration of between 10-9 and 10-11.V.B.Fetisov, G.A.Kozhina, A.J.Fishman, T.E.Kurennykh, V.B.Vykhodets: Journal ofApplied Physics, 1998, 83[11], 6876-8

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FeO: Ca DiffusionMeasurements were made of Ca impurity diffusion in single crystals, at temperaturesranging from 800 to 1150C. It was found that the results could be described by:

D (cm2/s) = 0.51R1.12exp[-{166.7 + 7.7(R-1)}(kJ/mol)/RT]where R was the CO2/CO partial pressure ratio. The activation energies ranged from 164,for R = 0.7, to 178kJ/mol for R = 2.5. Upon assuming that the diffusion mechanism wasthe same as that for self-diffusion, the activation enthalpies could be deduced to be equalto 124kJ/mol for R = 0.7 and to 141kJ/mol for R = 2.5. The difference (38.5kJ/mol) in theactivation energies for Ca and Fe was essentially independent of the R-value. It wasconcluded that there was a strong interaction between the Ca atoms and the main defectsin this oxide; which were single vacancies interacting with clusters.M.Labidi, H.Boussetta, C.J.A.Monty: Solid State Ionics, 1997, 104[1-2], 133-45

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Fe2(Ni,Zn,Cu)O4: Interfacial DiffusionThe interfacial diffusion between Ni-Zn-Cu ferrite and Ag during sintering, attemperatures ranging from 850 to 950C, was investigated by using the scanning electronmicroscopy energy-dispersive X-ray analysis technique. It was found that the diffusionincreased with increasing sintering temperature. Impurities in the ferrite, especially SiO2

and NH4Cl, promoted interfacial diffusion. The interfacial diffusion of Ag could beexplained by liquid-phase formation at grain boundaries in sintered ferrite.T.Nakamura, Y.Okano, S.Miura: Journal of Materials Science, 1998, 33[5], 1091-4

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109

Fe2O3 Oxides Fe3O4

Fe2O3: Surface ReconstructionSpin-density functional theory was used to investigate various possible structures for the(00•1) haematite surface. It was found that, depending upon the ambient O partialpressure, 2 geometries with (1 x 1) symmetry were particularly stable under thermalequilibrium conditions. One was terminated by Fe, and the other by O. Both of themexhibited enormous surface relaxations. This was -57% in the case of the Fe-terminatedsurface, and -79% in the case of the O-terminated surface. A large inward relaxation ofthe first layer of Fe-terminated surface was noted, as well as a huge contraction of theinterlayer spacing between the Fe sub-surface layers of the O3-terminated surface. Inaddition to relaxations along the [00•1] direction, the O layers of both surfaces exhibiteda plane rotational reconstruction. This appeared to be the first report of this type ofrelaxation, and of plane reconstruction of O layers, at a metal oxide surface. Scanningtunnelling microscopy revealed the presence of 2 different surface terminations whichcoexisted on monocrystalline (00•1) a-Fe2O3 films which had been prepared using an Opressure of 0.001mbar. An unusual electronic structure was predicted for the O3-terminated surface; with a noticeable incidence of states from the sub-surface Fe layer.This resulted in a magnetic polarization of the O.X.G.Wang, W.Weiss, S.K.Shaikhutdinov, M.Ritter, M.Petersen, F.Wagner, R.Schlögl,M.Scheffler: Physical Review Letters, 1998, 81[5], 1038-41

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Fe2O3/MgO: Surface ReconstructionAn investigation was made of the structural and compositional changes that wereintroduced by the segregation of substrate Mg to the surface of 1µ-thick (001) Fe3O4 filmson (001) MgO. The surfaces of the molecular-beam epitaxially grown samples were flat,and exhibited a (v2 x v2) R45º reconstruction with respect to the Fe3O4 surface unit cell.The onset of Mg segregation to the surface occurred at about 670K; with long narrowextensions of terraces growing along the [110] and [11̄0] directions. Heating in Oproduced a 1 x 4 surface reconstruction, and extremely long (100nm) wide terraces. Thisannealing stage was attributed to the formation of a MgFe2O4 surface phase whichexhibited a highly anisotropic surface diffusion and step formation energy.J.F.Anderson, M.Kuhn, U.Diebold, K.Shaw, P.Stoyanov, D.Lind: Physical Review B,1997, 56[15], 9902-9

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Fe3O4: Antiphase BoundariesSuch boundaries were observed in monocrystalline films that had been grown onto MgO.The antiphase boundaries were an intrinsic result of film nucleation and growth. The intrasub-lattice super-exchange coupling was greatly strengthened across an antiphaseboundary, while the inter sub-lattice super-exchange coupling was weakened; thusreversing the predominant interaction found in the bulk. The antiphase boundaries thusseparated oppositely magnetized regions; consistent with Lorentz microscopy

110

Fe3O4 Oxides Fe20(Sr,Bi)10Ox

measurements. The antiphase boundaries introduced very large saturation fields andalmost random magnetization distributions in zero field.D.T.Margulies, F.T.Parker, M.L.Rudee, F.E.Spada, J.N.Chapman, P.R.Aitchison,A.E.Berkowitz: Physical Review Letters, 1997, 79[25], 5162-5

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y Fe20(Sr,Bi)10Ox: O PermeationOxides with perovskite-like structures, Sr10-x/2BixFe20Oy, where x was equal to 4, 6, 8 or10, were synthesized. These oxides exhibited high O permeabilities. The O permeationrate at 1150K was equal to 0.90mlSTD/cm2min, for x = 10, and equal to 0.41mlSTD/cm2minfor x = 6 (figure 6).S.Li, W.Yang, L.Fang, L.Lin: Journal of Solid State Chemistry, 1997, 130[2], 316-8

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Figure 6Permeation of O in Sr10-x/2BixFe20Oy

1.0E-01

1.0E+00

1.0E+01

8 9 10 11

x = 10x = 6

104/T(K)

P (m

l/cm

2 min

)

111

Fe20(Sr,Bi)10Ox Oxides KNbO3

Fe20(Sr,Bi)10Ox: O PermeationA new series of mixed conducting oxides, Sr10-n/2BinFe20O, where n was equal to 4, 6, 8or 10, was synthesized by using a solid-state reaction method. They had high Opermeability, and the permeation rate at 1150K was 0.41mlSTD/cm2min for n = 6, and0.90mlSTD/cm2min for n = 10. This was twice as high as that for Sr0.5Bi0.5FeO3. In thecase of Sr1-xBixFeO3, where x was equal to 0.1, 0.3, or 0.5, the O flux increased withincreasing Bi content.S.Li, Y.Cong, L.Fang, W.Yang, L.Lin, J.Meng, Y.Ren: Materials Research Bulletin,1998, 33[2], 183-8

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GeO2-Rb2O-Ag2O: Ionic ConductionThe alternating-current conductivity of xAg2O-(1-x)Rb2O-4GeO2 glasses, where x rangedfrom 0 to 1, was measured at temperatures of between 5K and ambient. A typical low-temperature conductivity region was observed below 200K. The temperature andfrequency dependence of the conductivity in this region was explained in terms of thethermal excitation of asymmetrical double-well potential configurations. It was noted thatthe classical mixed mobile-ion effect was largely absent. This indicated that the effectwas associated with independent movements of the mobile ions.H.Jain, X.Lu: Journal of the American Ceramic Society, 1997, 80[2], 517-20

[446-164-111]

In2O3: Point DefectsThe electronic states of the oxide were calculated by using a molecular orbital method.The conduction band consisted mainly of O 2p and In 5p. The valence band consisted ofO 2p. When there was an O vacancy in the oxide, a localized orbital appeared at the Insite near to the O vacancy. This formed a vacancy level below the bottom of theconduction band. When there were many O vacancies in the oxide, the vacancy level wasnon-localized. When Sn atoms were substituted for In atoms, localized molecular orbitalsof Sn 5p and O 2p appeared in the high-energy region of the conduction band. On theother hand, when Sn atoms were inserted into interstitial positions, vacancy levelsappeared below the bottom of the conduction band. In either case, when there were manyO vacancies present, O vacancy levels appeared between the conduction band and thevalence band. A complex that was formed from In, Sn and O formed the doping level.M.Mizuno, T.Miyamoto, T.Ohnishi, H.Hayashi: Japanese Journal of Applied Physics - 1,1997, 36[6A], 3408-13

[446-164-111]

KNbO3: Point DefectsThe linear muffin-tin orbital method, combined with density functional theory and theintermediate neglect of differential overlap technique, was used to study F centers (Ovacancy with two electrons) in cubic and orthorhombic ferro-electric crystals.

112

KNbO3 Oxides LaAlO3

Calculations which were performed for 39-atom super-cells showed that the 2 electronswere considerably delocalized, even in the ground state of the defect. Their wavefunctions extended over the two Nb atoms which were closest to the O vacancy, and overother nearby atoms. The F center in this material therefore resembled the defects inpartially-covalent SiO2, rather than the usual F centers in ionic crystals such as MgO.Absorption energies were calculated, by means of the intermediate neglect of differentialoverlap technique, after relaxation of the atoms which surrounded the F center. In thecase of the orthorhombic phase, absorption bands were predicted to lie at 2.72, 3.04 and3.11eV. The first one was close to that observed under electron irradiation. In the case ofthe cubic phase, which was stable at temperatures above 708K, only 2 bands (2.73,2.97eV) were expected.R.I.Eglitis, N.E.Christensen, E.A.Kotomin, A.V.Postnikov, G.Borstel: Physical ReviewB, 1997, 56[14], 8599-604

[446-164-112]

KTaO3: Point DefectsPhoto-induced Ta4+ centers in nominally pure single crystals were studied by usingelectron-spin resonance techniques. Two of the centers (Ta4+-VO and Ta4+-VO-M4+) wererelated to O vacancies. A third center was associated with an OH- molecular ion (Ta4+-OH-). These attributions were decided on the basis of concentration measurements of thecorresponding centers after annealing in Ar, O, H or H2O atmospheres. It was shown thatthe Ta4+ centers were shallow donors. At temperatures above 30K, they were ionized andtransformed into ordinary VO and OH-; which were assumed to be the main lattice defectsbefore illumination. Their energy levels were deduced from the temperature dependenceof the relaxation rate of the light-induced non-equilibrium localized electron population.The energy levels of Ta4+-VO and Ta4+-VO-M4+ centers were situated at 0.026 and0.008eV below the bottom of the conduction band, respectively. The symmetry of thecenters was broken in the sense that the photo-electron was localized near to one of 2equivalent Ta5+ ions next to an O vacancy or OH-. The VO defects and OH- moleculesplayed a role in the nucleation of local polar clusters in nominally pure crystals at lowtemperatures.V.V.Laguta, M.I.Zaritskii, M.D.Glinchuk, I.P.Bykov, J.Rosa, L.Jastrabík: PhysicalReview B, 1998, 58[1], 156-63

[446-164-112]

LaAlO3: Point DefectsA sensitive dielectric resonator technique was used to measure the tangent loss andrelative permittivity of single crystals at 4 to 300K and 4 to 12GHz. At temperaturesgreater than 150K, the tangent loss was almost sample-independent, with a linearfrequency dependence and a monotonic temperature variation. The tangent loss below150K was characterized by a peak at about 70K. The height of the peak was frequency-and sample-dependent. The peak was attributed to defect dipole relaxation. The activationenergy for the relaxation process was deduced to be 0.031eV. This low value indicatedthat the defect dipoles were associated with interstitials.

113

LaAlO3 Oxides LaMnO2

C.Zuccaro, M.Winter, N.Klein, K.Urban: Journal of Applied Physics, 1997, 82[11],5695-704

[446-164-113]

(La,Ca)CrO3: Sr Grain Boundary DiffusionThe grain-boundary diffusivity of Sr in La0.9Ca0.13CrO3 was determined by means ofsecondary-ion mass spectrometry; using depth-profiling from the surface or line-scanningof fracture surfaces. Depth profiles which were sputtered using an O2+ primary ion beamgave 2 slopes for the Sr concentration profile. These corresponded to bulk and grain-boundary diffusion. The depth profiles were fitted to an appropriate equation thatfurnished bulk and grain-boundary diffusion coefficients of 6.5 x 10-20 1.6 x 10-15m2/s,respectively, at 1273K. The line-scanning measurements permitted the successfuldetermination of Sr concentration profiles around the grain boundary. However, the grain-boundary diffusivity which was deduced was equal to 6.0 x 10-13m2/s. This was a factorof 100 greater than that which was obtained using the depth-profiling method.T.Horita, N.Sakai, T.Kawada, H.Yokokawa, M.Dokiya: Journal of the American CeramicSociety, 1998, 81[2], 315-20

[446-164-113]

(La,Ca)MnO3: Grain BoundariesThe magneto-resistance of mechanically induced grain boundaries in La0.7Ca0.3MnO3 thinfilms was investigated. The boundaries were fabricated by mechanical deformation of theLaAlO3 substrate before film deposition. During deposition, the deformed substrateregion introduced growth disorder into the film and led to a wide grain boundary. Theresultant structures exhibited a reproducible and substantial magneto-resistance inmagnetic fields below 2kG.C.Srinitiwarawong, M.Ziese: Applied Physics Letters, 1998, 73[8], 1140-2

[446-164-113]

LaMnO3: Point DefectsIn order to clarify the role which was played by vacancies, resistivity and magnetizationmeasurements were carried out on a series of samples which were derived from thisparent compound, and La and Mn vacancies were introduced by systematically varyingthe O annealing conditions. The ferromagnetic transition temperature was found todecrease as the vacancy concentration, and Mn4+ concentration, increased. It wasconcluded that La and Mn vacancies played a significant role in determining the physicalproperties of the materials. At high vacancy concentrations, the magnetic properties wereanalogous to those of spin glasses.P.S.I.P.N.De Silva, F.M.Richards, L.F.Cohen, J.A.Alonso, M.J.Martínez-Lope,M.T.Casais, K.A.Thomas, J.L.MacManus-Driscoll: Journal of Applied Physics, 1998,83[1], 394-9

[446-164-113]

114

(La,Pb)FeO3 Oxides La8Sr8Co16O36

(La,Pb)FeO3: O Permeation and Ionic ConductionPerovskite-type solid solutions were found in the oxide system, La1-xPbxFeO3, where xranged from 0.1 to 0.3. Sintering of the ceramics at temperatures above 1300K led to acation deficiency, in the La sub-lattice, due to lead oxide evaporation. The electricalconductivity and O permeation fluxes increased with increasing x-value. The permeationfluxes through La1-xPbxFeO3 were shown to be limited by the bulk ionic conductivity.V.V.Kharton, A.P.Viskup, E.N.Naumovich, A.A.Tonoyan, O.P.Reut: Materials ResearchBulletin, 1998, 33[7], 1087-93

[446-164-114]

(La,Sr)CoO3: Point DefectsVacancy-related defects in laser-ablated thin films of La0.5Sr0.5CoO3 were detected byusing a variable-energy positron beam. The O non-stoichiometry was altered by varyingthe O partial pressure within the deposition chamber during cooling. Conductivitymeasurements confirmed the change in O content. An increased positron trapping atvacancy defects was observed with increasing non-stoichiometry. It was suggested thatvacancy clusters were present in films that were cooled in O (10-5 Torr).D.J.Keeble, A.Krishnan, T.Friessnegg, B.Nielsen, S.Madhukar, S.Aggarwal, R.Ramesh,E.H.Poindexter: Applied Physics Letters, 1998, 73[4], 508-10

[446-164-114]

(La,Sr)(Co,Fe)O3: O DiffusionSpecimens of La0.6Sr0.4Co0.6Fe0.4O3 were studied by using electrical conductivity andhigh-temperature coulometric titration experiments. The chemical diffusion coefficients,for temperatures of between 923 and 1255K and O partial pressures of 0.03 to 1bar,ranged from 10-6 to 5 x 10-5cm2/s. The associated activation energies ranged from 95 to117kJ/mol.J.E.ten Elshof, M.H.R.Lankhorst, H.J.M.Bouwmeester: Solid State Ionics, 1997, 99[1-2],15-22

[446-164-114]

La8Sr8Co16O36: Point DefectsThe crystal structure of an anion-deficient perovskite-type orthorhombic phase wasdetermined for the first time by the combined use of energy-dispersive X-rayspectroscopy, electron energy-loss spectroscopy, high-resolution transmission electronmicroscopy and electron diffraction techniques. The unit cell comprised 2 types offundamental module, and contained a total of 8 modules. Each module involved a c-axisstacking of anion-deficient SrCoO3 and LaCoO3 basic perovskite cells. The unit cellpreserved the characteristics of the perovskite framework, and had a superstructure whichwas induced by O vacancies. This work illustrated the correlation of anion deficiencywith the valence state of Co, and proved that O atom positions could be determined byusing a combination of transmission electron microscopy and other techniques.

115

La8Sr8Co16O36 Oxides (La,Sr)MnO3

Z.L.Wang, J.S.Yin: Philosophical Magazine B, 1998, 77[1], 49-65[446-164-115]

(La,Sr)2CuO4: Planar DefectsPlanar defects which extended through the entire film thickness were observed, by meansof transmission electron microscopy, in La1.85Sr0.15CuO4 thin-film samples. Plan-viewelectron diffraction observations showed that the defects were located on (101) or (011)-type planes. The defects were analyzed by using high-resolution lattice imaging, and theirdisplacement vectors were determined by using image processing. A comparison wasmade with pure shear defects that were due to non-stoichiometry. Evidence for thepresence of such pure shear defects, with a displacement vector of 1/6[031], was in factprovided by these experiments.A.Alimoussa, M.J.Casanove, J.L.Hutchison: Philosophical Magazine A, 1997, 76[5],907-19

[446-164-115]

(La,Sr)2CuO4: Point DefectsThe O non-stoichiometry of La2-xSrxCuO4, where x ranged from 0 to 0.3, was measuredas a function of Sr content, temperature (400 to 1000C) and O partial pressure (1 to 10-

10atm), using high-temperature gravimetric and coulometric titration techniques. The Onon-stoichiometry ranged from O excess to O deficiency, depending upon the partialpressure and the Sr content. An O excess was observed for specimens with an x-value ofless than 0.05. The dependence of O excess non-stoichiometry upon the O partialpressures could be explained by a model in which interstitials were the predominantdefects. The variations in partial molar enthalpy revealed that a strong interaction betweenO and vacancies occurred in O-deficient material. The experimentally determined Opartial molar entropy was compared with values which were calculated by using a so-called metal model, a hopping conduction model, and a narrow-band conduction model,where an increase in O vacancies hardly affected the carrier concentration, where theholes which were generated by specimen oxidation were trapped by Cu ions, and wheregenerated holes were itinerant, respectively. The variation in O partial molar entropy wasexplained well by the hopping model and by the narrow-band conduction model.H.Kanai, J.Mizusaki, H.Tagawa, S.Hoshiyama, K.Hirano, K.Fujita, M.Tezuka,T.Hashimoto: Journal of Solid State Chemistry, 1997, 131[1], 150-9

[446-164-115]

(La,Sr)MnO3: Point DefectsDefect-disorder models were derived for non-doped and Sr-doped material. A random-defect model and a cluster-defect model were both considered for regimes thatcorresponded to an O deficit and an O excess. The models were based upon publishedexperimental non-stoichiometry data. According to both models, the addition of Sr led toan increase in the concentration of electron holes and O non-stoichiometry. The defectclusters that were predicted by the cluster model exhibited a marked concentration only atvery low O partial pressures. Both models were checked against electrical conductivity

116

(La,Sr)MnO3 Oxides LiNbO3

data. Good agreement was found between the random-defect model and experimentaldata.J.Nowotny, M.Rekas: Journal of the American Ceramic Society, 1998, 81[1], 67-80

[446-164-116]

La0.85Sr0.15MnO3: Point DefectsSingle crystals were prepared which contained O-vacancy contents ranging from 0 to0.055. The b and c lattice parameters were found to increase with the deviation fromstoichiometry, while the a-parameter remained essentially constant. The O vacancies wereexpected to reduce the hole concentration, and therefore the Curie temperature andelectrical conductivity. However, the effects were slightly different to those observed inhole-doped (La,Sr)MnO3 systems, due to differing lattice distortions. The lattice changein the sample was reflected by a marked shift in the hyperfine field of Mn4+ ions; asdetected by using nuclear magnetic resonance. The O deficiency, and the resultantproperties changes, were reversible.A.M.De Léon-Guevara, P.Berthet, J.Berthon, F.Millot, A.Revcolevschi, A.Anane,C.Dupas, K.Le Dang, J.P.Renard, P.Veillet: Physical Review B, 1997, 56[10], 6031-5

[446-164-116]

LiAlSiO4: Ionic Conduction and Point DefectsThe environment around Li in glass samples was studied by using neutron scattering, withisotopic substitution of Li, and reverse Monte Carlo modelling. It was found that Li atomswere present in distorted tetrahedral sites that were edge-linked to (Si,Al)O4 tetrahedra;with a first Li-Li distance of about 0.5nm. An increase in the activation energy for ionicconduction, between aluminosilicate and silicate glasses, was attributed to structuraldifferences.L.Cormier, P.H.Gaskell, G.Calas, J.Zhao, A.K.Soper: Physical Review B, 1998, 57[14],R8067-70

[446-164-116]

Li2B4O7: TwinsGrowing twins were studied in single crystals, and twinning was found to occur accordingto Dauphiné's rule; with twinning axes of [100] and [010]. Twinning planes were absentfrom the single crystals, and only twinning axes were present. The [100] or [010]crystallographic directions were the axes of rotation, and the (100) and (010)crystallographic planes were the joining surfaces of twins. The visible trace of thesemeeting surfaces (the so-called twinning seam) was observed, on the (001) crystalsurface, in the form of steps which reversed sign during the transition from a (001)surface to a (001̄) surface. Such steps, involving twin seams, were not observed onpolished (100) or (010) surfaces. However, under careful examination, very weaklocalized ridges in the twinning region could be discerned. The twins were not revealedby polarized light but, taking account of the fact that the Z-axis was polar, the twinning(according to the Dauphiné rule) could be attributed to electrical twinning without having

117

LiNbO3 Oxides LiNbO3

any effect upon the optical properties. It was suggested that twinning might even have hada polysynthetic nature in some samples.Y.V.Burak: Journal of Crystal Growth, 1997, 186, 302-4

[446-164-117]

Figure 7Diffusivity of H in LiNbO3

Li0.5La0.5TiO3: Ionic ConductionNuclear magnetic resonance and electrical conductivity methods were used to study thedynamics of ionic diffusion in crystalline samples. It was found that the direct-currentconductivity exhibited a non-Arrhenius temperature dependence. Spin-lattice andconductivity relaxations were analyzed, in terms of a non-Arrhenius dependence of thecorrelation time, over the same ranges of frequency and temperature. Both relaxationscould be described by a function of the form, exp[-t0.4], over the whole temperature range.C.León, J.Santamaria, M.A.París, J.Sanz, J.Ibarra, L.M.Torres: Physical Review B, 1997,56[9], 5302-5

[446-164-117]

1.0E-02

1.0E-01

1.0E+00

1.0E+01

19 20 21 22 23

0mol%Li, X-cut0%molLi, Y-cut0mol%Li, Z-cut0.5mol%Li, X-cut0.5mol%Li, Y-cut0.5mol%Li, Z-cut1mol%Li, X-cut1mol%Li, Y-cut1mol%Li, Z-cut

104/T(K)

D (µ

2 /h)

118

LiNbO3 Oxides LiNbO3

y LiNbO3: H DiffusionDepth profiles of the proton concentration in proton-exchanged samples were measuredby means of secondary ion mass spectrometry. The concentration-dependent diffusioncoefficient for proton-Li counter-diffusion was then determined by using the Boltzmann-Matano method. The results (figure 7) obeyed Arrhenius relationships, and their analysisyielded activation energies of 78.6, 66.5, and 82.4kJ/mol for X-cut, Y-cut and Z-cutspecimens, respectively. The results also depended upon the Li content of the protonexchange source (which was present as the benzoate).K.Ito, K.Kawamoto: Japanese Journal of Applied Physics - 1, 1997, 36[11], 6775-80

[446-164-118]

LiNbO3: Point DefectsSpecimens of an Er-doped congruent phase were investigated by using electronparamagnetic resonance techniques. When the magnetic field was oriented parallel to thec-axis, the previously observed characteristic hyperfine split electron paramagneticresonance spectrum of Er3+ was observed. However, the angular dependence revealedline-splittings which had not been reported before. These were attributed to Er3+ defectsthat were located on Li+ sites which were next to Li+ vacancies, VLi

+. These tilted the g-tensor orientation away from the c-axis. Several differently oriented Er3+-VLi

+ complexeswere identified. The electron paramagnetic resonance spectra could be explained byassuming that a statistical distribution of VLi

+ existed in the neighborhood of Er3+-VLi+.

T.Nolte, T.Pawlik, J.M.Spaeth: Solid State Communications, 1997, 104[9], 535-9[446-164-118]

LiNbO3: TwinsThe relationship between cracking and mechanical twinning, and between ferro-electricpolarization and mechanical twinning, was investigated in stoichiometric material. Threesets of mechanical twins were found to cross each other, and the crossing points of thesetwins were preferred sites for crack nucleation. The mechanical twins exhibited a head-to-tail arrangement of ferro-electric polarization, and complicated polarization states wereobserved at the crossing points.B.M.Park, K.Kitamura, Y.Furukawa, Y.Ji: Journal of the American Ceramic Society,1997, 80[10], 2689-92

[446-164-118]

LiNbO3: TwinsTwinning in stoichiometric samples which had been grown by using the double crystalCzochralski technique was investigated. The geometry and structures of twin boundarieswere observed by using optical and transmission electron microscopy. The twinningwhich was observed was considered to be mechanical twinning. The mechanical twinswere suggested to form just after crystal growth, and had a doubly-twinned lamellarshape. They were coherent, and the composition plane was a {012} plane.

119

LiNbO3 Oxides MgO

B.M.Park, K.Kitamura, K.Terabe, Y.Furukawa, Y.Ji, E.Suzuki: Journal of CrystalGrowth, 1997, 180, 101-4

[446-164-119]

LiNbO3: TwinsThe geometry and atomic structure of the mechanical twin boundary were investigated. Apossible structural model for this twin was deduced and was compared with the observedstructure. The structure of the twin in stoichiometric material was coherent, and itscomposition plane was one of the {012} planes. This mechanical twin could be explainedin terms of a twin model for slip between two Nb-centered octahedral layers, with a slightdistortion of each octahedron.B.M.Park, K.Kitamura, K.Terabe, Y.Furukawa, E.Suzuki: Philosophical MagazineLetters, 1998, 77[1], 17-21

[446-164-119]

Li2O-(Al2O3,Ga2O3)-TiO2-P2O5: Ionic ConductionFast Li-ion conducting glass-ceramics were obtained by heat-treating glasses in thissystem. The glass-ceramics were composed mainly of LiTi2(PO4)3, in which the Ti4+ ionswere partially replaced by Al3+ or Ga3+ ions. Such substitution led to a considerableenhancement of the conductivity. The maximum room temperature conductivity whichwas obtained was 0.0013S/cm in the Al system and 0.0009S/cm in the Ga system.J.Fu: Journal of Materials Science, 1998, 33[10], 1549-53

[446-164-119]

Li2O-Al2O3-TiO2-SiO2-P2O5: Li Diffusion and Ionic ConductionFast Li-ion conducting glass-ceramics were prepared from the pseudo-binary system,2[Li1+xTi2SixP3-xO12]-AlPO4. The major phase which was present was LiTi2P3O12, inwhich Ti4+ and P5+ ions were partially replaced by Al3+ and Si4+ ions, respectively. It wasfound that increasing the x-value resulted in a considerable enhancement of theconductivity. Room-temperature conductivities above 0.001S/cm were measured over awide range of compositions.J.Fu: Journal of the American Ceramic Society, 1997, 80[7], 1901-3

[446-164-119]

y MgO: Ca, O, Zn Diffusion and Ion BombardmentThe radiation-enhanced diffusion and ion-beam mixing of 18O, Ca and Zn buried tracerlayers, in samples which had been grown by means of molecular beam epitaxy, weremeasured after bombardment with 2MeV Kr+ and 1MeV Ne+, He+ and H+ at temperaturesranging from 30 to 1500C (figure 8). The ion-beam mixing parameter varied between0.00001 and 0.00005nm5/eV for the various tracers, at 30C, and increased slowly withincreasing temperature. These results were consistent with ballistic mixing. In the highesttemperature range which was investigated (1350 to 1500C), the radiation-enhanceddiffusion coefficient for 18O was proportional to the square root of the irradiation flux,and

120

MgO Oxides MgO

indicated an apparent activation enthalpy of 1.2eV. These dependences upon flux andtemperature were indicative of recombination-limited kinetics, in which the measuredactivation enthalpy represented half of the migration enthalpy of anion vacancies. Anactivation enthalpy of 4.1eV was deduced for temperatures ranging from 1150 to 1350C.This unusually high value was attributed to the dissociation energy of small vacancyclusters. It was noted that measurements on the cation sub-lattice were limited totemperatures below 900C, due to an excessive thermal diffusivity that was associatedwith extrinsic vacancies which were linked to trivalent impurity-charge compensation.A.I.Van Sambeek, R.S.Averback, C.P.Flynn, M.H.Yang, W.Jäger: Journal of AppliedPhysics, 1998, 83[12], 7576-84

[446-164-120]

Figure 8Diffusivity of Ca, O and Zn in MgO

MgO: Pipe Diffusion, Dislocations, Grain Boundaries, and Point DefectsAtomistic simulations were used to evaluate the diffusion pathways and activationenergies for cation and anion vacancy migration in {410}/[001] symmetrical tilt grainboundaries. These could be considered as being an array of dislocation pipes. Theapproach used was based upon molecular dynamics, and it was found that the diffusion

1.0E-19

1.0E-18

1.0E-17

1.0E-16

1.0E-15

1.0E-14

1.0E-13

4 14 24

OCaZn

104/T(K)

D (c

m2 /s

)

121

MgO Oxides MgO

paths were anisotropic; with diffusion down the dislocation pipes being favored overdiffusion between the pipes. The lowest calculated activation energies for isolatedvacancies were equal to 1.05eV for Mg and to 1.01eV for O at 0GPa. The bulk activationenergies were equal to 1.94eV for Mg and to 2.12eV for O. The lower activation energies,coupled with increased defect concentrations at the interface, showed that the boundarieswere regions of high diffusivity. However, the concentrations of vacancy pairs at theinterface and the high binding energy of Mg-O pairs led to the prediction that a largefraction of the defects were bound. This in turn caused the activation energy for vacancymigration to approach that of the bulk. In this case, the higher boundary diffusivities werethe result of high defect concentrations at the boundary.D.J.Harris, G.W.Watson, S.C.Parker: Physical Review B, 1997, 56[18], 11477-84

[446-164-121]

MgO: Surface Diffusion and Electron IrradiationNanometre-scale electron-beam drilling was carried out on vacuum-depositedmonocrystalline films, in a 200kV high-resolution transmission electron microscope,using a thermionic gun. It was found that holes formed due to a surface mass-lossmechanism. Surface atomic diffusion and desorption were observed in situ at a spatialresolution of 0.2nm and a time resolution of 1/60s. It was found that most of the surfaceatoms first began to diffuse around steps on the surface, under the electron irradiation,and then desorbed into the vacuum at weakly-bonded positions such as steps, kinks andisolated positions on the surface.T.Kizuka, N.Tanaka: Philosophical Magazine Letters, 1997, 76[4], 289-97

[446-164-121]

MgO: Electron Irradiation, Dislocations, and Point DefectsMicrostructural evolution and radiation damage kinetics were studied in a wide variety ofsingle crystals by means of in situ observations during electron irradiation in a high-voltage electron microscope. The irradiation introduced interstitial dislocation loops, nodefect clusters, and bubbles into the crystals; depending upon the source and history ofthe specimen. Evidence was found that the presence of structural vacancies couldsuppress the nucleation and growth of dislocation loops during electron irradiation, thusenhancing the radiation resistance. The vacancies which appeared to promote radiationtolerance were related to the incorporation of OH- ions into the crystal lattice. Thesevacancies were in fact point-defect complexes which consisted of one or two OH- ionsthat resided on O lattice sites; together with a vacancy on a Mg site. These configurationsyielded either a fully charge-compensated point defect complex, (OH•VMg"HO•)X, or asingly-charged point defect complex, (OH•VMg")'.C.Kinoshita, T.Sonoda, A.Manabe: Philosophical Magazine A, 1998, 78[3], 657-70

[446-164-121]

MgO: Electron Irradiation and Grain BoundariesIrradiation-induced migration of a [001] S = 5 tilt boundary was monitored by means ofin situ high-resolution transmission electron microscopy, with a temporal resolution of

122

MgO Oxides MgO

1/60s. The migration process was clarified for the first time. It was found that the non-periodic boundary, which was made up of several structural units, migrated and became aperiodic structure along (310)A-(31̄0)B planes when the units transformed into ones whichwere surrounded by (310)A-(31̄0)B facets. The units which were surrounded by (120)A-(210)B facets grew and became aligned along the (120)A-(210)B planes. It was found thatmigration proceeded at every translation and transformation of 3 kinds of structural unitin the boundary, and that the migration speed was controlled by the time intervalsrequired for translation or transformation of the units. It was shown that the atomicarrangement of the boundary during translation and transformation was predicted by thecoincidence-site lattice model when the boundary was composed of stable structural units.T.Kizuka, M.Iijima, N.Tanaka: Philosophical Magazine A, 1998, 77[2], 413-22

[446-164-122]

MgO: Grain BoundariesA comparison of local-density approximations and classical potential calculations for the[001] S = 5 twist grain boundary showed that the complex structure which was observedin the defective grain boundary was caused by a simple electrostatic effect. Both methodspredicted complex structures of octagons and squares. The electronic structure indicatedthat there were no states in the gap; thus suggesting that the so-called vacancies in thestable relaxed structure should be considered to be structural elements rather than truedefects.J.H.Harding, C.Noguera: Philosophical Magazine Letters, 1998, 77[6], 315-20

[446-164-122]

MgO: Point DefectsThe free energies of vacancy pair formation and migration were computed usingmolecular dynamics methods involving free-energy integrations and a non-empirical ionicmodel with no adjustable parameters. The intrinsic diffusion constant was obtained forpressures ranging from 0 to 140GPa, and temperatures ranging from 1000 to 5000K.Excellent agreement was found, with zero-pressure diffusion data, to within experimentalerror. The homologous temperature model, which related diffusion to the melting curve,closely described the high-pressure results for this theoretical framework.J.Ita, R.E.Cohen: Physical Review Letters, 1997, 79[17], 3198-201

[446-164-122]

MgO: Point DefectsCathodoluminescence imaging of nano-indentations in crystals, made using a Berkovichdiamond triangular pyramid, was carried out at wavelengths ranging from 380 to 800nmby using a scanning electron microscope. It was shown that the centers of nano-indentations, with estimated residual depths ranging from about 570 to 325nm,luminesced quite strongly and that the spatial resolution of the image appeared to bebetter than 1µ; as revealed by the structure within the cathodoluminescence image. Thecenters of larger indents, with residual depths greater than about 1300nm, did not exhibitluminescence. It was suggested that most of the cathodoluminescence which arose from

123

MgO Oxides MgTiO3

the indentation centers was due to F and F+ colour centers which were produced byplastic flow; with emission peaks at 525 and 375nm. It was also suggested that therelative number density of F+ centers increased with increasing plastic strain; which, inturn, increased with increasing indentation size. Therefore, beyond some criticalindentation size, most of the cathodoluminescence from the indentation center was due tothe F+ centers. And since this luminescence was beyond the detection sensitivity of thecathodoluminescence system, the indentation center exhibited an absence ofcathodoluminescence.M.M.Chaudhri: Philosophical Magazine Letters, 1998, 77[1], 7-16

[446-164-123]

MgO/FeO: InterdiffusionAn investigation was made of the interdiffusion of cations in the MgO-FeO oxide systemat temperatures of between 1473 and 1603K. Diffusion-couple experiments were carriedout by using polycrystalline samples of the oxides. After heat-treatment, the concentrationprofiles were measured by means of energy-dispersive X-ray spectroscopy. It was foundthat the data could be described by:

D (m2/s) = 1.29 x 10-5 exp[-152(kJ/mol)/RT]exp[-12.4xMg]J.Bygden, A.Jakobsson, D.Sichen, S.Seethararaman: Physical Review Letters, 1997,88[5], 433-7

[446-164-123]

MgTiO3: Ion Bombardment and Point DefectsA study was made of the radiation response of geikielite; an ilmenite-group oxide.Oriented single crystals were bombarded with 200keV Ar2+ or 400keV Xe2+ ions.Rutherford back-scattering spectrometry ion channelling was then used to characterize theresultant radiation damage. In the case of 200keV Ar2+ ion bombardment, the sample washeld at 100K and a buried amorphous 55nm-thick layer formed beneath a 90nm-thickdefective crystalline layer after exposure to a fluence of 2 x 1015/cm2. More detailedexperiments used incremental 400keV Xe2+ ion bombardment and Rutherford back-scattering spectrometry ion channelling to determine the extent and rate of radiationdamage at temperatures of 170, 300 and 470K. This showed that there was a strongtemperature dependence of damage accumulation, and that the critical amorphizationfluence increased from 2 x 1015 Xe2+/cm2 (170K) to 6 x 1015 Xe2+/cm2 (300K) to morethan 2.5 x 1016 Xe2+/cm2 (470K). Damage appeared to accumulate in several stages. Arapid initial growth levelled out at an intermediate stage and was followed by an increaseup to eventual saturation. At 170 and 300K, the damage fraction saturated at 100%,whereas saturation occurred at about 80% at 470K. Rutherford back-scatteringspectrometry ion channelling data suggested the possible formation of a radiation-inducedmetastable phase, in the damaged region, which might be considered to be analogous topressure-induced or temperature-induced phase transformations in other ilmenite-groupoxides. The present results suggested that ionicity, composition and melting temperaturecould play important roles in the radiation response of ceramics; particularly with regardto determining the relative radiation tolerance of members of a solid-solution series.

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MgTiO3 Oxides Mn3O4

J.N.Mitchell, N.Yu, K.E.Sickafus, M.A.Nastasi, K.J.McClellan: Philosophical MagazineA, 1998, 78[3], 713-25

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Mn(La,Ba)O3: Point DefectsThermogravimetric methods were used to determine the conditions which were requiredin order to obtain samples of Ba-substituted LaMnO3 with a stoichiometric O content.Single-phase and vacancy-free samples of La1-xBaxMnO3, where x ranged from 0.12 to0.24, were prepared in air by quenching from 1100 to 1450C. Compositions with x-valuesof less than 0.12 required synthesis under a reduced O pressure. The structure was studiedby using neutron and powder diffraction methods, and was analyzed in terms of thetolerance factor concept. At room temperature, the transition from orthorhombic torhombohedral symmetry occurred at an x-value of about 0.13. Structural, magnetic andresistive measurements showed that the properties were similar to those of the Sr-substituted system, at slightly lower levels of substitution. The Curie temperatures werehigher than those for the Sr-substituted samples, for x-values of 0.10 to 0.14, due to abetter match of the Mn-O and La(Ba)-O bond lengths. They were slightly lower, for x-values of 0.20 to 0.24, due to an increased local distortion. The largest magnetoresistiveeffect at room temperature was found for x = 0.22. By comparing the Ba- and Sr-substituted systems, it was shown that the structural properties were governed by the ionicsizes whereas the physical properties were controlled by ionic size and hole doping.B.Dabrowski, K.Rogacki, X.Xiong, P.W.Klamut, R.Dybzinski, J.Shaffer, J.D.Jorgensen:Physical Review B, 1998, 58[5], 2716-23

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Mn3O4/Ag: DislocationsIt was observed that internal oxidation of Ag-3at%Mn produced Mn3O4 precipitateswhich had a so-called parallel topotaxy with respect to the metal matrix, and anoctahedral shape which was due to {111} facets. This orientation and shape did not givethe lowest strain-energy state, thus indicating a predominance of interfacial energy overstrain energy. On the basis of strain energy, the oxide precipitates should have had aplate-like shape with (001) as the dominant facet. Due to the tetragonality of the oxide,only a few planes and directions of the Ag and oxide could be parallel simultaneously.The precipitates exhibited a tendency to align their {111} planes parallel to the matrix forone pair of facets while, for another pair of planes, a tilt of 7.6º occurred which wasrelieved by the appearance of ledges in the Ag. An essentially 1-dimensional 10.4%resultant mismatch along <112> was usually accommodated by an array of dislocationswith a <110> line direction and alternating Burgers vectors of 1/6<112> and 1/3<112>.At parts of the parallel and tilted interfaces, dislocations were observed which stood offfrom the interface by 1 or 2 atomic spacings into the Ag. On the basis of atomisticcalculations, relaxation at the parallel {111} interfaces was attributed to the dissociationof unfavorable 1/3<112> Burgers vector into more favorable Shockley partials. At thetilted interfaces, ledges which corresponded to the tilt between the {111} planes couldcollapse and thus improve the parallelism of the {111} planes of the metal and oxide at

125

Mn3O4 Oxides Nb2(Pb,Ba)O6

the interface. The underlying mechanism of this collapse was the dissociation of a Frankpartial into a Shockley partial which was then emitted into the Ag along {111} and madea large angle with respect to the interface and to a stair-rod which remained at theinterface.B.J.Kooi, H.B.Groen, J.T.M.De Hosson: Acta Materialia, 1997, 45[9], 3587-607

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MnO/Cu, Mn2O3/Ag, Mn2O3/Cu: DislocationsParallel close-packed plane interfaces between MnO, or Mn3O4, and Ag or Cu, werestudied by means of high-resolution transmission electron microscopy and atomisticcalculations. The object was to determine what network of misfit dislocations existed atthe interfaces. Because of the projective nature of high-resolution transmission electronmicroscopy, an atomistic model which incorporated the displacements that wereassociated with possible misfit-dislocation networks was of appreciable value. It wasfound that, in the case of Ag/Mn3O4, there was an array of dislocations with a linedirection of [11̄0] and alternating Burgers vectors of 1/6[112̄] and 1/3[112̄]. In the case ofCu/MnO, there was a trigonal network of edge dislocations with a Burgers vector of1/6<112> type and a <110> line direction. In the case of Cu/Mn3O4, the Burgers vectorswere still of 1/6<112> type but, due to a slight change in line direction with respect to<110>, there was some screw character.B.J.Kooi, H.B.Groen, J.T.M.De Hosson: Acta Materialia, 1998, 46[1], 111-26

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Nb2(Pb,Ba)O6: H DiffusionSingle crystals of Pb1-xBaxNb2O6, where (1 - x) ranged from 0.5 to 0.6, were grown byusing the vertical Bridgman method. The OH- absorption spectra were measured bymeans of Fourier-transform infra-red spectrometry. The as-grown crystals exhibited astrong OH- absorption near to 3485/cm, thus revealing the existence of H+. Thepolarization dependence of the integrated OH- optical-absorption intensity showed that allof the O-H bonds were polarized either perpendicular or parallel to the c-axis; with agreater proportion being perpendicular. The proton concentration could be increased byheat treatment in a humid atmosphere, and was decreased by an ambient or O atmosphere.The diffusivity of protons perpendicular to the c-axis was measured in Pb0.5Ba0.5Nb2O6

by means of proton in-diffusion measurements. It was found that Do was equal to 2.19 x10-7cm2/s at 750C. This was comparable to the proton diffusivities in TiO2 (1.90 x 10-

6cm2/s) and LiNbO3 (1.93 x 10-7 to 2.51 x 10-6cm2/s) at the same temperature. The protonactivation energy was 0.95eV.M.Lee, R.S.Feigelson, A.Liu, L.Hesselink, R.K.Route: Physical Review B, 1997, 56[13],7898-904

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126

NiO Oxides (Pb,La)(Zr,Ti)O3

NiO/Fe2O3: InterdiffusionMulti-layer films of NiO on a-Fe2O3 were prepared by means of pulsed laser ablation.Interdiffusion led to the formation of NiFe2O4 at the interfaces. Its ferrimagnetismprovided a macroscopically measurable local probe signal for monitoring the admixture ofantiferromagnetic NiO and a-Fe2O3. Magnetization, torque and magneto-opticalmeasurements were performed on the multi-layer samples. Kinetically controlled (due tothe incoming particle energies) and thermally activated regimes of NiFe2O4 formationwere observed as a function of the substrate temperature. In the thermally activatedregime, the deposition duration per layer (and therefore the temperature-dependentdiffusion at the interfaces) played a crucial role in the formation of the ferrite.N.Keller, M.Guyot, A.Das, M.Porte, R.Krishnan: Solid State Communications, 1998,105[5], 333-7

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P2O5-ZrO2-SiO2: Ionic ConductionFast proton-conducting glasses were prepared via the hydrolysis of metal alkoxides. Theglasses which were obtained by heating at 150 to 400C were chemically stable andexhibited conductivities of about 104S/cm at room temperature. These conductivities werehigher, by about 4 orders of magnitude, than those of glasses which contained no P2O5.These high conductivities were attributed to fast proton transfer which was accelerated bymolecular water, bonded with POH groups.M.Nogami, R.Nagao, K.Makita, Y.Abe: Applied Physics Letters, 1997, 71[10], 1323-5

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(Pb,Ca)TiO3, (Pb,Sm)TiO3: Point DefectsThe low-frequency behaviour of the isotropic elastic modulus and mechanical losstangent of modified ferro-piezo-electric ceramics were studied by using the 3-pointbending technique. Measurements of these parameters, as a function of temperature, werecarried out at frequencies ranging from 0.2 to 20Hz. The results revealed anomalies, inboth parameters, at the transition temperature and below. The anomalies below thetransition temperature exhibited wide maxima of mechanical loss, whose temperaturesincreased when the measuring frequency was increased. Values for activation energieswere obtained which suggested the existence of mechanical relaxation mechanisms inwhich point defects and 90º ferro-electric domains were involved. The main species ofpoint defect was expected to be the O vacancy.B.Jiménez, J.M.Vicente: Journal of Physics D, 1998, 31[4], 446-52

[446-164-126]

(Pb,La)(Zr,Ti)O3: Point DefectsA study was made, of vacancy-related defects in ferroelectric capacitors, by using avariable-energy positron beam. Heterostructures of (Pb0.9La0.1)(Zr0.2Ti0.8)O3, withLa0.5Sr0.5CoO3 electrodes, were deposited by means of pulsed laser deposition and the

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(Pb,La)(Zr,Ti)O3 Oxides Pb2(Sn,W)2O6

effects of O deficiency were studied by using structures grown in O (760 or 10-5 Torr).The variable-energy positron beam depth profile revealed an increase in vacancy-relateddefects with increasing O non-stoichiometry. The formation of vacancy clusters in theLa0.5Sr0.5CoO3 top electrode, and VPb-VO defects in the (Pb0.9La0.1)(Zr0.2Ti0.8)O3 layer,with increasing O deficiency, was deduced.D.J.Keeble, B.Nielsen, A.Krishnan, K.G.Lynn, S.Madhukar, R.Ramesh, C.F.Young:Applied Physics Letters, 1998, 73[3], 318-20

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Pb[(Nb,Mg)0.9Ti0.1]O3: Point DefectsElectrical measurements were carried out upon various ceramics which exhibited aspontaneous relaxor-to-ferroelectric transition. Both the relaxor behavior and the phasetransition were found to depend strongly upon the Pb-vacancy content. This wasexplained in terms of an increase in the ordered domains, and reduced interactionbetween polar clusters.O.Bidault, E.Husson, A.Morell: Journal of Applied Physics, 1997, 82[11], 5674-9

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Pb(Nb,Ni)O3/Pb(Zr,Ti)O3: Nb, Ni, Ti, Zr DiffusionDiffusion reaction between these phases in gradient piezo-electric ceramics wasinvestigated as a function of the diffusion temperature and time. Ionic compositiondistribution profiles in the interdiffusion region were examined by means of electronmicroprobe analysis. On the basis of a diffusion model for the overlapped diffusionsolution from a thin slab, numerical simulations of the ionic composition distributionwere carried out. These were in agreement with the electron-probe micro-analysis results.The diffusion coefficients for Ni2+, Nb5+, Ti4+, and Zr4+ were deduced to be equal to 3.38x 10-11, 2.26 x 10-11, 1.08 x 10-11 and 9.9 x 10-12m2/s, respectively. The associatedactivation energies were equal to 94.4, 171.7, 257.5 and 325.8kJ/mol, respectively.X.Zhu, J.Xu, Z.Meng: Journal of Materials Science, 1998, 33[5], 1023-30

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Pb2(Sn1.5W0.5)O6.5, Pb2(Ti1.5W0.5)O6.5: Point DefectsThe bulk structures and surface compositions of these O defect pyrochlore oxides werestudied in order to determine the relationship between anion vacancies, electricalconductivity, cationic mobility, and surface segregation. The bulk structures were refinedby means of Rietveld analysis, combined with powder neutron and X-ray synchrotrondiffraction data. It was found that the oxides were isostructural (F4̄3m, a = 1.03501 [Ti] =1.06051nm [Sn]). The Ti (or Sn) and W atoms were randomly distributed on the 16e sites,and O-vacancy ordering occurred on the pyrochlore type-8b sites. In both cases, the Pbcation was surrounded by seven O atoms in a compressed scalenohedral arrangement inwhich the eighth vertex was occupied by the Pb 6s lone-pair electrons. In the Ti-containing pyroclore, the Pb atoms were displaced by 0.00328nm, along the [111]direction, towards the associated vacancy. In the larger Sn-containing pyroclore, the Pbcation was not displaced away from the ideal 16d site position. This difference was due to

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Pb2(Sn,W)2O6 Oxides PbWO2

the Pb-O bonding requirements. The surface compositions of both oxides were identicalto the bulk compositions. It was concluded that electrical conductivity was a criticalfactor which governed surface enrichment by potentially mobile A-type cations.Ismunandar, B.J.Kennedy, B.A.Hunter: Journal of Solid State Chemistry, 1997, 130[1],81-9

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PbTiO3: Point DefectsAn investigation was made of the nature of the atomic relaxations, around O-vacancydefects in this ferro-electric perovskite, via first-principles pseudopotential total-energycalculations. A tail-to-tail polarization was detected in the atomic relaxations around Ovacancies, and its stability was enhanced by charge trapping. It was noted that the Ovacancies in Ti-O-Ti chains along the polarization axis were more favorable than those inTi-O-Ti planes that were normal to the axis.C.H.Park, D.J.Chadi: Physical Review B, 1998, 57[22], R13961-4

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PbWO4: Neutron Irradiation and Point DefectsChanges in the absorption spectra, introduced by 60Co irradiation and high-temperatureannealing, were studied. On the basis of the radiation and annealing-induced absorptionspectra, 2 color centers were deduced to be responsible for shaping the absorptionspectrum of crystals below 3.6eV. These were: Pb3+, O-, F and F+. Thermoluminescenceglow curves were also studied above room temperature in order to determine the traplevels which resulted from irradiation.M.Nikl, K.Nitsch, S.Baccaro, A.Cecilia, M.Montecchi, B.Borgia, I.Dafinei, M.Diemoz,M.Martini, E.Rosetta, G.Spinolo, A.Vedda, M.Kobayashi, M.Ishii, Y.Usuki, O.Jarolimek,P.Reiche: Journal of Applied Physics, 1997, 82[11], 5758-62

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PbWO4: Point DefectsThe auto-localization of free electrons at WO4

2- complex anions was demonstrated bymeans of electron spin resonance measurements It was known that the WO4

3- center was ashallow donor, with its energy level situated some 0.05eV below the bottom of theconduction band. At temperatures of 50 to 60K, the captured electrons were de-trappedand the formation of other paramagnetic centers occurred. In La-doped material, anenhanced cross-section (with respect to that for undoped samples) for electron-trappingby MoO4

2- impurity centers was found. No intrinsic O hole and/or Pb3+ centers wererevealed.V.V.Laguta, J.Rosa, M.I.Zaritskii, M.Nikl, Y.Usuki: Journal of Physics - CondensedMatter, 1998, 10[32], 7293-302

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PbWO2 Oxides SiO2

PbWO4: Point DefectsThe frequency dependence of the dielectric loss factor in undoped and La-doped crystalswas investigated, using frequencies ranging from 10Hz to 10MHz, at temperaturesranging from 40 to 370C. Typical dielectric relaxation was observed in the case of La-doped samples. The results suggested that intrinsic mobile defects existed in the purecrystals and that Pb vacancies were the predominant mobile defects. In the case of La-doped crystals, La3+ ions were thought to be located at Pb2+ sites and to combine with Pbvacancies so as to form dipole complexes: [2(LaPb

3+)•-VPb"]. These were suggested to bethe cause of the dielectric relaxation in La-doped crystals.B.Han, X.Feng, G.Hu, P.Wang, Z.Yin: Journal of Applied Physics, 1998, 84[5], 2831-4

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Pb(Zr,Ti)O3: Antiphase BoundariesTransmission electron microscopy and neutron diffraction methods were used tocharacterize single crystals from the rhombohedral region of the Pb(Zr1-xTix)O3 phasediagram, where x ranged from 0.06 to 0.45. The electron diffraction patterns revealed theexistence of superlattice reflections of the form, ½{hkl}p, where h = k = l, and ½{hk0}p,which were not observed in neutron powder diffraction studies. An analysis of thesereflections revealed satellite spots, around the ½{hk0}p, which were associated withperiodic antiphase boundaries. The origin of these superlattice reflections was explainedin terms of the existence of local regions with anti-parallel cation displacements.J.Ricote, D.L.Corker, R.W.Whatmore, S.A.Impey, A.M.Glazer, J.Dec, K.Roleder:Journal of Physics - Condensed Matter, 1998, 10[8], 1767-86

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Pb(Zr,Ti)O3: Point DefectsStudies of the effects of defect distributions and mobility upon ferro-electric phasetransformations were performed by means of transmission electron microscopy, high-resolution electron microscopy and dielectric spectroscopy. Defects which could berandomly quenched-in from temperatures that were significantly above that of the ferro-electric phase transformation resulted in a relaxor ferro-electric behavior that wascharacterized by polar nano-domains. However, defects which remained mobile, up totemperatures that were significantly lower than the phase transformation, resulted in anormal ferro-electric state which was characterized by domain pinning. Mobile impuritiesand vacancies were shown to order into chain fragments or clusters which diffused todomain boundaries and resulted in pinning of the boundaries during cooling.Q.Tan, Z.Xu, J.F.Li, D.Viehland: Applied Physics Letters, 1997, 71[8], 1062-4

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SiOx: Point Defects and Defect AnnealingA strong photoluminescence of H-doped material, which had been prepared via plasma-enhanced chemical vapor deposition, was studied; together with infra-red and micro-

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SiOx Oxides SiO2

Raman spectra. It was found that each photoluminescence spectrum comprised 2Gaussian components: a main band, and a shoulder. The former was suggested to arisefrom amorphous Si clusters which were embedded in the SiOx network, and its red-shiftas a function of annealing temperature was attributed to expansion of the Si clusters. Theshoulder remained at about 835nm, regardless of the annealing temperature, and wastentatively attributed to luminescent defect centers. An enhanced photoluminescencespectrum, after annealing at 1170C, was attributed to the quantum confinement effects ofnanocrystalline Si; embedded in the oxide matrix.Z.X.Ma, X.B.Liao, J.He, W.C.Cheng, G.Z.Yue, Y.Q.Wang, G.L.Kong: Journal ofApplied Physics, 1998, 83[12], 7934-9

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SiO2: B DiffusionIt was noted that oxides which had been grown at low temperatures had a higher densitythan those which had been grown or annealed at high temperatures. An investigation wasmade here of the role which the oxide density played in affecting B diffusion throughthin-gate dielectrics. Capacitor structures were prepared by depositing polycrystalline Sionto 5.4 to 6nm of oxide which had been grown at 800C in dry O2. Some of these high-density oxides were then annealed at 1100C in order to relax grown-in stresses. After ionimplantation of B into the polysilicon, the structures were annealed in an inert ambient at950C in order to promote B penetration. Capacitance-voltage measurements revealed thatB penetration was greatly enhanced in films which had been subjected to high-temperature relaxation annealing. It was concluded that B diffusion in the oxide wasretarded by an increased density. The B diffusivity decreased from about 4 x 10-18cm2/s,at an excess density of 0.4%, to a diffusivity of about 10-18cm2/s at an excess density of4%.M.Navi, S.T.Dunham: Applied Physics Letters, 1998, 72[17], 2111-3

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SiO2: B DiffusionThe diffusivity of B in amorphous silica films was studied at temperatures ranging from1173 to 1373K. The films had been prepared by the thermal oxidation of Si substrates.The B was introduced from B vapor at pressures of between 5.5 x 10-18 and 7.2 x 10-13Pa.The coefficients were deduced from the B concentration profiles, as measured by usingsecondary ion mass spectrometry. It was found that the results could be described by:

D (m2/s) = 1.88 x 10-12exp[-205(kJ/mol)/RT]The diffusivity did not depend upon the B concentration at the film surface. It wassuggested that the B diffused through Si sites in the silica network..K.Hawagishi, M.Susa, T.Maruyama, K.Nagata: Journal of the Electrochemical Society,1997, 144[9], 3270-5

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SiO2 Oxides SiO2

SiO2: C DiffusionThe diffusion of C in oxide films, and its segregation at Si/SiO2 interfaces, wereinvestigated by using C-containing borophosphosilicate glass films and C-implanted SiO2films. It was found that C atoms diffused in oxide films at temperatures as low as 500C.The C atoms segregated at the Si/SiO2 interface and induced a positive charge. Thepositive charge density was proportional to the segregated C concentration. Fieldemission transmission electron microscopy and electron energy loss observations revealedthat C atoms existed on the SiO2 side of the interface, and that another C-rich phaseformed in SiO2.I.Mizushima, E.Kamiya, N.Arai, M.Sonoda, M.Yoshiki, S.Takagi, M.Wakamiya,S.Kambayashi, Y.Mikata, S.Mori, M.Kashiwagi: Japanese Journal of Applied Physics - 1,1997, 36[3B], 1465-8

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SiO2: D DiffusionThe diffusion of D in films was studied, at temperatures of between 485 and 655K, byusing a D2 gas-phase permeation technique. It was found that the diffusivity ranged from5.8 x 10-14 at 485K to 5.1 x 10-13cm2/s at 655K. The overall results could be described by:

D (cm2/s) = 1 x 10-10 exp[-0.34(eV)/kT]Grain-boundary diffusion through thin films was suggested to be the rate-limiting process.R.Checchetto, L.M.Gratton, A.Miotello, C.Tosello: Journal of Non-Crystalline Solids,1997, 216, 65-70

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SiO2: SiF4 DiffusionExtended Hückel calculations were used to estimate the change in electron density on Si-O back-bonds, and the total electronic energy during reaction at the oxide film surface. Itwas demonstrated that the desorption of SiF4 occurred via a series of steps. Firstly, HF2

-

ions dissociated into HF monomers plus F- ions near to the surface. Secondly, F- ionsattacked the -SiF3 surfaces and H ions attacked the O atoms of the back-bond. Thirdly,the O-Si-F bond-angle decreased and the Si-O bond strength was weakened. Finally, theH atoms passivated the O atoms after tetrahedral SiF4 molecules were generated and theSi-O back-bonds were broken. The activation energy for the desorption of SiF4 wasestimated to be 0.8eV.T.Oku, K.Sato, M.Otsubo: Japanese Journal of Applied Physics - 1, 1997, 36[3B], 1374-9

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SiO2: Electron Irradiation, Neutron Irradiation, and Point DefectsPhotoluminescence data, on amorphous samples, in the spectral range of the a (4.3eV), ß(3.1eV) and ? (2.7eV) emissions - as excited using synchrotron radiation - were reported.Neutron irradiation gave rise to photoluminescence features which were distinct fromthose observed in either non-irradiated Ge-doped silica or in samples which contained

132

SiO2 Oxides SiO2

other impurities. The present data suggested that the defect-formation processesdetermined the various photoluminescence patterns which were observed, and that theusually cited distinction between Si-like and Ge-like centers was irrelevant.F.Meinardi, A.Paleari: Physical Review B, 1998, 58[7], 3511-4

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SiO2: Electron Irradiation and Point DefectsThe irradiation-sensitive defect structures of pure anhydrous and hydrated fusedamorphous samples were investigated by using cathodoluminescence micro-analysis.Irradiation with a continuous stationary electron beam resulted in a sub-surface trapped-charge induced electric field which caused the electromigration of mobile charged defectspecies within the volume of the irradiated specimen. Cathodoluminescence emissions,observed between 300 and 900nm at temperatures ranging from 5 to 295K, werecorrelated with particular defect centers; including the non-bridging O-hole centers withstrained bond and/or non-bridging hydroxyl precursors (in hydrated specimens), the self-trapped exciton, O-deficient centers (such as the neutral O vacancy and/or the 2-foldcoordinated Si defect) and the charge-compensated substitutional Al center.M.A.S.Kalceff: Physical Review B, 1998, 57[10], 5674-83

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SiO2: Gamma Irradiation and Point DefectsFilms of Si-rich oxide, with a thickness of about 1.2µ, were deposited onto p-type Sisubstrates by means of radio-frequency magnetron sputtering. After annealing at varioustemperatures in a N2 atmosphere, the photoluminescence spectra of all of the samplesexhibited 2 main peaks; at about 710 and 800nm. After ?-irradiation, these 2photoluminescence peaks increased in intensity by 3 to 5 times. A strong new 580nmpeak also appeared. The positions of all 3 photoluminescence peaks underwent no evidentshift when the measurement temperature was increased from 10 to 300K. The resultswere explained in terms of a model which assumed that photo-emission occurred due toluminescence centers, rather than nm-sized Si particles, in the Si-rich oxide films.S.Y.Ma, B.R.Zhang, G.G.Qin, Z.C.Ma, W.H.Zong, X.T.Meng: Materials ResearchBulletin, 1997, 32[10], 1427-33

[446-164-132]

SiO2: Ion Bombardment and Point DefectsThe amorphization of a-quartz by bombardment with light and medium ions wasinvestigated by means of Rutherford back-scattering spectrometry channelling, resonantnuclear reaction analysis and mechanical surface profiling. Monocrystalline samples werebombarded with 10 to 100keV H+, N+, Ne+ or Na+ ions, at 77K, to fluences of up to1016/cm2. Resultant disordering of the quartz was found to occur via the 3-dimensionalnucleation and growth of defect agglomerates, and perhaps small spatially-separatedamorphous zones in the surrounding crystalline material. The nucleation rate was relatedto the energy density which was deposited in elastic collisions. Above a critical energydensity of 1.92eV/atom (corresponding to 0.04dpa), deposited by nuclear collisions, a

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SiO2 Oxides SiO2

coherent amorphous layer formed at the depth of maximum energy deposition. This thengrew towards the surface, and towards greater depths, during continued bombardment.Mechanical surface profiling revealed the presence of large compressive stresses (of theorder of 1.5GPa). These built up at low fluences and were then released at higher fluenceswhen the amorphous layer extended as far as the surface. Such stress release wasaccompanied by a significant (about 19%) reduction in the atomic density.F.Harbsmeier, W.Bolse: Journal of Applied Physics, 1998, 83[8], 4049-53

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SiO2: Point DefectsIt was pointed out that measurements of O-vacancy creation in Si/SiO2/Si structuresduring high-temperature annealing, which had suggested an activation energy of 1.5eVfor the process, had been interpreted in terms of a simple thermodynamic model. It wasdemonstrated here that this model was inconsistent with thermochemical calculationswhich had indicated that the energy required for this process was 4.5eV. Another process,which involved thermally-induced O out-diffusion at the SiO2/Si interface, had aneffective activation energy, for O-vacancy creation, of about 2.0eV. This was consideredto be more consistent with the experimental data.R.A.B.Devine, W.L.Warren, S.Karna: Journal of Applied Physics, 1998, 83[10], 5591-2

[446-164-133]

SiO2: Point DefectsDefect concentrations in amorphous material, which were created by the implantation of10MeV protons, were examined. The depth profiles of Si-Si bonds, E' centers, and peroxyradicals were close to that of the electronic energy loss. Interstitial O2 molecules wereidentified, and their concentration was found to be larger than that of the peroxy radicals.The total concentrations of Si-Si bonds and E' centers were comparable to those of theinterstitial O2 molecules and peroxy radicals. The present results provided experimentalevidence that O Frenkel defect formation occurred in the amorphous oxide due to denseelectronic excitation. The efficiency of Frenkel defect formation was estimated to beabout 5 x 10-7/eV.H.Hosono, H.Kawazoe, N.Matsunami: Physical Review Letters, 1998, 80[2], 317-20

[446-164-133]

SiO2: Point DefectsFirst-principles configuration interaction calculations were made of the optical transitionsof the E' center. This was a hole which was trapped at an O vacancy, (-O)3Si•+Si(O-)3, insilica. Two competing excitation mechanisms were found. These were the promotion ofan electron from an O(2p) valence band orbital into the singly-occupied Si dangling bond,and charge-transfer transitions from (-O)3Si• to +Si(O-)3. The 2 excitations occurred atsimilar energies, (about 5.9eV), but only the charge-transfer process had a high intensity.The excitation was followed by a complex non-radiative decay process. This wassuggested to explain the absence of luminescence from this center.

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SiO2 Oxides SiO2

G.Pacchioni, G.Ieranò, A.M.Márquez: Physical Review Letters, 1998, 81[2], 377-80[446-164-134]

SiO2: Point DefectsBy using synchrotron radiation as a photon source, photoluminescence spectra wereobtained from buried oxides that had been formed by O implantation. On the basis of thespectra, the oxide was deduced to contain relaxed and unrelaxed neutral O vacancies inconcentrations of 1.4 x 1020 and 1017/cm3, respectively.K.S.Seol, T.Futami, Y.Ohki: Journal of Applied Physics, 1998, 83[4], 2357-9

[446-164-134]

SiO2: Point DefectsThe defects which were generated during irradiation with energetic (10eV) photons werefound to trap electrons at a level which was 3.1eV below the oxide conduction band. Theelectron spin resonance data, and the behavior after H passivation, indicated that theoptically active state could be attributed to a H-complexed O vacancy. The observedinjection of electrons to these traps, from Si, suggested that the revealed defects were thepossible origin of a degradation-induced electrical conduction in thin SiO2 layers.V.V.Afanasev, A.Stesmans: Applied Physics Letters, 1997, 71[26], 3844-6

[446-164-134]

SiO2: Point DefectsTwo photoluminescence bands were detected from Si+-implanted films after rapid thermalannealing at temperatures above 950C. The photoluminescence band at 2.2eV wasobtained after rapid thermal annealing in dry N, and the other one (1.9eV) was obtainedafter rapid thermal annealing in wet N. The luminescence at 2.2eV disappeared when thefilms were re-annealed in an electric oven at temperatures above 600C. This was similarto the behavior of O-deficient and H-deficient structures, and this photoluminescenceband was therefore attributed to the Ed' center. The 1.9eV band, being closely related toSi-O-H structures and persisting after re-annealing at up to 800C, was attributed to theeffect of non-bridging O hole centers.S.T.Chou, J.H.Tsai, B.C.Sheu: Journal of Applied Physics, 1998, 83[10], 5394-8

[446-164-134]

SiO2: Point DefectsFirst-principles calculations, including electron correlation and cluster models, were madeof the formation energies of a series of defects in the bulk oxide and in silica glasseswhich contained Ge impurities. The defects which were considered included O vacancies,=Si-Si=, =Si-O-Si=, =Si-O-O•, =Si-O•, =Si•, E' centers, =Si•+Si=, double O vacancies, =Si-Si-Si=, Frenkel pairs, =Si-Si-O-O-Si=, et cetera. The corresponding analogues, in whichGe replaced network Si atoms, were also considered. It was found that the formation of asingle O vacancy, defined as being the energy required to remove and take to infinity aneutral O atom, was 8.5eV. This was consistent with recent thermodynamic

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SiO2 Oxides SiO2

estimates. It was shown that, in the presence of Ge impurities, the formation of O-deficient centers occurred at lower energies.G.Pacchioni, G.Ieranò: Physical Review B, 1997, 56[12], 7304-12. See also: PhysicalReview B, 1998, 57[2], 818-32

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SiO2: Point DefectsThe charge-state dependent structural stability of the O vacancy in a-quartz wasdetermined by using first-principles total-energy calculations. It was found that neutraland positively charged O vacancies exhibited bistability, and that a spontaneous structuraltransformation occurred for doubly positively and doubly negatively charged states. Thestructural transformation introduced a new electron trap, and could lead to thedegradation of thin films.A.Oshiyama: Japanese Journal of Applied Physics - 2, 1998, 37[2B], L232-4

[446-164-135]

SiO2: TwinsThe transient characteristics of x-axis inversion in AT-cut quartz plates were studied. Theresonance frequency of the quartz plate was monitored throughout the heat treatmentprocess which induced x-axis inversion. The inversion could be detected via a resonantfrequency change. It was observed that the critical temperature for x-axis inversiondecreased with film thickness. An expansion of the x-axis inverted area was observed attemperatures greater than the critical one.T.Uno: Japanese Journal of Applied Physics - 1, 1997, 36[5B], 3000-3

[446-164-135]

SiO2/Si: Interface DefectsThe passivation properties of local-oxidation-of-silicon oxide/Si interface defects wereinvestigated by means of reverse-current measurements and capacitance transientspectroscopy of pn junction diodes that had a large LOCOS-defined perimeter. TheLOCOS/Si interface defects had some properties which were similar to those of theSiO2/Si(100) interface states of metal-oxide-silicon diodes. However, there was asignificant difference between the 2 interfaces with regard to the levels of unpassivateddefects which remained after H2 annealing. The level was higher for a LOCOS/Siinterface than for a MOS interface.S.Fujieda, H.Nobusawa, M.Hamada, T.Tanigawa: Journal of Applied Physics, 1998,84[5], 2732-4

[446-164-135]

SiO2/Si: Interface DefectsThe correlation between detrimental electrically-active interface traps, and electron spinresonance-active point defects - Pb0 and Pb1 (unpaired Si orbitals) - was studied via

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SiO2 Oxides SiO2

controlled variations of the defect bath densities. No electrical activity of Pb1, as aninterface state, could be detected, while all of the Pb0 defects appeared to be active.A.Stesmans, V.V.Afanasev: Physical Review B, 1998, 57[16], 10030-4

[446-164-136]

SiO2/Si: Interface DefectsThe observed electron spin resonance hyperfine spectra which were associated with theunpaired electron of the Pb1 interface defect in thermal SiO2/(100)Si showed that thepredominant interaction arose from a single 29Si isotope. The hyperfine tensor exhibited aweakly monoclinic (nearly axial) symmetry, with the principal axes of the g andhyperfine tensors coinciding. A molecular orbital analysis indicated that the unpairedelectron resided, to the extent of some 58%, in a single unpaired Si hybrid orbital. Thiswas found to be 14% s-like and 86% p-like, with the p-orbital closely aligned with a<211> direction; at 35.26º to the [100] interface normal. It was noted that O did notconstitute an immediate part of the defect, and the results firmly established that a keypart of the Pb1 defect was a tilted (about 20° about <011>) Si3=Si• unit.A.Stesmans, B.Nouwen, V.V.Afanasev: Journal of Physics - Condensed Matter, 1998,10[27], L465-72

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SiO2/Si: Interface DefectsLow-energy cathodoluminescence spectroscopy was used to study optical transitions atdefect bonding arrangements in Si/SiO2 interfaces which had been prepared by means oflow-temperature plasma deposition. Variable-depth excitation, obtained by using variouselectron injection energies, provided a clear distinction between luminescence from thenear-interface region of the oxide film, from the Si/SiO2, and from the underlyingcrystalline Si substrate. Cathodoluminescence bands at about 0.8 and 1eV were attributedto interfacial Si atom dangling bonds with various numbers of back-bonded Si and Oatoms. Cathodoluminescence spectroscopy also revealed higher photon-energy features.These included bands, at about 1.9 and 2.7eV, which were attributed to sub-oxidebonding defects in the as-grown oxide films. There was also a substrate-related feature atabout 3.4eV. Hydrogenation at 400C, and/or rapid thermal annealing (900C), markedlyreduced the intensities of cathodoluminescence spectroscopy features which wereattributed to interfacial and sub-oxide bonding defects.J.Schäfer, A.P.Young, L.J.Brillson, H.Niimi, G.Lucovsky: Applied Physics Letters, 1998,73[6], 791-3

[446-164-136]

SiO2/Si, SiO2/SiC: Interface DefectsThe exposure of these interfaces to H, at temperatures ranging from 450 to 800C, wasfound to produce a considerable density (up to 1013/cm2) of positively charged centers.The absence of any correlation between the charging process, and the presence of Sidangling-bond centers in SiO2 (or at the Si/SiO2 interface), was assumed to indicate H

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SiO2 Oxides SiO2

bonding in a valence alternation state which was suggested to be an over-coordinated Ocenter, [Si2=OH]+, that was stabilized by SiO2 network rearrangement at the interface.V.V.Afanasev, A.Stesmans: Physical Review Letters, 1998, 80[23], 5176-9

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SiO2-GeO2: UV Irradiation and Point DefectsMicrostructural changes in glasses (5 or 13mol%GeO2) which had been irradiated with 5or 6.4eV light were investigated by means of Raman spectroscopy. The low-frequencyso-called boson bands at 50/cm, in both samples, shifted upwards during irradiation, buttheir intensities exhibited opposite changes. This indicated that the thermal damage to thesurface, caused by energetic ultra-violet photons, was related to the Ge content of theglass. The intensities of the D1 and D2 defect lines increased and shifted to higherfrequencies; thus reflecting a reduction in the Si-O-Si bond angle during irradiation. Thiswas suggested to be due to a change, in ring statistics, in favor of smaller rings. That is, 6-fold rings transformed to 3-fold and 4-fold rings during ultra-violet irradiation. Theopposite changes in intensity of the defect lines were attributed to variations in thenetwork structure.F.X.Liu, J.Y.Qian, X.L.Wang, L.Liu, H.Ming: Physical Review B, 1997, 56[6], 3066-71

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SiO2-GeO2-: Point DefectsThe photo-absorption of neutral O vacancy defects in silicate and germanosilicate glasseswas studied by using first-principles quantum-chemical techniques. The lowest singlet-singlet excitations in these defects occurred at about 7eV and involved the promotion ofan electron, from a bonding orbital between adjacent Si (or Ge) atoms, to a diffuseRydberg-type orbital. Such excitations were too high in energy to contribute significantlyto the 5eV absorption band of silicate and germanosilicate glasses.B.B.Stefanov, K.Raghavachari: Physical Review B, 1997, 56[9], 5035-8

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SiO2-GeO2: Point DefectsA 15.7mol%GeO2 glass was prepared by using a vapor axial deposition method, andoptical absorption at 4.92 and 5.08eV was observed between 300 and 600K. The resultswere then compared with electron spin resonance data for the same samples. Theabsorption at 4.92eV was found to increase with increasing temperature, while the opticalabsorption at 5.08eV decreased with increasing temperature. These changes in opticalabsorption indicated that the concentration of the neutral O mono-vacancy was reduced athigher temperatures. The changes were reversible. The concentration of the Ge E' center,as estimated from electron spin resonance data, increased with increasing temperature; incontrast to a decrease in the neutral O mono-vacancy. The results strongly suggested theexistence of thermal equilibrium reactions between neutral O mono-vacancies and Ge E'centers.

138

SiO2 Oxides (Sr,La)CoO3

M.Takahashi, T.Fujiwara, T.Kawachi, A.J.Ikushima: Applied Physics Letters, 1998,72[11], 1287-9

[446-164-138]

SiO2-GeO2: Point DefectsAn investigation was made of the formation of Ge-related defects in glass fiber which waspreformed by poling with ArF laser excitation. The electric field dependence of theinduced defect concentrations was measured by using optical absorption techniques.Color centers, such as Ge electron-trapped centers and Ge E' centers, were introduced.The concentrations of induced Ge electron-trapped centers and Ge E' centers increasedwith increasing electric field. The conversion efficiency of Ge electron-trapped centers toGe E' centers was found to be independent of the electric field. The results stronglysuggested that the ArF laser excitation was effective in forming Ge electron-trappedcenters.M.Takahashi, T.Fujiwara, T.Kawachi, A.J.Ikushima: Applied Physics Letters, 1997,71[8], 993-5

[446-164-138]

SiO2-Li2O-TiO2-BaO-La2O3: H DiffusionA commercial pH-responsive glass (SiO2-25Li2O-7TiO2-5BaO-3mol%La2O3) wassubjected to heat treatment in a vacuum furnace, and the pH response of the heat-treatedspecimen was compared with that of as-received glass. It was found that the heat-treatedspecimen, in which the content of so-called mobile H ions had been decreased, exhibiteda much lower Nernst slope than the ideal value (59mV). It was concluded that thegeneration of pH selectivity originated from a H-ion concentration cell between the testsolution and the glass electrode, and that an essential factor required for a glass specimento exhibit a pH response was that the glass should contain mobile H ions.Y.Abe, M.Maeda, N.Hayakawa, H.Umehara, M.Nogami: Materials Research Bulletin,1997, 32[11], 1535-42

[446-164-138]

(Sr,La)CoO3: O Diffusion and Ionic ConductionThe permeability of O in dense membranes of Sr0.7La0.3CoO3 was studied, attemperatures of between 750 and 1100C, under various O partial pressure gradients. Theresults indicated that, for thicknesses of 0.057 to 0.215cm, the O flux was controlledmainly by bulk diffusion across the membrane. The activation energy for this process wasestimated to be 60kJ/mol. The ionic conductivity was equal to 0.5S/cm at 1000C.C.H.Chen, H.J.M.Bouwmeester, R.H.E.Van Doorn, H.Kruidhof, A.J.Burggraaf: SolidState Ionics, 1997, 98[1-2], 7-13

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139

(Sr,La)TiO3 Oxides SrTiO3

(Sr,La)TiO3: Point DefectsPerovskite-type ceramics, which had been sintered in pure O at 1400C, were investigatedas a function of La contents ranging from 0 to 0.6. Optical and scanning electronmicroscopy, X-ray analysis, X-ray diffraction and pyconometry were used to determinethe mechanism that was responsible for the compensation of the electronic excess chargethat resulted from adding La. Pure Sr-vacancy compensation mechanism was observed atLa contents of up to 0.3. At La contents above 0.4, Ti vacancies also occurred but theirconcentration remained negligible when compared to that of the predominating Srvacancies. No indication of a reported solubility limit of La, at a content of 0.4, wasobserved. At La contents of 0.5 and 0.6, the lattice structure was slightly distorted:tetragonally and orthorhombically, respectively.R.Moos, T.Bischoff, W.Menesklou, K.H.Härdtl: Journal of Materials Science, 1997,32[8], 4247-52

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(Sr,La)3(VO4)2: Ba, Ca, Sr Diffusion and Point DefectsSerial-sectioning and radio-tracer techniques were used to monitor the diffusion of 90Sr,45Ca and 133Ba in Sr 3-3x La2xo x(VO4)2 solid solutions, where x ranged from 0 to 0.25. Itwas found that the diffusivities decreased in the order: Ca - Sr - Ba. An increase in thevacancy concentration, o , on the Sr sub-lattice led to an increase in the diffusivity of therare-earth cations. It was shown that doubly-charged cations migrated via Sr(2) positions.O.N.Leonidova, G.I.Dontsov, I.A.Leonidov, A.C.Zhukovskaya: Fizika Tverdogo Tela,1998, 40[2], 223-6 (Physics of the Solid State, 1998, 40[2], 200-3

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SrTiO3: H DiffusionThe effects of doped acceptor ions upon proton diffusion in Sc-doped samples werestudied by means of first-principles molecular dynamics simulations. It was found that theproton formed an O-H bond with the neighboring O ion, and that the frequency of the O-H stretching vibration depended upon the position of the proton in the crystal. Thefrequencies which were deduced from the simulations were consistent with experimentalresults obtained from infra-red transmission spectra. It was shown that the positionaldependence of the O-H stretching vibration was caused by the electron densitydistribution in the Sc-doped samples. Near to Sc ions, electrons tended to localize aroundeach ion and produced a higher frequency while, near to Ti ions, the electron densitybetween the Ti and O ions was larger than that in the undoped crystal; thus giving rise tolower frequencies.F.Shimojo, K.Hoshino, H.Okazaki: Journal of Physics - Condensed Matter, 1998, 10[2],285-94

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140

SrTiO3 Oxides SrTiO3

y SrTiO3: O DiffusionDiffusion profiles were recorded in situ and were evaluated by using an opticaltechnique. It was shown that theoretical calculations could reproduce the experimentaldata (figure 9), without using adjustable parameters, if account was taken of coupling ofthe diffusing species to internal redox changes in the dopants. Measurements on singlecrystals and on bicrystals, with and without crack formation, provided valuableinformation concerning the effect of free relaxed surfaces, freshly produced cracksurfaces, and grain boundaries upon diffusion. In particular, the inward diffusion of O viacrack surfaces was characterized by enhanced diffusion coefficients.I.Denk, F.Noll, J.Maier: Journal of the American Ceramic Society, 1997, 80[2], 279-85

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Figure 9Diffusivity of O in SrTiO3

SrTiO3: DislocationsTransmission electron microscopy was used to study structural defects in Verneuil-grownsingle crystals. The dislocation density was observed to decrease with increasing depth

1.0E-06

1.0E-05

1.0E-04

10 11 12 13 14 15

4.2E19Fe/cm36.9E18Fe/cm34.2E18Fe/cm3

104/T(K)

D (c

m2 /s

)

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SrTiO3 Oxides SrTiO3

from the original cut surface of the crystals. The high density of dislocations in the skinregion was suggested to be responsible for a second characteristic length scale which wasdetected when using various beam methods.R.Wang, Y.Zhu, S.M.Shapiro: Physical Review Letters, 1998, 80[11], 2370-3

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SrTiO3: Grain BoundariesThe effects of annealing upon the grain boundary of a 36.8º symmetrical [100] tiltbicrystal were studied. Scanning tunnelling microscopy and atomic force microscopywere used for the non-destructive observation of boundary structures. It was found thatannealing the bicrystalline substrates at temperatures as low as 780C led to the formationof grooves at their boundaries. This provided direct evidence that the thickness depressionof Cu3Ba2YO7 films at the bicrystal boundaries originated from the underlying groovedsubstrates. Hole-like defects, with diameters ranging from some 30 to 200nm, were alsoobserved.O.D.Jiang, Z.J.Huang, A.Brazdeikis, M.Dezaneti, C.L.Chen, P.Jin, C.W.Chu: AppliedPhysics Letters, 1998, 72[25], 3365-7

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SrTiO3: Grain BoundariesThe temperature dependence of the grain-boundary potential barrier height, and of theconductivity across the grain-boundary space-charge depletion layer in acceptor-dopedceramics, was investigated by using a numerical simulation technique. The underlyingmodel was that of a back-to-back double Schottky barrier at the grain boundary. Aninterpretation was offered in the form of a defect chemistry model for the bulk, and forthe space-charge depletion layer which surrounded the grain-boundary core on both sides.The temperature behavior of the potential barrier at the grain boundary could be dividedinto 3 different regimes. These were a linear regime, followed by saturation anddecreasing regimes. Two different spatial conductivity profiles at the grain boundarieswere identified with 2 different characteristic thermal activation energies for the effectivegrain-boundary conductivity. The barrier height itself was not equal to the thermalactivation energy of the effective grain-boundary conductivity. The electricalcharacteristics of the grain boundaries could be deliberately influenced by decorating theboundaries with suitable dopants.R.Hagenbeck, R.Waser: Journal of Applied Physics, 1998, 83[4], 2083-92

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SrTiO3: Grain Boundaries and Point DefectsA high-resolution transmission electron microscopic study was made of the atomisticstructure of a S = 3 (111) grain boundary. Quantitative evaluation of the high-resolutiontransmission electron microscopic images revealed that shear stresses, which arose fromprocessing or external loading, had an appreciable effect upon the translation state of the2 adjacent grains and upon the ionic positions at the grain boundary. Under a low shearstress, the boundary exhibited mirror symmetry with respect to the boundary plane, and a

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SrTiO3 Oxides Sr(Ti,Zr)O3

comparatively large free volume. The minimum spacing between Ti ions in the {111}layers which neighbored the boundary plane caused an expansion of 0.06nm with respectto the geometrical model. On the other hand, a high shear stress (of the order of 740MPa)transformed the structure into a so-called lock-in configuration which had no mirrorsymmetry and a smaller excess volume. This demonstrated that shear stresses could havean important effect upon the atomistic structure of the grain boundaries. Such a change instructure could alter the free energies which were associated with the formation of pointdefects, and thus influence charge accumulation on the boundary plane; as well as thecorresponding space-charge layers.O.Kienzle, F.Ernst: Journal of the American Ceramic Society, 1997, 80[7], 1639-44

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SrTiO3: Point DefectsThe electrical conductivity of donor-doped and undoped ceramics and of single crystalsof Sr1-xLaxTiO3, where x ranged from 0 to 0.1, was investigated at temperatures rangingfrom 1000 to 1400C; under O partial pressures of 10-21 to 1bar. By incorporatingpublished Hall-effect and thermopower data, a set of parameters for a defect-chemicalmodel was determined. This model precisely described the point-defect concentrationsand transport properties of these materials. It was shown that defects in the cation sub-lattice completely governed the electrical behavior of donor-doped and undoped material.In the latter case, frozen-in Sr vacancies acted as intrinsic acceptors.R.Moos, K.H.Härdtl: Journal of the American Ceramic Society, 1997, 80[10], 2549-62

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SrTiO3/LaAlO3: DislocationsSamples of Sr-deficient and stoichiometric epitaxial (001) film, deposited onto(110)rhombohedral LaAlO3 substrates by means of radio-frequency magnetron sputtering,were characterized by using high-resolution transmission electron microscopy.Subsequent heat treatment in O had a positive effect upon the dielectric properties. TheSr-deficiency had a large negative effect upon the microwave dielectric constant of thefilms. These changes were related to changes in lattice parameter. Misfit dislocationswere present at the film/substrate interface of all of the samples. The residual elasticstrain compressed the SrTiO3 unit cell in the substrate surface plane, and expanded it byan equal amount in the [001] direction. The use of X-ray diffraction revealed that atetragonal distortion due to mismatch strain was restricted to the narrow region which wasclosest to the film/substrate interface.L.Ryen, E.Olsson, L.D.Madsen, X.Wang, C.N.L.Edvardsson, S.N.Jacobsen,U.Helmersson, S.Rudner, L.D.Wernlund: Journal of Applied Physics, 1998, 83[9], 4884-90

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Sr(Ti,Zr)O3/SrTiO3, SrZrO3/SrTiO3, BaZrO3/SrTiO3, SrZrO3/LaAlO3: DislocationsCubic perovskite films with various compositions were grown onto perovskite substrates,using the chemical solution deposition method, in order to investigate the effect of a large

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Sr(Ti,Zr)O3 Oxides TiO2

lattice mismatch. The films were pyrolyzed in order to crystallize the perovskite with therequired composition, and were then heated to 1000C so as to promote epitaxial graingrowth. Pole figures were obtained by X-ray diffraction, and selected area diffractionstudies were made of transmission electron microscopy specimens with latticemismatches of 2.5% (SrTi0.5Zr0.5O3 on SrTiO3), 5% (SrZrO3 on SrTiO3), 7.4% (BaZrO3

on SrTiO3) and 8.2% (SrZrO3 on LaAlO3). This revealed only the orientation,[100](001)film||[100](001)substrate. The X-ray diffraction and transmission electronmicroscopy results also indicated an increasing polycrystallinity with increasing latticemismatch. High-resolution electron microscopy of films on SrTiO3 demonstrated that anarray of misfit dislocations was present at the interface. The misfit dislocations had linevectors of <100>-type and Burgers vectors of <010>-type. The dislocation separationdistances which were obtained by using high-resolution electron microscopy and X-raydiffraction lattice parameter measurements showed that the strain energy within the films,due to lattice mismatch, was almost fully relaxed.P.A.Langjahr, F.F.Lange, T.Wagner, M.Rühle: Acta Materialia, 1998, 46[3], 773-85

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Ta2O5: Point DefectsThe defect states which were responsible for leakage currents in ultra-thin (less than10nm) Ta2O5 films were measured by using a novel zero-bias thermally stimulatedcurrent technique. It was found that defect states, whose activation energy was estimatedto be about 0.2eV, could be more efficiently suppressed by using N2O rapid thermalannealing instead of O2 rapid thermal annealing for post-deposition treatment. Theleakage current was also smaller for samples given N2O rapid thermal annealing than forthose subjected to O2 rapid thermal post-deposition annealing. Hence, the above defectstates were expected to be important in causing leakage currents.W.S.Lau, L.Zhong, A.Lee, C.H.See, T.Han, N.P.Sandler, T.C.Chong: Applied PhysicsLetters, 1997, 71[4], 500-2

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TiO2: O, OH, Ti DiffusionA sputter-deposited Ti-20at%W alloy was anodized in order to determine the effect ofalloy species upon the structures of anodic films on Ti. It was of interest that anamorphous oxide film developed on the alloy at voltages of up to about 200V. Such filmformation, on relatively pure Ti, occurred at voltages of less than 20V. The amorphousfilms consisted of an outer TiO2 layer, and an inner layer which contained TiO2 and WO3.The layered films formed due to the migration of Ti4+ ions at a rate which was about 2.4times higher than that of W6+ ions. This was attributed to the relative strengths of W6+-Oand Ti4+-O bonds. Film growth proceeded via the migration of both O2- and/or OH- ionsand metal cations. This was typical of amorphous anodic films, with the transport numberof cations being determined, using a Si species marker, to be 0.39.

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TiO2 Oxides TiO2

H.Habazaki, K.Takahiro, S.Yamaguchi, K.Shimizu, P.Skeldon, G.E.Thompson,G.C.Wood: Philosophical Magazine A, 1998, 78[1], 171-87

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TiO2: Surface DiffusionThe diffusivity under a gas pressure of 100MPa was deduced by simulating a sinteringmodel which involved surface and volume diffusion. The value of the bulk diffusioncoefficient, as estimated by using the present simulation, was similar to the O diffusioncoefficient in the oxide. The present work showed that the surface diffusivity (2.2 x 10-

16m2/s at 850C) under high-pressure Ar gas was about 30 times higher than that under1atm.M.Nanko, K.Ishizaki: Physical Review B, 1997, 56[11], 6965-9

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TiO2: Dislocations and TwinsTransmission electron microscopy was used to study the crystal size distribution, grain-boundary disorder and defect structure in nanocrystalline specimens which had beenprepared by means of reactive sputtering. The average grain diameter, as determined bymeasurements of dark-field micrographs, was about 15nm. Evidence of both ordered anddisordered grain-boundary regions was found, and planar defects which were observedwithin grain interiors were identified as being (011) deformation twins. Thecrystallographic shear defects, that could occur as a result of the aggregation of Ovacancies in the hypo-stoichiometric oxide, were rarely present. No dislocations, apartfrom those required in order to accommodate misfit, were seen.D.G.Rickerby: Philosophical Magazine B, 1997, 76[4], 573-83

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TiO2: Grain BoundariesA detailed study was made of the atomic structure of a [001] tilt grain boundary, S = 5(210) 36.8º, in rutile and Z-contrast imaging was used to obtain a 2-dimensional map ofthe cation positions at the interface. Details of the charge state of the cations and theatomic structure around anion sites were then obtained by using electron energy-lossspectroscopy. The spectroscopic data for O were interpreted by using multiple scatteringtheory to obtain 3-dimensional structural information. Use of the combined techniquespermitted a unique grain boundary structure to be defined.D.J.Wallis, N.D.Browning, P.D.Nellist, S.J.Pennycook, I.Majid, Y.Liu, J.B.VanderSande: Journal of the American Ceramic Society, 1997, 80[2], 499-502

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TiO2: Point DefectsAnatase-rutile phase transformation in the presence of Fe2O3 was investigated in air or Aratmospheres by using X-ray diffraction and scanning electron microscopy techniques.Isothermal anatase-rutile transformation, as a function of time, was monitored between

145

TiO2 Oxides UO2

825 and 950C. The data could be closely fitted by using various rate laws. In the presenceof Fe2O3, the transition temperature was lowered and the transformation rate in air wasincreased. Transformation in the presence of Fe2O3, under an Ar atmosphere, was morerapid than it was in air. The enhancement effect of Fe2O3 was explained in terms of theformation of O vacancies. These were produced by Fe3+ diffusion in the TiO2 lattice. Itwas concluded that the enhancement effect was greater in an Ar ambient because Fe3+

diffusion in non-stoichiometric TiO2 was some 2 orders of magnitude faster than instoichiometric TiO2.F.C.Gennari, D.M.Pasquevich: Journal of Materials Science, 1998, 33[5], 1571-8

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TiO2: Surface DefectsAtomically resolved images of a (110)-(1 x 1) surface were obtained by means of non-contact atomic force microscopy under ultra-high vacuum. In contrast to previousscanning tunnelling microscopic studies, the outermost atoms of bridge-bound O ridgeswere seen as protruding rows. A high-resolution image of the surface revealed that thebridging O atoms ordered, with a (1 x 1) periodicity, on terraces. Point O-atom defectswere also seen as dark spots.K.Fukui, H.Onishi, Y.Iwasawa: Physical Review Letters, 1997, 79[21], 4202-5

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Ti(Sr,Ca)O3: Planar DefectsThe microstructure and dielectric properties of A-site excess and stoichiometric sampleswere investigated. Ruddlesden-Popper type planar faults in A-site excess specimens wereobserved by using high-resolution scanning transmission electron microscopy. Scanningtransmission electron microscopy micro-chemical analysis, calculation of high-resolutionimages using a multi-slice method, and structural energy calculations, revealed that Caions selectively occupied the cation sites of the faults. An observed shift of the peakdielectric constant to lower temperatures, and an expansion of the lattice parameter of theA-site excess specimens, were attributed to selective occupation of the cation sites of theplanar faults by Ca ions.M.Fujimoto, T.Suzuki, Y.Nishi, K.Arai, J.Tanaka: Journal of the American CeramicSociety, 1998, 81[1], 33-40

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UO2: Point DefectsThe formation energies of O and U vacancies, interstitials and Frenkel pairs, as well asSchottky trio defect energies, were calculated by using the linear muffin-tin orbital super-cell method in the atomic sphere approximation. These ab initio results were found to bein qualitative agreement with experiment, and they confirmed the results of previouscalculations that were based upon semi-empirical interaction models. A discrepancybetween certain experimental and theoretical values was not resolved.

146

UO2 Oxides YAG

T.Petit, C.Lemaignan, F.Jollet, B.Bigot, A.Pasturel: Philosophical Magazine B, 1998,77[3], 779-86

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WO3: Point DefectsModelling and simulations were used to explain the diffraction patterns of O-deficient Ti-doped WO3 thin films. A lack of 2-dimensional and 3-dimensional correlations which wasobserved in the structure was attributed to a random spacing, between {103} planes,which maintained the defect density of the WO2.9 ordered structure. The phase could beregarded as being an intermediate case between the crystalline and amorphous WO3

phases, due to freezing of the in-plane rotational degree of freedom of the [WO6]6-

structural unit.L.Sangaletti, L.E.Depero, E.Bontempi, R.Salari, G.Sberveglieri: Journal of Solid StateChemistry, 1997, 131[2], 215-20

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WO3: TwinsSheet conductivity was detected in O-reduced samples. The sheets were aligned along thetwin boundaries of the non-reduced starting material. The bulk transformed, during Oloss, to give an O-deficient phase with a tetragonal crystal structure: P4̄21m, a = 0.739, c= 0.388nm. The perovskite-like structure contained distorted octahedra with W-Odistances ranging from 0.17 to 0.218nm.A.Aird, M.C.Domeneghetti, F.Mazzi, V.Tazzoli, E.K.H.Salje: Journal of Physics -Condensed Matter, 1998, 10[33], L569-74

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WO3: TwinsSingle crystals which contained ferro-elastic twin walls were deoxygenated via gastransport reaction with Na vapor. The reaction product consisted of a non-superconducting matrix of tetragonal oxide, with superconducting twin boundaries. Thesuperconducting transition temperature was 3K. The upper critical field temperaturedependence exhibited BCS-type behavior below 2.5K, and a weak tail near to the criticaltemperature.A.Aird, E.K.H.Salje: Journal of Physics - Condensed Matter, 1998, 10[22], L377-80

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YAG: Dislocations and Point DefectsA perturbed F-center, with absorption bands at 370 and 625nm, which was associatedwith rare-earth ions was identified in Yb-doped samples. The dislocation density wasestimated to be lower than 100/cm2. Signs of Er3+ and Ho3+ were detected in theabsorption and emission spectra. Hydroxyl groups were also found.H.Yin, P.Deng, F.Gan: Journal of Applied Physics, 1998, 83[7], 3825-8

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147

Y2O3 Oxides YVO4

Y2O3: Point DefectsHigh-resolution X-ray and neutron scattering studies of powder samples of pure and Zr-doped (5at%) material revealed that mesoscopic compositional inhomogeneities existed,in the doped compound, which were unlikely to be revealed by conventional diffractionmethods. Even if structural refinement of the doped compound gave quite good reliabilityfactors within a single-phase model, a more careful analysis of the data proved theexistence of 2 phases. The main phase (90%) had compositions and structures which werevery close to those of the pure oxide. The phase which was richer in Zr exhibited siteselectivity. An abnormally high value (4.7Å2 at 294K) of the thermal displacementparameter of the site, as affected by Zr substitution, was explained in terms of a split-atom model which involved an off-site displacement of 0.38Å. Additional O atomsoccupied the vacancies in the structure, and provided charge balance for the Zr dopant.G.Baldinozzi, J.F.Bérar, M.Gautier-Soyer, G.Calvarin: Journal of Physics - CondensedMatter, 1997, 9[45], 9731-44

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Y2O3-Nb2O5: Ionic ConductionPhase equilibria in this system were studied, at temperatures of 1500 and 1700C, forcompositions ranging from 0 to 50mol%Nb2O5. The solubility limits of the C-type Y2O3

cubic phase and the YNbO4 monoclinic phase at 1700C were 2.5mol%Nb2O5 and0.2mol%Y2O3, respectively. The fluorite single phase existed between 20.1 and27.7mol%Nb2O5 at 1700C, and between 21.1 and 27.0mol%Nb2O5 at 1500C. Theconductivity of the system increased, as the Nb2O5 content was increased, up to amaximum at 20mol%Nb2O5. This was due to an increase in the fraction of fluorite phase.In the fluorite single-phase region, the conductivity decreased between 20 and25mol%Nb2O5. This was because of a decrease in the content of O vacancies. Theconductivity at 27mol%Nb2O5 was higher than that at 25mol%Nb2O5. The conductivitydecreased, between 27.5 and 50mol%Nb2O5, because of a decrease in the fraction of thefluorite phase. A 20mol%Nb2O5 sample exhibited the highest conductivity.J.H.Lee, M.Yashima, M.Kakihana, M.Yoshimura: Journal of the American CeramicSociety, 1998, 81[4], 894-900

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YVO4: Gamma Irradiation and Point DefectsDefects in powder samples of Eu3+-doped material, which had been prepared via solid-state diffusion, were studied by means of thermoluminescence and fluorescencemeasurements. The defects were generated by exposing the powder to ?-rays. The effectof radiation defects was studied by comparing the fluorescence spectra of non-irradiatedand ?-irradiated samples. The fluorescence output was found to decrease in the presenceof defects. Broad and weak characteristic bands at 425nm, due to Eu2+, were observed;together with Eu3+ emission lines at 575, 594 and 616nm. These were due to 5Do ? 7Fjtransitions, where j was equal to 0, 1 and 2, respectively. Two thermoluminescence peaks

148

YVO4 Oxides Zeolite

were found for doped powder, whereas only a single peak arose from pure samples. Thethermoluminescence emission was also greatly enhanced in the case of doped powder,and this increase was attributed to the presence of impurities. It was noted that defectaccumulation depended strongly upon the level of ?-ray exposure.R.B.Pode, S.J.Dhoble: Journal of Physics D, 1998, 31[1], 146-50

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YVO4: X-Irradiation and Point DefectsThe 3-dimensional thermoluminescence spectra of single crystals which had been grownfrom the melt by using the Czochralski technique, or pulled from a flux using the top-seeded solution growth method, were studied after irradiation. The thermoluminescencespectra were recorded for wavelengths ranging from 200 to 800nm, and temperaturesranging from 20 to 300K, before and after annealing in an O atmosphere. It was foundthat, in spite of the differing growth conditions, the 3-dimensional thermoluminescencespectra of the Czochralski samples exhibited similar thermoluminescence characteristics.The main thermoluminescence emission appeared around 450nm, at between 200 and250K. Further weak emissions were also detected at 570, 600, 650, and 710nm. Thesewere tentatively attributed to impurities. The thermoluminescence spectrum of the top-seeded solution grown crystal differed markedly from that of the Czochralski crystals.The major thermoluminescence peak appeared at lower temperatures here, whereas theemission spectrum exhibited a broader band around 500nm and the weak bands between500 and 700nm could not be detected. It was concluded that heat treatments, such asannealing in an O atmosphere, could modify the defect structures in 2 ways. Firstly, theycould compensate O vacancies and, secondly, they could produce additional precipitationdue to a limited solubility region in the quasi-congruent composition. Precipitates in theform of an Y or V excess further increased the defect content on anion and/or cation sitesin near-congruent crystals. Annealing in O increased most of the thermoluminescencepeak intensities (except for peaks at 210 and 241K). It was assumed that thecorresponding traps represented both cation and anion site defects.S.Erdei, L.Kovács, A.Peto, J.Vandlik, P.D.Townsend, F.W.Ainger: Journal of AppliedPhysics, 1997, 82[5], 2567-71

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Zeolite: Point DefectsEvidence for a body-centered cubic lattice of F centers was presented. A high electronspin density produced isotropic contact shifts, in the nuclear magnetic resonance spectraof the framework nuclei, whose magnitudes were a discrete function of the local electrondensity. Strong exchange coupling between unpaired electrons gave rise to anantiferromagnetic phase transition at 48K; thus providing the first example of an s-electron antiferromagnet.V.I.Srdanov, G.D.Stucky, E.Lippmaa, G.Engelhardt: Physical Review Letters, 1998,80[11], 2449-52

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149

ZnO Oxides ZnO

3 ZnO: O DiffusionIt was recalled that the models which were used to describe ZnO-based varistors reliedupon diffusion and defect data which had been obtained by using techniques that hadbeen superseded by methods such as secondary ion mass spectrometry. Values werereported here for O diffusivities in undoped monocrystalline ZnO as a function oftemperature and orientation. Evaporation was taken into consideration when analyzing theexperimental results. It was found that, to within experimental error, the energetics ofdiffusion were isotropic, but were slightly faster in the c-direction. The present results(table 5) could be described by:

a-direction: D (cm2/s) = 4.0 x 10-7exp[-2.22(eV)/kT]c-direction: D (cm2/s) = 9.0 x 10-6exp[-2.52(eV)/kT]

An analysis of previously published data suggested that the intrinsic activation energywas between 3.6 and 4.2eV.G.W.Tomlins, J.L.Routbort, T.O.Mason: Journal of the American Ceramic Society, 1998,81[4], 869-76

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Table 5Bulk Diffusivity of O in ZnO


925 8.20 x 10-17

995 2.62 x 10-15

1000 2.21 x 10-15

1040 5.48 x 10-15

1095 4.20 x 10-15

1100 6.16 x 10-15

1150 1.31 x 10-14

1175 1.97 x 10-14

1200 3.50 x 10-14

ZnO: Dislocations and Grain BoundariesBicrystals with large interfacial areas and controlled misorientations were prepared byusing the solid-phase inter-growth method. The structures of three [00•1] tilt boundaries,with misorientations of less than 1º, 17.8º or 31.5º, were studied by using high-resolutionelectron microscopy. The low-angle boundary comprised well-separated crystaldislocations, and the atomic structure of the large-angle boundaries could be described interms of sequences of [00•1] tunnels that were coordinated, 5-fold, 6-fold and 7-fold, byatomic columns. It was found that the 17.8º asymmetrical (9̄ 17 9̄ 0) boundary, S = 31,

150

ZnO Oxides ZnO

was planar and exhibited a relatively long-period repeating structure which wasoccasionally interrupted by interfacial dislocations. The Burgers vectors and step natureof these defects were investigated by using circuit mapping and were found to beconsistent with topological theory. The 31.5º interface was found to be extensivelyfacetted into (2̄ 7 5̄ 0) and (1 3 4̄ 0) symmetrical tilt boundaries. Mirror symmetry in theimmediate vicinity of the interface was suppressed by local relaxation. An angulardeviation, of -0.7º, from the periodic S = 13 system was seen to be accommodated byprimitive interfacial dislocations. Some of these defects exhibited compact cores; thusintroducing minimal disruption into the underlying periodic structure. Others exhibited amore complex reconstruction, leading to a reduction in interfacial area and defect energy.A.N.Kiselev, F.Sarrazit, E.A.Stepantsov, E.Olsson, T.Claeson, V.I.Bondarenko,R.C.Pond, N.A.Kiselev: Philosophical Magazine A, 1997, 76[3], 633-55

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ZnO: Grain BoundariesGrain boundaries in additive-free samples exhibited various types of conductive modecontrast, depending upon the electrical properties. Electron beam scattering patterns wereused to determine the misorientation of grain boundaries, showing certain types ofconductive mode contrast, in order to correlate electrical properties with grain-boundarystructure. No correlations were apparent when using the coincident-site lattice model, butother criteria suggested that grain-boundary electrical properties might becrystallographically controlled.J.D.Russell, D.C.Halls, C.Leach: Journal of Materials Science, 1997, 32[5], 4585-9

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ZnO: Grain BoundariesThe structure of a [00•1] 31.5° tilt boundary was investigated by means of high-resolutionelectron microscopy. In some regions, the interface structure could be best described byreferring to the periodic dichromatic pattern that corresponded to the misorientation,[00•1] 32.5° (S = 13). These regions exhibited arrangements of periodic facets, withlocalized interfacial dislocations (with primitive Burgers vectors) superposed on them.Other regions were better described by referring to the dichromatic pattern whichcorresponded to [00•1] 30°, which was periodic parallel to [00•1] alone. These segmentsof interface were parallel to (01̄•0) and (11•0) in the adjacent crystals, and their structurewas incommensurate perpendicular to [00•1]. No localized defects were found in suchsegments.F.Sarrazit, R.C.Pond, N.A.Kiselev: Philosophical Magazine Letters, 1998, 77[4], 191-8

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ZnO: Point DefectsNeutral-donor bound-exciton transitions were studied for isolated neutral donors thatwere made up of defect-pair complexes. The neutral-donor nature of these complexes wasdeduced from magnetic field measurements and from 2-electron transitions. Excited statesof the neutral-donor bound excitons were observed in the form of rotator states that were

151

ZnO Oxides ZnO

analogous to rotational states of the H2 molecule. Annealing experiments showed that thehigher-energy emission lines disappeared and that, at annealing temperatures near to800C, all of the emission intensity appeared in the lowest-energy donor bound-excitontransition. Integrated intensity measurements revealed that the total emission intensity wasapproximately conserved. It was suggested that the higher-energy emission lines were dueto neutral-donor bound-exciton transitions in which the pairs that made up the neutraldonors were more distantly spaced. These were the ones that were first to break up duringannealing. The conservation of emission intensity suggested that the pairs were noteliminated but took up a closer spacing. The measured activation energy of 3.6eV wasconsistent with this motion.D.C.Reynolds, D.C.Look, B.Jogai, C.W.Litton, T.C.Collins, W.Harsch, G.Cantwell:Physical Review B, 1998, 57[19], 12151-5

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ZnO: Point DefectsPowder samples were tribophysically activated by grinding in air, and were investigatedby using electron paramagnetic resonance methods to clarify the effect of the duration oftribophysical activation upon defect formation. Six electron paramagnetic resonancesignals were detected and it was shown that these were caused by the formation of VZn

-

:Zni0, VZn

-, (VZn-)2

-, F+, SDC, and interaction products of mill-material with ZnO. Thedependence of the formation of various centers upon grinding time suggested that thedefect structures in polycrystalline samples during grinding evolved in the order: Zn2+

lattice

? VZn-:Zni

+ ? VZn-:Zni

0 ? VZn- + Zni

-? (VZn-)2

-, and O2+lattice ? VO

+ + O-.N.G.Kakazey, T.V.Sreckovic, M.M.Ristic: Journal of Materials Science, 1997, 32[5],4619-22

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ZnO: TwinsTetrapod-like particles having 4 legs that were shorter than a micrometre were observedby means of transmission electron microscopy and electron diffractometry. The particleswere composed of 6 or more wurtzite-type elements. Four of them constituted legs, andothers formed (112̄2) twins with the legs. Particles which were composed of only 4wurtzite-type legs were also observed to form three (112̄4) twins among them. A growthmodel was proposed in which the tetrapod-like oxide particle grew from wurtzite-typemultiple twins which were introduced into a zincblende-type nucleus, and account wastaken of the combined slips which occurred during phase transformation.K.Nishio, T.Isshiki, M.Kitano, M.Shiojiri: Philosophical Magazine A, 1997, 76[4], 889-904

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ZnO/Al2O3, AlN/ZnO/Al2O3: Dislocations and Stacking FaultsAn investigation was made of the nature of the defects and hetero-interfaces in ZnO filmswhich had been grown onto (00•1) sapphire. High-quality epitaxial films were depositedby using pulsed lasers at temperatures ranging from 750 to 800C. The epitaxial

152

ZnO Oxides ZnO

relationship of the film with respect to the (00•1) substrate was found to be:(00•1)ZnO||(00•1)sapphire; with the in-plane orientational relationship:[01•0]ZnO||[1̄2•0]sapphire. The latter relationship corresponded to a 30º rotation of the ZnObasal planes with respect to the sapphire substrate. This was similar to the epitaxialgrowth characteristics of AlN and GaN on sapphire. The threading dislocations in ZnOwere found to have mainly 1/3<11•0> Burgers vectors. The planar defects, which weremainly I1 stacking faults, were found to lie in the basal plane and had a density of about105/cm. Epitaxial AlN films were grown at a temperature of about 770C by using aZnO/sapphire heterostructure as a substrate. This resulted in the formation of a thinreacted layer at the AlN/ZnO interface. These results had implications with regard to lowdefect contents in ZnO films, as compared to III-V nitrides, and to the use of ZnO filmsas buffer layers for III-V nitrides.J.Narayan, K.Dovidenko, A.K.Sharma, S.Oktyabrsky: Journal of Applied Physics, 1998,84[5], 2597-601

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ZnO-CuO: Grain BoundariesThe electrical properties of semiconducting ceramic composites which were made from n-type ZnO and p-type CuO were investigated. The conductivity increased with increasingCuO contents of between 1 and 95mol%. The impedance response exhibited 3semicircles, thus indicating that 3 resistive elements contributed to the total resistance. Anew model which was based upon the equivalent circuits was developed in order toexplain the contribution which grain boundaries made to the resistance of the composite.S.T.Jun, G.M.Choi: Journal of the American Ceramic Society, 1998, 81[3], 695-9

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ZnO-In2O3: Point Defects and Stacking FaultsThe incorporation of In2O3 into the ZnO wurtzite structure was investigated by using acombination of atomistic simulations and high-resolution electron microscopy. At lowconcentrations, the In ions were incorporated at Zn sites, with the creation of Znvacancies in order to maintain charge neutrality. Defect clustering which involved two Insubstitutional ions that surrounded a Zn vacancy was energetically favored. At higherconcentrations, In2O3 was accommodated via a series of inter-growth phases. Theseconsisted of extended planar defects which were stacked along the c-axis of thehexagonal wurtzite structure. A model which was proposed for the In2O3-(ZnO)n phasesconsisted of two In2O3 layers which were oriented along the (00•1) planes of ZnO andwere separated by a ZnO wurtzite layer that was n/2 unit cells thick. Each wurtzite regionwas displaced from the next by a translation of 1/3<01•0>. The associated solutionenergies of the various inter-growth phases were similar in magnitude, and were muchlower than the solution energies for either isolated defects or defect clusters.M.A.McCoy, R.W.Grimes; W.E.Lee: Philosophical Magazine A, 1997, 76[6], 1187-201

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153

ZnWO2 Oxides ZnWO2

ZnWO4: X-Irradiation and Point DefectsIn Li-doped single crystals, a hole-type defect which was attributed to an O- ion wasobserved, using electron spin resonance and electron nuclear double resonancespectroscopic techniques, after X-irradiation at 77K. The electron nuclear doubleresonance results proved that this O- ion associated with a nearby Li ion, to give an O--LiZn center. This defect was concluded to be a lower-symmetry version of the [Li]0 centerwhich was found in simple oxides.A.Watterich, A.Hofstaetter: Solid State Communications, 1998, 105[6], 357-62

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ZnWO4: Point DefectsElectron spin resonance and electron nuclear double resonance methods were used tocharacterize a new Mo5+-H center, with C1 symmetry, in reduced :Mo-doped singlecrystals. The electron nuclear double resonance results also confirmed that, for a recentlyreported W5+-H center in ultra-violet or ?-irradiated Li-doped crystals, an observedsuperhyperfine interaction resulted from a nearby H nucleus. A comparison of spin-Hamiltonian parameters revealed the similarity of the Mo5+-H and W5+-H centers in thismaterial.A.Watterich, A.Hofstaetter, R.Wuerz, A.Scharmann, O.R.Gilliam: Journal of Physics -Condensed Matter, 1998, 10[1], 205-13

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ZnWO4: Point DefectsParamagnetic centers in Tm-doped single crystals were characterized by means ofelectron spin resonance spectroscopy. Most of the dopant was found to be present in thecrystal as Tm3+ which substituted for Zn2+, and yielded no electron spin resonancespectrum. However, a low concentration of Tm2+ centers with the 4f13 electronconfiguration and S = 1/2 was detected, with C1 symmetry. The intensity of this centerwas greatly enhanced, by 366nm ultra-violet irradiation at 77K, due to electron capture.The low symmetry was attributed to an associated defect. This was probably a Znvacancy which had contributed to local charge compensation of the original Tm3+ ion,where the vacancy remained at the Tm ion even after electron capture at lowtemperatures. This center was described as being a Tm2+-VZn center. Some paramagnetichole-type defects were also observed, using electron spin resonance, after ultra-violetirradiation at 77K. One of these was suggested to a center which was near to a Tm3+ Zn-vacancy pair and was denoted as being a O-VZn-Tm3+ center.A.Watterich, L.A.Kappers, O.R.Gilliam: Solid State Communications, 1997, 104[11],683-8

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154

ZrO2 Oxides ZrO2

ZrO2: O DiffusionThe appearance of 2 distinct activation energies (at high and low temperatures), ananomalously high attempt frequency, and a stretched-exponential time correlationfunction at low temperatures, for O-ion diffusion in yttria-stabilized material, were shownto be quantitatively consistent with the predictions of the coupling model. The anomalies,which occurred in conductivity relaxation and quasi-elastic light-scattering data, wereattributed to interactions between diffusing O ions.K.L.Ngai: Philosophical Magazine B, 1998, 77[1], 187-95

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ZrO2: Zr DiffusionThe 96Zr tracer diffusivity was measured in Y2O3-stabilized monocrystalline samples, attemperatures ranging from 1200 to 1500C, by using secondary ion mass spectrometry. Itwas found that the activation energy for diffusion ranged from 4.4 to 5.0eV, and did notexhibit an appreciable dependence upon Y2O3 contents of between 8 and 32mol%. At agiven temperature, the diffusivity decreased with the Y2O3 content. This was explained interms of migration via vacancies in the cation sub-lattice; although a different mechanismwas suspected to operate at lower temperatures.M.Kilo, G.Borchardt, S.Weber, S.Scherrer, K.Tinschert: Berichte der Bunsengesellschaftfür Physikalische Chemie, 1997, 101[9], 1361-5

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ZrO2: DislocationsExperimental data were presented for cubic single crystals which were deformed in a<112> direction. These data included the flow stress, and its strain-rate sensitivity, as wellas the dislocation densities and the distances between glide obstacles along thedislocations. A model was proposed in order to describe the marked variation in plasticdeformation parameters below 1000K. It assumed that the mechanisms which controlleddislocation motion changed from pinning by localized obstacles (small precipitates orjogs), above a transition temperature, to the surmounting of Peierls relief at lowertemperatures. The model was in agreement with most of the experimental observations.B.Baufeld, B.V.Petukhov, M.Bartsch, U.Messerschmidt: Acta Materialia, 1998, 46[9],3077-85

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ZrO2: Grain BoundariesGrain boundary structures in superplastic SiO2-doped and undoped tetragonalpolycrystalline material were investigated by means of high-resolution electronmicroscopy, energy dispersive X-ray spectroscopy and electron energy loss spectroscopy,using a field emission-type transmission electron microscope. No amorphous phase wasobserved at any grain boundaries in either the doped or undoped samples. It was foundthat Y ions segregated over a width of 4 to 6nm at grain boundaries in both materials,

155

ZrO2 Oxides/Nitrides AlGaN

while Si ions segregated over a width of 5 to 8nm at grain boundaries in SiO2-dopedsamples. The average dihedral angle between the grain boundaries in doped materialcould be as high as 80º. This value agreed well with the fact that none of the grainboundaries contained a glass phase. Strain energy accumulated due to the dissolution ofSi ions in the tetragonal lattice. However, the grain boundary energy of the dopedmaterial was thought to be low enough to compensate for the increase in strain energynear to the grain boundaries. It was noted that O-K-edge electron energy loss spectra fromgrain boundaries in doped samples were shifted by 3 to 4eV, towards higher energies, ascompared with the signals from the grain interior. It was suggested that the chemicalbonding was strengthened at grain boundaries, due to the presence of solute Si.Y.Ikuhara, P.Thavorniti, T.Sakuma: Acta Materialia, 1997, 45[12], 5275-84

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ZrO2: Point DefectsThe effect of O vacancies upon the stability of phase structures and phase transformationsin zirconia at low temperatures was studied by using an electrochemical technique. It wassuggested that a decrease in O vacancy concentration decreased the stability of metastabletetragonal zirconia and promoted its transformation into the monoclinic phase.X.Lu, K.Liang, S.Gu, Y.Zheng, H.Fang: Journal of Materials Science, 1997, 32[6], 6653-6

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ZrO2-TiO2-Y2O3: Domain BoundariesDomain structures in tetragonal zirconia, which were formed by the diffusionless cubic-tetragonal transformation, were studied in ZrO2-3Y2O3-12mol%TiO2 samples whichexhibited a wide single tetragonal zirconia field. During annealing in the single-phasefield, domain boundaries with antiphase boundary-like contrast tended to become faceted;with a habit plane that was close to one of the {111} planes. Transmission electronmicroscopy energy-dispersive spectroscopic analysis showed that Y ions segregated to thedomain boundaries. Since the domain boundaries had a cubic structure, they werestabilized by the segregation of Y ions; which were known to stabilize cubic zirconia.H.Ogawa, A.Yasuda, N.Shibata, Y.Ikuhara, T.Sakuma: Philosophical Magazine Letters,1998, 77[4], 199-203

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Nitrides

AlGaN: Point DefectsThe transition from shallow to deep centers, as a function of pressure or alloying, wasinvestigated for O and Si donors on the basis of first-principles total-energy calculations.The stability of the localized deep state (DX center) was found to depend upon

156

AlGaN Nitrides AlN

interactions between the impurity, and third-nearest neighbor atoms, which occurred atdifferent positions in the zincblende and wurtzite phases. Such DX-center formation wassuppressed in the zincblende phase, as well as that of Si donors. The results supported theidentification of O as being the unintentional dopant in n-type GaN, and cast new light onthe driving force for DX formation.C.G.Van de Walle: Physical Review B, 1998, 57[4], R2033-6

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AlGaN, GaN: Point DefectsDoping in these materials was considered on the basis of state-of-the-art first-principlescalculations. In the case of n-type doping, it was found that N vacancies were too high inenergy to be incorporated during growth and were not responsible for unintentional n-type conductivity, whereas Si and O could be incorporated in large numbers and werelikely to cause unintentional n-type doping. The properties of O, including a propensity toDX-center formation, suggested that it was the main cause of unintentional n-typeconductivity. The DX transition did not occur in zincblende material. The Ga vacancieswere thought to be the likely source of yellow luminescence. In the case of p-type doping,it was found that the solubility of Mg was the main factor which limited the holeconcentration in GaN. The incorporation of Mg at interstitial sites or antisites was not aproblem. It was noted that H had a beneficial effect upon p-type doping because itsuppressed compensation and enhanced acceptor incorporation. The compensation ofacceptors by N vacancies could occur, and became increasingly marked as the Al-contentof AlGaN alloys was increased. No other acceptor impurity exhibited characteristics thatwere superior to those of Mg.C.G.Van de Walle, C.Stampfl, J.Neugebauer: Journal of Crystal Growth, 1998, 189-190,505-10

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AlGaN/GaN: Point DefectsDonor defect profiles in pseudomorphic AlGaN layers, grown onto GaN, were calculatedwhile taking account of the effects of a strain polarization field upon the defect formationenergy. Under certain conditions, the defect concentration could be increased by morethan an order of magnitude. These large concentrations, combined with the band-bendingeffects of the piezoelectric field, made charge transfer from the AlGaN barrier to the GaNwell extremely efficient; thus resulting in a 2-dimensional electron gas of very highdensity and low mobility.L.Hsu, W.Walukiewicz: Applied Physics Letters, 1998, 73[3], 339-41

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AlN: DislocationsSamples which had been plastically deformed, using high stresses at temperatures rangingfrom 20 to 800C, were studied by means of transmission electron microscopy. Theactivation of 1/3<11•0>(00•1) basal slip, as well as the first type of 1/3<11•0>{11̄•0}prismatic slip, was observed. Dislocation configurations in the basal plane revealed the

157

AlN Nitrides AlN

role that was played by Peierls forces acting along the screw orientation. It was easier fora dislocation to cross-slip in the prismatic plane than to lie, in the basal plane, along anorientation other than the screw direction. In the basal plane, as well as in the prismaticplane, no dislocation dissociation could be resolved under weak-beam imagingconditions. The occurrence of cross-slip was favoured by numerous pinning points onscrew dislocations; thus giving rise to dipole dragging. At temperatures ranging from 20to 500C, deformation proceeded via dislocation glide in the basal and prismatic planes;with frequent cross-slip between these planes. At 800C, a change in the mobility of edgedislocations on the prismatic plane was detected and was attributed to lattice frictionacting in the (00•1) orientation or to dislocation dissociation out of the glide plane.M.Audurier, J.L.Demenet, J.Rabier: Philosophical Magazine A, 1998, 77[4], 825-42

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AlN: Dislocations and Point DefectsThe structures and formation energies of neutral and charged edge dislocations wereinvestigated by means of density-functional calculations. Stoichiometric structures withfull or open cores were considered, as well as non-stoichiometric structures with Al or Nvacancies along the dislocation core. The formation energies were found to dependstrongly upon the Fermi level, due to the presence of defect levels in the band-gap, andupon the growth conditions; in the case of non-stoichiometric structures. A structure withAl vacancies along the dislocation core was predicted to be most stable in n-type materialthat was grown under N-rich conditions. On the other hand, a N-vacancy structure wasmost stable in p-type material that were grown under Al-rich conditions.A.F.Wright, J.Furthmüller: Applied Physics Letters, 1998, 72[26], 3467-9

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AlN: Dislocations and Stacking FaultsA study was made of the ability of dislocations to glide in one plane relative to another. Itwas noted that, according to the closest-packing criterion, the first type of prismatic planeappeared to be favoured. Line-energy calculations, which were performed underanisotropic elasticity conditions, did not permit the detection of preferential activation ofone of the glide planes. The crystallographic structures of basal and first-type prismaticplanes, the core structures of dislocations, and dissociation reactions were investigated,and configurations were proposed for stacking faults on prismatic planes.M.Audurier, J.L.Demenet, J.Rabier: Philosophical Magazine A, 1998, 77[4], 843-59

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AlN: Domain BoundariesPolytypoid structures in crystals which contained O were investigated by means ofelectron diffraction and high-resolution electron microscopy. The polytypoids consistedof flat and corrugated inversion domains. The flat inversion domain boundaries werefound to have 3 different interfacial structures. The distance between 2 neighboring flat

158

AlN Nitrides AlN

inversion domain boundaries was found to be either 5c or (5 ± 0.5)c, where c was thelattice parameter along the c-axis of the wurtzite 2H-type crystal.Y.Yan, M.Terauchi, M.Tanaka: Philosophical Magazine A, 1998, 77[4], 1027-40

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AlN: Point DefectsThe magneto-optical properties of the 0.9eV luminescence were investigated. Zero-phonon lines at 0.943eV were attributed to a transition, within a d2 configuration, becauseof a characteristic 3-fold ground-state splitting in magnetic fields. A g-value of 1.96 and azero-field splitting of 120µeV were determined for the 3A2(F) ground state. On the basisof the temperature dependence in a magnetic field, the 0.943eV zero-phonon line wasattributed to the 1E(D)-3A2(F) transition of isolated V3+.P.Thurian, I.Loa, P.Maxim, A.Hoffmann, C.Thomsen, K.Pressel: Applied PhysicsLetters, 1997, 71[20], 2993-5

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AlN: Stacking FaultsThe growth characteristics and atomic structures of defects were investigated in AlN thinfilms which were grown by metal-organic chemical vapor deposition onto the (10•2) r-plane of a-Al2O3. The AlN films were monocrystalline, and had an epitaxial relationship,(11•0)AlN||(10•2)sapphire, with an in-plane alignment of [00•1]AlN||[1̄0•1]sapphire. By usinghigh-resolution electron microscopy and multi-slice image simulation, the predominantdefects in the AlN thin films were found to be low-energy intrinsic stacking faults oftype-I1 which lay in the basal plane. This fault, with a 1/6[20•3] resultant displacementvector, could be formed by removing one (00•2) plane and then shearing the remaininghalf-crystal by a displacement of 1/3[10•0]. The faults appeared as a single face-centeredcubic stack, ABC, which was inserted into the normal, ... ABAB..., hexagonal sequence.K.Dovidenko, S.Oktyabrsky, J.Narayan: Journal of Applied Physics, 1997, 82[9], 4296-9

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AlN, GaN, InN: Stacking FaultsThe energies of basal-plane stacking faults in wurtzites were determined by using a 1-dimensional Ising-type model which incorporated effective layer-layer interactions thatwere based upon density-functional theory calculations. The stacking-fault energies werefound to be largest for AlN and smallest for GaN. This was consistent with density-functional results for wurtzite/zincblende energy differences. Estimates were alsoobtained for stacking-fault energies in the zincblende structure. The values were negative,in agreement with the observation that nominal zincblende films typically contained alarge number of stacking faults. A related result was that hexagonal structures withstacking sequences that repeated after 4 or 6 bilayers had lower energies than zincblende,for all 3 compounds.A.F.Wright: Journal of Applied Physics, 1997, 82[10], 5259-61

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159

AlN Nitrides AlN

AlN, GaN, InN: Stacking FaultsBasal-plane stacking faults in wurtzite-structured materials were studied by using density-functional pseudopotential calculations. The formation energies followed the trend thatwas exhibited by zincblende/wurtzite energy differences in the bulk materials. That is, thelowest energy was found for GaN and the highest was found for AlN. Type-I stackingfaults had the lowest energy, followed by type-II stacking faults, and finally extrinsicstacking faults. Another type of intrinsic stacking fault was identified that did not seem towere previously observed. Its energy was slightly lower than that of the type-II faults.There were no localized states in the band-gap. However, stacking faults could bound aquantum well-like region of zincblende material that was surrounded by the wurtzite host;thus giving rise to a luminescence line below the wurtzite band-gap.C.Stampfl, C.G.Van de Walle: Physical Review B, 1998, 57[24], R15052-5

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AlN, GaN: Surface DefectsAn investigation was made of the structure of the anion- and cation-terminated (00•1)surfaces of wurtzite-phase GaN and AlN, using ab initio local-orbital calculations thatwere based upon the local-density approximation and the pseudopotential method. All ofthe stable surface configurations which were studied differed, in atomic composition andperiodicity, from the ideal bulk-like termination. The calculated total energies werecompared for various p(2 x 2) geometries of GaN and AlN (00•1). Vacancy structureswere found to be the most stable configurations for the anion- and cation-terminatedsurfaces. Under metal-rich growth conditions, the predictions favoured the adsorption ofmetal atoms at the cation-terminated surface. Anion- and cation-derived dangling-bondstates appeared in the bulk band-gap as a result of the formation of vacancies or theadsorption of group-III atoms. Flat surfaces of both types were found to be stabilized by¾ of a monolayer of adsorbed H.J.Fritsch, O.F.Sankey, K.E.Schmidt, J.B.Page: Physical Review B, 1998, 57[24], 15360-71

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AlN/Si: DislocationsThe growth of hexagonal AlN directly onto (111)Si was studied, as a function of the filmthickness, by means of grazing incidence X-ray diffractometry and high-resolutionelectron microscopy. Two epitaxial relationships were observed. One was(00•1)[21̄•0]AlN||(111)[022̄]Si, which predominated at deposition temperatures greater than650C, and (00•1)[10•0]AlN||(111)[022̄]Si. The average in-plane crystallite size was 16.2nmfor a 4nm-thick layer, the in-plane rotation was about 2º, and the dislocations introducedan average strain distribution of 0.8%. The Si/AlN interface was very sharp, and completerelaxation (down to about 0.2%) occurred within one bilayer. No long-range order wasobserved at the interface; thus implying a low mobility of AlN species, on Si, whichinhibited any structural rearrangement. The in-plane rotations originated from early stages

160

AlN Nitrides BN

in the layer growth, and decreased with layer thickness; especially for thicknesses greaterthan 25nm.A.Bourret, A.Barski, J.L.Rouvière, G.Renaud, A.Barbier: Journal of Applied Physics,1998, 83[4], 2003-9

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AlN/SiC, GaN/SiC: Stacking FaultsThe polarity of GaN films and AlN buffer layers, grown onto the (00•1)Si surface of 6H-SiC using electron cyclotron resonance plasma-enhanced molecular beam epitaxy, wasinvestigated by means of convergent-beam electron diffraction and high-resolutionelectron microscopy. In the case of the AlN buffer layers, which contained a very highdensity of defects, the polarity was determined by using extensive high-resolutionelectron microscopy and image simulation. In both cases, the simulations were in goodagreement with the experimental results, and demonstrated that the free surfaces of theGaN and AlN layers were Ga- and Al-terminated, respectively. Moreover, (1̄2•0)prismatic planar defects which were observed in the AlN layers were identified as beingstacking faults, and analysis of various areas of the specimens confirmed that the layerswere unipolar.P.Vermaut, P.Ruterana, G.Nouet: Philosophical Magazine A, 1997, 76[6], 1215-34

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BN: H DiffusionPhoto-production and photo-bleaching of dangling bonds by ultra-violet irradiation andannealing were observed via electron spin resonance studies of amorphous hydrogenatedmaterial. A model which involved long-range H diffusion and H evolution was proposedin order to account for the electron spin resonance spectral line-shapes which wereobserved after ultra-violet irradiation and annealing. A Gaussian distribution of activationenergies for long-range H diffusion was used to explain the observed decrease in spinsusceptibility after thermal annealing. A best fit to the spin susceptibility data indicated anaverage activation energy of 0.0475eV and a half-width, of the Gaussian distribution, of0.027eV.I.M.Brown, S.H.Lin, B.J.Feldman: Physical Review B, 1997, 56[3], 994-6

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BN: Point DefectsThe electronic structures of Be, Mg and Si in zincblende BN were studied by using thetight-binding linearized muffin-tin orbitals technique. Calculations were performed byusing 64-atom super-cells which were centered on either a B or a N lattice site. While Beand Mg impurities were substituted only for B, substitution for both B and N wasconsidered in the case of Si. In each case, total-energy minimization was used to studylattice relaxation near to the impurity site, and the nature of the chemical bondingbetween the impurity and neighboring atoms of the host crystal was studied in detail. Itwas found that Be and Mg, when substituted for B, created delocalized levels whichmerged with states at the valence-band edge. These partially occupied levels could result

161

BN Nitrides BN

in p-type conductivity, as observed experimentally. In contrast to the behavior of isolatedBe and Mg impurities in crystalline BN, it was found that Si which was substituted at a Bsite introduced delocalized impurity states that overlapped with the conduction-band edgeof the host. These levels could contribute to the n-type conductivity of Si-dopedcrystalline BN. When Si was substituted into the anion sub-lattice it introduced sharp,partially occupied, and highly localized levels within the forbidden gap. Experimentallyobserved mixtures of n-type and p-type material could therefore be accounted for by thepresence of impurities of each type. The relaxation of the host lattice near to the Be andMg impurities was outward, as was the relaxation near to Si which was substituted at a Nsite. Inward relaxation was predicted to occur in the case of Si which was substituted forB. Total, orbital, and shell-projected densities of states for the impurity, and up to 6 of thecoordination shells nearest to the impurity, were analyzed in detail.V.A.Gubanov, E.A.Pentaleri, C.Y.Fong, B.M.Klein: Physical Review B, 1997, 56[20],13077-86

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BN: Point DefectsA study was made of the electronic and structural properties of native defects (N vacancy,B antisite, substitutional O), in the cubic phase, by using all-electron first-principles total-energy calculations. It was found that all of the defects introduced a deep state above themiddle of the energy-gap. These defects had C3v local symmetry. A Jahn-Teller distortionin the <111> direction was observed for the neutral isolated B antisite. Strong relaxationwas found in the case of the neutral N vacancy. A highly localized orbital in the vacancyregion (F-center precursor) and a resonant orbital in the conduction band appeared beforelattice relaxation. Such relaxation, plus distortion, split the resonant orbital and led to ahalf-occupied level close to the conduction band. The results suggested that a strongelectron paramagnetic resonance signal which was detected in cubic material could begenerated by the trapping of 2 electrons by acceptor levels which were introduced by Cdoping.P.Piquini, R.Mota, T.M.Schmidt, A.Fazzio: Physical Review B, 1997, 56[7], 3556-9

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BN: Point DefectsElectron paramagnetic resonance techniques were used to investigate defects in hexagonalBN, polycrystalline cubic BN, and pure BN or C-doped BN films which had beenproduced by reactive sputtering. The predominant paramagnetic center in the hexagonalmaterial was deduced to be the N vacancy. The zincblende samples exhibited a singleresonance line, and a model for the predominant paramagnetic defect could not beformulated. However, under the assumption that the N vacancy in this material was againresponsible for the observed resonance, an estimate of the electron density at the B nucleiof the first coordination shell was obtained from the experimental data. It was found thatfilms which were produced by using up to 10%N2 in Ar discharges exhibited single-lineelectron paramagnetic resonance signals with a g-value and line-width that wereconsistent with those observed for high-pressure cubic BN. The spin concentration was

162

BN Nitrides GaAlN

reduced upon increasing the N pressure in the discharge; thus supporting theidentification of the paramagnetic center as being a N vacancy. Films which wereproduced using N partial pressures in the Ar discharge that were greater than 10%exhibited an increased spin density. This was inconsistent with the idea that Nincorporation reduced the number of N vacancies. The BN films degraded in time, whenexposed to the atmosphere, and this was attributed to the presence of H3O+ or othercomplex defects which involved H and/or O. Doping with C increased the spin densityand the g-value, and this supported the suggestion that C stabilized the electron in the Nvacancy. The spin-lattice relaxation rates exhibited an unusual and almost lineartemperature dependence. It was concluded that theoretical analysis of the electronicstructure of the defects might clarify whether excited states of defects within the phononspectrum existed and permitted an Orbach relaxation process. This would then explain theobserved temperature dependence.M.Fanciulli: Philosophical Magazine B, 1997, 76[3], 363-81

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C3N4: Point DefectsBy using an ab initio super-cell approach, simulations were made of the electronicstructure around a N vacancy, and an investigation was made of its effect upon the bulkmodulus of the material. This was compared with a relaxed cubic structure with one N-atom removed. The ability of the super-lattice to model the electronic structure of thecrystalline material was used to deduce the properties of an amorphous structure. Theamorphous form had a bulk modulus which was appreciably lower than its crystallineequivalent, and a comparison of various structures suggested that the strength of thisnitride would ultimately depend upon the number of C-N bonds that could be formed.J.E.Lowther: Physical Review B, 1998, 57[10], 5724-7

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GaAlN: Point DefectsExperimental and theoretical evidence was presented for the metastability of O donors inthis material. As the Al content was increased, Hall effect measurements revealed anincrease, in the electron activation energy, which was consistent with the emergence of adeep DX level from the conduction band. After exposure to light, persistentphotoconductivity was observed in O-doped Al0.39Ga0.61N, at temperatures below 150K,with an optical threshold energy of 1.3eV. A configuration coordinate diagram wasdeduced from first-principles calculations, and yielded values for the capture barrier,emission barrier, and optical threshold which were in good agreement with theexperimental data.M.D.McCluskey, N.M.Johnson, C.G.Van de Walle, D.P.Bour, M.Kneissl,W.Walukiewicz: Physical Review Letters, 1998, 80[18], 4008-11

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163

GaInN Nitrides GaN

GaInN/GaN: V-DefectsIt was noted that, during the growth of InGaN/GaN multiple quantum-well structures, anew defect (the so-called V-defect) nucleated at threading dislocations in one of the firstquantum wells in a multiple quantum-well stack. This defect was common to almost allInGaN multiple quantum-well heterostructures. The nature of the V-defect was evaluatedhere by applying transmission electron microscopy, scanning transmission electronmicroscopy, and low-temperature cathodoluminescence techniques to a series ofGa0.80In0.20N/GaN multiple quantum-well samples. The structure of the V-defectcomprised buried side-wall quantum wells, on the {10•1} planes, and an open hexagonalinverted pyramid which was defined by the six {10•1} planes. This defect thereforeappeared as an open 'V' when seen in cross section. The formation of the V-defect wascontrolled kinetically by reduced Ga incorporation at the pyramid walls; the {10•1}planes. The V-defect was related to localized excitonic recombination centers that gaverise to a long-wavelength shoulder in photoluminescence and cathodoluminescencespectra.X.H.Wu, C.R.Elsass, A.Abare, M.Mack, S.Keller, P.M.Petroff, P.M.DenBaars,J.S.Specka, S.J.Rosner: Applied Physics Letters, 1998, 72[6], 692-4

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GaN: Ga, N DiffusionA kinetic model was used to explain the metal-organic vapor-phase epitaxial growth ofthis nitride. The model was based upon measured desorption rates, and upon assumedprecursor dissociation and sticking probabilities. The model showed how the growthtemperature and V/III ratio were linked with regard to the growth of high-quality films.By comparing reported growth conditions with the resultant quality of the material whichwas produced, it was deduced that optimum film growth occurred when the V/III ratiowas chosen so as to be slightly larger than the N/Ga desorption ratio. The relationshipsbetween the growth temperature, V/III ratio and sample quality were explained in termsof how the growth parameters influenced the incorporation of Ga and N atoms into thegrowing film. The Ga and N diffusion lengths at 1050C were estimated to be between 2and 20nm and less than 1nm, respectively, for reasonable metal-organic vapor-phaseepitaxial growth rates. The growth conditions which were required to give a smooth(00•1) surface morphology were described in terms of the present growth model, and thepossible origins of defect incorporation. As a result of the large N desorption rate, it wassuggested that N was incorporated, during growth, via an adsorption/desorption cycle.D.D.Koleske, A.E.Wickenden, R.L.Henry, W.J.DeSisto, R.J.Gorman: Journal of AppliedPhysics, 1998, 84[4], 1998-2004

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GaN: Ga, N Surface DiffusionThe diffusion of Ga and N adatoms on (00•̄1) and (00•1) surfaces was studied by usingdensity-functional theory. The calculations revealed very different diffusivities of Ga andN adatoms on the equilibrium surfaces. That is, whereas Ga was very mobile at typical

164

GaN Nitrides GaN

growth temperatures, the diffusion of N was slower by some orders of magnitude. It wasnoted that the desorption rate in excess N ambients could become smaller than theadsorption rate. Therefore, extensive regions could be formed in which the surface wascovered mainly by N atoms. These adatoms were likely to influence the migration pathand the diffusion barriers to Ga adatoms. The diffusion of Ga adatoms on the present, N-terminated, surfaces could thus be considered to be extreme cases of N coverage. Theenergetically favored binding sites were located at the face-centered cubic and hexagonalclose-packed positions, and the transition site was the bridge position. For both surfaceorientations, the diffusion barrier to Ga adatoms was strongly affected. On (00•1), themigration barrier increased from 0.4 to 1.8eV while, on (00•̄1), it increased from 0.2 to1.0eV. It was concluded that excess N at the surface, which could occur under N-richconditions, significantly reduced the mobility of Ga adatoms. The reason for this was theformation of strong Ga-N bonds, which had to be broken during adatom migration.T.Zywietz, J.Neugebauer, M.Scheffler: Applied Physics Letters, 1998, 73[4], 487-9

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y GaN: N DiffusionIsotopic heterostructures of the form, Ga14N/Ga15N/Ga14N, were used to study N self-diffusion by means of secondary-ion mass spectrometry and thermally activateddecomposition techniques. The diffusion profiles were measured after interdiffusion ofthe Ga14N and Ga15N layers at temperatures of between 770 and 970C. It was noted thatsuch isotopic heterostructures were especially suitable for self-diffusion studies becausediffusion took place at interfaces within the GaN crystal, and was therefore free frominterfering effects such as surface electric fields, mechanical stresses or chemicalpotential gradients. It was found that the temperature dependence of N self-diffusion(figure 10) in the hexagonal phase could be described by:

D (cm2/s) = 1600 exp[-4.1(eV)/kT]and by an entropy, for self-diffusion, of about 10k.O.Ambacher, F.Freudenberg, R.Dimitrov, H.Angerer, M.Stutzmann: Japanese Journal ofApplied Physics - 1, 1998, 37[5A], 2416-21

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GaN: Zn Diffusion and DislocationsThe optical activation of Zn ions which had been implanted into an epitaxial film wascarried out by annealing in N2 pressures of up to 1.6GPa. Such pressures avoideddecomposition, and could increase the annealing temperature to up to 1550C, instead ofthe 1000 to 1100C which could be used at ambient pressures. The Zn acceptor-relatedphotoluminescence intensity in implanted samples was maximized by annealing attemperatures above 1350C. The Zn photoluminescence intensity then exceeded, by afactor of 15, that of epitaxially Zn-doped GaN with a comparable Zn concentration. High-pressure annealing led to a significant diffusion of implanted Zn atoms in the films. It wasalso possible to diffuse Zn into the implanted/unimplanted layers from external sources.The presence of high dislocation densities strongly accelerated Zn diffusion.

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GaN Nitrides GaN

T.Suski, J.Jun, M.Leszczynski, H.Teisseyre, S.Strite, A.Rockett, A.Pelzmann, M.Kamp,K.J.Ebeling: Journal of Applied Physics, 1998, 84[2], 1155-7

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Figure 10Diffusivity of N in GaN

GaN: Electron Irradiation and Point DefectsHigh-energy (0.7 to 1MeV) electron irradiation, of samples which had been grown ontosapphire, produced shallow donors and deep or shallow acceptors at equal rates (1/cm).These data, in conjunction with theoretical considerations, were consistent only with theassumption that the shallow donor was the N vacancy, while the acceptor was the Ninterstitial. The N-vacancy donor energy was 0.064eV. This was much larger than thevalue (0.018eV) for the residual donor (probably Si) in this material. Hall-effectmeasurements also revealed a degenerate n-type layer at the GaN/sapphire interface. Thishad to be accounted for in order to determine the correct donor activation energy.D.C.Look, D.C.Reynolds, J.W.Hemsky, J.R.Sizelove, R.L.Jones, R.J.Molnar: PhysicalReview Letters, 1997, 79[12], 2273-6

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1.0E-18

1.0E-17

1.0E-16

1.0E-15

1.0E-14

1.0E-13

8 9 10

14N15N14N215N2

104/T(K)

D (c

m2 /s

)

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GaN Nitrides GaN

GaN: Electron Irradiation and Point DefectsDeep-level transient spectroscopic measurements of n-type epitaxial layers, which hadbeen irradiated with 1MeV electrons, revealed an irradiation-induced electron trap at Ec -0.18eV. The production rate was approximately 0.2/cm. This was lower than the rate(1/cm) which was found, for the N vacancy, by means of Hall-effect studies. The presentdefect trap could not be unambiguously identified.Z.Q.Fang, J.W.Hemsky, D.C.Look, M.P.Mack: Applied Physics Letters, 1998, 72[4],448-9

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GaN: Gamma Irradiation, Ion Bombardment, and Point DefectsTwo absorption bands, located at about 1730 and 2960/cm in the infra-red absorptionspectra, were observed in undoped samples which had been grown using low-pressuremetal-organic vapor phase epitaxy, irradiated with ?-rays, and exposed to a radio-frequency H plasma. Proton implantation, followed by ?-irradiation, could also activatethe infra-red band at around 1730/cm. On the basis of the experimental results, the1730/cm band was tentatively attributed to the local vibrational modes of Ga-Hcomplexes in the vicinity of N vacancies. The 2960/cm band was attributed to the localvibrational modes of N-H complexes in the vicinity of Ga vacancies, or of C-Hcomplexes.J.Q.Duan, B.R.Zhang, Y.X.Zhang, L.P.Wang, G.G.Qin, G.Y.Zhang, Y.Z.Tong, S.X.Jin,Z.J.Yang, X.Zhang, Z.H.Xu: Journal of Applied Physics, 1997, 82[11], 5745-7

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GaN: Ion Implantation and DefectsThe removal of Si and Te ion-implantation damage was studied as a function ofimplantation dose, implantation temperature and annealing temperature. Transmissionelectron microscopy revealed that amorphous layers, which could result from high-doseimplantation, recrystallized at temperatures of between 800 and 1100C to give a verydefective polycrystalline material. Lower doses (down to 5 x 1013/cm2) producedspecimens which were not amorphous, but were defective, after implantation. These alsoannealed poorly, at temperatures of up to 1100C, to leave a coarse network of extendeddefects. A high fraction of the Te was in substitutional positions after implantation andafter annealing. Although high-temperature implantation resulted in less disorder afterimplantation, this damage was still impossible to anneal out completely at up to 1100C.H.H.Tan, J.S.Williams, J.Zou, D.J.H.Cockayne, S.J.Pearton, J.C.Zolper, R.A.Stall:Applied Physics Letters, 1998, 72[10], 1190-2

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GaN: Ion Implantation and Point DefectsSeveral vibrational bands were observed near to 3100/cm in samples which had beenimplanted with H at room temperature and then annealed. The results indicated that thesebands were due to N dangling-bond defects which were created by implantation, and were

167

GaN Nitrides GaN

decorated with H. The frequencies were close to those predicted for VGa-Hn complexes.The new lines were therefore tentatively attributed to VGa defects that were decoratedwith various numbers of H atoms.M.G.Weinstein, C.Y.Song, M.Stavola, S.J.Pearton, R.G.Wilson, R.J.Shul, K.P.Killeen,M.J.Ludowise: Applied Physics Letters, 1998, 72[14], 1703-5

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GaN: DislocationsThe atomic structures and line energies of threading screw and threading edgedislocations in the wurtzite structure were calculated within the local-densityapproximation. It was noted that both types of dislocation were electrically inactive, witha band-gap that was free of deep levels. It was concluded that these features arose fromrelaxed core structures which were similar to (10•0) surfaces.J.Elsner, R.Jones, P.K.Sitch, V.D.Porezag, M.Elstner, T.Frauenheim, M.I.Heggie,S.Oberg, P.R.Briddon: Physical Review Letters, 1997, 79[19], 3672-5

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GaN: DislocationsThe first direct observations of the atomic structures of threading dislocation cores in thehexagonal phase of GaN were presented. By using atomic-resolution Z-contrast imaging,dislocations with edge character were found to exhibit an 8-fold ring core. The centralcolumn in the core of a pure edge dislocation had the same configuration as one row ofdimers on the {10-10} surface. On the basis of published theory, it was proposed thatedge dislocations did not have deep defect states in the band-gap, and did not contributeto cathodoluminescence dislocation contrast. On the other hand, both mixed and purescrew dislocations were found to have a full core, and full screw dislocation cores werepredicted to have states in the band-gap.Y.Xin, S.J.Pennycook, N.D.Browning, P.D.Nellist, S.Sivananthan, F.Omnès,B.Beaumont, J.P.Faurie, P.Gibart: Applied Physics Letters, 1998, 72[21], 2680-2

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GaN: DislocationsDefect structures in films which had been grown using hydride vapor-phase epitaxy werecharacterized by using transmission electron microscopy. Growth was obtained on SiO2-stripe patterned GaN layers that had been grown, by metal-organic vapor-phase epitaxy,onto sapphire substrates. Cross-sectional transmission electron microscopyunambiguously revealed that most of the dislocations, which originated from threadingdislocations that were aligned vertically in the metal-organic vapor-phase epitaxial layer,propagated laterally around the SiO2 mask, in the hydride vapor-phase epitaxial film,before the film thickness attained 5µ. This change in the propagation direction preventedthe dislocations from crossing the film, to the surface region, and thus led to a sharpreduction in the threading dislocation density of thicker films.

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GaN Nitrides GaN

A.Sakai, H.Sunakawa, A.Usui: Applied Physics Letters, 1997, 71[16], 2259-61[446-164-168]

GaN: DislocationsEpilayers which had been nitrided for various times were investigated by using light-scattering tomography and Raman scattering. In light-scattering tomographic images ofplan-view epilayers, the light-scattering defects were found to be distributed mainly in<11•0> directions. The density of the defects was lower in epilayers which had beennitrided for longer times. The defects were considered to be straight threading edgedislocations on {11̄•0} planes. The Raman shift of the E2 mode was larger in sampleswhich had been nitrided for longer times. The results demonstrated that misfit betweenthe GaN epilayer and an Al2O3 substrate was more unfavourably accommodated bythreading edge dislocations in epilayers which had been nitrided for longer times.J.Kang, T.Ogawa: Applied Physics Letters, 1997, 71[16], 2304-6

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GaN: DislocationsThe formation of homo-epitaxially and selectively grown GaN structures within thewindows of SiO2 masks were investigated by using transmission electron microscopy andscanning electron microscopy. The structures were produced by means of organometallicvapor-phase epitaxy. A GaN layer under the SiO2 mask provided a crystallographictemplate for the initial vertical growth of GaN hexagonal pyramids or striped pattern. TheSiO2 film provided an amorphous stage upon which lateral growth of GaN could occur, aswell as the possible accommodation of mismatch arising from thermal expansion duringcooling. Transmission electron microscopic observations revealed a substantial reductionin the dislocation density in areas of lateral growth of GaN which was deposited onto theSiO2 mask. No dislocations were observed in many of these areas.T.S.Zheleva, O.H.Nam, M.D.Bremser, R.F.Davis: Applied Physics Letters, 1997, 71[17],2472-4

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GaN: DislocationsA study was made of the organometallic vapor-phase lateral epitaxy and coalescence ofGaN layers which originated from GaN stripes that had been deposited within 3µ-widewindows, spaced 3µ apart, and contained within SiO2 masks on GaN/AlN/6H-SiC(00•1)substrates. The extent, and microstructural characteristics, of the lateral overgrowth werestrong functions of the stripe orientation. A high density of threading dislocations, whichoriginated from the interface of the underlying GaN with the AlN buffer layer, waspresent in GaN which was grown in the window regions. On the other hand, theovergrowth regions contained a very low density of dislocations.O.H.Nam, M.D.Bremser, T.S.Zheleva, R.F.Davis: Applied Physics Letters, 1997, 71[18],2638-40

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GaN: DislocationsImportant structural characteristics (correlation lengths of columnar crystallites,dislocation densities, angles of rotational disorder) were determined in the hexagonalphase, after its metal-organic chemical vapour deposition onto c-plane sapphire, by meansof transmission electron microscopy and triple-axis X-ray diffractometry. The filmsexhibited edge dislocation densities which were of the order of 1011/cm2, tilt and twistangles of 0.1º and 1.3º, respectively, and a columnar structure with lateral and verticalcorrelation lengths of 150 and 1000nm. It was shown that triple-axis X-ray diffractometrywas a good technique for distinguishing between defects, such as edge and screwdislocations, that led to a characteristic broadening of symmetrical and asymmetricalBragg reflections.T.Metzger, R.Höpler, E.Born, O.Ambacher, M.Stutzmann, R.Stömmer, M.Schuster,H.Göbel, S.Christiansen, M.Albrecht, H.P.Strunk: Philosophical Magazine A, 1998,77[4], 1013-25

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GaN: DislocationsOrganometallic vapor-phase epitaxial thin films, with the (001) plane parallel to thesurface of (110) sapphire substrates, were studied by using X-ray diffraction methods.The line profiles of the thin films along the [001] direction could be quantitativelyreproduced by assuming the existence of a strained lattice at the interface. Thedeformation and growth faults were determined to be equivalent, and each amounted to0.2%. Least-squares refinement of 42 independent peaks, after correcting for first-orderthermal diffuse scattering, furnished the Debye-Waller factors for Ga and N atoms. Thewurtzite positional parameter for the GaN thin film was found to be 0.3730. This was 1%smaller than that (0.377) for a strain-free single crystal, and the difference was attributedto strain effects.X.Xiong, S.C.Moss: Journal of Applied Physics, 1997, 82[5], 2308-11

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GaN: DislocationsA transmission electron microscopic study was made of films which had been producedby using an epitaxial lateral overgrowth technique which involved SiO2-mask/windows,stripe-patterned GaN layers, and hydride vapor-phase epitaxy. The regions which wereovergrown on SiO2 masks were carefully examined. Cross-sectional transmission electronmicroscopy clearly revealed the presence of characteristic defects along the [00•1]direction in the overgrown region. These consisted of arrays of dislocations which ran inthe mask stripe direction. These defects caused a crystallographic tilting, in the regionnear to the mask, with respect to other regions grown from the window area. The verticalre-propagation of dislocations, which had propagated laterally during epitaxial lateralovergrowth, was also observed at the coalesced site on the mask.A.Sakai, H.Sunakawa, A.Usui: Applied Physics Letters, 1998, 73[4], 481-3

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GaN: DislocationsThe effect of dislocations upon the electrical characteristics of p-n junctions was studiedby means of current-voltage measurements. Lateral epitaxial overgrowth was used toproduce areas of low dislocation density which were close to areas having the highdislocation density which was typical of growth on sapphire. A comparison of diodesfabricated in each region revealed that reverse-bias leakage currents were reduced by 3orders of magnitude on lateral epitaxially overgrown GaN. Temperature-dependentmeasurements indicated that the remaining leakage current in devices was associated witha deep trap level.P.Kozodoy, J.P.Ibbetson, H.Marchand, P.T.Fini, S.Keller, J.S.Speck, S.P.DenBaars,U.K.Mishra: Applied Physics Letters, 1998, 73[7], 975-7

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GaN: DislocationsLateral transport was investigated in films which had been produced by electron cyclotronresonance plasma-assisted molecular beam epitaxy, and were doped n-type by 1015 to1020/cm3Si. The room-temperature electron mobility, as a function of carrierconcentration, was found to describe a family of bell-shaped curves. This was consistentwith a proposed model for scattering by charged dislocations. The mechanism of thisscattering was investigated by studying the temperature dependences of the carrierconcentration and electron mobility. It was found that, in the low carrier concentrationregion (below 1017/cm3), the electron mobility was thermally activated. Overall, thetemperature dependence of the mobility in samples where dislocations played apredominant role showed that the mobility did not follow classical behavior but wasthermally activated; with an activation energy which was half of that predicted on thebasis of the temperature dependence of the carrier concentration.H.M.Ng, D.Doppalapudi, T.D.Moustakas, N.G.Weimann, L.F.Eastman: Applied PhysicsLetters, 1998, 73[6], 821-3

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GaN: DislocationsThick layers were grown by means of hydride vapor-phase epitaxy. Growth began asselective growth through openings in a SiO2 mask. Facets which consisted of {11̄•1}planes were formed in the early stages, and a continuous film developed via coalescenceof these facets on the SiO2 mask. As a result, GaN layers with dislocation densities as lowas 6 x 107/cm2 were grown onto 5cm-diameter sapphire wafers. The GaN layers werecrack-free and had a mirror-like surface.A.Usui, H.Sunakawa, A.Sakai, A.A.Yamaguchi: Japanese Journal of Applied Physics - 2,1997, 36[7B], L899-902

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GaN: DislocationsPlan-view transmission electron microscopy and cathodoluminescence images wereobtained from the same sample, at exactly the same location, of n-type material whichhad been grown onto sapphire substrates by means of metal-organic chemical vapordeposition. It was found that there was a clear one-to-one correspondence between thedark spots which were observed in cathodoluminescence images and the dislocations intransmission electron microscopy foils. This indicated that the dislocations were non-radiative recombination centers. The hole diffusion length in n-type material wasestimated to be about 50nm by comparing the diameters of the dark spots in thick sampleswhich were used for cathodoluminescence studies with samples which were thinned fortransmission electron microscopic observation. The efficiency of light emission was highwhen the minority carrier diffusion length was shorter than the dislocation spacing.T.Sugahara, H.Sato, M.Hao, Y.Naoi, S.Kurai, S.Tottori, K.Yamashita, K.Nishino,L.T.Romano, S.Sakai: Japanese Journal of Applied Physics - 2, 1998, 37[4A], L398-400

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GaN: DislocationsThe etch-pit density of organometallic vapor-phase epitaxial material, grown ontosapphire, was found to be markedly reduced by the insertion of a low-temperaturedeposited AlN buffer layer or GaN buffer layer between high-temperature grown GaN onsapphire. The insertion of a low-temperature deposited buffer layer between high-temperature grown GaN was effective in eradicating etch pits. Low-temperature depositedAlN and GaN buffer layers were both effective in eradicating etch pits. It was assumedthat the causes of the etch pits were micro-tubes because, if micro-tubes were present atthe surface, they would become enlarged due to the etching. The origin of the micro-tubeswas considered to be screw dislocations. However, further work was required in order toclarify the relationship between etch pits and dislocations. The insertion of a low-temperature deposited buffer layer might have the effect of stopping the threading ofscrew dislocations.M.Iwaya, T.Takeuchi, S.Yamaguchi, C.Wetzel, H.Amano, I.Akasaki: Japanese Journal ofApplied Physics - 2, 1998, 37[3B], L316-8

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GaN: DislocationsThe threading dislocations in films which had been grown onto Al2O3 substrates werestudied by means of plan-view transmission electron microscopy. A pure edgedislocation, and a mixed dislocation with various degrees of screw component, could bedistinguished by obtaining images near to, and far from, the <00•1> zone. Theexperimental results showed that some pure edge dislocations formed low-angleboundaries, while other pure edge dislocations and mixed dislocations were randomlydistributed in the films. The randomly distributed dislocations resulted from the reactionof partial dislocations. On the other hand, partial dislocations formed in order to eliminatestacking disorder near to the film/substrate interface.

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M.Hao, T.Sugahara, H.Sato, Y.Morishima, Y.Naoi, L.T.Romano, S.Sakai: JapaneseJournal of Applied Physics - 2, 1998, 37[3A], L291-3

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GaN: Dislocations and Stacking FaultsCathodoluminescence in the scanning electron microscope was used to study cross-sectional samples of epitaxial films which had been grown on sapphire. Increasedcathodoluminescence emissions, attributed to the presence of stacking faults anddecorated dislocations, were observed in a region of the buffer layer close to thefilm/substrate interface. A region of enhanced emission was also observed in theepilayers, which was partially caused by Si doping and in which structural defects wereinvolved. It was concluded that cross-sectional cathodoluminescence appeared to be auseful method for revealing features, of the spatial distribution of luminescence, whichwere not detectable by using plan-view measurements.M.H.Zaldivar, P.Fernández, J.Piqueras: Journal of Applied Physics, 1998, 83[5], 2796-9

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GaN: Domain BoundariesTransmission electron microscopy was used to investigate the structure of inversiondomain boundaries in a-phase (00•1) films which had been grown, via metal-organicchemical vapor deposition, onto sapphire substrates. Convergent-beam electrondiffraction was used to establish the existence of inversion domains with {10•0}boundaries. The displacement fringes, which were observed in 2-beam images that wereobtained from inclined inversion domain boundaries, were compared with dynamicsimulations. It was shown that the results were consistent with an atomic model in which4-fold bonding was preserved, with all bonds being of Ga-N type.D.Cherns, W.T.Young, M.Saunders, J.W.Steeds, F.A.Ponce, S.Nakamura: PhilosophicalMagazine A, 1998, 77[1], 273-86

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GaN: Nanotube DefectsA formation mechanism was proposed for so-called nanotube defects. It was shown that 2related types of defect were formed: nanotubes and pinholes. Both began with V-shapedfacets on {10•1} polar planes. Slow growth on these polar planes, and impurity poisoningof growth steps, were suggested to be responsible for the initiation of these defects. Theformation of a nanotube appeared to require the nucleation of a pinhole. The latterresulted from slow growth on polar planes, thus resulting in the observed V-shapedfeatures; with about 60º between the arms. This value was in good agreement with theangle between two {10•1} planes. It was possible that the formation of nanotubes andlarge pinholes could be eliminated by reducing the impurity levels below some criticalvalue, which was expected to depend upon the growth rate and the choice of substrateorientation.

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Z.Liliental-Weber, Y.Chen, S.Ruvimov, J.Washburn: Physical Review Letters, 1997,79[15], 2835-8

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GaN: Point DefectsPositron annihilation experiments were performed in order to identify native point defectsin n-type bulk crystals, as well as in epitaxial layers. The results showed that Gavacancies were present in concentrations of between 1017 and 1018/cm3; in both bulkcrystals and epitaxial layers. The Ga vacancies were negatively charged, and theirconcentration was related to the intensity of the yellow luminescence. It was concludedthat the Ga vacancies contributed to the electrical compensation of n-type material, andthat their acceptor levels were involved in the yellow luminescence transition.K.Saarinen, T.Laine, S.Kuisma, J.Nissilä, P.Hautojärvi, L.Dobrzynski, J.M.Baranowski,K.Pakula, R.Stepniewski, M.Wojdak, A.Wysmolek, T.Suski, M.Leszczynski, I.Grzegory,S.Porowski: Physical Review Letters, 1997, 79[16], 3030-3

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GaN: Point DefectsStrong defect-specific low-frequency peaks were detected in the low-temperature Ramanspectra of hexagonal material which had been grown, by using molecular beam epitaxy,onto sapphire substrates. The intensity of these peaks was found to be enhanced byexcitation in resonance with yellow luminescence transitions. Their attribution toelectronic Raman scattering was confirmed. The results implied that the observedelectronic Raman scattering peaks were related to shallow donors which were notnecessarily hydrogenic. A very low (0.0117eV) frequency Raman peak was insteadattributed to a pseudo-localized vibrational mode.D.S.Jiang, M.Ramsteiner, K.H.Ploog, H.Tews, A.Graber, R.Averbeck, H.Riechert:Applied Physics Letters, 1998, 72[3], 365-7

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GaN: Point DefectsDeep-level transient spectroscopy was used to characterize defects in n-type sampleswhich had been grown by using reactive molecular-beam epitaxy. Five deep-level defectswere observed, with activation energies of 0.234eV (E1), 0.578eV (E2), 0.657eV (E3)0.961eV (E4) and 0.240eV (E5). The E1, E2 and E3 levels corresponded to deep levelswhich had been previously reported for n-type material which had been grown by usingboth hydride vapor-phase epitaxy and metal-organic chemical vapor deposition. The E4

and E5 levels did not correspond to any previously reported defect levels, and werecharacterized for the first time here.C.D.Wang, L.S.Yu, S.S.Lau, E.T.Yu, W.Kim, A.E.Botchkarev, H.Morkoç: AppliedPhysics Letters, 1998, 72[10], 1211-3

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GaN: Point DefectsDeep centers in Si-doped n-type layers, which had been grown by means of reactivemolecular beam epitaxy, were studied as a function of the growth conditions by usingdeep-level transient spectroscopy. When Si-doped layers were grown onto a Si-doped n+-type contact layer at 800C, they exhibited a dominant trap, C1, with an activation energyof 0.44eV and a capture cross-section of 1.3 x 10-15cm2. However, samples which weregrown at 750C onto an undoped semi-insulating buffer exhibited prominent traps, D1 andE1, with activation energies and capture cross-sections of 0.20eV and 8.4 x 10-17cm2, and0.21eV and 1.6 x 10-14cm2, respectively. The E1 trap was believed to be related to a N-vacancy defect, since the Arrhenius signature for E1 was very similar to that of thepreviously reported trap, E, which was produced by 1MeV electron irradiation of GaNwhen this was prepared using metal-organic chemical-vapor deposition and hydridevapor-phase epitaxy.Z.Q.Fang, D.C.Look, W.Kim, Z.Fan, A.Botchkarev, H.Morkoç: Applied Physics Letters,1998, 72[18], 2277-9

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GaN: Point DefectsThe blue Mg-induced 2.8eV photoluminescence band in metal-organic chemical vapordeposited material was studied in samples having various Mg contents. It appeared near toa Mg concentration of 1019/cm3 and, at higher concentrations, dominated the room-temperature photoluminescence spectrum. The excitation power dependence of the 2.8eVband confirmed its donor-acceptor pair recombination nature. It was suggested that theacceptor was isolated MgGa, while the spatially separated deep donor (0.43eV) wasattributed to a nearest-neighbor associate of a MgGa acceptor with a N vacancy; formedby self-compensation.U.Kaufmann, M.Kunzer, M.Maier, H.Obloh, A.Ramakrishnan, B.Santic, P.Schlotter:Applied Physics Letters, 1998, 72[11], 1326-8

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GaN: Point DefectsDeep traps in undoped n-type layers which had been grown, by means of organometallicvapor-phase epitaxy, onto sapphire substrates were studied by using temperature-dependent conductivity, photo-induced current transient spectroscopy, thermallystimulated current, electron beam-induced current, and band edge cathodoluminescencemethods. Electron traps, with energy levels that were 0.1 to 0.2eV below the conductionband, were detected as well as hole traps with energy levels that were about 0.25, 0.5 and0.85eV above the valence band edge. The use of cathodoluminescence and electronbeam-induced current measurements showed that the deep recombination centers weredistributed inhomogeneously, with a well-defined cellular pattern. Both the carrierlifetime and the luminescence intensity increased at cell walls, thus reflecting a lowerdensity of recombination centers. On the other hand, the density of the main hole trap

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(0.85eV) increased. The photoconductivities of many samples revealed very long decaytimes at temperatures of between 100 and 300K. It was concluded that the effect wasprobably was not related to shallow donors, such as Si, but was instead associated withunidentified deep centers with a 0.2eV barrier to electron capture.A.Y.Polyakov, N.B.Smirnov, A.V.Govorkov, M.Shin, M.Skowronski, D.W.Greve:Journal of Applied Physics, 1998, 84[2], 870-6

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GaN: Point DefectsNominally undoped layers which had been grown by means of molecular beam epitaxywere investigated by using temperature-dependent and frequency-dependent admittancespectroscopy. Deep-level transient spectroscopy was inapplicable because the space-charge region, which was needed for the detection of deep defects, existed only at lowfrequencies. Two deep defect levels were identified in molecular beam epitaxially grownlayers. The thermal activation energies were 0.45 and 0.63eV, respectively. These deeptraps were familiar from deep-level transient spectroscopic and thermally stimulatedconductivity studies.A.Krtschil, H.Witte, M.Lisker, J.Christen, U.Birkle, S.Einfeldt, D.Hommel: Journal ofApplied Physics, 1998, 84[4], 2040-3

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GaN: Point DefectsThe ultra-violet to yellow luminescence ratio was investigated, as a function of dopantconcentration, for Si donor concentrations ranging from 5 x 1016 to 7 x 1018/cm3. Theexperimental results showed that this ratio was approximately constant in the low-excitation regime, and was independent of the dopant concentration. A theoretical modelthat was based upon rate equations was developed which permitted the prediction of theyellow luminescence defect concentration as a function of the dopant concentration. Acomparison of the model with experimental results revealed that the defect concentrationincreased almost linearly with dopant concentration. This dependence was consistent withthe existence of compensating centers such as acceptor impurities, like C, orcompensating native defects.E.F.Schubert, I.D.Goepfert, J.M.Redwing: Applied Physics Letters, 1997, 71[22], 3224-6

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GaN: Point DefectsOptically-active defects in undoped epilayers which had been grown onto sapphire, byusing metal-organic chemical vapor epitaxy, were investigated by means ofphotoluminescence techniques. A new metastable defect which emitted blue light wasfound, as well as the well-known yellow luminescence centers. Upon excitation by a325nm He-Cd laser, this metastable defect, at low temperatures, exhibited a luminescencefatigue effect, with a decay time of about 360s. When the temperature was increased toroom temperature, it recovered its optically-active state. The yellow band emissionincreased in intensity as the blue-band emission intensity decreased. Analysis showed that

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this metastable center was a hole trap, with the Ga vacancy being the most probablecandidate.S.J.Xu, G.Li, S.J.Chua, X.C.Wang, W.Wang: Applied Physics Letters, 1998, 72[19],2451-3

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GaN: Point DefectsA series of undoped or Si-doped epilayers, with a carrier concentration of between 4 x1017 and 1.6 x 1019/cm3, were grown onto (00•1) sapphire in order to investigate theeffect of Si incorporation upon stress relaxation. It was found that, as the Si doping wasincreased, the bound exciton peaks gradually shifted to lower energies, at the rate0.042eV/GPa; due to the relaxation of residual thermal stresses. The results showed thatthe full-width at half-maximum of double-crystal X-ray diffractometry, and the intensityratio of yellow luminescence to edge emission, increased gradually as the Si content wasincreased. It was proposed that the Si doping of GaN epilayers introduced widespreaddefects, and gave rise to stress relaxation during cooling. The yellow luminescence wasattributed to a complex of VGa with the Si-induced defects.I.H.Lee, I.H.Choi, C.R.Lee, S.K.Noh: Applied Physics Letters, 1997, 71[10], 1359-61

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GaN: Point DefectsExperimentally identified deep levels in p-type material, at about 0.9 to 1, 1.4, and 1.8 to2eV above the valence-band maximum, were attributed to Ga vacancies. On the basis offirst-principles calculations, it was concluded that it was necessary to consider thestructural modification, VGa ? Nanti + VN, which resulted from the transfer of a nearest-neighbor N atom to a Ga-vacancy site, in order to explain the levels at 1 and 2eV. IsolatedN antisite and N-vacancy defects were found to give rise to additional deep levels at 1.4and 0.8eV, respectively.D.J.Chadi: Applied Physics Letters, 1997, 71[20], 2970-1

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GaN: Point DefectsA thermal equilibrium method was used to calculate the background carrier concentrationthat was related to intrinsic defects in an ideal GaN crystal. The results showed that the Nvacancy concentration did not exceed 2 x 1017/cm3 in samples which had been grown attemperatures ranging from 800 to 1500K. It was concluded that the N vacancy was one ofmajor sources of carriers when the carrier concentration was less than 2 x 1017/cm3, butthat the major sources were other defects when the carrier concentration was greater than2 x 1017/cm3.G.Y.Zhang, Y.Z.Tong, Z.J.Yang, S.X.Jin, J.Li, Z.Z.Gan: Applied Physics Letters, 1997,71[23], 3376-8

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GaN: Point DefectsLayers of AlN and GaN were grown onto 4º off-axis 6H-SiC (00•1) substrates by usingHe supersonic beams that were seeded with NH3. The AlN films were used as bufferlayers for GaN growth at 800C. It was estimated that 39% of the NH3 molecules, whichimpinged on the substrate surface during GaN film growth, were incorporated. The highstructural quality of the epitaxial GaN layers was confirmed by transmission electronmicroscopy and electron channelling patterns. The GaN films, which had a thickness ofabout 105nm, had a defect density of about 2 x 1010/cm2.V.M.Torres, M.Stevens, J.L.Edwards, D.J.Smith, R.B.Doak, I.S.T.Tsong: AppliedPhysics Letters, 1997, 71[10], 1365-7

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GaN: Point DefectsNew lines in the photoluminescence spectrum of lightly Be-doped material were reported.The low-temperature photoluminescence spectrum of lightly-doped samples wasdominated by a transition at 3.385eV; with first and second LO phonon replicas. Power-resolved photoluminescence measurements showed that a peak at 3.385eV narrowed inwidth, and shifted to higher energies, with increasing excitation intensity. The transitionwas therefore attributed to donor-to-acceptor recombination which involved a Be acceptorwith an optical ionization energy of between 0.09 and 0.1eV. This was much shallowerthan the acceptor level (0.25eV) that was introduced by Mg doping. Increased dopingresulted in quenching of the band-edge luminescence, and in the appearance of a broadtransition that was centered around 2.4eV. This was attributed to a complex whichinvolved Be. All luminescence was quenched upon increasing the doping level evenfurther.D.J.Dewsnip, A.V.Andrianov, I.Harrison, J.W.Orton, D.E.Lacklison, G.B.Ren,S.E.Hooper, T.S.Cheng, C.T.Foxon: Semiconductor Science and Technology, 1998,13[5], 500-4

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GaN: Point DefectsOptical isothermal capacitance transient spectroscopy was used to characterize prominentmid-gap carrier traps in Si-doped n-type samples which had been grown by means ofmetal-organic vapor-phase epitaxy. Strong carrier photo-ionization was detected from 2deep levels to the conduction band. The first level photo-ionized over a broad range, frombelow 1.8 to over 2.3eV. It was seen in all n-type material, was believed to be defect-related and to be involved in yellow luminescence. The second and more dominant peakdeveloped with an incident photon energy of about 2.3eV. This was a previouslyunreported and unidentified impurity-related level.P.Hacke, H.Okushi: Applied Physics Letters, 1997, 71[4], 524-6

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GaN: Point DefectsSite-selective photoluminescence and photoluminescence excitation spectroscopicinvestigations were carried out, at 6K, on Nd3+ emissions from the 4F3/2 ? 4I11/2 transitionin Nd-implanted samples which had been grown by using metal-organic chemical vapordeposition. The site-selective photoluminescence excitation spectra which were detectedat emission wavelengths that were characteristic of each of the distinct Nd3+

photoluminescence bands included spectral features which were representative ofexcitation by above-gap absorption, by direct sharp-line Nd3+ 4f-shell absorption, and bybroad below-gap absorption bands. These could be attributed to defects and impuritiesand to a possible iso-electronic trap that was associated with one of the five Nd3+ sites. Itwas concluded that the Nd3+ site which was excited by direct sharp-line Nd3+ 4f-shellabsorption was the predominant or highest-concentration Nd3+ center. The excitationmechanisms for the other four Nd3+ sites all involved the non-radiative transfer of energy,from impurity-related or defect-related traps, to neighboring rare-earth atoms.S.Kim, S.J.Rhee, X.Li, J.J.Coleman, S.G.Bishop: Physical Review B, 1998, 57[23],14588-91

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GaN: Point DefectsThe atomic and electronic structures of H-vacancy complexes were investigated by meansof pseudopotential density-functional calculations. The calculated formation energiesprovided information as to the likelihood of incorporation of these complexes into n-typeor p-type material. The predicted binding energies provided a measure of the dissociationenergy. The estimated vibrational frequencies also yielded a signature, for the complex,that might facilitate experimental identification. It was suggested that an observed red-shift corresponded to a decrease in the intensity of the L1 line, accompanied by anincrease in intensity of a new line that was related to the N vacancy. The as-grownmaterial could contain a certain concentration of hydrogenated N vacancies which had alevel, near to the conduction band, that was thought to give rise to the L1 line. When thematerial was annealed, the hydrogenated vacancy complexes dissociated. The calculatedremoval energy of 1.56eV was consistent with complex dissociation at around 500C. Theresultant N vacancies had a level, near to the valence band, which was suggested to beresponsible for a line at around 2.9eV (420nm). It was noted that the +/3+ transition ofthe N vacancy was associated with a very large lattice relaxation.C.G.Van de Walle: Physical Review B, 1997, 56[16], R10020-3

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GaN, InGaN: Point DefectsBand structures and deep levels were calculated, for both cubic and hexagonal nitridesemiconductors, by using the sp3s* tight-binding formulation and the Green's functiontechnique. An anti-bonding s-like state which was produced by a N vacancy waspredicted to appear at 0.3eV below the conduction-band edge in GaN. This became

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shallower, and then resonant with the conduction band in InGaN, as the In content wasincreased.E.Yamaguchi, M.R.Junnarkar: Journal of Crystal Growth, 1998, 189-190, 570-4

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GaN: Stacking FaultsThe structure of a (12̄•0) prismatic stacking fault was calculated from first principles. Thefault vector was ½[10•1]. Such a fault was believed to occur when GaN was grown ontosapphire or SiC, and was thought to originate when basal-plane stacking faults foldedonto the prism plane. It was found that that the boundary was heavily reconstructed, butthat the atoms remained 4-fold coordinated, with no wrong bonds or dangling bondspresent. In spite of the existence of distorted 4-membered rings of bonds at the boundary,the fault did not introduce deep states into the gap. The calculated stacking fault energywas 0.072eV/Å2.J.E.Northrup: Applied Physics Letters, 1998, 72[18], 2316-8

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GaN: Stacking FaultsThe electronic structures of stacking faults were calculated within the framework of localempirical pseudopotential theory. The stacking faults, in both zincblende- and wurtzite-structured samples, were predicted to introduce electronic levels within the band gap;with an energy that was 0.13eV above the valence-band top. These levels were found tooriginate from interface states in hetero-crystalline wurtzite(00•1)/zincblende(111)interfaces.Z.Z.Bandic, T.C.McGill, Z.Ikonic: Physical Review B, 1997, 56[7], 3564-6

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GaN: Surface DefectsFilms which had been grown by means of metal-organic chemical vapor deposition, ontoc-sapphire substrates, were studied as a function of the post-growth rapid thermalannealing temperature (600C to 800C). Similar planar defects were observed in all of theheat-treated samples, and only their number density differed. It was demonstrated thatthere was a clear relationship between an improved crystalline quality, and post-growthhigh-temperature annealing. The number density of the near-surface defects wasdecreased by 61% as the annealing temperature was increased from 600 to 800C. Thesuppression of near-surface defects encouraged the development of a ß-W2N interfacialphase, and promoted interface smoothness.M.W.Cole, F.Ren, S.J.Pearton: Applied Physics Letters, 1997, 71[20], 3004-6

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GaN: Surface DefectsFirst-principles calculations were made of the formation energy of H-terminated (101̄0)surfaces. The calculations indicated that H adsorption on this surface would proceed viathe saturation of pairs of Ga and N dangling bonds, rather than via the exclusive

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occupation of only one type of bonding site. At absolute zero, the surface energy of thefully H-terminated surface was found to be 0.02eV/cell; as compared with 1.95eV/cell forthe bare surface. Results were presented for the N-H and Ga-H stretching and bendingeigen-frequencies. Dissociative adsorption of NH3 via the formation of N-H and Ga-NH2bonds was exothermic, and reduced the surface formation energy to a value which wasless than 0.1eV at absolute zero.J.E.Northrup, R.Di Felice, J.Neugebauer: Physical Review B, 1997, 56[8], R4325-8

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GaN: Surface Defects and Surface ReconstructionThe effect of growth conditions and surface polarity upon the morphology of (00•1)surfaces was studied. The decisive factors which favoured lowest-energy reconstructionswere the ionic nature of the Ga-N bond and the strength of the N-N bond. Under Ga-richconditions, Ga-adatom reconstructions were the most energetically favourable of the 2 x 2reconstructions to be studied; regardless of the polarity of the surface. Under N-richconditions, a N-adatom reconstruction was most stable on the Ga-terminated surface. Onthe N-terminated surface, a N atom spontaneously bonded to a surface N atom so as tocreate a N molecule and a vacancy substrate. The molecule was weakly bound, and wasexpected to evaporate at high growth temperatures. Since vacancy reconstruction wasunstable over the whole range of chemical potentials, evaporation of N2 molecules wasexpected to lead to a cascade of reconstructions of the substrate. For both polarities, ¾ ofa monolayer of H stabilized the relaxed ideally-cleaved surfaces. These reconstructionshad 2 x 2 symmetry. The very small relaxation of surface atoms on N-terminatedsurfaces, which was below the resolution of low-energy electron diffraction, caused thissurface to appear as 1 x 1.K.Rapcewicz, M.B.Nardelli, J.Bernholc: Physical Review B, 1997, 56[20], R12725-8

[446-164-180]

GaN: Surface ReconstructionTheoretical formation energies for various possible reconstructions, having 1 x 1 and 2 x2 periodicities on (00•1) and (00•̄1) surfaces, were presented. It was found that, duringmolecular beam epitaxial growth in the (00•1) direction, 2 x 2 structures became stableunder N-rich growth conditions while a Ga-rich environment was expected to yieldstructures having a 1 x 1 periodicity. With regard to molecular beam epitaxial growth on(00•̄1) surfaces, reconstructions with the 1 x 1 periodicity had low energies. Duringmetal-organic chemical vapor deposition, where H-terminated surfaces could occur, 1 x 1periodicities were found to be stable for both growth directions.J.Elsner, M.Haugk, G.Jungnickel, T.Frauenheim: Solid State Communications, 1998,106[11], 739-43

[446-164-180]

GaN: Surface ReconstructionAn investigation was made of clean, and As-covered, zincblende (001) surfaces by meansof first-principles total-energy calculations. In the case of clean surfaces, the results

181

GaN Nitrides GaN

revealed a novel surface structure which was very different from the well-establisheddimer structures which were usually observed on polar III-V (001) surfaces. Theenergetically most stable surface involved a Peierls distortion of the truncated (1 x 1)surface rather than the addition or removal of atoms. This surface exhibited a (1 x 4)reconstruction which consisted of linear Ga tetramers. It was also found that a sub-monolayer of As significantly lowered the surface energy; thus indicating that As mightbe a good surfactant.J.Neugebauer, T.Zywietz, M.Scheffler, J.E.Northrup, C.G.Van de Walle: PhysicalReview Letters, 1998, 80[14], 3097-100

[446-164-181]

GaN: Surface ReconstructionReconstructions of the (00•̄1) surface were studied for the first time. Four primarystructures, 1 x 1, 3 x 3, 6 x 6, and c(6 x 12) were observed by using scanning tunnellingmicroscopy and reflection high-energy electron diffraction. On the basis of first-principlescalculations, the 1 x 1 structure was shown to consist of a Ga monolayer that was bondedto a N-terminated GaN bilayer. By combining experimental data, and theory, it wasargued that the 3 x 3 structure was an adatom-on-adlayer structure with one additional Gaatom per 3 x 3 unit cell.A.R.Smith, R.M.Feenstra, D.W.Greve, J.Neugebauer, J.E.Northrup: Physical ReviewLetters, 1997, 79[20], 3934-7

[446-164-181]

GaN/Al2O3: Al, Ga InterdiffusionDistribution profiles of Ga and Al near to the interface of the n-GaN/sapphire systemwere determined by means of X-ray energy dispersive spectroscopy. It was found that Ga,diffusing into the sapphire substrate, obeyed the law-of-remainder probability function;giving a diffusivity of 2.30 x 10-13cm2/s. The diffusion was associated with GaN growthat high temperatures. When compared with the diffusion of Ga into the sapphire substrate,much less Al anti-diffusion from the substrate into the GaN film occurred; with adiffusivity of about 4.8 x 10-15cm2/s.S.Fung, X.Xu, Y.Zhao, W.Sun, X.Chen, N.Sun, T.Sun, C.Jiang: Journal of AppliedPhysics, 1998, 84[4], 2355-7

[446-164-181]

GaN/Al2O3: DislocationsA combination of atomic force microscopy and scanning capacitance microscopy wasused to investigate the relationship between the surface morphology and the near-surfaceelectrical properties of GaN films which had been grown onto c-axis sapphire substratesby means of metal-organic chemical vapor deposition. Local regions surrounding thesurface termination of threading dislocations exhibited a reduced change in capacitance,with applied voltage, relative to those regions that contained no dislocations. Capacitance-

182

GaN Nitrides GaN

voltage characteristics which were obtained for these regions indicated the presence ofnegative charge in the vicinity of dislocations.P.J.Hansen, Y.E.Strausser, A.N.Erickson, E.J.Tarsa, P.Kozodoy, E.G.Brazel,J.P.Ibbetson, U.Mishra, V.Narayanamurti, S.P.DenBaars, J.S.Speck: Applied PhysicsLetters, 1998, 72[18], 2247-9

[446-164-182]

GaN/GaAs: Dislocations and Stacking FaultsHigh-resolution electron microscopy was used to characterize the structure of ß-phaseepilayers which had been grown onto (001)GaAs substrates by means of plasma-assistedmolecular-beam epitaxy. A radio-frequency plasma source was used to producechemically active N. Exposure of the layer surface to the As flux during growth of thefirst few monolayers was shown to result in a markedly flat GaN/GaAs interface. Thebest-quality GaN layers were achieved by ensuring near-stoichiometric nucleation, usingan optimum Ga/N ratio. Deviation from these nucleation conditions led to interfaceroughening and to the formation of a wurtzite within the GaN layer. All of the layerscontained a high density of stacking faults, near to the interface, but this density sharplydecreased towards the surface. The stacking faults were anisotropically distributed withinthe GaN layer; probably due to the differing behavior of a, as compared with ß,dislocations in cubic GaN. Most of the stacking faults intersected the interface along linesthat were parallel to the major flat of the GaAs wafer. The stacking faults were oftenassociated with atomic steps at the GaN/GaAs interface.S.Ruvimov, Z.Liliental-Weber, J.Washburn, T.J.Drummond, M.Hafich, S.R.Lee: AppliedPhysics Letters, 1997, 71[20], 2931-3

[446-164-182]

GaN/InGaN: InterdiffusionThe effect of the growth parameters upon the optical quality of InGaN which had beengrown onto GaN was investigated. The photoluminescence spectrum of a sample with alow-temperature grown GaN cap layer, or a graded-temperature grown GaN cap layer,exhibited a shorter peak wavelength than that of a sample which had a GaN layer that hadbeen grown at normal temperatures. The shift in the peak wavelength increased, withincreasing layer thickness, for samples with caps that had been grown at lowtemperatures. This was because the defects that were present in a cap layer which hadbeen grown at low temperatures encouraged the out-diffusion of In atoms duringtemperature ramping. The narrower line-widths and higher intensities of thephotoluminescence spectra of InGaN after In out-diffusion were attributed to a reductionin strain, or in the dislocation or defect contents. The Raman and Auger electron spectraalso indicated that GaN which was grown at low temperatures contained a lot of defects.This reduced the phonon peak intensity, and led to the interdiffusion of In atoms duringthe growth of GaN/InGaN heterostructures.J.S.Tsang, J.D.Guo, S.H.Chan, M.S.Feng, C.Y.Chang: Japanese Journal of AppliedPhysics - 1, 1997, 36[3B], 1728-32

[446-164-182]

183

GaN Nitrides SiNx

GaN/SiC: Ion Implantation and Point DefectsIt was observed that He+ bombardment increased the relative intensity of the so-calledblue emission and resistivity of film samples, and decreased the intensity of the near-bandedge photoluminescence. Because the intensity of the main peak was sharplydecreased, the fine structure of the near-bandedge photoluminescence after He+

bombardment was observed. Upon comparing the observed sharp lines with thephotoluminescence peaks of O-doped material, it was concluded that O could produce acomplex which was characterized by a marked localization of free carriers and a largelattice distortion. The zero-phonon line of this defect had an energy which was close tothe band-gap energy of GaN.V.A.Joshkin, C.A.Parker, S.M.Bedair, L.Y.Krasnobaev, J.J.Cuomo, R.F.Davis,A.Suvkhanov: Applied Physics Letters, 1998, 72[22], 2838-40

[446-164-183]

GaN/SiC: Dislocations, Grain Boundaries, and Stacking FaultsTransmission electron microscopy and localized electron energy loss spectroscopy, in thebandgap-energy regime, were applied to this system. Although grown on a cubicsubstrate, the GaN film exhibited hexagonal phases having 2 different orientationalrelationships with respect to the substrate. One of these partially transformed into a cubicphase, via sequential faulting. The band-gap was determined in local regions of goodcrystallinity, and in the proximity of stacking faults, dislocations, grain boundaries, andhetero-interfaces. Extrapolation of the electron energy loss spectra gave a band-gap valueof 3.4eV for the unfaulted hexagonal close-packed phase. This was consistent with thepredicted value. The band-gap itself was heavily masked, in the electron energy lossspectra, by near-bandedge states around 3.2 and 3.3eV; especially in faulted grains. Localvariations in the mid-gap and near-bandedge density of states in the electron energy lossspectra were found to be related to microstructural defects in the GaN film.Photoluminescence measurements at 6K did not reveal the bound exciton at about3.47eV, but instead suggested the presence of donor acceptor pairs at 3.27eV.U.Bangert, A.Harvey, J.Davidson, R.Keyse, C.Dieker: Journal of Applied Physics, 1998,83[12], 7726-9

[446-164-183]

SiNx: Point Defects and Defect AnnealingNear-edge X-ray absorption fine-structure measurements at the N-K edge were used tomonitor the evolution of defect-related structures, in the spectra of buried nitride films, asa function of the implantation dose. The buried films were prepared by implanting 35keV14N+ ions into Si, to doses ranging from 2 x 1017 to 2 x 1018/cm2. Defect-relatedresonances, RL1 and RL2, appeared at 401.1 and 403.3eV, respectively. The RL1resonance was characteristic of a defect structure at low and intermediate implanteddoses, and could be annealed out with an activation energy of 0.5eV. The RL2 resonancewas the signature of excess N in N-rich films. It was attributed to the transitions of 1s-

184

SiNx Nitrides TiN

electrons to unfilled states with p-component at a defect site which contained a Ndangling bond. It could be annealed out only by prolonged annealing at 1150C.E.C.Paloura: Applied Physics Letters, 1997, 71[22], 3209-11

[446-164-184]

Si(O,N): Point DefectsBy using valence-band and Si 2p core-level photo-electron spectroscopy, it was shownthat short-range order in the amorphous oxynitride was governed by the Mott rule. Thatis, each Si atom was coordinated with four O and/or N atoms, each O atom (as in SiO2)was coordinated with two Si atoms, and each N atom (as in Si3N4) was coordinated withthree Si atoms. The effect of removing Si-Si bonds (hole traps) from the interface ofSiO2/Si by nitridation, and the cause of Si-Si bond creation near to the top surface of gateoxynitrides in semiconductor devices, could be understood for the first time by invokingthe Mott rule.V.A.Gritsenko, J.B.Xu, R.W.M.Kwok, Y.H.Ng, I.H.Wilson: Physical Review Letters,1998, 81[5], 1054-7

[446-164-184]

TaSiN: O DiffusionNon-crystalline layers were studied with regard to their barrier effect upon O diffusion. Itwas found that the penetration depth of O diffusion decreased markedly with increasingSi content of a TaSiN layer, and reached 20nm in a Ta22Si35N43 layer. However, theresistivity also increased. A good diffusion barrier layer with a low sheet resistance wasrepresented by Ta50Si16N34. Penetration depths of less than 40nm were observed in aslightly Si-rich Ta36Si27N37 layer during O2 annealing at 850C.T.Hara, M.Tanaka, K.Sakiyama, S.Onishi, K.Ishihara, J.Kudo: Japanese Journal ofApplied Physics - 2, 1997, 36[7B], L893-5

[446-164-184]

Ti(C,N): TwinsThe microstructures of samples which had been fabricated by hot pressing TiN and TiCpowders were studied by means of transmission electron microscopy. The carbonitridehad a face-centered cubic structure, but there was no evidence of ordering of theinterstitial atoms, and (111)-type twinning with a twin plane of Ti atoms was observed.Incoherent twin boundaries or steps were found to connect coherent twin boundarieswhich lay on (111). It was found that (a/6)<112̄>-type dislocations were found to beassociated with atomic steps at the twin boundary.K.Han, G.C.Weatherly: Philosophical Magazine Letters, 1997, 76[4], 247-58

[446-164-184]

TiN/B-C-N: Interdiffusion and Ion BombardmentDiffusion in multi-layers, during vacuum annealing at temperatures of up to 1000C and/orduring 300keV Ar bombardment, was studied. The effect of stress upon diffusion was

185

TiN Nitrides/Miscellaneous In16Fe8S32

proved by carrying out equivalent annealing or irradiation upon a multi-layer, with andwithout compressive stressing. Demixing or phase separation was observed duringthermal annealing. On the other hand, mixing occurred during irradiation. Bothphenomena were enhanced in the presence of a stress field.S.Fayeulle, M.Nastasi: Applied Physics Letters, 1998, 73[8], 1077-9

[446-164-185]

WSiN: P DiffusionThe diffusion-barrier behavior of films which were formed by ECR plasma nitridationwas investigated. It was noted that they exhibited good barrier characteristics when aradio-frequency bias was applied to the substrate during nitridation; even though the filmwas less than 6nm thick. Applying such a bias increased the N content. Also, Si atomswere preferentially sputtered, and local atomic ordering was decreased, because the effectof ion bombardment was markedly enhanced. It was concluded that these characteristicscontributed to the suppression of P diffusion via interstitial sites.A.Hirata, K.Machida, S.Maeyama, Y.Watanabe, H.Kyuragi: Japanese Journal of AppliedPhysics - 1, 1998, 37[3B], 1251-5

[446-164-185]

Miscellaneous

Ca5(PO4)3OH: Grain BoundariesThe electron structure of hydroxyapatite was investigated by using the tight-bindingrecursion method. The local density of states, partial density of states, structural energyand charge transfer were calculated, as well as the interatomic energy. The calculationsshowed that both the OH group and the PO4 group could be regarded as being individualstructural units, due to the strong interactions between the atoms. The atomic structure ofa grain boundary was also determined on the basis of the near-coincidence site latticemodel, while taking account of the bond lengths in the crystal, and of the latticedistortion. The electron structure and energy of a S = 7 grain boundary were calculated,and the results showed that a grain boundary with a slight lattice misfit had a lowerenergy.Q.Song, C.Wang, S.Wen: Philosophical Magazine A, 1998, 77[5], 1309-21

[446-164-185]

In16Fe8S32: Point DefectsCombined 57Fe Mössbauer and X-ray diffraction (Rietveld analysis) studies were made ofa spinel, before and after chemical Li insertion. In order to accommodate the newlyinserted cations, 2 mechanisms operated. Firstly, at low Li contents, cation migration tookplace from the usually occupied 8a and 16d sites towards the usually unoccupied 8b and16c sites; with distortion of the S coordination polyhedra. Secondly, at high Li contents,

186

In16Fe8S32 Miscellaneous KTiOPO4

when both 16c and 16d octahedral sites were occupied, a decrease in the distortion of thepolyhedra was observed. During insertion, no clear reduction in Fe atoms was observedbut there was an increase in the covalent character of the Fe-S bonds.C.Bousquet, A.Krämer, C.P.Vicente, J.L.Tirado, J.Olivier-Fourcade, J.C.Jumas: Journalof Solid State Chemistry, 1997, 134[2], 238-47

[446-164-186]

KH2PO4: DislocationsThe results of atomic force microscopic measurements of {101} faces showed that, forsupersaturations ranging from 3 to 30%, the terrace widths on vicinal growth hillocks -formed by dislocations - were nearly independent of both the supersaturation and thedislocation structure. This contradicted the predictions of simple Burton-Cabrera-Frankmodels. The data also showed that, for Burgers vectors greater than one unit step-height,the dislocations had hollow cores; in accordance with theoretical predictions. Bothanalytical and numerical analyses showed that a model which took account of the effectof these cores, upon the period of step rotation, predicted a dependence of the slope uponsupersaturation and Burgers vector. This conclusion was in good agreement withexperimental results. The calculations also showed that the effect of the core perimeterupon step transit-time predominated over the effect of a reduced step velocity that wasdue to stresses near to the core. A simple analytical expression could therefore be used todescribe the slope; even in the case of anisotropic step kinetics. The results were used toexplain the reproducible nature of the macroscopic growth rates, and to re-scale growth-rate data - for various temperatures and supersturations - so that they could be fitted by asingle curve.J.J.De Yoreo, T.A.Land, L.N.Rashkovich, T.A.Onischenko, J.D.Lee, O.V.Monovskii,N.P.Zaitseva: Journal of Crystal Growth, 1997, 182, 442-60

[446-164-186]

KTiOPO4: Er, Nb, Nd DiffusionConcentration profiles in (001) monocrystals were measured by means of Rutherfordback-scattering spectroscopy. It was found that the diffusion data which were deducedfrom the profiles could be described by:

Er: D (cm2/s) = 3.6 x 10-16exp[-6.4(eV)/kT]Nb: D (cm2/s) = 1.2 x 10-14exp[-3.8(eV)/kT]Nd: D (cm2/s) = 1.3 x 10-11exp[-5.3(eV)/kT]

M.J.Martin, C.Zaldo, M.F.Da Silva, J.C.Soares, F.Diaz, M.Aguilo: Journal of Physics -Condensed Matter, 1997, 9[34], L465-9

[446-164-186]

KTiOPO4: Point DefectsA careful analysis of the electron paramagnetic resonance spectra of Fe3+ ions in thismaterial was carried out in order to resolve the weaker lines. Eight additional sets of fine-structure lines were identified, and their respective spin-Hamiltonian parameters weredetermined. The 8 magnetically inequivalent Fe3+ sites were attributed to two Ti sites.

187

KTiOPO4 Miscellaneous MgCO3

Two different centers, which arose from Fe3+ ions at the above Ti sites, were explained interms of differing charge compensations. It was suggested that a vacancy of the first-neighbor O simultaneously played the role of a divalent positive charge compensator fortwo Fe3+ ions at the adjacent Ti sites.S.W.Ahn, S.H.Choh: Journal of Physics - Condensed Matter, 1998, 10[2], 341-8

[446-164-187]

MgCO3: Point DefectsThe low-temperature thermal depolarization spectrum of polycrystalline magnesiterevealed a unique broad dipolar relaxation peak which exhibited a maximum at about140K. The relaxation mechanism was suggested to be related to a dipole defectpopulation with a distribution of relaxation times. By using a partial heating technique,activation energies were found which ranged from 0.19 to 0.30eV. A review of dielectricrelaxation in calcite-family materials showed that each sub-lattice type (Ca and Mg)favored a certain defect structure. It was concluded that Mn2+ and Sr2+ impurities, as wellas H2O molecules or OH ions, were the most likely components of the dipole defectconfigurations.A.N.Papathanassiou, J.Grammatikakis: Physical Review B, 1997, 56[14], 8590-8

[446-164-187]

General

Twin Boundaries in d-Wave SuperconductorsTwin boundaries in orthorhombic d-wave superconductors were investigated numericallyby using the Bogoliubov-deGennes formalism within the context of an extended Hubbardmodel. The twin boundaries were represented by tetragonal regions of variable width,with a reduced chemical potential. For sufficiently large twin boundary widths andchanges in chemical potential, an induced s-wave component could break time-reversalsymmetry at a low temperature. This temperature, and the magnitude of the imaginarycomponent, were found to depend strongly upon the electron density.D.L.Feder, A.Beardsall, A.J.Berlinsky, C.Kallin: Physical Review B, 1997, 56[10],R5751-4

[446-164-187]

Enhancement of Superconductivity at Defects in High-Temperature SuperconductorsIt was shown that long-range strain fields around structural defects in high-temperaturesuperconductors could give rise to localized superconducting domains at temperaturesthat were appreciably higher than the bulk critical temperature; regardless of themicroscopic mechanism of superconductivity and the nanostructure of the defects. Thiseffect was attributed to the strongly non-monotonic dependence of the bulk critical

188

General

temperature upon pressure and hole concentration. The increase in critical temperaturewas calculated for edge dislocations, low-angle grain boundaries, and metastable lineardislocation arrays; while taking account of the anisotropic strain dependence of thecritical temperature in the ab-plane. It was suggested that the superconducting state on thegrain boundaries resulted from the proximity coupling of superconducting domains thatwere localized on a periodic chain of edge dislocations. In this case, the change in thecritical current decreased with the misorientation angle and vanished at a critical value ofthe latter. This angle was governed by competition between a strain-field enhancement ofthe critical temperature and the suppression of superconductivity in the dislocation cores.In the case of metastable dislocation arrays which were caused by plastic deformation, thestrain-induced increase in critical temperature was much more pronounced than that forgrain boundaries and occurred in macroscopic domains which were much larger than thecoherence length. The localized remanent strains in these domains could be high enoughto reveal the absolutely maximum critical temperature. It was expected that this might notbe attainable by using hydrostatic pressures.A.Gurevich, E.A.Pashitskii: Physical Review B, 1997, 56[10], 6213-25

[446-164-188]

Percolation Model for the Anomalous Conductivity of Fluorite-Related OxidesThe Monte Carlo simulation technique was used to study the anomalous conductivity ofaliovalently doped fluorite-related oxides. In order to investigate the effect of interactionsbetween dopants and O vacancies, 3 simple model interactions were tested. These were:nearest-neighbor attraction, nearest-neighbor repulsion, and a barrier model which wasbased upon the assumption of a reduced O mobility in the neighborhood of the dopants. Itwas found that percolation theory was essential for correlating local interactions and long-range mobilities of the ions. Numerical results showed that satisfactory agreement withexperiment could be obtained only in the case of the barrier model.M.Meyer, N.Nicoloso, V.Jaenisch: Physical Review B, 1997, 56[10], 5961-6

[446-164-188]

Point Defects in the Flux-Line Lattice of SuperconductorsThe self-energies and interaction energies of vacancies and interstitials in a triangularlattice of parallel Abrikosov vortices in type-II superconductors were calculated by usingLondon theory. Stable and metastable equilibrium configurations of the flux-line latticearound such point defects were investigated. It was again shown that the vacancy ofhighest (6-fold) symmetry usually did not exhibit the lowest energy. Due to relaxation ofthe surrounding vortex lattice, the defect energies were very small when compared withthe binding energy of a single vortex. The interaction of point defects was weak, andcould be repulsive or attractive; depending upon their type and distance and upon the ratioof the magnetic penetration depth to the average vortex spacing.E.Olive, E.H.Brandt: Physical Review B, 1998, 57[21], 13861-71

[446-164-188]

189

General

Low-Angle Grain Boundaries in High-Temperature SuperconductorsMechanisms were considered which could account for an observed rapid decrease in thecritical current density as a function of the misorientation angle of grain boundaries inhigh-Tc superconductors. It was shown that the angular dependence was governed mainlyby a decrease in the current-carrying cross-section, due to insulating dislocation cores andto a progressive local suppression of the superconducting order parameter near to thegrain boundaries as the angle increased. The insulating regions near to the dislocationcores resulted from a strain-induced local transition to the insulating antiferromagneticphase of the high-temperature superconductor. The structure of the non-superconductingcore regions and current channels in the grain boundaries was markedly affected by ananisotropy of the strain dependence of the critical temperature. A mechanism wasproposed for the progressive suppression of superconductivity, at grain boundaries, as afunction of the misorientation angle. This supposed that there was an excess ionic charge,on the grain boundaries, which shifted the chemical potential in the layer by an amountwhich was of the order of the screening length near to the grain boundaries. The localsuppression of the order parameter was magnified by the proximity, of any high-temperature superconductor, to a metal-insulator transition. This behavior was due, inturn, to their low carrier densities and extended saddle-point singularities in the electrondensity of states near to the Fermi surface. By taking account of these mechanisms, theangular variation of the critical current was calculated analytically by solving theGinzburg-Landau equation. The model described well the observed quasi-exponentialdecrease in critical current, with misorientation, for many high-temperaturesuperconductors. The d-wave symmetry of the order parameter weakly affected thecritical current density, at small angles, and could not account (by several orders ofmagnitude) for the fall in the critical density as the angle increased from 20 to 40º.A.Gurevich, E.A.Pashitskii: Physical Review B, 1998, 57[21], 13878-93

[446-164-189]

191

Author Index

Abare, A........................................ [163]Abe, H.............................................. [86]Abe, K........................................ [73, 75]Abe, Y.................................... [126, 138]Adriaenssens, G.J. ............................ [69]Afanasev, V.V.................[134, 136, 137]Aggarwal, S. .................................. [114]Aguilo, M....................................... [186]Ahlgren, T........................................ [68]Ahn, C.C. ......................................... [82]Ahn, S.W. ...................................... [187]Ainger, F.W. .................................. [148]Aird, A. .......................................... [146]Aitchison, P.R. ............................... [110]Ajayan, P.M. .................................... [65]Akagi, K........................................... [71]Akasaki, I. ...................................... [171]Albrecht, M.................................... [169]Alimoussa, A.................................. [115]Alonso, J.A. ................................... [113]Amano, H....................................... [171]Ambacher, O.......................... [164, 169]Anane, A. ....................................... [116]Anderson, J.F. ................................ [109]Andrén, H.O................................... [100]Andrianov, A.V.............................. [177]Angerer, H. .................................... [164]Anwand, W. ..................................... [73]

Aoki, Y. ..................................... [74, 75]Arai, K. .......................................... [145]Arai, N. .......................................... [131]Aswal, D.K. ................................... [103]Auciello, O..................................... [102]Audurier, M. .................................. [157]Augustine, G. ................................... [77]Averback, R.S. ............................... [120]Averbeck, R. .................................. [173]

Baccaro, S. .................................... [128]Balakrishna, V. ................................ [77]Baldinozzi, G. ................................ [147]Bandic, Z.Z. ................................... [179]Bangert, U...................................... [183]Baranov, P.G.................................... [78]Baranowski, J.M. ........................... [173]Barbier, A. ..................................... [160]Barrett, D.L...................................... [77]Barski, A........................................ [160]Barsoum, M.W........................... [83, 84]Bartsch, M. .................................... [154]Baufeld, B...................................... [154]Beardsall, A. .................................. [187]Beaumont, B. ................................. [167]Bechstedt, F. .................................... [80]Bedair, S.M.................................... [183]Beling, C.D. ..................................... [80]

Author Index

192

Bérar, J.F........................................ [147]Berastegui, P. ................................. [100]Berghuis, P..................................... [104]Berkowitz, A.E............................... [110]Berlinsky, A.J................................. [187]Bernaerts, D. .................................... [69]Bernasik, A. ..................................... [98]Bernhardt, J...................................... [80]Bernholc, J. .................................... [180]Berthet, P. ...................................... [116]Berthon, J....................................... [116]Bharuth-Ram, K. .............................. [69]Bidault, O....................................... [127]Bigot, B.......................................... [146]Birkle, U. ....................................... [175]Bischoff, T. .................................... [139]Bishop, S.G. ................................... [178]Bleichner, H. .................................... [76]Bolse, W. ....................................... [133]Bondarenko, V.I. ............................ [150]Bonn, D.A...................................... [101]Bontempi, E. .................................. [146]Borchardt, G................................... [154]Borgia, B........................................ [128]Börjesson, L. .................................... [93]Born, E........................................... [169]Borstel, G....................................... [112]Botchkarev, A. ............................... [174]Botchkarev, A.E............................. [173]Boudreau, M. ................................... [80]Bour, D.P. ...................................... [162]Bourret, A. ..................................... [160]Bousquet, C.................................... [186]Boussetta, H. .................................. [108]Bouwmeester, H.J.M.........[95, 114, 138]Brandon, D.G................................... [90]Brandt, E.H. ................................... [188]

Brauer, G. .................................. [73, 80]Brazdeikis, A. ................................ [141]Brazel, E.G. ................................... [182]Bremser, M.D. ............................... [168]Briddon, P.R. ............................[70, 167]Briggs, G.A.D. ................................. [98]Brillson, L.J. .................................. [136]Bronsveld, P.M. ............................... [89]Brook, R.J. ....................................... [93]Brötz, J........................................... [106]Brown, I.M. ................................... [160]Browning, N.D........................[144, 167]Burak, Y.V..................................... [117]Burggraaf, A.J................................ [138]Bygden, J. ...................................... [123]Bykov, I.P. ..................................... [112]Byun, J.D. ........................................ [93]

Calas, G. ....................................... [116]Calcagno, L...................................... [73]Calvarin, G................................[79, 147]Campillo, J.M. ................................. [73]Cantwell, G.................................... [151]Cardwell, D.A. ............................... [103]Carim, A.H. ..................................... [89]Carolan, J.F.................................... [101]Carter, C.B....................................... [97]Casais, M.T.................................... [113]Casanove, M.J................................ [115]Castaing, J........................................ [88]Cawley, J.D...................................... [81]Cecilia, A. ...................................... [128]Chadi, D.J. ..............................[128, 176]Chai, B.H.T...................................... [92]Chan, S.H....................................... [182]Chang, C.Y. ................................... [182]Chao, Y.C. ....................................... [81]

Author Index

193

Chapman, J.N................................. [110]Charlier, J.C. .................................... [65]Chaudhri, M.M............................... [123]Checchetto, R................................. [131]Chen, C.H. ..................................... [138]Chen, C.L....................................... [141]Chen, J.H. ........................................ [69]Chen, S.J. ......................................... [87]Chen, X.......................................... [181]Chen, Y.......................................... [173]Cheng, T.S. .................................... [177]Cheng, W.C.................................... [130]Cherns, D. ...................................... [172]Ching, W.Y. ..................................... [90]Cho, W.J. ......................................... [95]Cho, W.S.......................................... [95]Choh, S.H....................................... [187]Choi, G.M. ..................................... [152]Choi, I.H. ....................................... [176]Choi, S.S. ......................................... [95]Choi, W.K........................................ [78]Chong, T.C..................................... [143]Chou, S.T. ...................................... [134]Christen, J. ..................................... [175]Christensen, N.E. ........................... [112]Christiansen, S. .............................. [169]Chu, C.W. ...................................... [141]Chu, P. ............................................. [94]Chua, S.J. ....................................... [176]Citrin, P.H........................................ [85]Claeson, T. ............................. [105, 150]Cockayne, D.J.H. ........................... [166]Cohen, L.F. .................................... [113]Cohen, R.E..................................... [122]Cole, M.W. .................................... [179]Coleman, J.J................................... [178]Coleman, P.G................................... [73]

Collins, T.C.................................... [151]Cong, Y.......................................... [111]Corker, D.L.................................... [129]Cormier, L. .................................... [116]Cremades, A. ................................... [96]Croke, E.T........................................ [82]Cuomo, J.J. .................................... [183]

Da Silva, M.F. ............................... [186]Dabrowski, B. ................................ [124]Dafinei, I........................................ [128]Daidouh, A....................................... [85]Das, A............................................ [126]Davidson, J. ................................... [183]Davis, R.F. ..............................[168, 183]De Hosson, J.T.M. ....................[89, 125]De Léon-Guevara, A.M.................. [116]De Silva, P.S.I.P.N......................... [113]De Yoreo, J.J.................................. [186]Dec, J. ............................................ [129]Dediu, V. ......................................... [99]Dehm, G. ......................................... [83]Demenet, J.L. ................................. [157]den Otter, M.W. ............................... [95]Denanot, M.F. ................................ [105]DenBaars, P.M............................... [163]DenBaars, S.P. ........................[170, 182]Deng, P. ......................................... [146]Denk, I. .......................................... [140]Depero, L.E.................................... [146]DeSisto, W.J. ................................. [163]Devine, R.A.B................................ [133]Dewhurst, J.K. ................................. [69]Dewsnip, D.J.................................. [177]Dezaneti, M. .................................. [141]Dhoble, S.J..................................... [148]Di Felice, R.................................... [180]

Author Index

194

Díaz, A........................................... [104]Diaz, F. .......................................... [186]Diebold, U...................................... [109]Diéguez, E........................................ [96]Dieker, C........................................ [183]Diemoz, M. .................................... [128]Dimitrov, R. ................................... [164]Djuricic, B........................................ [89]Doak, R.B. ..................................... [177]Dobrzynski, L. ............................... [173]Dokiya, M. ..................................... [113]Domeneghetti, M.C........................ [146]Donlon, W.T. ................................... [92]Dontsov, G.I................................... [139]Doppalapudi, D. ............................. [170]Dorsch, W. ....................................... [77]Dovidenko, K......................... [152, 158]Doyle, B.P........................................ [69]Doyle, J.P......................................... [79]Drummond, T.J. ............................. [182]Duan, J.Q. ...................................... [166]Duijn-Arnold, A.V. .......................... [78]Dunham, S.T. ................................. [130]Dunn, B............................................ [97]Dupas, C. ....................................... [116]Duval, D.J. ....................................... [85]Dybzinski, R. ................................. [124]

Eastman, L.F. ............................... [170]Ebeling, K.J.................................... [165]Eckstein, R. ...................................... [77]Edvardsson, C.N.L. ........................ [142]Edwards, J.L. ................................. [177]Egi, T. ............................................ [100]Eglitis, R.I. ..................................... [112]Einfeldt, S. ..................................... [175]El-Raghy, T................................ [83, 84]

Elsass, C.R..................................... [163]Elsner, J. .................................[167, 180]Elstner, M. ..................................... [167]Engelhardt, G................................. [148]Erdei, S. ......................................... [148]Erickson, A.N. ............................... [182]Ernst, F. ......................................... [142]Esaka, T. .......................................... [97]Evetts, J.E. ..................................... [104]

Fan, Z............................................ [174]Fanciulli, M. .................................. [162]Fang, H. ......................................... [155]Fang, L....................................[110, 111]Fang, Z.Q................................[166, 174]Farber, L. ......................................... [84]Faurie, J.P. ..................................... [167]Fayeulle, S. .................................... [185]Fazzio, A........................................ [161]Feder, D.L...................................... [187]Feenstra, R.M................................. [181]Feigelson, R.S. ............................... [125]Feldman, B.J. ................................. [160]Feng, C.N....................................... [106]Feng, M.S. ..................................... [182]Feng, X. ......................................... [129]Fernández, P. ................................. [172]Fetisov, V.B. .................................. [108]Fini, P.T......................................... [170]Fischer, G.M. ................................. [105]Fisher, C.A.J. ................................... [93]Fisher, D. ......................................... [70]Fishman, A.J. ................................. [108]Flynn, C.P...................................... [120]Fong, C.Y. ..................................... [161]Foxon, C.T..................................... [177]Franke, M......................................... [80]

Author Index

195

Frauenheim, T. ....................... [167, 180]Freudenberg, F. .............................. [164]Friessnegg, T............................ [80, 114]Fritsch, J......................................... [159]Fu, C.L........................................... [104]Fu, J. .............................................. [119]Fuess, H. ........................................ [106]Fujieda, S. ...................................... [135]Fujimoto, M. .................................. [145]Fujimura, T. ..................................... [91]Fujishima, A..................................... [67]Fujita, K. ........................................ [115]Fujiwara, T..................................... [138]Fukui, K. ........................................ [145]Fung, S..................................... [80, 181]Furthmüller, J........................... [80, 157]Furukawa, Y......................[72, 118, 119]Futami, T........................................ [134]

Gan, F. .......................................... [146]Gan, Z.Z........................................ [176 ]García, J.A. ...................................... [96]Gardner, M.W. ............................... [101]Garratt-Reed, A.J. ............................ [87]Gaskell, P.H. .................................. [116]Gautier-Soyer, M. .......................... [147]Gemming, T. .................................... [88]Gennari, F.C................................... [145]Ghamnia, M. .................................... [91]Gibart, P......................................... [167]Gillespie, S....................................... [94]Gilliam, O.R................................... [153]Giocondi, J. ...................................... [77]Glazer, A.M. .................................. [129]Glinchuk, M.D. .............................. [112]Göbel, H......................................... [169]Goepfert, I.D. ................................. [175]

Gong, M........................................... [80]Goonewardene, A.D....................... [106]Gorman, R.J. .................................. [163]Goss, J.P. ......................................... [70]Gottstein, G...................................... [86]Govorkov, A.V. ............................. [175]Govorkov, S.A. .............................. [101]Graber, A. ...................................... [173]Grammatikakis, J. .......................... [187]Grattepain, C.................................... [86]Gratton, L.M. ................................. [131]Greer, A.J....................................... [106]Greve, D.W.............................[175, 181]Grimaldi, M.G. ................................ [73]Grimes, R.W. ................................. [152]Gritsenko, V.A............................... [184]Groen, H.B..................................... [125]Grzegory, I..................................... [173]Gu, S.............................................. [155]Gu, Z.Q............................................ [90]Gubanov, V.A. ............................... [161]Guedj, C........................................... [79]Guo, J.D......................................... [182]Gupta, S.K. .................................... [103]Gurevich, A.............................[188, 189]Guyot, M........................................ [126]

Haage, T........................................ [106]Habazaki, H. .................................. [144]Hacke, P......................................... [177]Hafich, M....................................... [182]Hagenbeck, R................................. [141]Hahn, T.S......................................... [95]Halls, D.C. ..................................... [150]Hamada, E. ...................................... [95]Hamada, M. ................................... [135]Han, B............................................ [129]

Author Index

196

Han, K............................................ [184]Han, T. ........................................... [143]Hansen, P.J..................................... [182]Hao, M................................... [171, 172]Hara, S. ............................................ [81]Hara, T........................................... [184]Harbsmeier, F................................. [133]Harding, J.H................................... [122]Härdtl, K.H. ........................... [139, 142]Hardy, W.N.................................... [101]Harris, D.J...................................... [121]Harrison, I. ..................................... [177]Harsch, W. ..................................... [151]Harshman, D.R............................... [106]Harvey, A....................................... [183]Hashimoto, K. .................................. [67]Hashimoto, T. ................................ [115]Haugk, M. ...................................... [180]Hautojärvi, P. ................................. [173]Hawagishi, K. ................................ [130]Hayakawa, N.................................. [138]Hayashi, H. .................................... [111]He, J............................................... [130]Heera, V........................................... [73]Heggie, M.I. ................................... [167]Heindl, J........................................... [77]Heinz, K........................................... [80]Helmersson, U................................ [142]Hemmingsson, C.............................. [72]Hemsky, J.W.......................... [165, 166]Henry, R.L. .................................... [163]Hesselink, L. .................................. [125]Heuer, A.H................................. [88, 92]Hirano, K. ...................................... [115]Hirata, A. ....................................... [185]Hobbs, L.W...................................... [87]Hobgood, H.M. ................................ [77]

Hoffmann, A. ................................. [158]Hoffmann, S..................................... [94]Hofmann, D. .................................... [77]Hofstaetter, A............................[75, 153]Hommel, D. ................................... [175]Hooper, S.E.................................... [177]Hopkins, R.H. .................................. [77]Höpler, R. ...................................... [169]Horita, T. ....................................... [113]Hoshino, K..................................... [139]Hoshiyama, S. ................................ [115]Hosono, H...................................... [133]Howells, W.S. .................................. [93]Howitt, D.G. .................................... [87]Hsu, L. ........................................... [156]Hu, G. ............................................ [129]Hu, Q.H. ........................................ [100]Hu, W. ............................................. [86]Huang, Z.J...................................... [141]Huebener, R.P. ............................... [102]Hulett, L.D....................................... [89]Hung, K.C...................................... [104]Hunt, D.C......................................... [70]Hunter, A.T...................................... [82]Hunter, B.A.................................... [128]Huntz, A.M. ..................................... [99]Husson, E....................................... [127]Hutchison, J.L. ............................... [115]Huvey, N.......................................... [76]

Iakoubovskij, K. ............................. [69]Ibarra, J. ......................................... [117]Ibbetson, J.P............................[170, 182]Ide, N. .............................................. [67]Ieranò, G. ................................[134, 135]Ihara, S............................................. [71]Iijima, M. ....................................... [122]

Author Index

197

Iijima, S. .......................................... [65]Ikoma, T........................................... [78]Ikonic, Z......................................... [179]Ikuhara, Y. ..................................... [155]Ikushima, A.J. ................................ [138]Impey, S.A. .................................... [129]Irene, E.A....................................... [102]Ishihara, K...................................... [184]Ishii, M........................................... [128]Ishizaki, K...................................... [144]Islam, M.S........................................ [93]Ismunandar .................................... [128]Isoya, J. ............................................ [69]Isshiki, T. ....................................... [151]Ita, J. .............................................. [122]Ito, K.............................................. [118]Itoh, H.................................. [73, 74, 75]Itoh, S............................................... [71]Ivanov, Z.G. ................................... [105]Iwasawa, Y..................................... [145]Iwaya, M. ....................................... [171]Iyoda, T............................................ [67]Izumi, F............................................ [72]

Jacobsen, S.N. ............................... [142]Jaenisch, V..................................... [188]Jäger, W. ........................................ [120]Jain, H............................................ [111]Jakobsson, A. ................................. [123]Jang, J.W.......................................... [95]Janzén, E.......................................... [72]Jardin, C........................................... [91]Jarolimek, O................................... [128]Jastrabík, L..................................... [112]Ji, Y. ...................................... [118, 119]Jia, C.L............................................. [94]Jiang, B. ........................................... [94]

Jiang, C. ......................................... [181]Jiang, D.S....................................... [173]Jiang, J.C. ........................................ [71]Jiang, L. ........................................... [67]Jiang, O.D. ..................................... [141]Jiménez, B. .................................... [126]Jin, P. ............................................. [141]Jin, S.X. ..................................[166, 176]Jogai, B. ......................................... [151]Johansson, L.G............................... [100]Johnson, N.M................................. [162]Jollet, F. ......................................... [146]Jones, R. ...................................[70, 167]Jones, R.E. ....................................... [94]Jones, R.L. ..................................... [165]Jorgensen, J.D................................ [124]Joshkin, V.A. ................................. [183]Jumas, J.C. ..................................... [186]Jun, J.............................................. [165]Jun, S.T.......................................... [152]Jungnickel, G. ................................ [180]Junnarkar, M.R. ............................. [179]

Käckell, P........................................ [80]Kagi, H. ........................................... [69]Kakazey, N.G................................. [151]Kakihana, M. ................................. [147]Kalceff, M.A.S............................... [132]Kallin, C. ....................................... [187]Kalnacs, J......................................... [67]Kalogeras, I.M. .................................. [1]Kambayashi, S. .............................. [131]Kamiya, E. ..................................... [131]Kamiya, N........................................ [76]Kamp, M........................................ [165]Kanai, H......................................... [115]Kanda, H.......................................... [69]

Author Index

198

Kaneko, K. ....................................... [88]Kang, J. .......................................... [168]Kappers, L.A.................................. [153]Karduck, P. ...................................... [86]Karna, S. ........................................ [133]Kashino, S........................................ [72]Kashiwagi, M................................. [131]Kaufmann, U.................................. [174]Kawachi, T..................................... [138]Kawada, T...................................... [113]Kawamoto, K. ................................ [118]Kawasuso, A. ................................... [73]Kawazoe, H.................................... [133]Kebbede, A. ..................................... [89]Keeble, D.J............................. [114, 127]Keinonen, J. ..................................... [68]Keller, N. ....................................... [126]Keller, S. ................................ [163, 170]Kennedy, B.J.................................. [128]Keskitalo, N. .................................... [79]Keyse, R......................................... [183]Kharton, V.V............................ [93, 114]Kienzle, O. ..................................... [142]Killeen, K.P. .................................. [167]Kilo, M. ......................................... [154]Kim, S............................................ [178]Kim, W. ................................. [173, 174]Kim, Y. ............................................ [67]Kim, Y.K. ...................................... [107]Kinoshita, C. ............................ [86, 121]Kiselev, A.N. ................................. [150]Kiselev, N.A................................... [150]Kitamura, K. .......................... [118, 119]Kitano, M....................................... [151]Kittelberger, S. ............................... [102]Kizuka, T. .............................. [121, 122]Klamut, P.W. ................................. [124]

Klein, B.M. .................................... [161]Klein, N. ........................................ [113]Klein, S.P....................................... [102]Kneissl, M...................................... [162]Knights, A........................................ [80]Kobayashi, M................................. [128]Kobayashi, N. .................................. [71]Kohno, N. ........................................ [97]Kojima, K. ....................................... [67]Koleske, D.D. ................................ [163]Kong, G.L. ..................................... [130]Konstantinov, A.O. .......................... [76]Kooi, B.J........................................ [125]Kordina, O. ...................................... [72]Kossler, W.J................................... [106]Koster, E. ....................................... [106]Kotomin, E.A................................. [112]Kovács, L....................................... [148]Koyama, S. .................................... [100]Kozhina, G.A. ................................ [108]Kozodoy, P. ............................[170, 182]Krämer, A. ..................................... [186]Krasnobaev, L.Y. ........................... [183]Krauss, A.R.................................... [102]Krishna, N.M. ................................ [106]Krishnan, A.............................[114, 127]Krishnan, R.................................... [126]Kristoffel, N................................... [100]Krtschil, A. .................................... [175]Kruidhof, H.................................... [138]Kubozono, Y.................................... [72]Kudo, J........................................... [184]Kuhn, M......................................... [109]Kuisma, S....................................... [173]Kunzer, M...................................... [174]Kurai, S.......................................... [171]Kurennykh, T.E.............................. [108]

Author Index

199

Kutami, H. ..................................... [100]Kwok, R.W.M................................ [184]Kyuragi, H. .................................... [185]

Labidi, M. ..................................... [108]Lacklison, D.E. .............................. [177]Lagerlöf, K.P.D................................ [88]Laguta, V.V............................ [112, 128]Laine, T.......................................... [173]Lam, C.C........................................ [104]Land, T.A....................................... [186]Lange, F.F...................................... [143]Langjahr, P.A. ................................ [143]Lankhorst, M.H.R. ......................... [114]Lau, S.S.......................................... [173]Lau, W.S. ....................................... [143]Laursen, T. ....................................... [82]Lawson, S.C..................................... [70]Le Dang, K..................................... [116]Leach, C......................................... [150]Lee, A. ........................................... [143]Lee, B.J. ........................................... [91]Lee, C.R. ........................................ [176]Lee, H.J............................................ [93]Lee, I.H. ......................................... [176]Lee, J.D.......................................... [186]Lee, J.H.................................... [95, 147]Lee, M............................................ [125]Lee, S.R. ........................................ [182]Lee, W.E. ....................................... [152]Lemaignan, C................................. [146]Lengauer, W..................................... [82]León, C. ......................................... [117]Leonidov, I.A. ................................ [139]Leonidova, O.N.............................. [139]Lesage, B. .................................. [86, 99]Leszczynski, M. ..................... [165, 173]

Levin, I. ..................................... [84, 90]Li, G. ............................................. [176]Li, J...........................................[99, 176]Li, J.F............................................. [129]Li, S. .......................................[110, 111]Li, X. ............................................. [178]Liang, K......................................... [155]Liang, R. ........................................ [101]Liao, X.B. ...................................... [130]Likonen, J. ....................................... [68]Liliental-Weber, Z...................[173, 182]Lin, L. .....................................[110, 111]Lin, S.H. ...................................[71, 160]Lind, D........................................... [109]Lindström, J.L............................ [72, 79]Linnarsson, M.K. ............................. [79]Lippmaa, E..................................... [148]Lisker, M. ...................................... [175]Litton, C.W. ................................... [151]Liu, A............................................. [125]Liu, F.X. ........................................ [137]Liu, J.Z. ......................................... [106]Liu, L. ............................................ [137]Liu, Y............................................. [144]Lo, W............................................. [103]Loa, I. ............................................ [158]Loh, F.C........................................... [78]Lombardi, F. .................................. [105]Look, D.C. .............. [151, 165, 166, 174]Loudjani, M.K. .......................... [86, 99]Lovett, D.R. ................................... [106]Lowther, J.E..............................[69, 162]Loyalka, S.K. ................................... [67]Lu, X.......................................[111, 155]Lucovsky, G................................... [136]Ludowise, M.J. .............................. [167]Lynn, K.G. ..................................... [127]

Author Index

200

Ma, S.Y. ........................................ [132]Ma, Z.C.......................................... [132]Ma, Z.X. ........................................ [130]Machida, K. ................................... [185]Mack, M......................................... [163]Mack, M.P. .................................... [166]Mackenzie, J.D................................. [97]MacManus-Driscoll, J.L. ................ [113]Madhukar, S........................... [114, 127]Madsen, L.D. ................................. [142]Maeda, H. ........................................ [72]Maeda, M....................................... [138]Maeyama, S. .................................. [185]Maier, J. ......................................... [140]Maier, M. ....................................... [174]Majid, I. ......................................... [144]Makita, K. ...................................... [126]Mallamaci, M.P................................ [97]Manabe, A...................................... [121]Manika, I.......................................... [67]Maniks, J.......................................... [67]Marchand, H. ................................. [170]Margulies, D.T. .............................. [110]Márquez, A.M................................ [134]Martin, M.J. ................................... [186]Martínez-Lope, M.J........................ [113]Martini, M...................................... [128]Maruyama, T.................................. [130]Mascher, P. ...................................... [80]Mason, T.O. ................................... [149]Matacotta, F.C.................................. [99]Matsunami, N................................. [133]Mattsson, M.S. ................................. [57]Maxim, P........................................ [158]Mayer, J. .......................................... [83]Mayer, J.W....................................... [82]Mazzi, F. ........................................ [146]

McClellan, K.J. .............................. [124]McCluskey, M.D............................ [162]McCoy, M.A. ................................. [152]McGarry, D...................................... [89]McGill, T.C.................................... [179]Mechin, L....................................... [104]Mei, L. ............................................. [66]Meinardi, F. ................................... [132]Menesklou, W................................ [139]Meng, J. ......................................... [111]Meng, X.T...................................... [132]Meng, Z. ........................................ [127]Mertin, M......................................... [94]Messerschmidt, U. ......................... [154]Messing, G.L. .................................. [89]Metzger, T. .................................... [169]Meyer, B.K. ..................................... [75]Meyer, M. ...................................... [188]Michaelis, A................................... [102]Mikado, T. ....................................... [74]Mikata, Y. ...................................... [131]Millot, F......................................... [116]Ming, H.......................................... [137]Miotello, A..................................... [131]Mishra, P.K.................................... [103]Mishra, U....................................... [182]Mishra, U.K. .................................. [170]Mitchell, J.N. ................................. [124]Mitchell, T.E.................................... [92]Miura, S. ........................................ [108]Miyamoto, T. ................................. [111]Mizuno, M. .................................... [111]Mizusaki, J..................................... [115]Mizushima, I. ................................. [131]Mochiji, K........................................ [71]Mokhov, E.N. .................................. [78]Molnar, R.J. ................................... [165]

Author Index

201

Monovskii, O.V. ............................ [186]Montecchi, M................................. [128]Monty, C.J.A.................................. [108]Moos, R. ................................ [139, 142]Morell, A........................................ [127]Mori, S. .......................................... [131]Morisaki, R. ..................................... [86]Morishima, Y. ................................ [172]Morita, Y. ........................................ [69]Morkoç, H.............................. [173, 174]Morris, R.C. ..................................... [92]Moss, S.C....................................... [169]Mota, R. ......................................... [161]Moustakas, T.D.............................. [170]Moya, E.G........................................ [86]Moya, F............................................ [86]Mu, Y............................................... [66]Müllejans, H..................................... [88]Müller, S.G. ..................................... [77]Murakami, Y.................................... [72]Muto, S. ........................................... [71]

Nagao, R........................................ [126]Nagata, K. ...................................... [130]Nahm, S. .......................................... [93]Nakagawa, Z. ................................. [100]Nakamura, S................................... [172]Nakamura, T. ................................. [108]Nam, O.H....................................... [168]Nanko, M. ...................................... [144]Naoi, Y. ................................. [171, 172]Narayan, J. ............................. [152, 158]Narayanamurti, V........................... [182]Nardelli, M.B. ................................ [180]Nashiyama, I. ............................. [74, 75]Nastasi, M. ..................................... [185]Nastasi, M.A. ................................. [124]

Naumovich, E.N. ......................[93, 114]Navi, M.......................................... [130]Nellist, P.D. ............................[144, 167]Nesládek, M..................................... [69]Neugebauer, J. ........ [156, 164, 180, 181]Newton, M.E.................................... [70]Ng, H.M......................................... [170]Ng, Y.H. ........................................ [184]Ngai, K.L. ...................................... [154]Nguyen, N.Q.................................... [95]Nicoloso, N. ................................... [188]Nielsen, B. ..............................[114, 127]Nieminen, R.M. ............................... [75]Niimi, H......................................... [136]Nikl, M. ......................................... [128]Ning, X.J.......................................... [76]Nishi, Y. ........................................ [145]Nishino, K...................................... [171]Nishio, K........................................ [151]Nissilä, J. ....................................... [173]Nitsch, K........................................ [128]Nobusawa, H.................................. [135]Nogami, M..............................[126, 138]Noguera, C..................................... [122]Noh, S.K. ....................................... [176]Noll, F............................................ [140]Nolte, T.......................................... [118]Nordell, N. ....................................... [79]Nörenberg, H. .................................. [98]Northrup, J.E...................[179, 180, 181]Nouet, G. ....................................... [160]Nouwen, B. .................................... [136]Nowotny, J................................[98, 116]

Oberg, S. ..................................[70, 167]Obloh, H. ....................................... [174]Obradors, X. .................................. [105]

Author Index

202

Ogawa, H. ...................................... [155]Ogawa, T........................................ [168]Ohdaira, T. ................................. [74, 89]Ohki, Y. ......................................... [134]Ohnishi, T. ..................................... [111]Ohshima, T. ......................... [73, 74, 75]Ohta, T............................................. [72]Ohtani, S. ......................................... [71]Oka, A............................................ [100]Okada, I............................................ [67]Okada, S........................................... [73]Okamoto, A...................................... [76]Okano, Y........................................ [108]Okazaki, H. .................................... [139]Oktyabrsky, S......................... [152, 158]Oku, T............................................ [131]Okumura, H. .............................. [74, 75]Okumura, K. .................................... [76]Okushi, H....................................... [177]Olive, E.......................................... [188]Olivier-Fourcade, J......................... [186]Olsson, E................. [100, 105, 142, 150]Omnès, F........................................ [167]Ong, T.Y. ......................................... [78]Onischenko, T.A. ........................... [186]Onishi, H........................................ [145]Onishi, S. ....................................... [184]Orton, J.W...................................... [177]Oshiyama, A. ................................. [135]Otsubo, M. ..................................... [131]Overhof, H. ...................................... [75]

Pacaud, Y. ....................................... [73]Pacchioni, G........................... [134, 135]Page, J.B. ....................................... [159]Paik, J.H........................................... [93]Pakula, K........................................ [173]

Paleari, A. ...................................... [132]Paloura, E.C................................... [184]Papathanassiou, A.N. ..................... [187]Parikh, A.S..................................... [105]París, M.A...................................... [117]Park, B.M................................[118, 119]Park, C.H. ...................................... [128]Park, H.M. ....................................... [93]Parker, C.A. ................................... [183]Parker, F.T. .................................... [110]Parker, S.C..................................... [121]Pashitskii, E.A.........................[188, 189]Pasquevich, D.M............................ [145]Pasturel, A. .................................... [146]Pawlik, T........................................ [118]Pearton, S.J. ....................[166, 167, 179]Pellegrino, P..................................... [79]Pelzmann, A................................... [165]Pennycook, S.J. .......................[144, 167]Pensl, G. .......................................... [77]Pentaleri, E.A................................. [161]Petersen, M. ................................... [109]Petit, T. .......................................... [146]Peto, A. .......................................... [148]Petroff, P.M. .................................. [163]Petukhov, B.V................................ [154]Phillips, B.L. .................................... [85]Pickering, S...................................... [89]Pico, C. ............................................ [85]Pint, B.A. ................................... [87, 89]Piqueras, J.................................[96, 172]Piquini, P. ...................................... [161]Piriou, B........................................... [79]Pirouz, P. ..............................[76, 78, 81]Plazaola, F. ...................................... [73]Ploog, K.H. .................................... [173]Pode, R.B....................................... [148]

Author Index

203

Poindexter, E.H.............................. [114]Poluektov, O.G................................. [78]Polyakov, A.Y................................ [175]Ponce, F.A. .................................... [172]Pond, R.C....................................... [150]Popovici, G. ..................................... [67]Porezag, V.D.................................. [167]Porowski, S. ................................... [173]Porte, M. ........................................ [126]Postnikov, A.V............................... [112]Pöykkö, S. ........................................ [75]Prelas, M.A. ..................................... [67]Pressel, K. ...................................... [158]Puff, W. ........................................... [80]Pujar, V.V. ....................................... [81]Püsche, W. ....................................... [83]

Qian, J.Y. ...................................... [137]Qin, G.G. ............................... [132, 166]

Rabe, K.M....................................... [85]Rabier, J. ................................ [105, 157]Rafaja, D. ......................................... [82]Raghavachari, K............................. [137]Ramakrishnan, A............................ [174]Ramesh, R.............................. [114, 127]Ramsteiner, M................................ [173]Rapcewicz, K. ................................ [180]Rashkovich, L.N. ........................... [186]Ravikumar, V................................... [65]Redlich, P......................................... [66]Redwing, J.M................................. [175]Reiche, P. ....................................... [128]Rekas, M. ....................................... [116]Remón, A......................................... [96]Ren, F. ........................................... [179]Ren, G.B. ....................................... [177]

Ren, Y............................................ [111]Renard, J.P..................................... [116]Renaud, G. ..................................... [160]Reut, O.P. ...................................... [114]Reuter, K.......................................... [80]Revcolevschi, A. ............................ [116]Reynolds, D.C.........................[151, 165]Rhee, S.J. ....................................... [178]Richards, F.M. ............................... [113]Rickerby, D.G. ............................... [144]Ricote, J. ........................................ [129]Riechert, H..................................... [173]Risbud, S.H...................................... [85]Ristic, M.M.................................... [151]Ritter, M. ....................................... [109]Robertson, J. .................................... [66]Rockett, A. ..................................... [165]Rogacki, K. .................................... [124]Rohrer, G.S. ..................................... [77]Roleder, K...................................... [129]Romano, L.T...........................[171, 172]Rosa, J. .............................[69, 112, 128]Rosetta, E....................................... [128]Rosner, S.J. .................................... [163]Routbort, J.L. ................................. [149]Route, R.K. .................................... [125]Rouvière, J.L.................................. [160]Rowley, A.T................................... [103]Rubin, P. ........................................ [100]Rudee, M.L. ................................... [110]Rudner, S. ...................................... [142]Rühle, M. ....................................... [143]Russell, J.D. ................................... [150]Ruterana, P..................................... [160]Ruvimov, S. ............................[173, 182]Ryen, L. ......................................... [142]Ryu, H. ............................................ [93]

Author Index

204

Saarinen, K. .................................. [173]Sahni, V.C...................................... [103]Sakai, A. .........................[168, 169, 170]Sakai, N. ........................................ [113]Sakai, S. ................................. [171, 172]Sakiyama, K................................... [184]Sakuma, T. ..................................... [155]Salama, K....................................... [105]Salari, R. ........................................ [146]Salje, E.K.H. .................................. [146]Samant, A.V..................................... [81]Samokhval, V.V............................... [93]Sandiumenge, F.............................. [105]Sandler, N.P. .................................. [143]Sangaletti, L. .................................. [146]Sankey, O.F.................................... [159]Santamaria, J. ................................. [117]Santic, B......................................... [174]Santos, M.T...................................... [96]Sanz, J............................................ [117]Sarrazit, F....................................... [150]Sartain, K.B...................................... [97]Sathyamurthy, S. ............................ [105]Sato, H. .................................. [171, 172]Sato, K. .......................................... [131]Sauer, C. .......................................... [83]Saunders, M. .................................. [172]Sberveglieri, G. .............................. [146]Schäfer, J........................................ [136]Schardt, J.......................................... [80]Scharmann, A........................... [75, 153]Scheffler, M. ...................[109, 164, 181]Scherrer, S................................ [98, 154]Schlögl, R. ..................................... [109]Schlotter, P..................................... [174]Schmidt, J. ....................................... [78]Schmidt, K.E.................................. [159]

Schmidt, T.M. ................................ [161]Schöner, A. ...................................... [79]Schubert, E.F. ................................ [175]Schulze, N........................................ [77]Schuster, M.................................... [169]See, C.H......................................... [143]Seethararaman, S............................ [123]Seguchi, T........................................ [71]Seo, J.W........................................... [69]Seol, K.S. ....................................... [134]Shaffer, J........................................ [124]Shaikhutdinov, S.K. ....................... [109]Shapiro, S.M. ................................. [141]Sharma, A.K. ................................. [152]Shaw, K. ........................................ [109]Shelton, R.N................................... [106]Shen, G.J........................................ [104]Sheu, B.C....................................... [134]Shibata, N. ..................................... [155]Shimizu, H. ...................................... [71]Shimizu, K. .................................... [144]Shimojo, F. .................................... [139]Shin, M. ......................................... [175]Shiohara, Y. ................................... [100]Shiojiri, M...................................... [151]Shrivastava, K.N. ........................... [106]Shul, R.J......................................... [167]Sichen, D. ...................................... [123]Sickafus, K.E. ................................ [124]Sigle, W. .......................................... [66]Simner, S.P. ..................................... [97]Simpson, P.J. ................................... [80]Sitch, P.K....................................... [167]Sivananthan, S. .............................. [167]Sizelove, J.R. ................................. [165]Skeldon, P...................................... [144]Skorupa, W. ............................... [73, 80]

Author Index

205

Skowronski, M......................... [77, 175]Sleight, A.W. ................................. [102]Smirnov, N.B. ................................ [175]Smith, A.R. .................................... [181]Smith, D.J. ..................................... [177]Soares, J.C. .................................... [186]Somieski, B...................................... [89]Son, N.T........................................... [72]Song, C.Y....................................... [167]Song, Q. ......................................... [185]Sonoda, M...................................... [131]Sonoda, T....................................... [121]Soper, A.K. .................................... [116]Spada, F.E...................................... [110]Spaeth, J.M. ................................... [118]Speck, J.S........................[163, 170, 182]Specka, J.S. .................................... [163]Spinolo, G. ..................................... [128]Srdanov, V.I. .................................. [148]Sreckovic, T.V. .............................. [151]Srinitiwarawong, C. ....................... [113]Stall, R.A. ...................................... [166]Stals, L.M......................................... [69]Stampfl, C. ............................. [156, 159]Starke, U. ......................................... [80]Stavola, M...................................... [167]Steeds, J.W..................................... [172]Stefanov, B.B. ................................ [137]Stein, B.L. ........................................ [82]Stepantsov, E.A.............................. [150]Stepniewski, R. .............................. [173]Stesmans, A. ............. [69, 134, 136, 137]Stevens, M. .................................... [177]Stiller, K......................................... [100]Stoll, O.M. ..................................... [102]Stömmer, R. ................................... [169]Störmer, J. ........................................ [73]

Stoyanov, P. ................................... [109]Strausser, Y.E. ............................... [182]Strite, S. ......................................... [165]Strunk, H.P. ..............................[77, 169]Stucky, G.D. .................................. [148]Stutzmann, M..........................[164, 169]Suarez-Sandoval, D.......................... [97]Sugahara, T.............................[171, 172]Sugiyama, N. ................................... [76]Sun, N............................................ [181]Sun, T. ........................................... [181]Sun, W. .......................................... [181]Sunakawa, H. ..................[168, 169, 170]Sung, T. ........................................... [67]Susa, M. ......................................... [130]Suski, T...................................[165, 173]Suvkhanov, A................................. [183]Suzuki, E........................................ [119]Suzuki, R. .................................. [74, 89]Suzuki, T........................................ [145]Svensson, B.G.................................. [79]Swenson, J. ...................................... [93]

Tagawa, H..................................... [115]Takagi, S........................................ [131]Takahashi, M. ................................ [138]Takahiro, K. ................................... [144]Takai, S............................................ [97]Takayanagi, K.................................. [95]Takeuchi, T. ................................... [171]Tamaki, S......................................... [67]Tamura, R. ....................................... [71]Tan, C. ............................................. [66]Tan, H.H. ....................................... [166]Tan, K.L........................................... [78]Tan, L.S. .......................................... [78]Tan, Q............................................ [129]

Author Index

206

Tanabe, T. ........................................ [71]Tanaka, I. ......................................... [88]Tanaka, J. ....................................... [145]Tanaka, M. ............................. [158, 184]Tanaka, N............................... [121, 122]Tanaka, S.I. ...................................... [91]Tani, T. ............................................ [76]Tanigawa, S. .............................. [74, 75]Tanigawa, T. .................................. [135]Tarsa, E.J. ...................................... [182]Taylor, D.......................................... [94]Tazzoli, V. ..................................... [146]Teisseyre, H. .................................. [165]ten Elshof, J.E. ......................... [95, 114]Terabe, K. ...................................... [119]Terada, S. ......................................... [81]Terauchi, M.................................... [158]Terjak, M.J.E. .................................. [85]Tews, H.......................................... [173]Tezuka, M. ..................................... [115]Thavorniti, P. ................................. [155]Thomas, J. ........................................ [83]Thomas, K.A.................................. [113]Thompson, G.E. ............................. [144]Thomsen, C.................................... [158]Thurian, P. ..................................... [158]Tinschert, K. .................................. [154]Tirado, J.L...................................... [186]Tomlins, G.W. ............................... [149]Tong, Y.Z............................... [166, 176]Tonoyan, A.A................................. [114]Torpo, L. .......................................... [75]Torres, L.M. ................................... [117]Torres, V.M. .................................. [177]Tortorelli, P.F................................... [89]Tosello, C....................................... [131]Tottori, S. ....................................... [171]

Townsend, P.D............................... [148]Tsai, J.H......................................... [134]Tsai, M.H......................................... [71]Tsang, J.S....................................... [182]Tsong, I.S.T. .................................. [177]Tsukada, M. ..................................... [71]

Uedono, A. ................................ [74, 75]Uhrberg, R.I.G. ................................ [81]Umehara, H.................................... [138]Uno, T............................................ [135]Urakawa, T. ..................................... [72]Urban, K. ..................................[94, 113]Usui, A............................[168, 169, 170]Usuki, Y......................................... [128]

Vainonen, E. ................................... [68]Van de Walle, C.G. ................. [156, 159,

.................................... 162, 178, 181]Van Doorn, R.H.E.......................... [138]Van Sambeek, A.I. ......................... [120]Van Tendeloo, G.............................. [69]Vander Sande, J.B.......................... [144]Vandlik, J....................................... [148]Vanecek, M...................................... [69]Vassilikou-Dova, A............................ [1]Vázquez-Navarro, D. ..................... [103]Vedda, A........................................ [128]Veiga, M.L....................................... [85]Veillet, P. ....................................... [116]Vermaut, P. .................................... [160]Vicente, C.P. .................................. [186]Vicente, J.M................................... [126]Viehland, D.................................... [129]Vilalta, N. ...................................... [105]Viskup, A.P...............................[93, 114]Vykhodets, V.B.............................. [108]

Author Index

207

Wagner, F. .................................... [109]Wagner, T. ..................................... [143]Wakamiya, M................................. [131]Wallis, D.J. .................................... [144]Walukiewicz, W..................... [156, 162]Wang, C. ........................................ [185]Wang, C.D. .................................... [173]Wang, L.P. ..................................... [166]Wang, P. ................................ [129, 166]Wang, R. ........................................ [141]Wang, R.P...................................... [102]Wang, W. ....................................... [176]Wang, X......................................... [142]Wang, X.C. .................................... [176]Wang, X.G. .................................... [109]Wang, X.L...................................... [137]Wang, Y.G. ...................................... [89]Wang, Y.Q. .................................... [130]Wang, Z.L. ..................................... [115]Warren, W.L. ................................. [133]Warren, W.W................................. [102]Waser, R. ................................. [94, 141]Washburn, J. .......................... [173, 182]Watanabe, Y................................... [185]Watson, G.W.................................. [121]Watterich, A................................... [153]Weatherly, G.C. ............................. [184]Weber, S. ................................. [98, 154]Wedler, H......................................... [80]Wei, X.L. ......................................... [81]Weimann, N.G. .............................. [170]Weinstein, M.G.............................. [167]Weiss, W........................................ [109]Weissgärber, T. ................................ [83]Wen, J.G. ....................................... [100]Wen, S. .......................................... [185]Wernlund, L.D. .............................. [142]

Wetzel, C. ...................................... [171]Whatmore, R.W. ............................ [129]White, B........................................... [94]Wickenden, A.E. ............................ [163]Wiesenberger, H. ............................. [82]Williams, D.L. ............................... [106]Williams, J.S. ................................. [166]Wilson, I.H. ................................... [184]Wilson, R.G. .............................[67, 167]Wimbauer, T. ................................... [75]Winnacker, A. .................................. [77]Winter, M. ..................................... [113]Wirth, H........................................... [80]Witte, H. ........................................ [175]Wojdak, M. .................................... [173]Wollschläger, J................................. [37]Wolverton, C.................................... [98]Wood, G.C..................................... [144]Wright, A.F.............................[157, 158]Wroe, R. ........................................ [103]Wu, X.H. ....................................... [163]Wuerz, R........................................ [153]Wysmolek, A. ................................ [173]

Xia, Y. ............................................. [66]Xin, Y. ........................................... [167]Xing, Y. ........................................... [66]Xiong, X. ................................[124, 169]Xu, H. ............................................ [108]Xu, J. ........................................[89, 127]Xu, J.B........................................... [184]Xu, S.J. .......................................... [176]Xu, X. ............................................ [181]Xu, Y.N. .......................................... [90]Xu, Z.............................................. [129]Xu, Z.H.......................................... [166]

Author Index

208

Yamada, K. ..................................... [72]Yamada, Y. .................................... [100]Yamaguchi, A.A............................. [170]Yamaguchi, E................................. [179]Yamaguchi, S......................... [144, 171]Yamamoto, S.................................... [71]Yamashita, K. ................................ [171]Yan, Y............................................ [158]Yang, M.H. .................................... [120]Yang, W................................. [110, 111]Yang, Z.J................................ [166, 176]Yaremchenko, A.A........................... [93]Yashima, M.................................... [147]Yasuda, A. ..................................... [155]Yasuda, K. ....................................... [86]Yeom, H.W. ..................................... [81]Yin, H. ........................................... [146]Yin, J.S. ......................................... [115]Yin, Z............................................. [129]Yokokawa, H. ................................ [113]Yoshida, S............................ [74, 75, 81]Yoshida, Y. ...................................... [72]Yoshikawa, M. ........................... [74, 75]Yoshiki, M. .................................... [131]Yoshimura, M. ............................... [147]Yoshiya, M. ..................................... [88]Young, A.P. ................................... [136]Young, C.F..................................... [127]Young, W.T. .................................. [172]Yu, E.T. ................................... [82, 173]Yu, L.S........................................... [173]Yu, N. ............................................ [124]Yue, G.Z. ....................................... [130]

Zafar, S. .......................................... [94]Zaitseva, N.P.................................. [186]

Zaldivar, M.H. ............................... [172]Zaldo, C. ........................................ [186]Zaritskii, M.I. ..........................[112, 128]Zavaliangos, A. ................................ [84]Zegenhagen, J. ............................... [106]Zhang, B.R..............................[132, 166]Zhang, G.Y. ............................[166, 176]Zhang, X. ....................................... [166]Zhang, Y. ......................................... [65]Zhang, Y.X. ................................... [166]Zhao, J. .......................................... [116]Zhao, Y.......................................... [181]Zheleva, T.S................................... [168]Zheng, Y. ....................................... [155]Zhong, L. ....................................... [143]Zhong, X.F....................................... [90]Zhu, X............................................ [127]Zhu, Y............................................ [141]Zhukovskaya, A.C.......................... [139]Ziese, M......................................... [113]Zolper, J.C. .................................... [166]Zong, W.H. .................................... [132]Zou, J. ............................................ [166]Zuccaro, C. .................................... [113]Zunger, A......................................... [98]Zywietz, T...............................[164, 181]

211

Keyword Index

2-dimensional ........... 106, 144, 146, 1563-dimensional ..... 80, 132, 144, 146, 148adatom........................................80, 164

Ag2O-B2O3-SiO2-AgIIonic Conduction ................................ 85

Ag2VP2O8Ionic Conduction ................................ 85

AlGaNPoint Defects .............................155, 156

AlGaN/GaNPoint Defects .................................... 156

alignment............. 77, 97, 122, 136, 146,..................................................158, 167

Al2MgO4Dislocations........................................ 85Electron Irradiation............................. 85Ion Bombardment............................... 85

AlNDislocations...............................156, 157Domain Boundaries .......................... 157Point Defects .............................157, 158Stacking Faults ..................157, 158, 159Surface Defects................................. 159

AlN/SiDislocations...................................... 159

AlN/SiCStacking Faults ................................. 160

AlN/ZnO/Al2O3

Dislocations.......................................151Stacking Faults ..................................151

Al2O3Dislocations.........................................87Grain Boundaries.................................88Grain Boundary Diffusion ...................86Ni Diffusion ........................................86Point Defects .................................89, 87Y Diffusion .........................................86

Al2O3/CuO Interdiffusion ...................................91

Al2O3/TiInterdiffusion.......................................91

Al2O3-MgODislocations.........................................91Point Defects .......................................91

AlPO4Grain Boundaries.................................92

ambipolar ...........................................95amorphization..........................123, 132amorphous .65, 66, 78, 80, 88, 123, 130,.......... 131, 132, 133, 143, 146, 154, 160,.................................. 162, 166, 168, 184anions .........90, 114, 120, 127, 128, 144,.......................................... 148, 159, 161anisotropy ......84, 88, 90, 109, 121, 157,.......................................... 186, 188, 189annealing ........69, 72, 73, 74, 78, 79, 84,................95, 98, 99, 102, 109, 112, 113,

Keyword Index

212

..........128, 130, 132, 133, 134, 135, 136,

..................141, 143, 148, 151, 155, 160,

.......................... 164, 166, 179, 184, 185annihilation ....................73, 74, 80, 173antiphase .................. 109, 110, 129, 155antisites....................... 75, 156, 161, 176arrays ........................... 84, 92, 169, 188

BNH Diffusion....................................... 160Point Defects .............................160, 161

B2O3-Li2O-NaBrIonic Conduction ................................ 92

B2O3-Li2O-NaClIonic Conduction ................................ 92

B2O3-Na2O-NaBrIonic Conduction ................................ 92

B2O3-Na2O-NaClIonic Conduction ................................ 92

Ba(Bi,La)O3O Permeation...................................... 93

back-scattering........... 73, 123, 132, 186

Ba2In2O5Point Defects ...................................... 93

(Ba,La)(Mg,Nb)O3Domain Boundaries ............................ 93

Ba0.5Sr0.5TiO3Point Defects ...................................... 94

BaTiO3Dislocations........................................ 94Twins ............................................94, 95

BaZrO3/SrTiO3Dislocations...................................... 142

bicrystals ......................................... 140

(Bi,Er)2O3

O Permeation.......................................95

Bi12GeO20Point Defects .......................................95

bilayer ......................................159, 181

Bi2O3-Gd2O3Ionic Conduction .................................96

Bi2O3-Y2O3Ionic Conduction .................................96

Bi12SiO20Point Defects .......................................95

Bi2(V,Cu)O5.35

Ionic Conduction .................................97

bombardment ......... 119, 123, 132, 133,.......................................... 183, 184, 185boundaries......68, 84, 86, 87, 88, 90, 91,.. 92, 93, 97, 99, 100, 101, 104, 105, 107,.. 108, 109, 110, 113, 118, 119, 120, 121,.. 122, 129, 140, 141, 142, 144, 146, 149,.. 150, 152, 154, 155, 157, 158, 171, 172,........... 179, 183, 184, 185, 187, 188, 189

CElectron Irradiation..............................65Point Defects .......................................65Surface Reconstruction........................65

C60Dislocations...................................66, 67Electron Irradiation..............................66

C (Diamond)B Diffusion..........................................67D Diffusion .........................................67Dislocations.........................................68Grain Boundaries.................................68Point Defects .................................69, 70Surface Reconstruction........................70

Keyword Index

213

C (Graphite)Dislocations........................................ 71Electron Irradiation............................. 71Ion Bombardment............................... 71Point Defects ...................................... 71

C3N4Point Defects .................................... 162

Calcium Hexaluminate/Al2O3Dislocations........................................ 97

capacitance .......................135, 177, 181

Ca5(PO4)3OHGrain Boundaries.............................. 185

carrier................ 72, 115, 170, 171, 174,..........................................176, 177, 189cascade..................................73, 85, 180cathodoluminescence 95, 122, 123, 132,...................136, 163, 167, 171, 172, 174cations...................... 84, 86, 87, 92, 114,.........................................120, 127, 129,...................142, 144, 145, 148, 154, 185cationic ............................................ 127ceramic catalysts ................................. 1channelling ................. 73, 123, 132, 177cleavage ........................................... 107climb .................................................. 92

CeO2Surface Defects................................... 97

coalescence ........... 77, 94, 104, 168, 170coincidence-site ............................... 122

CoLixO2Point Defects ...................................... 98

columnar ......................................... 169convergent-beam........................77, 160

CoOCr Diffusion ....................................... 98

Cr2O3Y Diffusion .........................................98

Cu3Ba2GdO6O Diffusion .........................................99Point Defects .......................................99

Cu3Ba2NdO7Twins ..................................................99

Cu3Ba2YO6Point Defects .....................................100Twins ................................................100

Cu3Ba2YO6.5

Twins ................................................100

Cu3Ba2YO7Dislocations.......................................104Electron Irradiation............................104Grain Boundaries.......................104, 105O Diffusion ....................... 101, 102, 103Planar Defects ...................................105Point Defects .............................104, 106Stacking Faults ..................................104Twins ................................................106

Cu3Ba2YO7/SrTiO3Twins ................................................106

Cu3Ba2YOxTwins ................................................100

Cu2Y2O5Stacking Faults ..................................106

Czochralski ..............................118, 148dangling-bond ...... 65, 66, 136, 159, 166deep-level ..72, 75, 79, 81, 173, 174, 175defect-free ..........................................74deformation-induced .........................81deposition .........72, 76, 80, 94, 106, 113,.......... 114, 126, 129, 133, 136, 137, 142,.......... 158, 159, 169, 171, 172, 173, 174,.................................. 178, 179, 180, 181desorption ...........80, 121, 131, 163, 164dielectric relaxation ............................ 1

Keyword Index

214

diffusion............................................. 57dipole ........................ 112, 129, 157, 187dislocations.....67, 68, 71, 76, 77, 83, 84,..........86, 87, 88, 94, 104, 120, 121, 140,..........143, 146, 154, 157, 164, 167, 168,...........169, 170, 171, 182, 186, 188, 189donors.80, 112, 128, 150, 151, 155, 156,...................162, 165, 173, 174, 175, 183doping.............87, 88, 92, 103, 106, 111,...........124, 147, 153, 156, 161, 172, 175..................................................176, 177DX.....................................155, 156, 162electroluminescence .......................... 76ellipsometry..................................... 102epitaxy ........37, 78, 79, 81, 97, 119, 160,..........166, 167, 168, 169, 170, 173, 174,..........................................175, 177, 182etching ........................................76, 171exciton ..............................132, 176, 183faults .........81, 84, 87, 92, 104, 145, 152,..................157, 158, 159, 160, 169, 172,..........................................179, 182, 183

(Fe,Cr)2CuO4Diffusion .......................................... 106

FeldsparTwins ............................................... 107

(Fe,Ni)2O4O Diffusion....................................... 108Point Defects .................................... 108

Fe2(Ni,Zn,Cu)O4Interfacial Diffusion ......................... 108

FeOCa Diffusion ..................................... 108

Fe2O3Surface Reconstruction ..................... 109

Fe2O3/MgOSurface Reconstruction ..................... 109

Fe3O4Antiphase Boundaries....................... 109

Fe20(Sr,Bi)10Ox

O Permeation.............................110, 111

films ........66, 67, 78, 80, 89, 94, 99, 102,.......... 104, 106, 109, 113, 114, 121, 126,.......... 130, 131, 132, 134, 135, 136, 141,.......... 142, 143, 146, 151, 152, 158, 160,.......... 161, 162, 163, 164, 167, 169, 170,...................171, 172, 177, 181, 183, 185

GaAlNPoint Defects .....................................162

GaInN/GaNV-Defects ..........................................163

GaNDislocations.............. 164, 167, 168, 169,.......................................... 170, 171, 172Domain Boundaries ...........................172Electron Irradiation....................165, 166Ga Diffusion......................................163Gamma Irradiation ............................166Ion Bombardment..............................166Ion Implantation ................................166Micropipes ..........................................78N Diffusion ...............................163, 164Nanopipes ...........................................78Nanotube Defects ..............................172Point Defects .....156, 165, 166, 173, 174,.................................. 175, 176, 177, 178Stacking Faults .......... 158, 159, 172, 179Surface Defects ................. 159, 179, 180Surface Diffusion ..............................163Surface Reconstruction..............180, 181Zn Diffusion......................................164

GaN/Al2O3

Al Diffusion ......................................181Dislocations.......................................181Ga Diffusion......................................181Interdiffusion.....................................181

GaN/GaAsDislocations.......................................182Stacking Faults ..................................182

Keyword Index

215

GaN/InGaNInterdiffusion.................................... 182

GaN/SiCDislocations...................................... 183Grain Boundaries.............................. 183Ion Implantation ............................... 183Point Defects .................................... 183Stacking Faults ................................. 160

GeO2-Rb2O-Ag2OIonic Conduction .............................. 111

growth ............................................... 37impedance spectroscopy ................... 57

In16Fe8S32Point Defects .................................... 185

InGaNPoint Defects .................................... 178

InNStacking Faults ..........................158, 159

In2O3Point Defects .................................... 111

intercalation ...................................... 57interdiffusion............ 123, 127, 164, 182interstitials ........ 69, 71, 82, 90, 93, 111,...................121, 133, 156, 165, 184, 185ionic conductors .................................. 1

KH2PO4Dislocations...................................... 186

KNbO3Point Defects .................................... 111

K2RbC60Point Defects ...................................... 72

KTaO3Point Defects .................................... 112

KTiOPO4Er Diffusion.......................................186Nb Diffusion .....................................186Nd Diffusion .....................................186Point Defects .....................................186

LaAlO3Point Defects .....................................112

(La,Ca)CrO3

Grain Boundary Diffusion .................113Sr Diffusion.......................................113

(La,Ca)MnO3

Grain Boundaries...............................113

LaMnO3Point Defects .....................................113

(La,Pb)FeO3Ionic Conduction ...............................114O Permeation.....................................114

(La,Sr)(Co,Fe)O3

O Diffusion .......................................114

(La,Sr)CoO3Point Defects .....................................114

La8Sr8Co16O36Point Defects .....................................114

(La,Sr)2CuO4Planar Defects ...................................115Point Defects .....................................115

(La,Sr)MnO3

Point Defects .....................................115

La0.85Sr0.15MnO3Point Defects .....................................116

LiAlSiO4Ionic Conduction ...............................116Point Defects .....................................116

Keyword Index

216

Li2B4O7Twins ............................................... 116

Li0.5La0.5TiO3

Ionic Conduction .............................. 117

LiNbO3H Diffusion....................................... 118Point Defects .................................... 118Twins ........................................118, 119

Li2O-(Al2O3,Ga2O3)-TiO2-P2O5Ionic Conduction .............................. 119

Li2O-Al2O3-TiO2-SiO2-P2O5

Ionic Conduction .............................. 119Li Diffusion ...................................... 119

low-energy electron diffraction ........ 37luminescence ....... 76, 95, 122, 123, 132,..........133, 134, 136, 156, 158, 159, 172,.......................... 173, 174, 175, 176, 177

MgCO3Point Defects .................................... 187

MgOCa Diffusion ..................................... 119Dislocations...............................120, 121Electron Irradiation........................... 121Grain Boundaries...............120, 121, 122Ion Bombardment............................. 119O Diffusion....................................... 119Pipe Diffusion .................................. 120Point Defects .....................120, 121, 122Surface Diffusion.............................. 121Zn Diffusion ..................................... 119

MgO/FeOInterdiffusion.................................... 123

MgTiO3Ion Bombardment............................. 123Point Defects .................................... 123

micro-twin ......................................... 87micro-voids........................................ 89

mid-gap ....................................177, 183migration..........66, 73, 74, 93, 107, 120,...................121, 122, 143, 154, 164, 185misfit83, 94, 97, 125, 143, 144, 168, 185mismatch .................. 124, 142, 143, 168misorientation .......... 105, 150, 188, 189

Mn(La,Ba)O3

Point Defects .....................................124

MnO/CuDislocations.......................................125

Mn2O3/AgDislocations.......................................125

Mn2O3/CuDislocations.......................................125

Mn3O4/AgDislocations.......................................124

mobility ............65, 67, 85, 94, 127, 129,...................156, 157, 159, 164, 170, 188models . 93, 102, 115, 134, 145, 149, 186molecular beam epitaxy ....................37monolayer .................. 81, 159, 180, 181Monte Carlo .......................92, 116, 188mosaic...............................................106Mössbauer................................106, 185multi-layers ...................... 126, 185, 184muon.................................................106nanocrystalline.........................130, 144native ................................ 161, 173, 175

Nb2(Pb,Ba)O6

H Diffusion .......................................125

near-coincidence ..............................185neutron ..75, 92, 116, 124, 127, 129, 147

NiO/Fe2O3Interdiffusion.....................................126

non-equilibrium ...............................112non-stoichiometry ........... 91, 92, 97, 99,...........................114, 115, 127, 145, 157

Keyword Index

217

n-type............79, 81, 152, 156, 157, 161,... 165, 166, 170, 171, 173, 174, 177, 178nucleation .......................................... 37organometallic .................168, 171, 174out-diffusion ............. 102, 103, 133, 182oxidation.............. 37, 69, 86, 87, 90, 98,..........................................115, 124, 130oxide films ......................................... 37

P2O5-ZrO2-SiO2


passivated ........................................ 131passivation................................134, 135patterning...................................79, 167

(Pb,Ca)TiO3

Point Defects .................................... 126

(Pb,La)(Zr,Ti)O3Point Defects .................................... 126

Pb[(Nb,Mg)0.9Ti0.1]O3Point Defects .................................... 127

Pb(Nb,Ni)O3/Pb(Zr,Ti)O3Nb Diffusion..................................... 127Ni Diffusion ..................................... 127Ti Diffusion...................................... 127Zr Diffusion...................................... 127

(Pb,Sm)TiO3Point Defects .................................... 126

Pb2(Sn1.5W0.5)O6.5

Point Defects .................................... 127

PbTiO3Point Defects .................................... 128

Pb2(Ti1.5W0.5)O6.5Point Defects .................................... 127

PbWO4Neutron Irradiation ........................... 128Point Defects .............................128, 129

Pb(Zr,Ti)O3

Antiphase Boundaries........................129Point Defects .....................................129

Peierls ......................... 92, 154, 157, 181permeation ..........95, 110, 111, 114, 131photoconductivity ............................162photo-electron ....................78, 112, 184photo-emission ...........................81, 132photo-induced ..........................100, 174photo-ionization ...............................177photoluminescence........ 65, 74, 80, 129,.......... 130, 131, 132, 134, 163, 164, 174,...........................175, 177, 178, 182, 183positrons.................... 72, 73, 74, 80, 89,.......................................... 114, 126, 127protons ..............118, 125, 126, 133, 139p-type..................65, 132, 152, 156, 157,.......................................... 161, 176, 178radiation-enhanced..........................119radiation-induced ............................123Raman .........79, 130, 137, 168, 173, 182

Rb6C60Point Defects .......................................72

Rietveld ........................ 72, 85, 127, 185roughening .................................99, 182Rutherford ................. 73, 123, 132, 186scanning tunnelling microscopy ........37Schottky ...................................141, 145self-diffusion.............................108, 164self-interstitials ..................................66Shockley .............................67, 124, 125

SiCDefect Annealing...........................73, 74Dislocations.............................75, 76, 77Electron Irradiation..............................72Ion Implantation ............................73, 74Micropipes ....................................77, 78Nanopipes ...........................................78Neutron Irradiation ..............................75Point Defects ....72, 73, 74, 75, 78, 79, 80Stacking Faults ..............................80, 81Surface Reconstruction........................81

Keyword Index

218

(Si,Ge)C/SiPoint Defects ...................................... 81

simulations .....65, 66, 71, 79, 81, 89, 91,..............92, 93, 120, 127, 139, 141, 144,...........146, 152, 158, 160, 162, 172, 188

SiNxPoint Defects .................................... 183Defect Annealing.............................. 183

sintering...88, 89, 97, 103, 108, 139, 144

SiO2B Diffusion....................................... 130C Diffusion....................................... 131D Diffusion....................................... 131Electron Irradiation....................131, 132Gamma Irradiation............................ 132Ion Bombardment............................. 132Neutron Irradiation ........................... 131Point Defects ..... 131, 132, 133, 134, 135SiF4 Diffusion .................................. 131Twins ............................................... 135

SiO2/SiInterface Defects........................135, 136

SiO2/SiCInterface Defects............................... 136

SiO2-GeO2

Point Defects .............................137, 138UV Irradiation .................................. 137

SiO2-Li2O-TiO2-BaO-La2O3

H Diffusion....................................... 138

SiOxDefect Annealing.............................. 129Point Defects .................................... 129

Si(O,N)Point Defects .................................... 184

(Sr,La)CoO3


O Diffusion .......................................138

(Sr,La)TiO3Point Defects .....................................139

(Sr,La)3(VO4)2Ba Diffusion......................................139Ca Diffusion......................................139Point Defects .....................................139Sr Diffusion.......................................139

SrTiO3Dislocations.......................................140Grain Boundaries...............................141H Diffusion .......................................139O Diffusion .......................................140Point Defects .............................141, 142

SrTiO3/LaAlO3Dislocations.......................................142

Sr(Ti,Zr)O3/SrTiO3

Dislocations.......................................142

SrZrO3/LaAlO3

Dislocations.......................................142

SrZrO3/SrTiO3Dislocations.......................................142

stacking faults ................ 75, 76, 87, 158stoichiometric.............. 91, 92, 118, 119,.......................................... 124, 142, 145super-dislocations ..............................77superlattice.................................93, 129superstructure..................................114surface defects....................................37surface morphology ...........................37

TaCC, N Diffusion.....................................82

Ta2CC, N Diffusion.....................................82

TaNC, N Diffusion.....................................82

Keyword Index

219

Ta2NC, N Diffusion .................................... 82

Ta2O5Point Defects .................................... 143

TaSiNO Diffusion....................................... 184

temperature-dependent ...126, 174, 175thermoluminescence ..........90, 147, 148

TiC/CuDislocations........................................ 82

Ti(C,N)Twins ............................................... 184

TiN/B-C-NInterdiffusion.................................... 184Ion Bombardment............................. 184

TiO2Dislocations...................................... 144Grain Boundaries.............................. 144O Diffusion....................................... 143Point Defects .................................... 144Surface Defects................................. 145Surface Diffusion.............................. 144Ti Diffusion...................................... 143Twins ............................................... 144

Ti3SiC2C Diffusion......................................... 83Si Diffusion ........................................ 83Dislocations........................................ 84Stacking Faults ................................... 84

Ti(Sr,Ca)O3

Planar Defects .................................. 145

tracers...............................108, 119, 154transients....72, 75, 79, 81, 94, 135, 166,.................................. 173, 174, 175, 177transitions.......... 70, 89, 90, 93, 99, 100,..........102, 105, 113, 116, 124, 126, 127,..........145, 146, 148, 151, 154, 155, 156,

...................158, 164, 173, 177, 178, 189traps .................134, 135, 148, 174, 175,.......................................... 177, 178, 184tunnelling 71, 80, 97, 109, 141, 145, 181twinning ......87, 106, 116, 117, 118, 184twins ... 81, 87, 94, 95, 99, 100, 101, 106,..................107, 116, 118, 119, 144, 146,.......................................... 151, 184, 187ultra-thin ..........................................143

UO2Point Defects .....................................145

vacancies ........66, 69, 70, 72, 73, 74, 75,............ 78, 84, 85, 89, 90, 92, 93, 94, 97,..............98, 99, 100, 104, 106, 108, 111,.......... 112, 113, 114, 115, 116, 118, 120,.......... 121, 122, 126, 127, 128, 129, 132,.......... 134, 133, 134, 135, 137, 139, 142,.......... 144, 145, 147, 148, 152, 153, 154,.......... 155, 156, 157, 159, 161, 162, 165,... 166, 173, 174, 176, 178, 180, 187, 188vacancy-type ................................73, 74wafers .................................83, 170, 182weak-beam .......................................157

WO3Point Defects .....................................146Twins ................................................146

WSiNP Diffusion ........................................185

wurtzite ............152, 156, 158, 159, 167,.......................................... 169, 179, 182X-irradiation ....... 72, 78, 79, 81, 85, 86,.......... 88, 90, 91, 93, 106, 108, 114, 123,.......... 127, 139, 142, 143, 144, 147, 153,........... 154, 159, 169, 176, 181, 183, 185

YAGDislocations.......................................146Point Defects .....................................146

Y2O3Point Defects .....................................147

Keyword Index

220

Y2O3-Nb2O5


YVO4Gamma Irradiation............................ 147Point Defects .............................147, 148X-Irradiation..................................... 148

Yb2.75C60Point Defects ...................................... 84

ZeolitePoint Defects .................................... 148

zeolites ................................................. 1

zig-zag................................................ 89

ZnODislocations...................................... 149Grain Boundaries.......................149, 150O Diffusion....................................... 149Point Defects .............................150, 151Twins ............................................... 151

ZnO/Al2O3

Dislocations.......................................151Stacking Faults ..................................151

ZnO-CuOGrain Boundaries...............................152

ZnO-In2O3

Point Defects .....................................152Stacking Faults ..................................152

ZnWO4Point Defects .....................................153X-Irradiation......................................153

ZrO2Dislocations.......................................154Grain Boundaries...............................154O Diffusion .......................................154Point Defects .....................................155Zr Diffusion.......................................154

ZrO2-TiO2-Y2O3Domain Boundaries ...........................155

Defects and Diffusion in Ceramicspearton.mse.ufl.edu/semic_properties/data/5037.pdf65 Defects and...

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