Nitrogen in germanium · Nitrogen in germanium I. Chambouleyrona) Instituto de Fı´sica ‘‘Gleb...

Nitrogen in germaniumI. Chambouleyron and A. R. Zanatta Citation: Journal of Applied Physics 84, 1 (1998); doi: 10.1063/1.368612 View online: http://dx.doi.org/10.1063/1.368612 View Table of Contents: http://scitation.aip.org/content/aip/journal/jap/84/1?ver=pdfcov Published by the AIP Publishing Articles you may be interested in Dopant effects on solid phase epitaxy in silicon and germanium J. Appl. Phys. 111, 034906 (2012); 10.1063/1.3682532 Electron transport and band structure in phosphorus-doped polycrystalline silicon films J. Appl. Phys. 105, 033715 (2009); 10.1063/1.3068349 Improvement on electron field emission properties of nanocrystalline diamond films by co-doping of boron andnitrogen J. Vac. Sci. Technol. B 21, 1074 (2003); 10.1116/1.1576396 Valence band spectra of nitrogen incorporated amorphous carbon films J. Appl. Phys. 89, 2414 (2001); 10.1063/1.1337602 Defects, doping, and conduction mechanisms in nitrogen-doped tetrahedral amorphous carbon J. Appl. Phys. 81, 1289 (1997); 10.1063/1.363907

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Nitrogen in germaniumI. Chambouleyrona)

Instituto de Fı´sica ‘‘Gleb Wataghin’’, Universidade Estadual de Campinas, P.O. Box 6165,13083-970 Campinas, SP, Brazil

A. R. ZanattaInstituto de Fı´sica de Saõ Carlos, Universidade de Saõ Paulo, P.O. Box 369, 13560-970 Saõ Carlos, SP,Brazil

~Received 8 July 1997; accepted for publication 23 March 1998!

The known properties of nitrogen as an impurity in, and as an alloy element of, the germaniumnetwork are reviewed in this article. Amorphous and crystalline germanium–nitrogen alloys areinteresting materials with potential applications for protective coatings and window layers for solarconversion devices. They may also act as effective diffusion masks for III-V electronic devices. Theexisting data are compared with similar properties of other group IV nitrides, in particular withsilicon nitride. To a certain extent, the general picture mirrors the one found in Si–N systems, asexpected from the similar valence structure of both elemental semiconductors. However, importantdifferences appear in the deposition methods and alloy composition, the optical properties of asgrown films, and the electrical behavior of nitrogen-doped amorphous layers. Structural studies arereviewed, including band structure calculations and the energies of nitrogen-related defects, whichare compared with experimental data. Many important aspects of the electronic structure of Ge–Nalloys are not yet completely understood and deserve a more careful investigation, in particular thestructure of defects associated with N inclusion. The N doping of thea-Ge:H network appears to bevery effective, the activation energy of the most effectively doped samples becoming around 120meV. This is not the case with N-dopeda-Si:H, the reasons for the difference remaining an openquestion. The lack of data on stoichiometricb-Ge3N4 prevents any reasonable assessment on thepossible uses of the alloy in electronic and ceramic applications. ©1998 American Institute ofPhysics.@S0021-8979~98!07213-2#

TABLE OF CONTENTS

I. INTRODUCTION. . . . . . . . . . . . . . . . . . . . . . . . . . . . 2II. PREPARATION OF COLUMN IV NITRIDES. . . 3

A. Thin film deposition. . . . . . . . . . . . . . . . . . . . . . . 3B. Column IV nitrides. . . . . . . . . . . . . . . . . . . . . . . . 3

1. Silicon nitride. .. . . . . . . . . . . . . . . . . . . . . . . 32. Germanium nitride. . . . . . . . . . . . . . . . . . . . . 43. Carbon nitride. . . . . . . . . . . . . . . . . . . . . . . . . 44. Tin nitride. . . . . . . . . . . . . . . . . . . . . . . . . . . . 4

C. Other metal and semiconductor nitrides. . . . . . . 4D. Deposition conditions, particle bombardment,

hydrogenation and structure of group IVamorphous thin films. . . . . . . . . . . . . . . . . . . . . . 5

III. STRUCTURE OFa-Ge~N! ALLOYS. . . . . . . . . . 7A. Structural studies by EXAFS. . . . . . . . . . . . . . . . 7B. Electronic structure. . . . . . . . . . . . . . . . . . . . . . . . 8

1. Theoretical approaches and coordinationdefects. . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . 8

2. Experimental reports. . . . . . . . . . . . . . . . . . . 8C. Structural studies by optical techniques. .. . . . . 10

1. Infrared spectroscopy. . . . . . . . . . . . . . . . . . . 102. Raman spectroscopy. . . . . . . . . . . . . . . . . . . . 12

D. Hydrogenation, structure and stability ofa-GeN films. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13

E. The Ge–N bond: randomness, charge transferand electronegativity. . . . . . . . . . . . . . . . . . . . . . 14

IV. OPTICAL PROPERTIES OFa-Ge ALLOYS. . . . 16A. Optical properties. . . . . . . . . . . . . . . . . . . . . . . . . 16B. Optical absorption in amorphous

semiconductors. . . . . . . . . . . . . . . . . . . . . .. . . . . 161. Optical band gap~E04 andETauc!. . . . . . . . . 162. Band-gap widening. . . . . . . . . . . . . . . . . . . . . 17

C. Electronic versus structural disorder in

amorphous semiconductors. . . . . . . . . . . . . . . . . 17

1. Optical absorption and Urbach edge. .. . . . . 17

2. Structural disorder. . . . . . . . . . . . . . . . . . . . . 19

3. Composition, structural disorder and

optical band gap. . . . . . . . . . . . . . . . . . . . . . . 19

V. TRANSPORT PROPERTIES. . . . . . . . . . . . . . . . . . 20

A. H-free Ge12xNx alloys: crystalline and

amorphous. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

B. Hydrogenateda-Ge12xNx films (x>0.01). . . . . 21

C. Nitrogen as an impurity in Ge (x<0.01). . . . . . 22

1. N in crystalline germanium. . . . . . . . . . . . . . 22

2. N as an active dopant in amorphous Ge

~and Si!. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23

D. Photoconductivity in N-dopeda-Ge:H films. . . 26

VI. CONCLUDING REMARKS. . . . . . . . . . . . . . . . . . 27a!Electronic mail: [email protected]

JOURNAL OF APPLIED PHYSICS VOLUME 84, NUMBER 1 1 JULY 1998

10021-8979/98/84(1)/1/30/$15.00 © 1998 American Institute of Physics


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I. INTRODUCTION

The nitrogen atom, 1s22s22px2py2pz , can complete itsvalence shell in one of the following ways:1

~i! electron gain to form the nitride ion N32;

~ii ! electron-pair bonds~single or multiple!, like in mo-lecular N2;

~iii ! electron-pair bonds with electron gain, as in NH22;

~iv! electron-pair bonds with electron loss, as in tetrahe-dral ammonium.

The molecules NR3 are pyramidal; the bonding involvessp3 orbitals so that the lone pair occupies the fourth position.The propensity of nitrogen, like carbon, to formpp –ppmultiple bonds is a feature that distinguishes it from phos-phorus and the other Group V B elements. Thus nitrogen asthe element is dinitrogen, N2, with a very high bond strengthand a short internuclear distance~1.094 Å!, whereas phos-phorus forms P4 molecules or infinite layer structures inwhich there are only single bonds.1

The numerous covalent nitrides and their properties, aswell as their potential applications, vary greatly dependingon the element with which nitrogen is combined. For ex-ample, in recent decades an increasing interest developed inthe wurtzite polytypes of Group III nitrides: GaN, AlN, andInN, which form a continuous alloy system, the band gapsranging from ;2 eV for InN, to ;3.5 eV for GaN, to;6.5 eV for AlN ~Fig. 1!.2–6 Thus Group III N alloys couldpotentially be fabricated into active optical devices at wave-

lengths going from the red well into the ultraviolet, in thesame way as the highly successful As-based and P-basedGroup III alloys used in the infrared, red and green regionsof the spectrum.

Of particular relevance to the present review is the roleof nitrogen, from doping concentrations to the alloy phase, inGroup IV elements: tetrahedral C, Si, Ge anda-Sn. Thedriving idea behind this presentation is to review all reportedproperties of nitrogen in germanium, either crystalline oramorphous. As the understanding of a rather new materialimproves when its properties are discussed in the context ofthose of similar~and better known! systems, we will fre-quently refer to nitrogen in silicon and, whenever appropri-ate, to nitrogen in tetrahedral carbon anda-Sn.

By far the most studied among the above systems issilicon nitride.7 The research on Si3N4 has been largely fu-eled by its use in microelectronics technology, where it actsas an effective insulating material and a diffusion mask forimpurities. These applications require the etching of thefilms, which proceeds isotropically for amorphous andhighly microcrystalline insulator compounds. Etchants oftechnical importance are ammonium fluoride, buffered hy-drofluoric acid at room temperature, hot 85% phosphoricacid for pattern etching with metal masks, and miscellaneousother etchants, usual strong mineral acid or bases.8 Table Iindicates the etching rate of different solutions ona-Si anda-Ge based nitrides. The electronic applications of siliconnitride have been complemented with ceramic uses, the im-portance of which has not ceased to grow in recent years.9 Atthe other end of the research efforts on Group IV nitrides liesthe SnN system, the properties of which are poorly under-stood at present.10

Nitrogen in diamondlike carbon and carbon nitride CNalloys are the subject of intense research. The interest stemsfrom at least two reasons. First, the possibility of preparingb-C3N4, a hypothetical material predicted to have a hardnesscomparable to that of diamond.5 Up to now, however, theefforts made to synthesizeb-C3N4 have not been verysuccessful.11 Second, nitrogen is a possible donor in dia-mond. Substitutional nitrogen gives, besides a variety ofother not yet fully understood nitrogen related defects, a do-nor level at 1.7 eV below the conduction band.12 Controllingthe conductivity of diamondlike films by chemical dopingwould open the possibility for high temperature solid-stateelectronics, a matter of interest for military and other appli-cations.

Germanium nitride films13,14 were prepared in the 1960sand 1970s but the subject did not develop consistently.In 198515 the Laboratory of Photovoltaic Research at

FIG. 1. Optical gap of column III and IV nitrides as a function of bondlength~see Refs. 2–6!. Note the different bond length and gap energy scales.The nitrides of heavy isocore elements display similar band-gap variations.

TABLE I. Etch rates of group IV nitrides produced by strong acids and bases.

MaterialDeposition or

anneal temperature

Etch rate at RT

HF BHF H3PO4 H2SO4 NaOH

a-Si3N4 CVD: 1000–1300 K medium low medium nil nila-SixNyHz PACVD: 500 K very high medium high nil nila-Ge3N4 CVD: 700–1000 K very high high medium low low

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UNICAMP, Campinas, started a program aimed at a system-atic study of the structural, compositional, electrical and op-tical properties of GeN alloys, and of N as an impurity in thea-Ge:H network. Other research groups, particularly in Japanand Germany, have contributed to unravel many puzzlingquestions. In the present review, we report results on N inGe, with particular emphasis on our own research at Campi-nas.

II. PREPARATION OF COLUMN IV NITRIDES

A. Thin film deposition

The chemical vapor deposition~CVD! of thin films hasbecome one of the most important methods of film formationand now constitutes a powerful and versatile tool in moderntechnologies, such as those employed in solid-state electron-ics. The reasons for the rapidly growing importance of CVDin the last few years lay primarily in its versatility for depos-iting a very large variety of elements and compounds at rela-tive low temperatures, in the form of both vitreous and crys-talline layers having a high degree of perfection and purity.Other unique advantage of CVD over other methods of filmformation is the relative ease of creating materials of a widerange of accurately controllable composition and layer struc-tures that are difficult or impossible to attain by othertechniques.16

Roughly speaking, CVD can be defined as a materialsynthesis method in which the constituents~in the vaporphase! react to form a solid film at some surface.17 Thus theoccurrence of chemical reactions is an essential characteristicof CVD. In order to understand CVD processes, one mustknow which chemical reactions occur in the reactor and towhat extent. Furthermore, the effects of process variablessuch as temperature, pressure, input concentrations, and flowrates on these reactions must be understood. The basic con-figurations of CVD systems have evolved along the years,the most important ones being the so-called plasma-assistedCVD method.

The uniqueness of the plasma to generate chemicallyreactive species at low temperatures is due to the nonequi-librium nature of the plasma state. By nonequilibrium, wemean a gas plasma typically sustained at;0.1 to severalTorr, which exhibits temperatures of the free electrons oftens of thousands of degrees Kelvin, while the temperatureof the translational and rotational modes of free atoms, radi-cals, or molecules will be only hundreds of degrees Kelvin.Illuminating discussions about the fundamental aspects andcharacteristics of low temperature plasma species can befound in the specialized literature.8,18,19

One of the prime motivating factors in utilizing plasmadeposition processes is that the substrate temperature can bekept relatively low, typically 500 K or lower. ConventionalCVD processes require temperatures substantially higher thatmay be inappropriate for certain substrate materials or devicestructures. The films that are deposited by plasma reactionsare usually amorphous in nature, with short-range structuralordering only. The composition of the films can be varied ina controlled way by changing some key plasma parameters,such as reactant gas flow ratios~and/or target composition if

a sputtering system is employed!. It is to be expected, andindeed observed, that electrical, mechanical, and chemicalproperties vary with composition and deposition conditions.

Both the dc- and rf-driven CVD methods depend largelyon the plasma chemistry. As a result, the significant param-eters that can affect the characteristics and properties of adeposited film are more varied than in evaporated or sput-tered material. The literature ona-SiN-based materials de-posited by CVD-like plasma-assisted methods is quite ampleand very interesting reviews can be easily found.20–23 Re-ports on CVD-like depositeda-GeN materials, on the otherhand, are relatively scarce.24–28

In the physical vapor deposition~PVD! of thin layers,more energetic phenomena usually occur. Within the variousPVD-like methods, the sputtering technique is by far themost widely used. Sputtering is the plasma-assisted vapor-ization of a material~called target! by bombarding it withhigh energy particles~normally Ar1 ions!. Due to these en-ergetic collisions atoms or fragments are ejected from thetarget’s surface and contribute to film formation. In additionto the particles from the target, most of the gaseous speciespresent in the plasma undergo physical/chemical reactionsand determine important properties of the thin films beingdeposited. In current practice sputtering discharges aredriven by high frequency~13.56 MHz! power supplies, andstrictly speaking, there is no cathode or anode since the netflow of charged species to each electrode is zero. Due to thevery different mobility of electrons and ions and differencein electrode area, however, a negative dc bias develops onthe powered electrode. The election for rf-instead of dc-driven systems is determined by the convenience of the lowpressure at which rf plasmas operate as well as by the kind ofmaterial ~e.g., insulator! to be produced. This advantagecombined with capacitive coupling allows the reactive depo-sition of insulating and semiconducting films~a-SiN anda-GeN, for example! under highly controlled conditions.Most of the pioneering work on the reactive sputtering ofa-SiN:~H! thin films used gaseous Ar1N2

29–36 mixtures in-stead of Ar1NH3.

37 The sputter deposition ofa-GeN:~H!films started just a few years ago.15,38–40

B. Column IV nitrides

1. Silicon nitride

Thin Group IV nitride films can be deposited by CVDfrom a variety of precursor gases. In thermal processes,Si3N4 films have been prepared at high substrate tempera-tures~1000–1200 K! from silane SiH4 and ammonia NH3 ina hydrogen atmosphere41,42 and from dichlorosilane SiCl4

and ammonia at low pressure.43,44 Plasma deposition ofSi3N4 films from SiH4, N2 and/or NH3 has been studied ex-tensively and is now at the stage of production applicationsin semiconductor device manufacture, mainly for the passi-vation of device surfaces. Si3N4 films 3000–10 000 Å thickare excellent diffusion masks for alkali ion contaminants andother impurities. The recognition of the potential of low-temperature deposited nitrides in semiconductor applicationsis not new, and the delay in the widespread use of this ma-terial has resulted, in part, from the lack of equipment ad-

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equate to the film uniformity and production requirements ofthe industry. The chemistry of plasma Si3N4 deposition isextremely complex and the detailed reaction kinetics is notfully understood, largely because such deposition reactionsare difficult to diagnose, as opposed to processes takingplace exclusively in the gas phase. Plasma-assisted siliconnitride produced via plasma is different from silicon nitrideproduced by conventional CVD or PVD techniques. In theformer, the composition can be controlled varying the ratioof the reactant gas flow, the power level, the substrate tem-perature during deposition, and the pressure in the reactorvessel. Remote and direct plasma-enhanced CVD techniquesuse SiH4 and either NH3 or N2 as the precursor gases todeposit silicon nitride.45–49 Disilanyl amine50 and cyclopen-tadienyl substituted silanes~Cp SiH3 and Cp2 SiH2! in com-bination with N2/NH3 have also been employed.51 Finally,electron cyclotron resonance~ECR! plasmas~which producea higher degree of ionization than normal rf or microwaveplasmas! have recently been used to deposit silicon nitride.In ECR deposition processes, SiH4 and N2/NH3 orSi~NMe2!3H and N2 are used as precursors and films aredeposited at substrate temperatures of typically 350–600 Kwith growth rates of 50– 200 Å min21.52–54

2. Germanium nitride

Crystalline germanium nitride layers produced by ther-mally activated chemical methods were found to containGe3N2 and Ge3N4 as constituents.55 Ge3N2 is obtained by thefollowing chemical reactions:56

GeI2 ——→NH3

~GeNH!n1NH4I, ~1!

and

3~GeNH! ——→550– 600 K

Ge3N21NH3. ~2!

Ge3N2 can also be formed when germane is decomposed byactive nitrogen:57

3GeH41N2* ——→400– 600 K

Ge3N216H2. ~3!

Similarly, crystalline Ge3N4 can be prepared from el-emental Ge reacting with NH3 or starting from Ge~NH!2

groups by one of the following processes:58,59

Ge1NH3 ——→1000 K

Ge3N4 ←700 K

@Ge~NH!2#n ~4!

in which case,c-Ge3N4 has the hexagonal structure ofphenacite (b-Ge3N4), the lattice parameters being:58,60 a50.8038 nm;c50.3074 nm. There is also a rhombic struc-ture attributed to c-Ge3N4 with lattice parameters61 a51.384, b50.906 andc50.818 nm. The density14,62 ofc-Ge3N4 is r;5.3 g cm23 and its heat of formation63

2DH;65 kJ mol21. Other possibilities for the preparationof nonstoichiometric GeN thin films, using thermally acti-vated CVD methods, are: GeH4 and NH3 ~at ;800 K!;64

GeCl4 and NH3 ~with temperatures in the 700–900 Krange!13,65 and GeH4 and N2H4 ~at ;700 K!,66 to mentionjust a few of them.

Stoichiometric (Ge3N4) germanium nitride films havebeen mostly prepared by thermally activated CVD of GeCl4

and NH3 at 650–850 K,13,65 from GeH4 and NH3 at 900 K,64

and GeH4 and hydrazine at 600 K.66 Plasma-assisted CVDhas also been used to deposit GeN films at substrate tempera-tures of 500–650 K using GeH4/NH3 or GeH4/N2.

27,28How-ever, most of the present understanding of GeN alloys de-rives from films deposited by the rf sputtering technique.

3. Carbon nitride

The interest in hydrogenated diamondlike carbon films(a-C:H! stems from their unique electrical, chemical andmechanical properties. In the past few years the effect of theincorporation of dopants in these films has been studied. Inparticular, since it was proposed that the bulk modulus of thehypothetical materialb-C3N4, structurally analogous tob-Si3N4, may be similar to that of diamond,67 the incorpo-ration of nitrogen into carbon films has received special at-tention. After an early study on rf sputtered CN films, carriedout by Cuomoet al. in 1979,68 attempts to incorporate nitro-gen into carbon films were made using several techniques: rfor dc sputtering,69,70 laser ablation,71 and ion-beam assisteddeposition.72 These films are amorphous, and for some par-ticular deposition conditions the presence ofb-C3N4 nanoc-rystals embedded into the amorphous matrix wasdetermined.70,71 The incorporation of nitrogen improves themechanical properties of diamondlikea-C films, in particularthe tribological ones like wear and friction.73 Nitrogen incor-poration up to 20 at % ina-C:H films was achieved throughplasma-assisted CVD of different gases mixtures.74–78It wasshown thata-CN:H films deposited by rf plasma decompo-sition of CH4–N2 mixtures can be as hard asa-C:H films,and that the incorporation of nitrogen causes a reduction ofthe internal stress without any noticeable change in theirhardness.79,80 A possible explanation for this effect is thereduction of the mean atomic coordination number, and thusa reduced over-constraint due to the presence of threefold Nsubstituting fourfold coordinated C atoms.

4. Tin nitride

There has been little mention of tin nitride films in theliterature. Remy and Hantzpergue used reactive cathodicsputtering to prepare the first nitride films.10 Crystalline SnNfilms have been recently prepared by reactive sputtering ofSn in a N2 plasma,81 and by magnetron sputtering of a pureSn target in a gas mixture of Ar and N2 at roomtemperature.82 The only CVD SnN reported to date is thedeposition of polycrystalline material by atmospheric pres-sure CVD from Sn~NMe2!4 and NH3 at substrate tempera-tures of 500–700 K.83

C. Other metal and semiconductor nitrides

Nitrides of a number of metals are of technical impor-tance, particularly those of Ti, Zr, Hf, Nb, Ta, and Be~andmore recently B, Al, and V!, which form hard and highlystable refractory materials with very high melting points.Widely used, the applications of several of these coatings~especially TiN! include the improvement of the wear resis-tance of cemented carbide tools, and protective coatings.84–86

The chemical systems for producing CVD refractory nitrides



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are based on reacting volatile metal halides with N2, with orwithout H2, at high temperatures, i.e., 2700–3000 K for theZrCl4–N2 or HfCl4–N2 systems, and 1300–2100 K for theTiCl4–N2–H2 system.87 Temperatures in the range of 1100–1400 K, however, can be used to deposit TiN films by thelatter reaction.88 Films of hafnium nitride HfN, from the re-action of chlorides of Hf with N2–H2, have been depositedon W, also at relatively low temperatures in the range of1200–1600 K, and thermodynamic and kinetic studiesshowed that the process is mass transport limited at low H2

concentrations.89

Boron nitride~BN!, prepared under optimal CVD condi-tions, is an excellent insulator andp-type dopant source forjunction formation in Si device fabrication. Several CVDmethods have been investigated for depositing thin BN filmsof both microcrystalline and amorphous structure.90,91 Clearvitreous BN films are deposited on a variety of substrates at850–1250 K by reacting B2H6 and NH3.

90 The electricalconductivity of these films shows, however, that BN depos-ited at 1050 K on Si is similar to Si3N4 but is not as good aNa1 barrier and not as stable chemically.90

Another refractory nitride of interest is aluminum nitride~AlN !, which has potential as a good dielectric for active andpassive components in semiconductor devices because of itslarge energy gap and high thermal stability.92 Aluminum ni-tride films have been deposited by reacting NH3 with AlCl3

or Al~CH3!3.92–95

Films of gallium nitride~GaN! have been grown origi-nally by the ammonolysis of gallium monochloride GaCl.96

Pyrolysis of the GaBr3–NH3 complex,97 and ammonolysis ofGa~CH3!3 have also been employed.95 As for AlN, the maininterest in GaN stems from the possibility of high-temperature electronics and short-wavelength optical

applications.6 In addition to the above mentioned thin filmnitrides and respective methods of deposition, a great effortis made nowadays to prepare N-based new materials withmethods other than CVD.6

D. Deposition conditions, particle bombardment,hydrogenation and structure of group IV amorphousthin films

Regardless of the chemical/physical processes involved,surface adatom mobility ranks among the most importantdeposition parameters in plasma-assisted deposition meth-ods. Besides the thermal energy associated with the substratetemperature, extra energy can be supplied to the growingfilm by bombarding it with energetic particles. The processenhances adatoms mobility and may induce the removal ofloosely bonded particles, producing denser films with no co-lumnar microstructure.98 In other words, under certain con-ditions, particle bombardment may cause film etching, re-moving preferentially weakly bonded atoms and/orprotrusions that develop into columns. Ion, neutral speciesand/or electron bombardment of the growing surface add en-ergy to adatoms and also heat the substrate. The precise iden-tification and the role played by these different bombardingspecies is not an easy task and depends very much on thedeposition method and conditions. Table II reviews brieflythe ongoing debate on the nature and importance of the bom-barding species affecting most of the structural properties ofrf sputtereda-Si:H films in an Ar atmosphere.99–103

A systematic study on the effects of particle bombard-ment on the properties of rf sputtereda-Ge:H does not exist.However, the importance of particle bombardment to thestructural and electronic properties of this semiconducting

TABLE II. Particle bombardment in the formation of thin films according to different authors. Depending onthe adopted deposition conditions either electrons, charged or neutral particles can be associated to the bom-bardment of the thin film surface during growth.

Author Event Particle Consequences Observation

power delivered:Brodiea general rf electrons substrate heating Wion;30 mW cm22

sputtering (sp) We2;500 mW cm22

PAr determines bombardment removesAndersonb rf sp a-Si:H neutral Si species substrate loosely bonded species

bombardment

PAr determines Ar1 ions areRossc rf sp a-Si:H Ar1 ions structural accelerated by the

features floating potential

PAr determines electron bombardmentMoustakasd rf sp a-Si:H electrons or Ar1 substrate leads to high quality

bombardment films

Ar and Si species hydrogen atoms canTardye rf sp a-Si:H hydrogen atoms thermalize in the reach energies of

plasma 1.5 keV

aSee Ref. 99.bSee Ref. 100.cSee Ref. 101.dSee Ref. 102.eSee Ref. 103.



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alloy has been known for a decade.104–107 In particular rfsputtering was one of the first methods used to producea-Ge:H films of improved quality.104 Good qualitya-Ge:Hsamples are currently deposited under varied PVD-likeconditions.106,107

Just to illustrate the importance of the adatom mobilityon the optical properties of thea-GeN:H films consider Figs.2~a! and 2~b!, where the effects of the substrate temperatureand of the rf power delivered to the plasma during depositionare presented.108 The influence of the dc bias and substratetemperatures~adatom mobility! on theE04 optical gap~pho-ton energy at which the absorption coefficient reaches104 cm21! of the films is apparent in the figure.

The electronic properties ofa-semiconductors dependlargely on hydrogenation. Device qualitya-Si:H ~a-Ge:H!films have a density of dangling bond on the order of1015 cm23 (1017 cm23), as determined from electron spinresonance~ESR! data. In nonhydrogenateda-Si and a-Gethis density is;1019 cm23. The role of hydrogen, however,is not only to passivate dangling bonds, but also to reducethe strain of the Si~Ge! network improving the optoelec-tronic properties. Bonded hydrogen in thesea-structures alsodetermines the optical gap because the valence band recedeswith hydrogenation. Roughly speaking, the optical gap scaleswith the hydrogenation. The method/condition of depositionmay determine some upper limits~optical gap or hydrogencontent values, for example!. Figures 3~a! and 3~b! show theoptical gapE04 for a-Si:H anda-Ge:H samples deposited inthe same rf sputtering system, in an Ar1H2 atmosphere.Both sample series were deposited under similar conditionsbut possess quite different degrees of hydrogenation.

The incorporation of nitrogen ina-Si anda-Ge occurs ina way similar to that of hydrogen. Most samples consideredin this review were prepared by the rf sputtering method. Inthem the nitrogen concentration is proportional to the N2 orNH3 partial pressure during deposition. Figure 4 shows thenitrogen content ofa-GeN anda-SiN alloys, as determinedfrom a deuteron induced nuclear reaction analysis@14N(d,p)15N#, vs the N2 or NH3 partial pressure duringdeposition. Most of the features present in Fig. 4 are associ-ated with the deposition conditions~substrate bombardment

and temperature, for example! while others reflect chemicalaspects.

Closing this brief introduction on plasma-assisted depo-sition of N-containing thin films, it is worth mentioning thatdevice qualitya-Si:H ~and alloy! films are best produced byusing CVD-like or soft deposition methods,109 like the an-odic glow-discharge of the SiH4 technique. The sputteringtechnique has been much less used to preparea-Si:H, al-

FIG. 2. Optical gap (E04) vs ~a! the target bias and~b! the substrate tem-perature of rf sputtereda-GeN:H thin films. As seen, the optical gap~ornitrogen content! is influenced by both the target bias and substrate tempera-ture ~see Ref. 108!.

FIG. 3. Optical gaps (E04) of rf sputtereda-Si:H and a-Ge:H thin films~deposited under similar conditions and in the same deposition system! as afunction of hydrogen content. Note the band-gap widening mechanism in-duced by hydrogenation.

FIG. 4. Nitrogen concentration~as determined from nuclear reaction analy-sis! in rf sputtered group IV amorphous thin films vs the N2 or NH3 partialpressure. Within the studied range the nitrogen concentration scales linearlywith the N2 and NH3 partial pressures.



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though a-Si:H films of electronic quality have been alsoreported.110–112Contrary to the case ofa-Si:H films, PVD-like or hard deposition methods involving more energeticprocesses during deposition, like sputtering deposition tech-niques and cathodic glow-discharge, proved to be the mostappropriate for the deposition of optimizeda-Ge:H. The rea-sons for this difference are not yet understood, but it is clearby now that they may be at the origin of the difficultiesfound in the preparation of optimizeda-SiGe:H alloys. Yet,it is important to keep in mind that each deposition methodand condition produces thin films with quite specific charac-teristics, some of them being directly correlated to the physi-cal processes occurring during growth.

III. STRUCTURE OF a-GE„N… ALLOYS

It is convenient to report the structural properties of newbinary alloys in the broad context of well characterizedmodel systems of known stoichiometry and similar valencecoordination. For GeN alloys, the obvious choice isc-Si3N4,an intensively studied electronic material. Two structures ofc-Si3N4 have been identified by x-ray diffraction.113,114Theyhave been nameda and b phase, both having a hexagonalcrystal lattice. The coordination numbers for Si and N are 4and 3, respectively. Figure 5 shows the phenacite structure,as it is called, of crystallineb-Si3N4. It involves covalentbonds between planar bonded N and tetrahedral bondedSi.115 This local structure is consistent with Sisp3 hybridorbitals while N bonding is explained in terms of a linearcombination ofp orbitals, the planar geometry being givenby a strong repulsion of nonbonded Si atoms. According tothis picture thes electrons of N do not participate in bondingand the top of the valence band ofb-Si3N4 is filled by non-bonding Npz electrons.

Off-stoichiometric SiN alloys can be prepared by plasmaassisted chemical vapor deposition~PACVD! adjusting thepartial pressure of N2 or NH3 in the reaction chamber. Thecomposition of off-stoichiometric SiN alloys is given in theliterature with two different notations. Some authors preferto use Si12xNx , whereas others indicate composition usingSiNx . In the former notation 0<x<0.57 (4/7), whereas inthe latter 0<x<1.33 (4/3). The notation Si12xNx

(Ge12xNx) has been retained in the present review. For thesake of clarity, the data from the literature expressed as0<x<1.33 will be transformed into the 0<x<0.57 nota-tion.

Electron-diffraction measurements ona-Si12xNx :H ob-tained by the glow discharge of SiH4 and NH3 mixtures116

yield radial distribution functions~RDFs!, the peaks ofwhich are a combination of those expected for the first- andsecond-nearest neighbors inc-Si3N4 and Si. More structuralinformation is obtained by x-ray diffraction117 and neutronscattering118 on nearly stoichiometric samples deposited byCVD. The RDF of thea-phase was found to resemble that ofthe b-phase c-Si3N4 and bond angles of;109.8° and;121° were found around Si and N, respectively.116 More-over, analysis of small-angle scattering pointed to the exis-tence of voids, the presence of which reduces slightly thecoordination numbers of Si and N. Extended x-ray absorp-tion fine structure ~EXAFS! studies indicate that off-stoichiometric a-Si12xNx compounds are chemicallyordered.119

A. Structural studies by EXAFS

The structure of GeN compounds has been much lessstudied than that of SiN alloys. As expected from the similarvalence structure of Ge and Si, the structure ofc-Ge3N4 issimilar to c-Si3N4. Stoichiometricb-Ge3N4 has a densityr55.3 g cm23 and a nitrogen concentration of4.631022 atoms cm23.14,120 Boscherini et al.121 studied byEXAFS ~Ge K-edge! the local order ofc-Ge3N4 and of fivea-Ge12xNx :H samples with N concentrations ranging fromx50 up to x537 at. %. Figure 6 shows the Fourier trans-form of EXAFS signals corresponding toa-Ge:H;a-Ge12xNx :H andc-Ge3N4 samples. The arrows at the bot-tom of Fig. 6 indicate the expected position of features dueto first and second shell configurations, once phase shifts aretaken into account. Quantitative data analyses performed inkspace indicate that the full EXAFS signal ina-Ge12xNx :H

FIG. 5. A basal projection of the phenacite structure.b-phase germaniumnitride is hexagonal with two molecular units (2Ge3N4) in the unit cell. Thespace group isP63 /m. The fractions indicate the distance to the basal planein c-axis units.

FIG. 6. Magnitude of the Fourier transforms of EXAFS fora-Ge12xNx :Handc-Ge3N4 samples~Ref. 121!. The arrows indicate expected positions offeatures due to possible first- and second-shell configurations, once phaseshifts are taken into account. The arrow positions correspond to Ge–N,Ge–Ge, Ge–N–N, Ge–N–Ge, Ge–Ge–N, and Ge–Ge–Ge,respectively,from lower to higher interatomic distance. The right-hand side panel showsthe four possible second-shell configurations ina-Ge12xNx :H and the in-teratomic distances expected.



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can be completely simulated by a first shell component dueto Ge and N and a second shell component of Ge in a Ge–Nconfiguration, the relative importance of these contributionsvarying according to the N concentration. An interestingconclusion of this work is that the presence of Ge–N–N canbe excluded, an indication that when N is inserted in thea-Ge:H network its bonding to Ge is determined mostly byGe–N interactions.121 The constant value found for theGe–Ge second shell distance, equal to the average value inc-Ge3N4, is a further confirmation of this picture. Anotherimportant feature is that, irrespective of the amount of N inthe samples, interatomic distances remain constant:RGe–N51.835 Å, RGe–Ge52.44 Å, andRGe–N–Ge53.19 Å.The root-mean-square N bond-angle fluctuation was found tobe ,6° in all samples, a value smaller than the minimumfluctuation for bond angles on Ge;9°, suggesting thatbonds on N are more rigid than on Ge. Finally, the fact thatthe second shell EXAFS signal can be fitted with Ge in aGe–N–Geconfiguration gives further evidence that when Nenters thea-Ge network it exclusively form bonds with Geand thus induces a totally chemically ordered network, inclose agreement with the findings ona-Si12xNx :H.122

B. Electronic structure

1. Theoretical approaches and coordination defects

The electronic structure of stoichiometric H-freea-Si12xNx anda-Si12xNx :H alloys has been studied exten-sively from a theoretical and from an experimental point ofview.123–125The only calculation on the electronic structureof germanium nitride has been done by Makleret al.126

Compared toa-Si12xNx , only a small experimental researcheffort has been made on the electronic structure ofa-Ge12xNx alloys.127–129The available results indicate thatthe gross features of the electronic structure of both elemen-tal semiconductor nitrides is very much alike, as expectedfrom the similar valence configuration of Si and Ge. How-ever, some differences appear concerning the band gap wid-ening mechanism with increasing nitrogen, the nature andthe energy level of defects, and the structure of the valenceband.

Let us first summarize the electronic structure ofa-Si12xNx .123–125Tight binding calculations ofc-Si3N4 in-dicate that the top of the valence band~VB! consists of N2pp nonbonding states, the conduction band~CB! beinggiven by Si–N anti-bonding orbitals.123,126The main valenceband consists of three peaks separated by a gap from a deepfourth peak of Ns states. The three main peaks of the VB areassigned~increasing binding energy! to N pp, a mixture ofN p and Sip bonding states, and Np and Sis states, re-spectively. The results of calculations are in good overallagreement with the photoemisison data of Ka¨rcher et al.124

Theory predicts that the band gap ofa-Si12xNx should in-crease slowly with increasing N up tox50.42, with the bandedges opening almost symmetrically about mid-gap, andthen rapidly untilx50.57 ~stoichiometric composition!. Hy-drogenation is found effective in widening the band gap inSi-rich alloys (x,0.42) only, i.e., the N concentration atwhich Si–Si bonds fail to form continuous percolation paths

across the network. Ina-Si3N4, the silicon dangling bond(SiBD) gives rise to states 3 eV above the VB edge, whereasthe nitrogen dangling bond (NDB) produces a strongly local-ized state close to the VB edge. Large N concentrations in-crease the disorder and broaden the VB tail, which is due tovariations in second neighbor N–N interactions.123

There is a clear difference between the behavior of de-fects of Si-rich and of N-richa-Si12xNx :H alloys. In Si-richmaterial the neutral SiDB produces a half-filled level nearmid-gap, as ina-Si:H, whereas the energy level of NDB isembedded deep in the VB. Inc- or a-Si3N4, however, ahighly localizedp-like NDB appears just above the VB maxi-mum. The roughlysp3 hybridized neutral dangling bond of atrivalent Si site~Si surrounded by 3 N, calledK0 center! islocated around 3 eV above the VB maximum.130 The simul-taneous presence of both types of DBs should cause a chargetransfer from SiDB to NDB giving: K01N05K11N2. Atroom temperature a smallK0 signal is measured by ESRsuggesting that SiDBs are in excess. In Si-rich alloys(x.0.47) theK0 center is more stable in its diamagneticconfiguration and there is strong experimental evidence infavor of the K0 center having a negative correlationenergy.131

Makler et al.126 calculated the electronic structure ofa-Ge12xNx :H alloys, in particular the structure of the VB,the band-gap widening mechanism for increasing N concen-trations, and the energy levels of gap states due to GeDB andto NDB. Besides the expected smaller band gap for similar Nconcentration, the calculations ona-Ge12xNx :H indicate nomajor qualitative differences with respect to the electronicstructure ofa-Si12xNx :H. Tight binding calculations126 pre-dict a preferential attachment of H to N ina-Ge12xNx :Halloys, in agreement with experimental evidence.132,133 Re-garding the energy levels of coordination defects Makleret al.126 find the GeDB at the center of thea-Ge:H gap, whereit remains for N concentration increasing up tox50.4.Above x50.4 the GeDB disappears from the band gap. Thecalculations do not find NDB states in the gap ofa-Ge12xNx :H.126

ESR is a powerful tool to obtain structural informationon coordination defects. The strength of the ESR absorptionband is proportional to the density of paramagnetic electronsand the ESR signal structure gives information about localbonding. The most useful information comes from theg fac-tor of the paramagnetic center and the hyperfine interaction,which depend on the nuclear magnetic properties of the at-oms around the center. The neutral~paramagnetic! danglingbonds in hydrogenated amorphous Si and Ge haveg52.0055 andg;2.019 signatures, respectively.134,135

The nature of coordination defects in H-free and hydro-genateda-Ge12xNx alloys has been studied by ESR andlight-induced ESR in rf magnetron sputtered samples.136–138

It has been found that both theg value and the linewidth ofthe GeDB decrease with increasing N concentration, as shownin Fig. 7. The effects have been attributed to the large an-isotropy of theg value for Ge in H-free samples,137 and inhydrogenated samples to the existence of separated Ge-richand N-rich phases in films with largex.138 Note that this lastinterpretation is at odds with EXAFS data which indicate a



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chemically ordered~stricto sensu! material.121 The origin ofthe satellites in the resonance spectra is due to hyperfineinteraction with N nuclei, as confirmed by the different hy-perfine structure of GeN samples made using15N ~nuclearspin I 51/2! and 14N (I 51).137 These preliminary ESR re-ports indicate the need for more detailed investigations, inorder to gain a deeper understanding of coordination defectsof a-Ge12xNx alloys.

Let us consider here now the main differences betweenthe electronic structure of both elemental semiconductor ni-trides.

~a! Makler et al.126 find that the band gapEg ofa-Ge12xNx remains unaltered up to N concentrations ofx;0.5, and that forx.0.5, Eg grows quickly to reach thevalue Eg54.7 eV (x50.57). This is qualitatively differentfrom Robertson’s123 slow linear increase ofEg in a-Si12xNx

for 0,x,0.4, and a faster increase rate at largerx values.Note that neither fora-Ge12xNx , nor in the case ofa-Si12xNx samples does the calculated gap opening rate cor-respond to the variations measured by photoemission.124,129

~b! Nitrogen is a group V element and, as an impurity, itmay be an active dopant in thea-Si:H anda-Ge:H networks.Makler et al.126 find that a N impurity is a deep donor ina-Si:H, the anti-bonding state laying at 0.45 eV below thebottom of the CB. The calculation performed for N impurityin the a-Ge:H network gives a too small energy differencebetween the bottom of the CB and the N donor level and theauthors cannot conclude whether N impurity is a shallowdonor or not. Robertson’s130 calculation locates the N4

1 statejust below the CB minimum. The calculations, however, donot agree with the experimental results on N-dopeda-Si:Handa-Ge:H.139–141We discuss the problem in more detail inthe section on electronic transport and N doping.

2. Experimental reports

The structure of the VB and its evolution as a function ofN content (0,x,0.36) was investigated in H-freea-Ge12xNx films129 by ultraviolet photoelectron spectros-copy ~UPS!.142 Photoelectrons were analyzed with an energyresolution of 0.2 eV, the light source being HeI ~21.2 eV!and HeII ~40.8 eV! photons. Figure 8 shows the photoelec-

tron spectra ofa-Ge anda-Ge12xNx ~x50.11, 0.31 and0.36! samples. A recession of the VB maximum is apparentin Fig. 8 ~left!. Two N-related features can be recognizedclearly in the figure. The peak at binding energy~BE!;5.2 eV is close to the value determined for the nonbondingN 2pp lone pair observed fora-Si12xNx alloys,124 and hasprobably the same origin. As discussed above, two additionalpeaks have been reported fora-Si12xNx , one with a BE;7.5 eV and the other at energies ranging between 11.4 and12.4 eV, which were attributed to bonding N 2px,y orbitalshaving some contributions from Si 3s and 3p orbitals.124

Following this interpretation the feature at BE;10.9 eV inFig. 8 is probably due to bonding N 2px,y states with contri-butions from either Ge 4s or Ge 4p states. Comediet al.,129

however, do not provide a conclusive interpretation of thisexperimental finding which does not correspond to the threeVB peaks predicted theoretically by Makleret al.126

The Fermi energy and the VB maximum have beenfound to experience129 a small upward shift in energy withincreasingx up to x;0.25 ~see Fig. 9!, a behavior differentfrom that reported fora-Si12xNx ,124 in which the VB maxi-mum recedes linearly with increasingx. The reason for thisdifference is not yet understood. The sudden asymmetricalwidening of the band gap ina-Ge12xNx samples forx.0.22 deduced from Fig. 9 is similar to that reported forsilicon nitride compounds.124,125In both elemental semicon-ductor nitrides the band-gap opening starts at a N concentra-tion when N 2pp states dominate the VB maximum. It wasfound experimentally that, similar toa-Si:H, hydrogenationdoes not control the optical gap ofa-Ge12xNx :H alloysabove a well defined N concentration~see Fig. 9!. Note,however, that the N concentration at which the sudden open-ing of the band gap occurs is significantly lower in GeN thanin SiN alloys, i.e.,x;0.25 andx;0.5, respectively. Thisdifferent behavior, related to the nature of the dominant or-

FIG. 7. g value and linewidth of the resonant ESR broad line ofa-Ge12xNx

alloys as a function of increasing N2 partial pressure in the deposition cham-ber ~Ref. 137!.

FIG. 8. Ultraviolet photoelectron spectra of purea-Ge anda-Ge12xNx

samples~x50.11, 0.31, and 0.36!. The spectra have been obtained using HeI ~21.2 eV! and HeII ~40.8 eV! photons. A recession of the VB maximumand new electronic states are apparent for samples with large N contents~Ref. 129!.



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bital at the VB maximum, has not yet been explained. Tosome extent, it might be the consequence of different meth-ods used to estimate the concentration of nitrogen.

C. Structural studies by optical techniques

1. Infrared spectroscopy

The analysis of lattice or network vibration modes hasproved to be an efficient tool in structural studies.The structure of a-Ge12xNx and a-Ge12xNx :H alloyshas been studied with infrared~IR! and Ramanspectroscopies.24–26,39,40,143–146Next we discuss the IR spec-tra of H-freea-Ge12xNx films of varying composition.

Crystalline Si and Ge lattices have no first-order IR ab-sorption. In amorphous Si and Ge, however, the lack of long-range order relaxes the crystal momentum and the symmetryrules prohibiting first-order absorption and, in principle, allvibration modes can contribute to the first-order absorptionprocess. In other words, within a disordered network all vi-bration modes can be active in the IR. In the case understudy, the introduction of N atoms into thea-Ge networkenhances the disorder and induces a considerable chargetransfer from Ge to N producing stronger IR absorptionbands related to the Ge–N bond structure.

The group formed by a planar bonded N atom and itsthree Ge neighbors is the skeletal Ge3N group. Its normalvibration modes are:147 a breathing mode, an out-of-planestretching mode, the symmetric and asymmetric in-planestretching mode, and an in-plane bending mode. The asym-metric stretching vibration involves the displacement of theN atom and of the three Ge neighbors and is strongly IRactive. Its strength has been found to be proportional to Nconcentration.39,148Figure 10 shows the IR absorption spec-tra of six a-GeN samples as a function of wave number inthe 200– 1600 cm21 range. The main absorption band peak-ing at ;700 cm21 is the in-plane asymmetric stretchingvibration.132,133 The analysis of this absorption band as afunction of N concentration indicates that:

~a! The integrated absorptionA(vst)5*st@a(v)/v#dvis proportional to the N concentration, i.e.,@N#5KGe–N

A(vst), with KGe–N5531018 atoms cm22. A deuteron-induced nuclear reaction@14N(d,p)15N# was used by thepresent authors148 to determine the absolute N concentrationin a series ofa-Ge12xNx samples. The method allows thedetermination of N concentrations as low as1019 atoms cm23 in samples of typical thickness 1024 cm~see Fig. 11!. The existence of a calibration constant facili-tates the rapid determination of alloy composition from op-tical measurements.

~b! The presence of back impurity atoms of a differentelectronegativity induces a shift of the peak energy of dipolevibrations.149,150It is worth mentioning here that the absenceof back atom~s! @or dangling bond~s!# also influences thecharge distribution around the Ge–N dipole. Ina-Ge12xNx

alloys the frequency shifts are associated with changes in theGe–N interatomic distance and their analysis sheds light onthe structural modifications produced by the incorporation ofvarying amounts of N, or of other light impurities, like hy-

FIG. 9. ~a! Energy position of the valence band maximumEV , the FermienergyEF and the conduction band edgeEC ~resulting from the addition ofE04 to the VBM! as a function of N content ina-Ge12xNx ~see Ref. 129!.~b! Same kind of representation fora-Si12xNx ~see Ref. 124!.

FIG. 10. Absorption IR bands vs wave number ofa-Ge12xNx samples. Notethe appearance of new absorption bands as the N content in the film in-creases~see Ref. 39!.

FIG. 11. Integrated absorption of the main in-plane Ge–N stretching vibra-tion mode vs N content fora-Ge12xNx samples~see Ref. 39!. The slope ofthe straight-line fit isKGe–N55.031018 cm22. The filled circles indicate theintegrated area of the absorption bands centered at about 690 cm21. Theopen triangles represent the integrated area of all stretching vibrations~690,870, and 1100 cm21!.



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drogen or deuterium~to be discussed in a coming section!.Inductive effects due to increasing nitrogen are detected inthe main stretching vibration mode. The appearance of newabsorption bands at;870 and;1100 cm21 ~see Fig. 10!have been ascribed to Ge–N stretching vibrations in an en-vironment including N atoms and dangling bonds in differentbonding configurations. Figure 12 sketches the different con-figurations which, according to our interpretation,39 give riseto the three bands just discussed.

~c! The absorption features appearing at;450 and300 cm21 ~see Fig. 10! have been associated with the sym-metric stretching and breathing modes involving the disor-dereda-Ge3N skeletal group, respectively.

The above mentioned vibration modes were identifiedcomparing the IR features ofa-Ge12xNx alloys with those ofwell known gaseous compounds and investigating theirstrength and energy position dependence on alloying. Acompound that allows us to establish a useful analogy withthe a-Ge3N skeleton of GeN alloys is trigermylamine~GeH3!3N, an extremely unstable and reactive gaseousspecies.151,152The structure of~GeH3!3N is shown in Fig. 13,which also displays its main vibration modes and the corre-sponding energy peaks. Note that, besides the absence of aconnective network, the main difference between the Ge3Nskeleton ofa-Ge12xNx films and~GeH3!3N is the presenceof H atoms as Ge bond terminators. Briefly, the vibrationproperties of trigermylamine indicate that:~i! the Ge atomsare probably noncoplanar with N, a valence force fieldcalculation151 giving a bond angle of 116° at N.~Note thatthe vibrations of the pyramidal site for a threefold-coordinated N are qualitatively similar to the vibrations of Nin a planar configuration, both having four characteristic vi-bration modes involving the displacement of N and its threeGe neighbors!; ~ii ! there are two IR absorption bands in~GeH3!3N which do not move on deuteration:151,152 a verystrong and broad band at;850 cm21, attributed to the asym-metric stretching mode of the Ge3N structure@Fig. 13~b!#and a weak band at;370 cm21 originating from the sym-metric stretching mode of trigermylamine@Fig. 13~a!#.

The vibration frequencies of the skeletal Ge3N structurein solid a-GeN and gaseous~GeH3!3N differ because the

trigermylamine molecule possesses three H atoms bonded toeach Ge. A charge redistribution around the Ge atoms isexpected because H is more electronegative than Ge.153 As aconsequence, the stretching vibrations should be more ener-getic in the molecule than in an isolated Ge–N dipole. As theN concentration increases ina-Ge12xNx , however, the pres-ence of N back atoms leads to a charge redistribution aroundGe atoms which goes in the direction of the trigermylaminestructure. This fact explains the reasonable agreement ofpeak energies of the asymmetric in-plane stretching vibra-tions in both the gaseous and the solid phases~850 and870 cm21, respectively! of the Ge3N skeletal group.

The absorption band at;450 cm21 associated with thein-plane symmetric stretching mode ofa-Ge12xNx is foundat 370 cm21 in trigermylamine. The energy difference mayoriginate from the nonplanar configuration in~GeH3!3Nwhich involves the out-of-plane motion of N. Finally, thedependence of the absorption strength of the 300 cm21 fea-ture on N concentration ina-Ge12xNx lead us to attribute the300 cm21 vibration to a disorder-induced breathing mode ofthe a-Ge network, absent in trigermylamine.

A linear relationship between absorption strength andcomposition with no extra bands at high alloy ratios, as wellas a constant band shape, means that just the density of di-pole oscillators is varying with composition. Similarly, anequal shift of the absorption peak energy with compositionfor two different absorption bands, or a linear relationshipbetween maximum absorption coefficients for a couple ofbands, indicate that both have a common origin. Figure 14shows the dependence of the absorption strength on alloyingof the asymmetric in-plane Ge–N stretching mode(;700 cm21), the Ge–N symmetric stretching mode(;450 cm21) and the Ge–N breathing (;300 cm21)modes. As can be seen, the overall absorption increases onalloying, an indication that N is associated with the threemodes.

FIG. 12. Different chemical environments of Ge atoms suggested to inducefrequency shifts of the main Ge–N dipole vibration~see Ref. 39!.

FIG. 13. Simplified skeletal representation for the trigermylamine~GeH3!3Nmolecule~top!. The possible vibrations of the Ge3N group in thea-Ge hostare: ~a! symmetric in-plane stretching vibration mode (;370 cm21); ~b!asymmetricin-planestretching vibration mode (;850 cm21); ~c! symmet-ric ~pyramidal configuration! stretching mode; and~d! symmetric nearin-plane~pyramidal configuration! stretching mode. Vibrations~c! and~d! arecharacteristic of gaseous trigermylamine~see Refs. 39 and 152!.



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All the inductive effects measured in IR absorptionoriginate from charge transferred from Ge to N atoms. X-rayphotoelectron spectroscopy~XPS! measurements performedon the samea-Ge12xNx samples128,154corroborate the over-all picture of charge transfer deduced from IR spectra.

2. Raman spectroscopy

Raman spectroscopy is a well developed probe tech-nique to study the structural features ofa-semiconductor al-loys. In a Raman process the incident light frequency isshifted by an amount equal to the natural frequency of thevibration mode involved in the inelastic scattering. The ab-sence of long-range order favors the observation of all vibra-tions usually forbidden in crystalline lattices by selectionrules. As the short-range order ina-semiconductors is similarto that in the crystalline parents, the phonon density of statesis expected to be similar to that of the crystalline phase withsome broadening due to disorder in bond length and bondangle.155 In addition to structural information, Raman scat-tering measurements ina-semiconductors can give insightinto the bond type and the nature of disorder. A Raman studyof the structural changes produced by different N concentra-tions in thea-Ge network was recently reported.145 Figure 15shows the reduced Raman spectra of purea-Ge and of twoa-Ge12xNx samples. In the case of purea-Ge the spectrumis dominated by the contributions of the transverse acoustic~TA!-like mode at;80 cm21 and from the transverse optic~TO!-like mode at;270 cm21. The incorporation of N up toaround 1022 atoms cm23 does not affect essentially the shapeof the spectrum, except for a slight broadening of the TO-like mode, related to increased topological disorder. Forlarger N concentrations, however, radical changes are de-tected. The TA- and the TO-like modes broaden consider-ably and are also shifted to higher energies. In addition, abroad structure appears at;700 cm21 corresponding to thein-plane asymmetric stretching vibration referred to in theprevious paragraph. Xuet al.26 reported Raman spectra ofa-Ge12xNx :H films deposited by plasma-enhanced chemicalvapor deposition~PECVD!. These authors found that whenx

reaches 0.37, i.e., just above the maximum N concentrationof the samples reported in Ref. 145, the Ge–Ge TO modedisappears while the Ge–N absorption band becomes verystrong.

There is wide agreement on the fact that the full width athalf maximum ~FWHM! height of the TO-like vibrationmode (GTO) and theI TA /I TO amplitude ratio are useful indi-cators of structural disorder. Changes of the measuredGTO

correspond to modifications of the short-range order, particu-larly of the bond angle distribution.156 In the study of Zanattaet al.,145 the laser energy~2.6 eV! used to record the Ramanspectra is coupled mainly to valence and conduction statesassociated with Ge–Ge bonds, the Ge–N bonds having amuch higher energy separation~the band gap widens as@N#increases, as discussed in the following section!. Thus thechanges in frequency and bandwidth of the TO-like modereflect basically the distortion of the Ge–Ge bonds with in-creasing N concentration. For small nitrogen contents,@N#,1022 atoms cm23, the symmetry of the bonds is deter-mined mostly by thesp3-like character of the Ge–Ge bond.The changes of GTO become more important for@N#.1022 atoms cm23 and are attributed to the drastic bondangle distortions due to changes in the character of the domi-nant bond, from essentiallysp3-like ~Ge–Ge! to some com-bination betweensp3- andsp2- ~or p-! like Ge–N orbitals.The structural changes in thea-Ge12xNx network as the Ncontent increases were also probed by photothermal deflec-tion spectroscopy~PDS!, which indicates a linear relation-ship between electronic and structural disorder.~The pointwill be discussed in detail in Sec. IV.! The interpretation hasbeen confirmed also by photoelectron emission studies madeon the same samples.129

Figure 16 shows the shifts of the IR vibration and Ra-man scattering signal induced by nitrogen. In this figure,bothv IR andvRamanare represented as a function of theE04

optical gap, which scales with N content. The changes can beunderstood on the basis of a lightening and a stiffening of thenetwork. Roughly, v IR and vRaman are proportional to

FIG. 14. Absorption coefficient at peak energy vs N concentration in thesolid phase of the stretching and breathing modes ofa-Ge12xNx alloys.Note the dependence of all vibration modes on nitrogen concentration~seeRef. 39!.

FIG. 15. Reduced Raman spectra of somea-Ge~N! samples with increasingN concentrations. The features corresponding to the TA- (80 cm21) andTO-like (270 cm21) vibration modes are indicated~see Ref. 145!.



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(k/m)1/2, wherek is an effective force constant associatedwith Ge–Ge~and Ge–N! bonds, andm is the correspondingreduced mass. As N is added to thea-Ge host, heavy Geatoms are substituted by light N. In addition, Ge–N bondsare stronger than Ge–Ge bonds~EGe–N;2.65 eV andEGe–Ge;1.94 eV!,1 producing a stiffening of the network.Note that the interpretation is consistent with results ofEXAFS,121 which indicate that root mean square N bondangle fluctuations are smaller than the minimum fluctuationfor bond angles on Ge, an indication that Ge–N bonds aremore rigid than Ge–Ge.

D. Hydrogenation, structure and stability of a-GeNfilms

Hydrogen atoms reduce significantly the density of dan-gling bonds of group IVa-semiconductors and, thus, play afundamental role in their structural and electronic properties.Another effect of H is to break weak bonds during the depo-sition process. Thus, from the structural point of view, theinclusion of monovalent H reduces the coordination of Si orGe and helps in relaxing the over-coordinated network. To alesser extent, a similar relaxation process is induced by ni-trogen, N having a coordination number smaller than Ge.Hydrogenation ofa-Ge12xNx has several consequences, ex-tending from growth rate to structure and electronic trans-port. Of particular relevance to the present review is thebonding of H to N40,132,133and the likely influence of N–Hbonds on deposition rate, network relaxation, and electronictransport. The absorption bands associated with hydrogenand deuterium bonded to Ge and N have been identified.40

These results suggest that ‘‘preferential attachment’’157 of Hto N132,133 takes place when H atoms are adsorbed on thegrowing surface, an indication that the different bonding en-ergies of N–H~4.5 eV! and Ge–H~3.8 eV! measured inmolecules hold qualitatively in thea-network.

X-ray inspection shows that as-deposited hydrogenatedGe12xNx films are amorphous. Honmaet al.133 investigatedthe structural stability ofa-Ge12xNx :H films by studyingcrystallization processes by differential thermal analysis~DTA! and differential thermal gravimetry~DTG!. Heatingup the films under controlled conditions induces the crystal-lization of the alloy, revealed by the Ge~111! diffractionpeak. The crystallization temperatureTc increases with Nconcentration and varies from;635 K for a-Ge:H up toabove 925 K for a film containing;40 at. %N. There is aweight decrease on heating, attributed to the out-diffusion ofH and~probably! Ar, always present in rf-sputtered material.Similar qualitative results were obtained by Xuet al.26 ona-Ge12xNx :H films deposited by PACVD using GeH4 andNH3 gas mixtures. Thermal annealing of the films up to 875K shows that the inclusion of N improves the structural sta-bility and increases the crystallization temperature. For filmscorresponding tox50.44 no crystallization signal can be de-tected by Raman spectroscopy after 30 min anneal at 875K.26 Summarizing, nitrogen improves considerably the ther-mal stability ofa-Ge12xNx films.

The mechanical properties ofa-Ge12xNx :H films, in-cluding stress, thermal expansion, and elasticity as a functionof temperature and N concentration, have been investigatedby Marqueset al.158 Figure 17~a! indicates that as-depositedfilms have compressive stress ranging from approximately;1.0 to ;7.0 kbar as theE04 optical gap increases fromabout 1.0 to 3.0 eV~@N# ;30 at. %, by NRA!. The compres-sive stress, as found in sputtereda-Ge12xNx :H films, is as-sociated with the presence of impurities, to hydrogenation,and to Ar inclusion. It gives stable and compact films, inagreement with the studies of DTA.133 Tensile stress, on the

FIG. 16. Frequency shifts of the Ge–Nin-planeasymmetric stretchingv IR

and of the TO-like (vRaman) vibration modes as a function of the opticalparameterE04 ~which depends on N concentration! ~see Ref. 145!. Theasymmetric stretching vibration and TO-like modes display a similar behav-ior. Also represented are theoretical data from Barrioet al. ~see Ref. 146!showing the expected increase of the TO-like mode frequency. Large Ncontents are responsible for a lightening and a stiffening of the network. Thedashed straight lines are linear regression fits of experimental data.

FIG. 17. ~a! Intrinsic compressive stress vs optical gapE04 ~}N content! ofa-Ge12xNx :H films deposited by rf sputtering, and~b! thermal expansioncoefficient vs optical gapE04 of some of thea-Ge12xNx :H films of ~a! ~seeRef. 158!.



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other hand, is characteristic of unstable films with structuraldefects, such as columnar structure, voids, etc. The thermalexpansion coefficient decreases with nitrogen@see Fig.17~b!#, a consequence of the structural changes of the alloy,which evolves from a diamondlike toward a hexagonal struc-ture ~phenacite! of smaller overall coordination.

E. The Ge–N bond: randomness, charge transfer andelectronegativity

Charge transfer is an important effect to be considered inSiN and GeN alloys. Electronegativity was originally pro-posed by Pauling, as a parameter that would allow predictionof the approximate polarity of a partially covalent bond.153

The organic chemistry of Ge shows many similarities to thatof Si, in some chemical reactions Ge appearing more elec-tronegative than Si. Nitrogen has an electron affinity and anelectronegativity greater than Si and Ge.153 Therefore a Natom in the Si~Ge! network gets an extra electronic chargecoming from Si~Ge! neighbors, the charge of which will beless than 4 by an amount which depends on the number of Nneighbors. Since Pauling’s original proposal a number ofnew definitions of electronegativity and the ways to evaluatethem have appeared. Sanderson’s electronegativity scale,159

which besides the atom’s ‘‘ability’’ to hold their own elec-trons tightly together taking into account the electronic con-figuration, gives: xH53.55; xC53.79; xN54.49;xSi52.84; xGe53.59. These figures indicate that Ge is theelement having the closest electronegativity to that of H. Asa consequence, the Ge–H bond is less polar than the Si–Hbond, which is consistent with the experimentally foundpreferential attachment of H to Si ina-SiGe:H alloys.157

X-ray photoelectron spectroscopy~XPS! is a very sensi-tive probe of the atomic environment of a selected element,the binding energy of its core electrons depending on theatomic surroundings. The chemical shifts induced by N inthe Ge 3d core levels have been studied by XPS and byx-ray excited Auger electron spectroscopy~XAES!.128,154

The surface analysis techniques probe the outermost atomiclayers and provide useful information on the atomic environ-ment of a given species and its influence on the macroscopicproperties.142 Consider Fig. 18 where the Ge 3d core levelsignals of H-freea-Ge12xNx samples containing differentamounts of N are represented.127 On these ex situ samples nochemical or physical cleaning processes were used in orderto avoid possible preferential sputter etching of species. Thesignal associated with a surface oxide Ge–O, apparent inFig. 18, does not shift as the N concentration increases in thealloy.

The bands associated with the presence of N are a maxi-mum of five, corresponding to Ge units having 0, 1, 2, 3, and4 bonded N atoms. The additivity of the chemical shifts hasbeen assumed, the line shapes being the same for each com-ponent. The fitting procedure is illustrated in Fig. 19 fora-Ge anda-Ge0.89N0.11. The present authors127 suggestedthat it was possible to analyze the subsurface information ofthe Ge 3d despite the presence of a thin oxide surface layer.Similar to the case of the parent SiN alloys160,161the analysisof the Ge–N bond distribution indicates a fair agreementwith the theoretical prediction of the random bonding

FIG. 18. Ge 3d core level signals of samples containing increasing amountsof nitrogenx, as indicated in the figure. The centroid of the main contribu-tion of the Ge–Ge and Ge–N signals, as determined from fitting procedures,is indicated by vertical lines. Note that, whereas the Ge 3d core levelsassociated with Ge–N bonds experience an apparent chemical shift~as Nincreases!, the signal associated with Ge–O bonds of the surface oxide doesnot shift significantly~see Ref. 127!.

FIG. 19. Experimental photoelectron spectra of the Ge 3d core level~pluses! and the corresponding fitting with Gaussian curves representing thedifferent Ge–N components~solid lines! for a couple of samples~see Ref.127!. ~a! Purea-Ge with a native oxide overlay. The deconvolution repre-sents the Ge 3d5/2 (EB;29.2 eV) and the Ge 3d3/2 (EB;29.8 eV) corelevels corresponding to bulk Ge–Ge bonds, and the contribution of nativeGe–O bonds (EB

Ox;33 eV). ~b! The deconvolution of thea-Ge0.89N0.11

sample displays the main Ge–N and Ge–O bonding components. Neither in~a! nor in ~b! was the core level signal associated with the Ge oxide de-convoluted in its spin orbit split components.



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model.162 A chemical shift per bondDEb bond21;0.560.1 eV for the Ge–N bond is found, a value intermediatebetween the shifts induced by H and O on Ge. A similar UPSstudy by Comediet al.129 gives a slightly smaller chemicalshift per bondDEb bond21;0.3060.08 eV. These valueshave to be compared with the binding energy shift of;0.6 eV measured per Si–N bond ina-Si12xNx

alloys.124,163

Figure 20 compares the position and shape of the Si 2pand Ge 3d core levels ofa-SiN anda-GeN alloys as a func-tion of nitrogen concentration. The shifts reported in the fig-ure correspond to the centroid of the core level signal andscale in a way similar for both Si and Ge alloys. Figure 20also displays the FWHM of the XPS signal of the Si 2p andGe 3d core level signals. Note the asymmetric bell-likeshape in both cases, which merely reflects the Si–N andGe–N bonding distribution. However, two differences ap-pear:~i! the maximum FWHM is smaller fora-GeN than fora-SiN alloys and,~ii ! this maximum occurs at different Nconcentrations. A different FWHM maximum may be theconsequence of the different chemical shift per bond of Si–Nand Ge–N bonds, but why the maximum FWHM occurs atdifferent composition is not yet understood.

Using experimental data of XPS and XAES core levelshifts, Zanattaet al.128 determined the modified Auger pa-rameter shiftDa8,164–166 which is exempt from problemsinherent to the interpretation of XPS and XAES shifts.Da8values were used to estimate the charge transferred on alloy-ing (DnGe). A proportionality betweenDa8 and DnGe wasfound, reminiscent of the idea of the local electronegativity

concept.159 Among the various electronegativity definitionsand scales being reported in the literature, the one promotedby Sanderson,159 which has been successful in explaining theIR transmission spectra ofa-Si alloys,150 was retained. Thepreceding considerations suggest the possible use ofDa8 todefine an electronegativity scale. In order to test this idea,Zanattaet al.128 plot ~as in Fig. 21! the XPS shift versusDa8of somea-Si- and a-Ge-based alloys. The resulting linearcorrelation indicates first, the possibility of calibrating XPSshifts in terms of partial charges and second, thatDa8 mightfurnish a reasonable scale of electronegativities in such non-stoichiometric compounds.

Finally, an interesting correlation is found consideringthe chemical shift induced in the elemental semiconductorcore levels by different impurities. Figure 22 displays thechemical shift per bond (DEB bond21) for Si and Gehosts154 as a function of the electronegativity of the foreignatom~Sanderson’s159 scale!. Note the important direct corre-spondence found between transferred charge and electrone-gativity for a collection of quite different systems and from

FIG. 20. Ge 3d and Si 2p core level binding energies and full width at halfmaximum~FWHM! height of the XPS signal as a function of N concentra-tion in a-Si12xNx anda-Ge12xNx samples. The data of FWHM correspondto the whole signal.~Open squares! Ref. 127; ~filled triangles! Ref. 124;~open triangles! Ref. 161. Note that bothEB and FWHM exhibit similarbehavior in Ge and Si hosts, except for the values of the maximum FWHM,and the N concentration corresponding to this maximum~from Ref. 127!.

FIG. 21. XPS shift of Si 2p and Ge 3d core levels as a function of thecorresponding modified Auger parameterDa8. The possibility of calibrat-ing XPS shifts in terms of the modified Auger parameter is apparent fromthe figure~see Ref. 128!.

FIG. 22. Chemical shift per bond (DEB bond21) as a function of the foreignatom’s electronegativity~Sanderson’s scale!. The error bars are associatedwith some dispersion of reported data and impurity content. The straightlines are linear regression fits of the available data.~a! SiH, see Ref. 252,~b!SiN, see Ref. 124,~c! SiN, see Ref. 253,~d! SiN, see Ref. 161,~e! SiO, seeRef. 173,~f! SiO, see Ref. 254,~g! SiF, see Ref. 255,~h! GeH, see Ref. 256,~i! GeN, see Ref. 154,~j! GeO, see Ref. 257,~k! GeO ~see Ref. 258!.



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data of different authors. The figure suggests the possibilityof a new electronegativity scale, with electronegativityvalues154 of Si and Ge~corresponding toDEB bond2150! ofxSi;2.6 andxGe;3.1 in agreement with the idea that, insome cases, Ge atoms are more electronegative than Si.

IV. OPTICAL PROPERTIES OF a-Ge ALLOYS

A. Optical properties

The introduction of N atoms in a group IVa-networkleads to a new class of materials having quite new optical,electronic and structural properties with potential technologi-cal applications. Large impurity contents~alloying range! in-duce important changes of the optical, electronic and struc-tural properties of group IV amorphous semiconductors.Hydrogen, for example, relaxes the network and provokes aband-gap widening, a consequence of the recession of thetop of the VB.167 On the other hand, small concentrations ofcolumn III and V impurities~doping range,1 at %! provokeimportant changes in the electronic properties of group IVa-semiconductors. Figure 23 shows the transmission spectrain the near infrared-visible~NIR-VIS! and mid-infrared~mid-IR! energy ranges of somea-GeN thin films with dif-ferent N and H contents. As seen increasing amounts of N inthe a-Ge network widen the pseudo-gap. The effects of Ndoping on the transport properties ofa-Ge:H anda-Si:H arediscussed in Sec. V.

B. Optical absorption in amorphous semiconductors

The breakdown of the electronk wave vector selectionrule for optical transitions in amorphous solids entails anoptical response which is free of fine features, like van Hovesingularities and well-defined band edges. Moreover, withdisorder the abrupt band edges of crystals broaden into tails

of localized states extending into the pseudo-gap. The widthof the tails determines the most important optoelectroniccharacteristics of ana-semiconductor and depends on thedegree of disorder as well as on the bonding character of themaster states.

The optical absorption ina-semiconductors is usuallyreported in three different energy regions:~i! a high photonenergy region involving transitions between extended elec-tron states,~ii ! an intermediate photon energy region corre-sponding to transitions from or to localized states, and ex-hibiting an exponential-like dependence on photon energy.@The tailing of these localized states~or band edges! is char-acterized by a parameter called the Urbach energyE0 ,168

usually determined from PDS data169#, and ~iii ! a sub-gapabsorption region which is associated with:~1! transitionsbetween deep defect states and the VB or CB,~2! local vi-brations involving lighter atoms; and~3! the resonant modesof the host network.

1. Optical band gap (E 04 and E Tauc)

The presence of localized states at energies between va-lence and conduction bands, makes the optical gapEg anill-defined parameter ina-semiconductors and, consequently,several ways are currently in use to defineEg . The simplestone is to considerEg as the energy corresponding to an ab-sorption coefficienta5104 cm21 (E04). Another usual defi-nition of the optical gap in ana-semiconductor is the so-called Tauc’s gapETauc,

170 given by the energy where the(anE)1/2 or (aE)1/2 vs E plot goes to zero, wheren is theindex of refraction andE the photon energy, i.e.:

~anE!1/25B1/2~E2ETauc!. ~5!

TheB1/2 Tauc parameter includes information on the convo-lution of the VB and CB states, and on the matrix element ofoptical transition, which reflects not only the relaxedk se-lection rule but also the disorder-induced spatial correlationof optical transitions between the VB and CB. Formally,B1/2

FIG. 23. Transmittance spectra of somea-Ge-based thin films in the NIR-VIS and mid-IR energy regions. Note that both the presence~and concen-tration! of hydrogen and nitrogen atoms drastically affect the optical trans-parency of these compounds.

FIG. 24. ~a! Photothermal deflection spectroscopy~PDS! curves of somea-Ge-based thin films~the same represented in Fig. 23!. The optical bandgap E04 and the Urbach energyE0 are indicated for each sample.~b!(anE)1/2 vs photon energy for a series ofa-Ge-based compounds.



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depends on the product of the oscillator strength of the opti-cal transition, the deformation potential, and the mean devia-tion of the atomic coordinates.171

Figures 24~a! and 24~b! display the optical gapsE04 andETauc of a-GeN compounds. The optical gap and the Urbachenergy depend on the presence~or not! of H as well as on N.The optical gapsE04 andETauc and theB1/2 Tauc parameterfor a-GeN samples with different N contents are representedin Fig. 25. As seen, important changes in the optical dataoccur for N concentrations higher than;1022 atoms cm23.The optical gap ofa-GeN alloys does not depend linearly onthe nitrogen concentration.

2. Band-gap widening

The mechanism of band-gap widening of group IVa-semiconductors due to alloying with lighter atoms~H, C,N, O, etc.! is well reported in the literature, withparticular emphasis on a-SiN- and a-SiO-basedalloys.20,147,124,130,172,173

The overall effect of N atoms on the optical band gap ofa-Ge:~H! and a-Si:~H! compounds are represented in Fig.26. According to the figure, there is a threshold N concen-tration ~which is different for Si and Gea-hosts! abovewhich important changes of the optical band gap occur. Thephenomenon was considered when discussing the electronicstructure ofa-Ge~N! alloys and is related to the substitutionof Si–Si ~or Ge–Ge! bonds by stronger Si–N~or Ge–N!bonds. As the N content increases, a N lone-pair band devel-ops and dominates the VB maximum as stoichiometry isapproached. The largest optical gap is obtained for the sto-ichiometric compound, as expected. On the contrary, forsmaller N contents Si–Si~or Ge–Ge! bonds dominate theVB maximum. A similar optical gap widening mechanism isfound in Si and Ge carbides and oxides.

C. Electronic versus structural disorder inamorphous semiconductors

All disordered solids display an exponential-like absorp-tion edge, i.e., lna}E ~photon energy!. The microscopic ori-gin of the optical absorption tail ina-semiconductors, usu-

ally associated with static and thermal structural disorder, isnot yet well understood despite extensive experimental174–179

and theoretical174,180–198investigations. Part of the problem,in fact, lies in the precise identification furnished by the dif-ferent experimental techniques used to probe it. Codyet al.176 have proposed a model fora-Si:H assuming theequivalence of thermal and structural disorder. For the al-loys, however, the characteristic energy of the exponentialabsorption edge increases to large values and additionalterms, which take into account the compositional disorderdue to alloying and hydrogenation, must be considered.199,200

Today, the absorption tail is considered to reflect the VB andCB edge joint density of states~DOS!. It represents then, thedisorder-induced broadening of the bands. The physical ori-gin of localized states in group IVa-semiconductors is at-tributed to the existence of strained bonds, always present inover-coordinated materials. According to the picture, thea-network must be strained locally in order to accommodatethe atoms in a nonperiodic array. To these strained bondscorrespond electronic states near the band edges, the densityof which fall off exponentially with energy. Optical transi-tions to and from localized tail states would originate theUrbach edge in the absorption spectrum. In spite of the ap-peal of this qualitative description, it is difficult to makepredictions based solely upon it.

1. Optical absorption and Urbach edge

In a recent publication the present authors201 discuss therelationship between the characteristic energy of the expo-nential absorption edgeE0 and the Tauc parameterB1/2 ofa-SiN- and a-GeN-based alloys~Fig. 27!. No correspon-dence was experimentally found betweenB1/2 and E0 forsmall N concentrations~less than a few at. %!. Note that, inall cases, chemical doping provokes a noticeable broadeningof the Urbach energy. In contrast, in the GeN and SiN alloyregime a correlation betweenB1/2 andE0 is found. This ex-perimental finding reflects the structural changes induced byelements of different atomic coordination, i.e., the characterof the bonding orbital at the top of the VB as the N contentincreases. A connection between the exponential absorption

FIG. 25. Optical band gaps~E04 and ETauc! and Tauc parameterB1/2 of aseries ofa-GeN alloys. The values corresponding to ana-Ge:H sample hasbeen also represented. Note the band-gap widening apparent for nitrogenconcentrations higher than;1022 atoms cm23 ~see Ref. 39!.

FIG. 26. Optical band gap (ETauc) vs the nitrogen content ofa-GeN:~H!~Refs. 39 and 108! and a-SiN:~H! ~Ref. 259! thin films. Qualitativelysamples exhibit the same behavior, i.e., a sharp optical band-gap widening atlarger N contents.



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tail and the short-range potential fluctuations produced bycharged impurities is suggested in Fig. 27. Photoemissionspectra, much more sensitive than optical techniques al-though not susceptible to selection rules, indicate that this isthe case: N concentrations higher than;10 at. % give riseto important changes in the VB extrema ofa-SiN anda-GeNcompounds corroborating the above picture.124,129

The energy region probed by photons around the absorp-tion edge corresponds to band-to-band and tail-to-band~orband-to-tail! electronic transitions in which the probability ofa photon of energyE being absorbed is proportional to theproduct of the initial~valence! and final~conduction! statesseparated by that energy. This is the joint DOS defined as

DOSJ~E!5E NV~E8!NC~E81E!dE8, ~6!

where NV and NC are the VB and CB density of states func-tions, respectively. The initial and final states, however, havea wave nature and transitions between them are possible onlyto the extent that the corresponding wave functions overlap.In crystals, the periodicity of the lattice is reflected in theperiodicity of the electron wave functions which leads to thequantization of momentum. Ina-materials, the phase of theelectron wave function is random, i.e., not well defined withrespect to atomic position. Momentum is not a good quan-tum number and allowed transitions occur between any twostates for which energy conservation applies and wave func-tion overlapping exists. The matrix elements for optical tran-sitions between any two such states is the dipole matrix ele-ment R. The imaginary part of the dielectric constant and,

consequently, the absorption coefficient depend on both thejoint DOS and the transition matrix elements:

e2~E!}R2~E!~DOSJ!, ~7!

and

a~E!}e2~E!/n, ~8!

where R2(E) is the normalized dipole matrix elementsquared averaged over all transitions separated byE, andnis the index of refraction. Jacksonet al.202 have reportedmeasurements ona-Si:H indicating thatR2(E);10 Å2 overthe photon energy range 1.5 eV,E,3.0 eV, and slightlygreater than 10 Å2 for 0.6 eV,E,1.5 eV. Strictly speaking,however, in the sub-gap regionR2(E) should depend on thenature of defects and, in particular, on the distribution ofcharge around it, which modifies the dipole moment of thebound electron. So the constancy ofR2(E) measured in theband tail energy range of intrinsica-Si:H is not expected tohold when different impurities are added or defects of a newkind created. Potential fluctuations produced by impuritieshaving an electronegativity different than the atoms of thehost network, like N ina-Si or a-Ge, for example, are ex-pected to induce an increased absorption via changes of thedipole moment of localized electrons and holes. The distanceto where the perturbation extends will depend on free carrierscreening. In singly doped materials the concentration of freecarriers increases rapidly with doping, the potential fluctua-tions arising from random charged defects will be effectively

FIG. 27. UrbachE0 energy vs theB1/2 Tauc parameter fora-Si- ~Ref. 260!anda-Ge-based~Ref. 201! thin films under different impurities concentra-tions. There is no clear relationship betweenE0 andB1/2 at low impuritiesconcentration. In contrast,E0 andB1/2 scale linearly in the alloying regime.

FIG. 28. Urbach energyE0 as a function of the FWHM of the TO-likevibration modeGTO ~Raman scattering! of N-baseda-Si ~Refs. 207–209!anda-Ge ~Ref. 145! thin films. Even for the lowest N concentration~dopingregime! electronic disorder is present in thesea-hosts. At higher N contentsE0 andGTO increase.



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screened, inducing important dipole moment changes. Incompensateda-semiconductors the carrier concentration isvery low and charges are very poorly screened. The compen-sated material is therefore anticipated to have long-range po-tential fluctuations. The optical transition is spatially local-ized and the absorption modified when the period of thefluctuation is comparable with the wave function extent ofthe electron and the hole. Much less impressive changes inthe absorption edge of compensateda-semiconductors areexpected, in agreement with experimental evidence.109

The picture suggests that the Urbach slope changes mea-sured in singly doped samples may originate from a broad-ening of the joint DOS~increased disorder! and/or fromchanges of the dipole matrix element of the optical transi-tion, induced by the dopant species.203

2. Structural disorder

The relationship between Raman linewidth and struc-tural disorder ina-semiconductors is abundantly reported inthe literature.156,204–206There is agreement that the FWHMof the TO-like vibration mode (GTO) is a useful indicator ofstructural disorder: the larger the disorder the larger theGTO

value. Roughly speaking, changes inGTO correspond tomodifications of the short-range order~SRO!, particularly ofthe bond angle distribution. Figure 28 shows the UrbachslopeEo as a function of the FWHM of the TO-like Ramanvibration modeGTO measured on different N-containinga-Siand a-Ge samples. SinceGTO is associated with the bondangle distribution~or SRO! a relationship is expected in thealloy regime, and experimentally found~Fig. 28!. At high NconcentrationsGTO increases indicating the occurrence of

structural and/or electronic disorder.145,207–209In the small Nconcentration range~doping regime!, however, botha-Si anda-Ge hosts do not present appreciable changes in their struc-ture, as probed by Raman spectroscopy.

The dependence of bothEo and GTO on the N concen-tration can be better appreciated in Fig. 29 where a series ofN-dopeda-Ge:H samples~using either N2, or NH3 as dopingsources! anda-GeN alloys are represented. According to thefigure, while Eo experiences important changes~as deter-mined from PDS! in the whole range of@N# the RamanGTO

signal is affected just for@N#>531021 atoms/cm3. Remem-ber that, at this concentration level, considerable structuraland optoelectronic modifications are detected. Another inter-esting feature of Fig. 29 relates to the use of N2 and NH3 asdoping sources. As can be seen, while the structure is essen-tially the same in both series of N-doped samples the use ofNH3 leads to a broader Urbach tail, a result linked to theexistence of NHn complexes which might originate fromlarge electrostatic fluctuations in thea-Ge network at voidsurfaces without affecting significantly the surroundings ofmost host atoms.

3. Composition, structural disorder and optical bandgap

Impurities may influence the properties of tetrahedrallycoordinateda-semiconductors in a number of ways. Figure30~a! shows the Urbach slope of N-containinga-Si:H filmsas a function of their optical gap. It can be seen that in bothdoping and alloying regimes, the characteristic energy of the

FIG. 29. UrbachE0 energy and TO-like FWHM Raman signalGTO as afunction of the nitrogen content ofa-Si- and a-Ge-based compounds.WhereasE0 ~electronic disorder! is largely influenced by N concentrationsas small as 1019 atoms cm23, theGTO values~structural disorder! do presentimportant changes only for@N#>531021 atoms cm23.

FIG. 30. Urbach energy (E0) vs optical band gap (Eg) of various amor-phous thin films. At low impurities content~doping regime! there is nochange inEg while the electronic disorder increases for increasing amountsof impurities. In the alloying regime bothE0 andEg scale according to theN concentration.



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Urbach tail increases with the introduction of nitrogen in thehost network. The same phenomenon is observed ina-Ge-based alloys@Fig. 30~b!#. The optical gap changesoriginate from long-range order modifications occurring onlyat large nitrogen concentrations. These optical gap changesare always accompanied by a broadening of the Urbach tail,the converse not being always true. The Urbach energy mayincrease considerably without any measurable changes of theoptical gap, as in the case of chemical doping for example.

A brief summary of the above mentioned optical prop-erties in a-semiconductors follows:~i! in group IVa-semiconductors, short-range order potential fluctuationsinduced by chemical doping seem to contribute to the broad-ening of the exponential absorption tails; and~ii ! in botha-Si:H anda-Ge:H samples no structural disorder associatedwith doping levels could be observed by Raman spectros-copy. These findings suggest that the increased sub gap ab-sorption of N-doped samples does not originate only fromchanges of the joint density of states, and may also reflectchanges of the dipole matrix element for optical transitions.The valence structure and the electronegativity of the impu-rities, as well as the existence of impurity clusters, may thusplay a role in the Urbach tail broadening process.

V. TRANSPORT PROPERTIES

Crystalline and amorphousb-Ge3N4 are insulating wideband-gap semiconductors with potential technologicalapplications.14,133 The band gap of germanium nitride is;4.5 eV and its refractive index 2.0560.05.13,14,65

Yashiro210 used Ge3N4 as a passivating layer and a diffusionmask to make low leakage Gep-n junctions. Asc-Ge3N4 isan insulator, the discussion of the electronic transport in GeNalloys will refer to:~a! the conductivity of H-free and hydro-genated off-stoichiometric amorphous GeN alloys and,~b! Nas an impurity in thec-Ge lattice and in thea-Ge network.Along the presentation these topics will be compared withsimilar studies on N in Si, particularly when notably differ-ences appear between the Ge and the Si host.

A. H-free Ge 12xNx alloys: Crystalline and amorphous

The insulating properties of Si and Ge nitrides deterio-rate notably as the N concentration decreases to belowx;0.57. To the present authors’ knowledge there are no re-ports on the electronic properties of off-stoichiometricc-Ge12xNx . In contrast, the transport properties of H-freeand hydrogenated off-stoichiometric amorphous GeN alloysas a function of N content have been addressed by severalauthors.15,26,39,113,211,212

One of the first reports on the effects of N on the opto-electronic properties ofa-Ge was the study by Takanoet al.213 on rf sputtered H-freea-Ge12xNx samples depositedat ;350 K. The conductivity was found to depend on thetemperature and on the N2 partial pressure during growth.Concomitant changes of the optical properties of the filmswere detected simultaneously and attributed by Takanoet al.213 to the incorporation of N in thea-Ge network.Whereas some of the changes measured by Takanoet al.could be due to nitrogen incorporation, a detailed analysis of

the data leads us to believe that most of these changes re-sulted from film contamination upon exposure to theatmosphere.214 The large scatter of the conductivity data, theappearance of the Ge–O related absorption band~;800cm21) in the N-free films, and the very low room tempera-ture dark conductivity@sd(RT)# values for H-free material,indicate that Takanoet al.213 were in fact studying Ge oxyni-trides, as those reported a few years later by Bagratishviliet al.215 We come back to the subject when discussing thedopant properties of N in thea-Ge network.

Figure 31 shows the dependence on temperature of thedc dark conductivity@sd(T)# of somea-Ge12xNx samples.39

It is apparent from the figure that the addition of N affectsthe conductivity ofa-Ge films. As expected for H-free amor-phous material having a large density of defects in thepseudo-gap, the conductivity is dominated by a variablerange hopping mechanism,216 with a characteristicsd(T)5s0* exp[2(T0 /T1/4)] behavior. Figure 31 also showsthe conductivity of an intrinsica-Ge film deposited under thesame conditions as the alloys. Table III gives the electricalcharacteristics of the above series of samples for N varyingin the 0,x,0.36 range.39 The N content of these samples

FIG. 31. Logarithm of the dark conductivity of some selecteda-Ge12xNx

samples vsT21/4. Note the straight-line behavior in all cases, an indicationof variable range hopping transport~Ref. 39!.

TABLE III. Compositional and electrical characteristics ofa-Ge12xNx

samples deposited by the rf sputtering technique.a Note thatsRT does notchange with nitrogen up to a concentration of nearly 4 at %. For larger@N#the band gap widens andsRT drops very fast.

SampleNo.

@N#(cm23)

x~%!

sRT

(V cm)21N(EF)

(eV21 cm23)

1 ,1018 — 9.031023 1.731018

2 nab — 1.531022 2.331018

3 na — 1.831022 2.431018

4 1.631020 0.002 1.831022 1.631018

5 6.031020 0.007 1.731022 1.731018

6 9.131020 0.011 1.531022 1.131018

7 3.031021 0.037 1.131022 1.231018

8 8.831021 0.108 3.531023 —9 1.831022 0.221 2.531025 —

10 2.531022 0.306 ,1027 —11 2.931022 0.355 ,1029 —

aSee Ref. 39.bNot available.



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was determined from a deuteron-induced nuclear reaction,and the concentration evaluated by interpolating betweenpure Ge and stoichiometricb-Ge3N4, which has4.631022 N atoms cm23 and a density of 5.3 g cm23.14,120

Two facts at least emerge from Fig. 31 and Table III:~i! theabsolute value ofsd and its dependence onT do not changeappreciably up to N;331021 cm23 (x;0.04) and ~ii !higher N concentrations decreasesd at all temperatures andthe samples display a stronger dependence onT.

The relative independence ofsd on @N# for x,0.04 isthe consequence of the large density of deep defects (GeDB)around the Fermi energyN(EF) which dominate the trans-port. Assuming the wave function localization parameterg;0.1 Å21 determined by Knoteket al.217 for a-Ge, and theconductivity data of Table III, the density of states forx,0.04 samples at the Fermi energyN(EF) is of the orderof 1018 eV21 cm23 or higher39 ~see Table III!, a defect den-sity compatible with transitions between states 100 Å or lessfar apart. Hence, in Ge-rich material (x,0.04) EF is pinnedby GeDBs and hopping between them dominates conductionat all temperatures. Honmaet al.133 have also reported theconductivity of ana-Ge12xNx film (x50.1) against tempera-ture, and found that logsd becomes a straight line whenplotted againstT21/4, an independent indication that variablerange hopping prevails in highly defective materials, as forsamples of Fig. 31.

As the N concentration into the samples increases~x.0.04, Table III! a progressive decrease of the conductiv-ity is expected because nitrogen provokes the widening ofthe pseudo-gap, inducing a stronger localization of deep de-fects ~see electronic structure Sec. III B, and optical proper-ties, Sec. IV B!. As a consequence,sd is expected not onlyto decrease with increasing N, but to have a stronger depen-dence on temperature, in agreement with experimental re-sults ~Fig. 31!. The lack of reliable data on the defect struc-ture in N-rich a-Ge12xNx samples prevents a quantitativeanalysis along the line used for the Ge-rich alloy.

The optical and transport properties ofa-Ge0.76N0.17O0.07

films formed by nitriding Ge in hydrazine N2H4 vapor atT;1000 K have been reported by Bagratishviliet al.218 Anycomparison with the present authors results39 becomes diffi-cult due to the large quantity of oxygen present in Bagratish-vili et al.218 samples~the hydrazine contained 1%–2% impu-rity water!. For H-free GeN alloys,x;0.17, a Tauc gap170 ofapproximately 0.9 eV is expected,39 whereasETauc51.2 eVis reported for the samples containing oxygen,218 a clear in-dication of the important effects of oxygen contamination onthe optical properties. From PDS data of Ge oxynitridesamples, Bagratishviliet al.218 estimated a defect density atthe Fermi energyN(EF);1.431018 eV21 cm23, a valuecomparable to those reported in Table III forx,0.04 O-freesamples. Contrary to the case ofa-Ge12xNx reported inTable III, a temperature activated dark conductivity at RTand above is found for Ge oxynitride samples.218

B. Hydrogenated a-Ge12xNx films „x>0.01…

Hydrogen atoms play a fundamental role on transportproperties.a-Ge12xNx :H films, prepared by rf reactive sput-

tering of ac-Ge target in a gaseous mixture of Ar1N21H2,were reported by Honmaet al.143 and by Chambouleyronet al.132 These preliminary reports established the role of Hin cleaning the pseudo-gap via the passivation of GeDB’s andthe preference of H to bond to N instead of to Ge atoms.More extensive research was done by Marques,211 who stud-ied several series of samples deposited either by the rf sput-tering technique~Ar1N21H2, Ar1NH3 and Ar1NH31H2

gaseous mixtures!, or by the rf glow discharge~GD! ofGeH41NH3. The Tauc170 optical gap of these sample seriescovered the 1.0–3.5 eV range, the wider gap samples beingthose prepared by rf GD.211,212The transport properties of rfglow discharge depositeda-Ge12xNx :H samples have alsobeen reported by Honmaet al.133 and by Xuet al.26

The most striking effect of hydrogenation on the elec-tronic transport of amorphous semiconductors is the appear-ance of a thermally activated conductivity. Figure 32 showsthe logarithm of the dc dark conductivity of a series ofa-Ge12xNx :H alloys, rf sputtered in an Ar1NH3 atmo-sphere, against inverse temperature.211,212 The conductivityof a nominally undopeda-Ge:H sample deposited under thesame conditions is also shown in Fig. 32 for comparison.The activation energyEa is obtained fitting the data pointswith the expressionsd(T)5s0 exp(2Ea /kT), wherek is theBoltzmann constant. The figure shows that the conductivitydepends strongly on N content. Small NH3 partial pressuresin the chamber cause an increase in the dc conductivity and asimultaneous reduction of the activation energy. This effectwas interpreted as originating from N active chemical doping

FIG. 32. Dark conductivity ofa-Ge12xNx :H films prepared by rf sputteringa c-Ge target in an Ar1NH3 atmosphere as a function of inverse tempera-ture. The conductivity of an intrinsica-Ge:H sample is also shown forcomparison. Note the active doping effects of small NH3 partial pressures inthe deposition chamber and, for high NH3 partial pressures, a decreasingconductivity due to the band-gap widening mechanism induced by nitrogen.The dark conductivity is thermally activated for all sample series~see Refs.211 and 212!.



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of thea-Ge:H network, a subject we will discuss in detail ina coming section. Increasing the partial pressure of NH3

above 531024 mbar during deposition reverses the trend,i.e., Ea increases again andsd decreases very fast, reachings;10212 (V cm)21 at room temperature for thex;0.37sample. The trend toward an insulating material is concomi-tant with the band-gap opening mechanism induced by trigo-nal N bonding, referred to in the previous sections on thestructural and the optical properties of the alloys. Similarlyto the crystalline case, activated transport ina-semiconductors corresponds to thermally excited electronsin the neighborhood of the CB mobility edge. It dominates inthe samples under analysis because hydrogenation has re-duced drastically the density of deep defects, preventing anysignificant contribution from phonon assisted tunneling be-tween localized GeDBs.

The conductivity of a series ofa-Ge12xNx :H samplesprepared by the rf GD of 5% H2 diluted (GeH41NH3) gas-eous mixtures was recently reported by Xuet al.26 Figure 33shows the room temperature conductivity of Xuet al.’ssamples as a function of N content~as determined fromXPS! and those of Marqueset al.211,212deposited by rf sput-tering. Note the differentsRT values between the two series.The difference may be due to the use of different depositionmethods, for it is well known that rf sputtereda-Ge:H pos-sesses a smaller density of deep defects than anodic glowdischarge samples104 ~see Fig. 33!. The large DOS in thepseudo-gap of GD deposited samples may hinder the effectsof nitrogenation at small NH3 partial pressures. Othersources of disagreement might stem from different methodsused to estimate the N content of the two series, as well asthe total amount of bonded hydrogen atoms.

C. Nitrogen as an impurity in Ge „x<0.01…

1. N in crystalline germanium

With the exception of N, the elements of groups III andV of the Periodic Table are known to produce localized elec-tronic levels near the band edges of the elemental semicon-

ductors Si and Ge. The simple model of chemical dopingassumes that four electrons are needed to fill the bondingorbitals of the tetrahedral neighbors. The extra electron~do-nor case! is bound to the impurity by a Coulomb-like attrac-tive potential produced by the extra opposite charge of theimpurity atom. The attractive force is modified by the dielec-tric constant of the crystal, the lattice potential being takeninto account through the use of an appropriate effectivemass.219 A similar argument may be used to give account forthe electric activity of shallow acceptor states.

Within this context, N appears to be distinctly differentfrom other group-V elements, and its electrical activity in thecrystalline lattice of Si and Ge is still a matter of study. Letus note at this point that the properties of N in diamond arevery different than in Si or Ge, N being highly soluble indiamond220 but having a very low solubility in Si and Ge.221

The difficulty of studying the electrical activity of substitu-tional N in crystalline elemental semiconductors partly de-rives from:222 ~i! the rather high dissociation energy of theN2 molecule which prevents the normal mechanism ofatomic diffusion and~ii ! the fact that the N2 molecule mayenter the tetrahedral interstice of Si and Ge. Moreover, bycooling melted Si or Ge in an atmosphere of pure N2 or NH3

the reaction with the liquid phase forms the correspondingnitride phase, with planar bonded N and tetrahedral bondedSi or Ge.

As a consequence of the above, the manifestation of do-nor properties of substitutional N inc-Si andc-Ge can bedetermined only if it is introduced in the atomic state at

FIG. 33. Room-temperature dark conductivity ofa-Ge12xNx :H samples asa function of N content.~Dots! samples from a plasma enhanced chemicalvapor deposition system using diluted GeH4 ~5% in H2! and NH3 gaseousmixtures.~Ref. 26! ~Squares! rf sputtered samples in an Ar1NH31H2 at-mosphere~see Ref. 261!.

FIG. 34. Energy defect levels associated with N inc-Si as found by variouscharacterization and implantation techniques. The donor levels produced bysubstitutional phosphorus and arsenic are also shown for comparison.~a! W.Kohn, Ref. 219;~b! See Ref. 264;~c! See Ref. 265;~d! A. H. Clark, J. D.MacDougall, K. E. Manchester, P. E. Roughan, and F. W. Anderson, Bull.Am. Phys. Soc.13, 376 ~1968!; ~e! A. G. Milnes, in Deep Impurities inSemiconductors~J. Wiley, New York, 1973!; ~f! W. J. Kleinfelder, Techn.Rep. ~Stanford Electronics Labs, 1967!; ~g! Y. Tokumaro, H. Okushi, T.Masui, and T. Abe, Jpn. J. Appl. Phys. 221, L443 ~1982!; ~h! S. T. Pan-telides and C. T. Sah, Phys. Rev. B10, 638 ~1974!; ~i! K. L. Brower, Ref.225; ~j! K. Nauka, M. S. Goorsky, H. C. Gatos, and J. Lagowski, Appl.Phys. Lett.47, 1341 ~1985!; ~k! L. C. Kimerling and J. L. Benton, Appl.Phys. Lett.39, 410 ~1981!; ~l! D. Wruck and P. Gaworzewski, Phys. StatusSolidi ~A! 56, 557 ~1979!. Singly and double charged Si divacancies giverise to levels at 0.42 and 0.56 eV below the CB, respectively@J. Stein, inRadiation Effects Semiconductors, edited by J. Corbett and G. Watkins~Gordon & Breach, New York, 1971!#.



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relatively low doses. Ion implantation has been used for sucha purpose.223–227 The discussion around experimental find-ings involves the subtleties of annealing processes, alwaysnecessary to relax the radiation damage produced by the ionbombardment. Figure 34 shows schematically the energylevels associated with N and with N-induced defects reportedfor c-Si, together with the donor levels of substitutional Pand As.

Probably due to nitrogen’s tendency to form only threebonds and leave two electrons in a lone pair, substitutional Nin diamond and Si is distorted in the~111! direction. Thisgives rise to a deep level in the gap.228,229A study of N inc-Ge by Campbellet al.224 reports that;85% of implantedN is located in nonsubstitutional sites, and that N does notoutdiffuse fromc-Ge for anneals as high as;1000 K. More-over, electrical measurements indicate that substitutional Nis not a donor inc-Ge. The dominant center observed inN-dopedc-Ge orc-Si is not substitutional N but a nitrogenpair,226,227which consists of two neighboring~100!-orientedGe~Si!–N split interstitials, arranged in an antiparallel con-figuration and with the four bonds forming a square lying on$011%, as shown in Fig. 35.

The defect structure of N as a substitutional impurity incrystalline elemental semiconductors has also been ad-dressed theoretically230,231 with two different approaches:perturbation methods and cluster methods, which stress thelocal environment of the defect. Atomic relaxation aroundthe defect and electronic correlation are seldom considered.As both can influence calculations by a factor as large as thegap energy, the defect levels with respect to band edges arenot reliably located.

Summarizing, N is a poor donor in crystalline group-IVsemiconductors, partly because of its low solubility, andpartly because of the tendency of N to form pairs in im-planted as well as in grown materials.

2. N as an active dopant in amorphous Ge (and Si)

The problems related to the inclusion of substitutional Natoms in the crystalline Si and Ge lattices may be partiallyovercome ina-networks. The deposition conditions~fastcondensation from the vapor phase! allow the incorporationof atomic components, or radicals of complex molecules,

dissociated by an rf or a dc plasma. Spear and LeComber232

established for the first time that, in tetrahedrally coordinateda-semiconductors, the doping properties of group III and Velements are similar to those found in the crystalline semi-conductor parents. The doping process, however, is differentin an a-semiconductor compared with a crystal, in the sensethat in the crystal constraints of symmetry force an impurityatom of a different valence to have the same coordinationnumber as its host in a substitutional site. Ina-networks thelocal minimization of free energy may also be attained withthe impurity atom having a first coordination number equalto its chemical valence~self compensated site!. In this case,the first coordination number is determined by purely chemi-cal parameters and not by the impurity’senvironment.216,233,234Both active and inactive doping sitescoexist ina-networks and their relative stability results fromthe combined short-range chemical contributions and the lat-tice strain term, for it is known that tetrahedrally coordinatedamorphous networks are topologically constrained.235

The electrical activity of N impurity ina-Si:H anda-Ge:H has been a matter of debate for many years. It hasbeen argued theoretically that over-coordination of nitrogen~N4

1[positively charged, tetrahedrally coordinated N! shouldbe impeded in thea-Si anda-Ge networks by its small size,and should only be possible if one or more neighbors arehydrogen.123 Another proposal236 to explain the large densityof charged Si dangling bonds normally found in N-dopeda-Si:H considers the existence of N charged dangling bonds(N2

2) and positively charged SiDBs (Si31). This configuration,

however, does not give rise ton-type doping, disagreeingwith the electrical properties found experimentally in N-dopeda-Si:H anda-Ge:H networks.

The nature of these defects was analyzed in detail byShimizu et al.237 by means of several spectroscopic tech-niques. Measurements of ESR, light-induced electron spinresonance~LESR!, and sub-gap optical absorption using theconstant photocurrent method~CPM!, were carried out fora-Si:H, a-Si12xCx :H, a-Si12xOx :H and a-Si12xNx :Hfilms as a function of film thickness. From the combinedinformation provided by these different characterizationtechniques and from samples of different thickness, both thesurface density of dangling bonds in the disordered surfacelayer and the density of dangling bonds in the bulk regionwere obtained, discriminating neutral from charged defects.The main conclusions of Shimizuet al.,237 agreeing with ourmicroscopic picture of N doping ina-Ge:H films,139,238canbe summarized as follows:~i! the density of dangling bondsdeduced from CPM agrees fairly well with the sum of neu-tral and negatively charged dangling bonds deduced fromESR and LESR, both neutral and negatively charged dan-gling bonds contributing to CPM signal;~ii ! the energy po-sition and the width of the defect level distribution are al-most unchanged regardless of the fraction of charged defects,suggesting that both positively and negatively charged de-fects are paired through a Coulombic interaction,~iii ! manycharged dangling bonds exist in the bulk ofa-Si:H, theirdensity being more than five times that of the neutral dan-gling bond. Although some charged dangling bonds existeven in highly purea-Si:H films, the origin of these charged

FIG. 35. The antiparallel model of the N pair defect inc-Si andc-Ge ~seeRef. 227!. Filled circles: Ge atoms; open circles: N atoms.



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dangling bonds appears to be unintentionally incorporated N;~iv! finally, the incorporation of N largely increases the den-sity of dangling bonds. The charged dangling bonds arelikely to be intimate pairs of N4

1 and Si32 rather than N2

21Si31

pairs as suggested by Robertson and Powell.236 It is worthmentioning here that a recent calculation of defect levels ofN impurity in a-Si finds the N4

1 center just belowEc .130

The active doping properties of N in rf sputtereda-Ge:Hfilms using N2 as a doping source were independently stud-ied at Campinas139,140,211,212,238–241~see Fig. 32, for example,for the first evidence!211 and at Kaiserslauten.242,243 Figure36 shows the changes of the dark conductivity ofa-Ge:Hsamples when small amounts of N2 are allowed to flow intothe reaction chamber during deposition.139 Experimental dataindicate that the temperature dependence ofsd can be wellrepresented by either one exponential function, or by the sumof two exponential functions. The single activation energyEa which dominates the electric transport of undoped andlightly dopeda-Ge:H in the 170–400 K temperature rangecorresponds, under the assumption of a linear temperaturedependence of the pseudo-gap, to the energy difference, atT50 K, between the Fermi level and the conduction bandmobility edge. A second exponential, which appears only insamples deposited under N2 partial pressures higher than1025 mbar, dominates the transport at low temperatures. Theanalysis ofsd(T) indicates that transport through band taillocalized states~hopping! dominates the conductivity at lowtemperatures. The most effectively N-dopeda-Ge:H samplepossesses a high temperature activation energyEa;0.12 eV, a remarkably low value for a dopeda-semiconductor, and a low temperature activation energyEH;10 meV, consistent with a thermally activated nearestneighbor hopping at the Fermi energy~see Fig. 36, samplewith @N# ;0.1 at. %!. These experimental findings indicatethat active N doping is a very efficient process ina-Ge:Hsamples of improved quality. Qualitative similar results wereobtained by Dru¨sedauet al.242 The doping efficiency of N ina-Ge:H, defined as the fraction of all incorporated nitrogen

atoms which are electrically active, was quantified by thepresent authors238 and found to be of the same order of mag-nitude as the doping efficiency of P ina-Si:H films, asshown in Fig. 37. It is interesting at this point to compare theabove results with experimental data of N-dopeda-Si:Hfilms. In what follows we show that important differencesappear between both materials.

One of the first studies on N as a dopant ofa-Si:Hreported244 an eight orders of magnitude increase in conduc-tivity upon incorporation of N, with an activation energy assmall as 20 meV for the most effectively doped sample.These results, however, could never be reproduced by otherresearch groups. Most studies on N-dopeda-Si:H140,141,245,246report the following overall picture. N-dopeda-Si:H films possess a thermally activated conductiv-ity, the activation energyEa decreases down to a minimumvalue of around 0.5 eV with doping and a simultaneous 3–5orders of magnitude increase ofsd is measured~see Fig. 38!.The experimental data suggest that, either the doping effi-ciency of N ina-Si:H is very small, or the energy level of thedefect center giving rise ton-type doping is relatively deepin the pseudo-gap, or both. Meaningful comparisons are of-

FIG. 36. Logarithm of the dark conductivity of N-dopeda-Ge:H as a func-tion of inverse temperature. Intrinsic and lightly N-doped samples possess asingle thermally activated conduction regime down to low temperatures. Inheavier N-dopeda-Ge:H films, thermally activated hopping through local-ized states near the conduction-band edge dominates the conductivity~seeRef. 139!.

FIG. 37. The doping efficiency of nitrogen atoms vs the concentration of Nin the solid phase for N-doped rf sputtereda-Ge:H samples~solid circles! ascompared with P-doped glow dischargea-Si:H films ~see Refs. 238 and262!.

FIG. 38. Room-temperature dark conductivity vs activation energy for N-dopeda-Si:H films. The experimental data labeled bya-Si:H~N! correspondto rf sputtered samples while the other points were obtained from rf glowdischarge (SiH41N2) deposited samples~see Refs. 141, 246, and 263!.



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ten difficult ~or impossible! to establish between samples de-posited in different systems under different conditions ormethods, or deposited using different gaseous dopantsources. In spite of the difficulty, let us speculate on thenature of the N donor impurity under the light of availableexperimental results.

A comparative study of N doping in rf sputtereda-Si:Handa-Ge:H has been made by the present authors.140,240Fig-ure 39 shows the incorporation of N atoms ina-Si:H anda-Ge:H samples prepared in the same deposition system as afunction of the N2 partial pressure (PN2

) in the chamber.Figure 39 indicates that for similar N2 partial pressures;~i!the rate of N incorporation with varyingPN2

is the same forboth elements and,~ii ! under similarPN2

almost a one orderof magnitude larger fraction of N is incorporated ina-Si:H.The comparison only considers samples having the same de-gree of hydrogenation and containing at most 1 at % N at-oms, which corresponds to a true doping concentrationrange. rf sputtering, however, does not lead to device qualitya-Si:H and its DOS is one order of magnitude larger thanthat of a-Ge:H.

Figure 40~a! showssRT andEa for N-dopeda-Si:H as afunction of PN2

for the sample series under consideration.Important conductivity changes appear upon N incorpora-tion, a clear indication of chemical doping. Note that themeasured changes occur at relatively low N2 partial pres-sures, a consequence of the relatively large rf power fed intothe plasma during the sputtering process, which dissociates aconsiderable fraction of N2 molecules. Compare this picturewith that obtained by Zhouet al.141 @Fig. 40~b!# who foundan almost identical dependence ofEa and sRT on N-dopeda-Si:H samples deposited by the rf glow discharge methodusing a (SiH41N2) gaseous mixture. In the case of GDsamples, however, a much higher N2 pressure is required inthe chamber~see Figs. 1 and 2 of Ref. 141! to detect anyconductivity change, a consequence of the low rf power fedinto the GD plasma, which is insufficient to break a sufficientfraction of N2 molecules. Zhouet al.141 compare N dopingwith P doping ofa-Si:H under the same deposition condi-tions, and conclude that the doping efficiency of N is at least

three orders of magnitude smaller than that of P. We believethis conclusion to be erroneous because the dissociation en-ergy of PH3 and N2 are very different. Under the same dis-charge conditions the dissociation of PH3 molecules will beaccomplished to a much larger extent than for N2. The re-sults of Zhou et al.141 and those of the present authorssuggest,140 on the contrary, that the N doping efficiencymight not necessarily be smaller ina-Si:H than ina-Ge:H.Instead, experimental results suggest that the donor levelproduced by N ina-Si:H may be deeper than that of N ina-Ge:H, as suggested by the calculations of Makleret al.126

Note that, because of a much larger Ge dielectric constant,this is the case for all group III and V elements inc-Si ascompared withc-Ge. The limited energy range swept byEa

might then be the consequence of a deep donor level and notof a poor doping efficiency. To the present authors knowl-edge, the doping efficiency of N ina-Si:H has not yet beenclearly established.

Why is N a highly efficient dopant ina-Ge? The size ofthe nitrogen atom would hinder the N4

1 configuration, unlessa considerable atomic relaxation occurs around the impurityatom, as suggested by the present authors.139,239 FollowingRobertson and Powell’s ideas Dru¨sedauet al.242 suggest thatthe N-related donor state originates in the presence of hydro-gen complexes, such as→Ge–H2–Ge← for example, andnot from an over-coordinated N atom. The reason behind theH-complex configuration should be, according to the Kaiser-slautern group, the preferential attachment of hydrogen tonitrogen.242 To investigate the likely role of H in the dopingprocess, N-doped rf sputtereda-Ge:H samples using NH3and N2 as a gaseous dopant source were compared atCampinas.139,239 The analysis of the transport data indicatethat, although the doping effect looks similar with the use ofeither of the two gases, equal amounts of N in the networkdo not induce equal activation energy and conductivitychanges ina-Ge:H. Figure 41 shows the variations of theactivation energy in N2- and NH3-dopeda-Ge:H as a func-tion of N concentration in the solid phase.239 When NH3 isusedEa vs @N# flattens and shifts toward larger N concentra-tions. This is an indication that the N-active doping effi-

FIG. 39. Nitrogen incorporation ina-Si:H anda-Ge:H thin films as a func-tion of N2 partial pressure during deposition. The N concentration ina-Siwas determined from IR absorption data, whereas the impurity content ina-Ge was measured by nuclear reaction techniques~NRA! ~see Ref. 140!.

FIG. 40. Dark conductivity at room temperature and activation energy ofN-dopeda-Si:H films as a function of N2 partial pressure in the chamberduring deposition.~a! data from rf sputtered samples and~b! data from rfGD-deposited films. Note the important changes ofsRT and ofEa inducedby small N2 partial pressures~see Refs. 140 and 141!.



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ciency is smaller for NH3 than for N2. Moreover, the use ofNH3 induces a large network disorder~see optical properties,Sec. IV C! due to the incorporation of NH and NH2 groupsas Ge dangling bond terminators.239 Tsu et al.247 preparedsilicon nitride films from (SiH41N2) and (SiH41NH3). Inthe latter, they found evidence of the incorporation of NHand NH2 groups into the films. Samples deposited using ND3

instead of NH3 indicate that the source of the H in the NHand NH2 groups is NH3 and not SiH4, suggesting that thedonor center induced by nitrogen ina-Si:H is also over-coordinated N atoms, in agreement with the results ofShimizu et al.237 Infrared studies and chemical bondingmodeling lead Lucovskyet al.248 to propose that fourfoldcoordinated N with second neighbor H atoms, as inN1–Si–H linkages, is the active donor configuration of N inthe a-Si:H network. These results, however, should be con-sidered with caution, in the sense that the N-doped samplesstudied by Lucovskyet al.248 contain up to 15 at. % nitrogen,i.e., a N content well above the normal doping concentrationrange and typical of off-stoichiometric silicon nitride.

The question of N dopinga-Ge has been recently ad-dressed theoretically by Venezuela and Fazzio.249 The struc-ture of a-Ge is generated using the continuous-space MonteCarlo method on a large number of atoms. The correspond-ing electronic and structural configurations are obtainedwithin the framework of the density–functional theory usingthe local density approximation. The results are illustrativeof the importance of atomic relaxation, which appears to bevery large in thea-Ge network. In all N atom inclusion casesstudied by Venezuelaet al.249 ~trigonal or tetrahedral coor-dinated N!, a strong inward relaxation of the Ge atoms in thevicinity of the impurity is found. The calculations also showthat, of all possible charge states of tetrahedral N, the moststable is N4

1, in agreement with the predictions of the experi-mental studies at Campinas.139,238 As with other group Vactive dopant elements in Si and Ge,234 the N4

1 impurity is inthe center of an almost perfect tetrahedron of Ge atoms, inwhich the Ge–N bond length is approximately 2.1 Å. Thisfigure is to be compared with 1.83 Å, the Ge–N bond length

determined from EXAFS data of trigonal N in thea-Genetwork.121 Let us note here that the above calculation wereperformed in a H-freea-Ge network. The presence of H mayaffect considerably the kind of most favorable impurity con-figuration. In addition, Venezuela and Fazzio249 madeab ini-tio calculations of defect formation energies, the values ofwhich, however, appear to be at odds with chemical intuitionand the available experimental data. It is theoretically found,for example, that active electronic doping originates fromN3

1 and not from N41, as expected from the standard theory

of chemical doping ofa-semiconductors.234 The two mostsurprising results of the defect energy calculations are:~i!The neutral N3

0 center cannot exist because the reactionN3

0→N311D2 is exothermic, the donor electron becoming

localized at deep Ge dangling bond centers. In other words,the Ge–N bond of the Ge3N skeleton would be shallowerthan thesp3 Ge dangling bond. This is at odds with therecession of the VB ofa-GeN alloys as the N concentrationincreases, as measured by photoelectron spectroscopies.129 Itis also contrary to chemical intuition considering that N ismore electronegative than Ge.~ii ! The ab initio calculationsalso indicate that the relative concentration of@N3#/@N4# de-creases as the N concentration increases in the network. Thisprediction is not compatible with the known phenacite struc-ture of stoichiometrica-Ge3N4, in which Ge atoms are tet-rahedrally bonded to N atoms, and N is trigonally bonded toGe in a planar configuration.

Summarizing, the experiments done at Campinas usingN2 and NH3 as dopant sources and the structural calculationsof Venezuela and Fazzio,249 referring to threefold and four-fold coordinated N ina-Ge, give support to the idea thatactive N doping originates from over-coordinated N and notfrom NH complexes. Yet, there are no reliable theoreticalvalues for the energy level of the N4

1 antibonding state. Thequestion of the different activity of N in the Si and Gea-networks remains to be answered. The available informa-tion suggests that the doping mechanisms are pretty muchthe same in both semiconductors, the less efficient N dopingof a-Si:H being probably associated with a deep donor level,say 300–350 meV below the CB edge. The different energyfor the N donor level betweena-Ge:H ~50–60 meV! anda-Si:H might originate from the smaller dielectric constantof Si and a less effective atomic relaxation around the impu-rity.

D. Photoconductivity in N-doped a-Ge:H films

The photoconductivity, defined as the change in conduc-tivity upon illumination, can be expressed as:Ds5e(mnDn1mpDp), with Dn5hGt, where m is thecarrier mobility,G the photogeneration rate of electron-holepairs, h the quantum efficiency for free-carrier productionandt the lifetime of excess carriers. Depending on recombi-nation kinetics,t is usually a function of the generation rate.For strongly absorbed photonsDs}(aF)g, wherea andFare the absorption coefficient and the photon flux, respec-tively, g is found to vary between 0.5 and 1.0~bimolecularand monomolecular recombination kinetics, respectively!and may also depend on the generation rate and the photonenergy.

FIG. 41. High temperature (T.300 K) activation energy of the dark con-ductivity vs nitrogen concentration in N-dopeda-Ge:H films using N2~circles! and NH3 ~squares! as doping sources. Both dopant sources inducesimilar changes of the activation energy. Note, however, the flattening of thecurve and its shift to larger N concentrations for samples doped from NH3

source~see Ref. 239!.



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Marcanoet al.250 reported on the secondary photocon-ductivity of N-dopeda-Ge:H samples. Figure 42 shows themeasuredhmt product andg factor as a function of N dop-ing. Three different behavior regions are apparent from thefigure. The intrinsic and lightly doped samples(N,1019 cm23) do not display important changes in thehmtproduct and theg exponent remains close to 0.75, an indica-tion that light N doping has not substantially altered the re-combination kinetics. It is expected that a fraction of neutralGeDB will be converted toD2, i.e., doubly occupied DBswhich are ineffective as recombination centers. TheEF shift~see shift ofEa in Fig. 41! is not sufficient to transform allD0’s into D2’s. The density of remainingD0’s is sufficientto keep thehmt product in the 1027 cm2 V21 range.

The second region corresponds to three effectivelydoped samples possessing an activation energyEa<0.2 eVand ahmt product distinctly higher than the samples of thefirst group (hmt.1026 cm2 V21). The small activation en-ergy for these samples means that most GeDB’s are doublyoccupied and thehmt product is almost independent of dop-ing. The small increase ing may originate from the GeDB

density increase with N doping. The remaining samples arethe most N contaminated. The active doping process is veryineffective at large N doping concentrations~see Fig. 37!,most N entering in thea-Ge network with trigonal coordina-tion. A very different coordination and atom size induce theappearance of a large density of coordination defects. As aconsequence, an importanthmt product drop is expected aswell as a pure monomolecular recombination kinetics, asmeasured.

Masudaet al.251 studied the ESR, light-induced ESR andthe decay time of the photoconductivity of N-dopeda-Si:Hthin films. From the behavior of the decay time they con-clude that, up to a N concentration of approximately 2 at %,the conductivity variations are due to the incorporation oftetrahedral N ~effective dopant!. For N concentrations.2 at. % the decay time of the photoconductivity decreasesand the density of deep defects continues to increase, sug-gesting that most of charged SiDB’s found in the samples do

not originate from N donors but from structural randomnesscaused by the incorporation of nitrogen.

VI. CONCLUDING REMARKS

The properties of nitrogen impurity in germanium and ofGe–N alloys have been reviewed. Amorphous and crystal-line germanium–nitrogen alloys are interesting materialswith potential applications for protective coatings and win-dow layers for solar conversion devices. They may also actas effective diffusion masks for III-V electronic devices. Thescarce existing information on crystallineb-Ge3N4 indicatesthat its lattice and electronic structure are similar tob-Si3N4,by far the most studied column IV nitride. In particular, theband-gap widening mechanism induced by nitrogen closelyresembles both alloys. The top of the valence band is domi-nated ina-Ge3N4 by a N 2p nonbonding orbital, the conduc-tion band being given by Ge–Ge antibonding orbitals. Themechanical properties of crystallineb-Ge3N4 are very poorlyknown, as well as the possible ceramic usesc-GeN alloys.

The review focuses on the experimental and theoreticalproperties of amorphous germanium–nitrogen alloy films, onwhich a more abundant literature exists. A general corre-spondence has been found between these and the structuraland optical properties ofa-SiN, as expected from the similarvalence structure of both elements. Some differences appear,however, regarding the nitrogen concentration at which theband-gap widening induced by nitrogen becomes effective,the structure of the valence band as determined from electronphotoemission spectra, and the properties of N impurity inthe amorphous network.

Nitrogen is not an active dopant in crystal column IVelements, but an active dopant in the corresponding amor-phous networks. The doping efficiency of N ina-Ge:H anda-Si:H has been discussed. Some important differences be-tween the two cases have been highlighted. Nitrogen is avery effective active dopant in thea-Ge:H network, inducinga donor level at around 50–60 meV below the CB edge. Itsdoping efficiency is similar to that of phosphorus ina-Ge:Hor in a-Si:H. The conductivity variations ofa-Si:H inducedby N doping are not as impressive. The discussion suggeststhat this poor performance may stem from a N-induced deepdonor level, and not from a reduced doping efficiency. Theshallowness of the defect level energy ina-Ge:H may be theconsequence of a larger atomic relaxation around the impu-rity, as found theoretically froma-Ge structural simulationsusing a continuous-space Monte Carlo method on a largenumber of atoms. More theoretical and experimental workon the nature and energy of N-induced defects would allowus to understand the role of filledd states in the microscopicorigin of N doping efficiency.

ACKNOWLEDGMENTS

Most of the work on N in Ge done at UNICAMP waspartially financed by the Brazilian agencies Fundaçao deAmparo aPesquisa do Estado de Saõ Paulo~FAPESP! andConselho Nacional de Desenvolvimento Cientıćo e Tecno-logico ~CNPq!. The authors are indebted to their colleaguesat the Laboratory of Photovoltaic Research for fruitful dis-cussions.

FIG. 42. hmt product andg factor for intrinsic and N-dopeda-Ge:H thinfilms as a function of N concentration in the solid phase~N2 has been useda dopant source gas!. ~See Ref. 250!. Note the plateau at N concentrationsbetween 1019 and 1020 cm23. See Fig. 41 for the corresponding activationenergies.



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1E. A. Cotton, G. Wilkinson, and P. L. Gaus,Basic Inorganic Chemistry,2nd ed.~J. Wiley, New York, 1987!.

2W. A. Harrison,Electronic Structure and The Properties of Solids~W. H.Freeman & Company, San Francisco, CA, 1980!.

3S. M. Sze,Physics of Semiconductor Devices~J. Wiley, New York,1981!.

4S. Strike and H. Morkoc¸, J. Vac. Sci. Technol. B10, 1237~1992!.5A. Y. Liu and M. L. Cohen, Phys. Rev. B41, 10727~1990!.6H. Morkoc, S. Strike, G. B. Gao, M. E. Lin, B. Sverdlov, and M. Burns,J. Appl. Phys.76, 1363~1994!. See also J. T. Torriak, J. I. Pan Kove, E.Iliopoulos, H. M. Ng, and T. D. Moustakas, Appl. Phys. Lett.72, 244~1998!.

7R. N. Katz, Science208, 841 ~1980!.8Thin Film Processes, edited by J. L. Vossen and W. Kern~Academic,New York, 1978!.

9Various authors, in ‘‘Silicon Based Ceramics,’’ Mater. Res. Bull.20,~1995!.

10J. C. Remy and J. J. Hantzpergue, Thin Solid Films30, 197 ~1975!; 30,205 ~1975!.

11D. Martok, K. J. Boyd, and J. W. Rabalais, Int. J. Mod. Phys.9, 3527~1995!.

12R. G. Farrer, Solid State Commun.7, 685 ~1969!.13H. Nagai and T. Niimi, J. Electrochem. Soc.115, 671 ~1968!.14G. D. Bagratishvili, R. B. Dzhanelidze, N. I. Kurdiani, and D. V. Saksa-

ganskii, Phys. Status Solidi A36, 73 ~1976!.15I. Chambouleyron, Appl. Phys. Lett.47, 117 ~1985!.16Handbook of Thin Film Technology, edited by L. I. Maissel and R. Glang

~McGraw-Hill, New York, 1970!.17See, for example,Thin Films from Free Atoms and Particles, edited by

K. J. Klabunde~Academic, New York, 1985!.18F. K. McTaggart,Plasma Chemistry in Electrical Discharges~Elsevier,

Amsterdam, 1967!.19Techniques and Applications of Plasma Chemistry, edited by J. R. Hol-

lahan and A. T. Bell~Wiley Interscience, New York, 1974!.20See, for example, C.-E. Morosanu, Thin Solid Films65, 171 ~1980!.21Silicon Nitride Thin Insulating Films, edited by V. Kapoor and H. Stein,

Proceedings of the Electrochemical Society Vol. 83–8~The Electro-chemical Society, Pennington, NJ, 1983!.

22A. C. Adams, inPlasma Deposited Thin Films, edited by J. Mort and F.Jansen~CRC, Boca Raton, FL, 1988!, Chap. 5.

23See, for example, Proc. Int Conf. Amorphous Semiconductors, J. Non-Cryst. Solids198&200 ~1996!.

24F. C. Marques, I. Chambouleyron, and F. Evangelisti, J. Non-Cryst. Sol-ids 114, 561 ~1989!.

25J. Xu, S. Miyazaki, and M. Hirose, Jpn. J. Appl. Phys., Part 135, 2043~1996!.

26J. Xu, K. Chen, D. Feng, S. Miyasaki, and M. Hirose, J. Appl. Phys.80,4703 ~1996!.

27D. B. Alford and L. G. Meiners, J. Electrochem. Soc.134, 979 ~1987!.28G. A. Johnson and V. J. Kapoor, J. Appl. Phys.69, 3616~1991!.29C. Weissmantel, Thin Solid Films32, 11 ~1976!.30C. J. Mogab and E. Lugujjo, J. Appl. Phys.47, 1302~1976!.31A. W. Stephens, J. L. Vossen, and W. Kern, J. Electrochem. Soc.123,

303 ~1976!.32G. J. Kominiak, J. Electrochem. Soc.122, 1271~1975!.33P. C. Y. Chen, Thin Solid Films21, 245 ~1974!.34W. Rothemund and C. R. Fritzche, Thin Solid Films15, 199 ~1973!.35L. F. Cordes, Appl. Phys. Lett.11, 383 ~1967!.36S. M. Hu, D. R. Kerr, and L. V. Gregor, Appl. Phys. Lett.10, 97 ~1967!.37J. Pompei,Proceedings of the Symposium on Deposition Thin Films

Sputter, 2nd ed.~University of Rochester Press, Rochester, NY, 1967!, p.127.

38J. J. Hantzpergue, Y. Doucet, Y. Pauleau, and J. C. Remy, Ann. Chim.~Paris! 10, 211 ~1975!.

39A. R. Zanatta and I. Chambouleyron, Phys. Rev. B48, 4560~1993!.40J. Vilcarromero and F. C. Marques, J. Appl. Phys.76, 615 ~1994!.41V. Y. Doo, D. R. Nichols, and G. A. Silvey, J. Electrochem. Soc.113,

1279 ~1966!.42V. Y. Doo, D. R. Kerr, and D. R. Nichols, J. Electrochem. Soc.115, 61

~1968!.43F. Roenigk and K. F. Jensen, J. Chem. Soc.134, 1777~1987!.44S. L. Zhang, J. T. Wang, W. Kaplan, and M. O¨ stling, Thin Solid Films

213, 182 ~1992!.45N. S. Zhou, S. Fujita, and A. Sasaki, J. Electron. Mater.14, 55 ~1985!.

46J. A. Thiel, S. V. Hattangady, and G. Lucovsky, J. Vac. Sci. Technol. A10, 719 ~1992!.

47K. R. Lee, K. B. Sundaram, and D. C. Malocha, J. Mater. Sci.4, 283~1993!.

48G. M. Samuelson and K. M. Mar, J. Electrochem. Soc.129, 1773~1982!.49V. S. Nguyen, W. A. Lanford, and A. L. Reiger, J. Electrochem. Soc.

133, 970 ~1986!.50H. Schuh, T. Schlosser, P. Bissinger, and H. Schmidbaur, Z. Anorg. Allg.

Chem.619, 1347~1993!.51J. Dahlhaus, P. Jutzi, H. J. Frenck, and W. Kulisch, Adv. Mater.5, 377

~1993!.52J. R. Flemish and R. L. Pfeffer, J. Appl. Phys.74, 3277~1993!.53M. Boudreau, M. Boumerzoug, P. Mascher, and P. E. Jessop, Appl. Phys.

Lett. 63, 3014~1993!.54J. Ahn and K. Suzuki, Appl. Phys. Lett.64, 3249~1994!.55F. Glockling,The Chemistry of Ge~Academic, London, 1969!.56W. C. Johnson, G. H. Morey, and A. E. Kott, J. Am. Chem. Soc.54, 4278

~1932!.57R. Storr, A. N. Wright, and C. A. Winkler, Can. J. Chem.40, 1296

~1962!.58W. C. Johnson, J. Am. Chem. Soc.52, 5160~1930!.59R. Schwarz and P. W. Shenk, Chem. Ber.63, 296 ~1930!.60R. Juza and H. Hahn, Z. Anorg. Allg. Chem.241, 32 ~1939!.61W. Leslie, K. G. Carroll, and R. M. Fisher J. Met.4, 204 ~1952!.62O. Johnson, Usp. Khim.25, 1 ~1956!.63R. Juza and H. Hahn, Z. Anorg. Allg. Chem.244, 124 ~1940!.64V. A. Gritzenko, Prog. Surf. Sci.28, 387 ~1975!.65A. B. Young, J. J. Rosenberg, and I. Szendro, J. Electrochem. Soc.134,

2867 ~1987!.66K. P. Pande and C. C. Shen, Appl. Phys. A: Solids Surf.28, 123 ~1982!.67A. Y. Liu and M. L. Cohen, Science245, 841 ~1989!.68J. J. Cuomo, P. A. Leary, D. Yu, W. Reuter, and M. Frisch, J. Vac. Sci.

Technol.16, 229 ~1979!.69C. J. Torn, J. M. Silverstein, J. H. Judy, and C. Chang, J. Mater. Res.5,

2490 ~1993!.70M. Y. Chen, D. Lin, X. Lin, V. P. David, Y. W. Chung, M. S. Wong, and

W. D. Sproul, J. Vac. Sci. Technol. A11, 521 ~1993!.71C. Niu, Y. Z. Lu, and C. M. Lieber, Science261, 335 ~1993!.72F. Fujimoto and K. Ogata, Jpn. J. Appl. Phys., Part 232, L420 ~1993!.73T. A. Yel, C. L. Lin, J. M. Silverstein, and J. H. Judy, J. Magn. Magn.

Mater.120, 314 ~1993!.74J. H. Kaufman, S. Metin, and D. D. Saperstein, Phys. Rev. B39, 13053

~1989!.75O. Amir and R. Kalish, J. Appl. Phys.70, 4958~1991!.76D. F. Franceschini, C. A. Achete, F. L. Freire, Jr., and G. Mariotto, in

Novel Forms of Carbon, edited by C. L. Renschler, J. J. Pouch, and D. M.Cox ~MRS, Pittsburg, PA, 1992!, p. 481.

77A. Mansour and D. Ugolini, Phys. Rev. B47, 10201~1993!.78J. Seth, R. Padiyath, and S. V. Babu, Diamond Relat. Mater.3, 210

~1994!.79D. F. Franceschini, C. A. Achete, and F. L. Freire, Jr., Appl. Phys. Lett.

60, 3229~1992!.80F. L. Freire, Jr., D. F. Franceschini, and C. A. Achete, Phys. Status Solidi

B 192, 493 ~1995!.81L. Maya, J. Vac. Sci. Technol. A11, 604 ~1993!.82R. S. Lima, P. H. Dionisio, and W. H. Schreiner, Solid State Commun.

79, 395 ~1991!.83R. G. Gordon, D. M. Hoffman, and U. Riaz, Chem. Mater.4, 68 ~1992!.84R. Kieffer, D. Fister, H. Schoof, and K. Mauer, Powder Metall.5, 1

~1973!.85R. Kieffer and P. Ettmayer, High Temp.-High Press.6, 253 ~1974!.86W. Schintlmeister and O. Pacher, Metall.28, 690 ~1974!.87Vapor Deposition, edited by C. F. Powell, J. H. Oxley and J. M. Blocher,

Jr. ~Wiley, New York, 1966!, Chap. 11.88T. Takahashi and H. Itoh, J. Electrochem. Soc.124, 797 ~1977!.89Chemical Vapor Deposition, 5th International Conference, edited by G.

F. Wakefield and J. M. Blocher, Jr.~Electrochemical Society, Penning-ton, NJ, 1975!, p. 634.

90M. J. Rand and J. F. Roberts, J. Electrochem. Soc.115, 423 ~1968!.91M. Hirayama and K. Shohno, J. Electrochem. Soc.122, 1671~1975!.92T. L. Chu and R. W. Kelm, Jr., J. Electrochem. Soc.122, 995 ~1975!.93D. W. Lewis, J. Electrochem. Soc.117, 978 ~1970!.94W. M. Kim, E. J. Stofko, P. J. Zanzucchi, J. I. Pankove, N. Ettenberg, and

S. L. Gilbert, J. Appl. Phys.44, 292 ~1973!.



143.106.108.149 On: Thu, 18 Jun 2015 18:49:29

95H. M. Manasevit, F. M. Erdmann, and W. I. Simpson, J. Electrochem.Soc.118, 1864~1971!.

96N. Ilegems, J. Cryst. Growth13/14, 360 ~1970!.97T. L. Chu, J. Electrochem. Soc.118, 1200~1971!.98M. J. Thompson,The Physics of Hydrogenated Amorphous Silicon I-

Topics in Applied Physics, edited by J. D. Joannopoulos and G. Lucovsky~Springer, New York, 1984!, Vol. 55, and references therein.

99I. Brodie, L. T. Lamont, Jr., and D. O. Myers, J. Vac. Sci. Technol.6, 124~1968!.

100D. A. Anderson, G. Moddel, M. A. Paesler, and W. Paul, J. Vac. Sci.Technol.16, 906 ~1979!.

101R. C. Ross and R. Messier, J. Appl. Phys.52, 5329~1981!.102T. D. Moustakas, Sol. Energy Mater.8, 187 ~1982!.103J. Tardy and R. Meaudre, Philos. Mag. B48, 73 ~1983!.104F. C. Marques and I. Chambouleyron,Proceedings of the 9th European

Photovoltaic Solar Energy Conference, edited by W. Palz, G. T. Wrixon,and P. Helm~Kluwer Academic, Dordrecht, 1989!, p. 1042.

105A. Antoine, B. Drevillon, and P. Roca, J. Non-Cryst. Solids77&78, 769~1985!.

106F. H. Karg, H. Bohm, and K. Pierz, J. Non-Cryst. Solids114, 477~1989!.107W. A. Turner, S. J. Jones, D. Pang, B. F. Bateman, J. H. Chen, Y. M. Li,

F. C. Marques, A. A. Wetsel, P. Wickboldt, W. Paul, J. Bodart, R. E.Norberg, I. ElZawawi, and M. L. Theye, J. Appl. Phys.67, 7430~1990!.

108J. Vilcarromero and F. C. Marques, Phys. Status Solidi B192, 543~1995!.

109See, for example, R. A. Street,Hydrogenated Amorphous Silicon~Cam-bridge University Press, Cambridge, 1991!, and references therein.

110W. Paul and D. A. Anderson, Sol. Energy Mater.5, 229 ~1981!.111T. D. Moustakas, H. P. Maruska, R. Friedman, and M. Hicks, Appl. Phys.

Lett. 43, 368 ~1983!.112T. D. Moustakas and H. P. Maruska, Appl. Phys. Lett.43, 1037~1983!.113D. Hardie and K. H. Jack, Nature~London! 180, 332 ~1957!.114R. Grun, Acta Crystallogr., Sect. B: Struct. Crystallogr. Cryst. Chem.35,

800 ~1979!.115J. Robertson, Philos. Mag. B44, 215 ~1981!.116M. V. Coleman and D. J. Thomas, Phys. Status Solidi25, 241 ~1968!.117T. Aiyama, T. Fukunaga, K. Nihara, T. Hirai, and K. Suzuki, J. Non-

Cryst. Solids33, 131 ~1979!.118M. Misawa, T. Fukunawa, K. Nihara, T. Hirai, and K. Suzuki, J. Non-

Cryst. Solids33, 135 ~1979!.119S. Mobilio and A. Filipponi, J. Non-Cryst. Solids97&98, 365 ~1987!.120S. N. Ruddlesden and P. Popper, Appl. Sci. Res.11, 465 ~1958!.121F. Boscherini, A. Filipponi, S. Pascarelli, F. Evangelisti, S. Mobilio, F. C.

Marques, and I. Chambouleyron, Phys. Rev. B39, 8364~1989!.122A. Filipponi, P. Fiorini, F. Evangelisti, and S. Mobilio,MRS Symposium

Proceedings, edited by A. Madan, M. Thompson, D. Adler, and Y. Ha-makawa~MRS, Pittsburgh, PA, 1987!, Vol. 95, p. 305.

123J. Robertson, Philos. Mag. B63, 47 ~1991!. Note that in this paper thenotationa-SiNx is used instead ofa-Si12xNx . A review of theoreticalcalculations on the electronic structure ofa-Si12xNx is found in thisarticle.

124R. Karcher, L. Ley, and R. L. Johnson, Phys. Rev. B30, 1896~1984!.125M. M. Guraya, H. Ascolani, G. Zampieri, J. H. Dias da Silva, M. P.

Cantao, and J. I. Cisneros, Phys. Rev. B49, 13446~1994!.126S. Makler, G. M. Rocha, and E. V. Anda, Phys. Rev. B41, 5857~1990!.127A. R. Zanatta and I. Chambouleyron, Phys. Status Solidi B193, 399

~1996!.128A. R. Zanatta, R. Landers, S. G. C. de Castro, G. G. Kleiman, I. Cham-

bouleyron, and M. L. Grilli, Appl. Phys. Lett.66, 1258~1995!.129D. Comedi, A. R. Zanatta, F. Alvarez, and I. Chambouleyron, J. Non-

Cryst. Solids198–200, 136 ~1996!.130J. Robertson, Philos. Mag. B69, 307 ~1994!.131W. L. Warren, J. Kanicki, J. Robertson, E. H. Poindexter, and P. J.

McWhorter, J. Appl. Phys.74, 4034~1993!.132I. Chambouleyron, F. C. Marques, J. I. Cisneros, F. Alvarez, S.

Moehlecke, W. Losch, and I. Pereyra, J. Non-Cryst. Solids77&78, 1309~1985!.

133I. Honma, H. Kawai, H. Komiyama, and K. Tanaka, J. Appl. Phys.65,1074 ~1989!.

134M. H. Brodsky and R. S. Title, Phys. Rev. Lett.23, 581 ~1969!.135M. Stutzmann, J. Stuke, and H. Dersch, Phys. Status Solidi B105, 265

~1983!.136G. Chen, F. Zhang, W. Jia, and W. Chen, Phys. Status Solidi A114, 1170

~1989!.

137H. Yokomichi, Materials Science Forum~Trans. Tech. Publications,Switzerland, 1995!, Vols. 196–201, p. 1291.

138H. Min, S. Ueda, N. Ishii, M. Kumada, and T. Shimizu, J. Non-Cryst.Solids198–200, 375 ~1996!.

139A. R. Zanatta and I. Chambouleyron, Phys. Rev. B46, 2119~1992!.140A. R. Zanatta and I. Chambouleyron, Braz. J. Phys.24, 434 ~1994!.141J-H. Zhou, K. Yamaguchi, Y. Yamamoto, and T. Shimizu, J. Appl. Phys.

74, 5086~1993!.142M. Cardona and L. Ley, inPhotoemission in Solids I, General Principles,

edited by M. Cardona and L. Ley~Springer, Berlin, 1978!, Chap. 1.143I. Honma, H. Kawai, H. Komiyama, and K. Tanaka, Appl. Phys. Lett.50,

276 ~1986!.144R. Liu Chun, Q. Sun Zhao, Z. Xu Jing, and X. Zheng, J. Phys. D24, 2215

~1991!.145A. R. Zanatta, I. Chambouleyron, and P. V. Santos, J. Appl. Phys.79,

433 ~1996!.146R. A. Barrio, A. S. Carric¸o, F. C. Marques, J. Sanjurjo, and I. Chambou-

leyron, J. Phys.: Condens. Matter1, 69 ~1989!.147G. Lucovsky, J. Yang, S. S. Chao, J. E. Tyler, and W. Czubatyj, Phys.

Rev. B28, 3234~1983!.148A. R. Zanatta, F. L. Freire, Jr., and I. Chambouleyron, J. Phys.: Condens.

Matter 5, A313 ~1993!.149A. L. Smith and N. C. Angelotti, Spectrochem. Acta15, 412 ~1959!.150G. Lucovsky, Solid State Commun.29, 571 ~1979!.151D. W. H. Rankin, J. Chem. Soc. A , 1926 ~1969!.152C. Gridewell, D. W. H. Rankin, and A. G. Robiette, J. Chem. Soc. A ,

2935 ~1970!.153L. Pauling,The Nature of the Chemical Bond, 3rd ed.~Cornell University

Press, New York, 1960!.154A. R. Zanatta and I. Chambouleyron, Solid State Commun.95, 270

~1995!.155J. E. Smith, Jr., M. H. Brodsky, B. L. Crowder, M. I. Natham, and A.

Pinczuk, Phys. Rev. Lett.26, 642 ~1971!.156J. S. Lannin, J. Non-Cryst. Solids97&98, 39 ~1987!.157W. Paul, D. K. Paul, B. von Roedern, J. Blake, and S. Oguz, Phys. Rev.

Lett. 46, 1016~1981!.158F. C. Marques, R. G. Lacerda, M. M. Lima, Jr., and J. Vilcarromero,

Phys. Status Solidi B192, 549 ~1995!.159R. T. Sanderson,Chemical Bonds and Bond Energy~Academic, New

York, 1976!, Chap. 2.160S. Hasegawa, L. He, T. Inokuma, and K. Kurata, Phys. Rev. B46, 12478

~1992!.161G. M. Ingo, N. Zacchetti, D. della Sala, and C. Coluzza, J. Vac. Sci.

Technol. A7, 3048~1989!.162Z. Yin and F. W. Smith, Phys. Rev. B42, 3658~1990!.163D. Bolmont, J. L. Bischoff, F. Lutz, and L. Kubler, Appl. Phys. Lett.59,

2742 ~1991!.164C. D. Wagner, Faraday Discuss. Chem. Soc.60, 291 ~1975!.165C. D. Wagner in,Handbook of X-Ray and Ultraviolet Photoemission

Spectroscopy, edited by D. Briggs~Heyden & Son Ltd., London, 1978!.166A. R. Williams and N. D. Lang, Phys. Rev. Lett.40, 954 ~1978!.167L. Ley, R. Karcher, and R. L. Johnson, Phys. Rev. Lett.53, 710 ~1984!.168F. Urbach, Phys. Rev.92, 1324~1953!.169See, for instance, H. Curtins and M. Favre, inAmorphous Silicon and

Related Materials, edited by H. Fritzsche~World Scientific, Singapore,1988!, p. 329.

170J. Tauc, R. Grigorovici, and A. Vancu, Phys. Status Solidi15, 627~1966!.

171R. Tsu, P. Menna, and H. Mahan, Sol. Cells21, 189 ~1987!.172S. Hasegawa, T. Tsukao, and P. C. Zalm, J. Appl. Phys.61, 2916~1987!.173F. G. Bell and L. Ley, Phys. Rev. B37, 8383~1988!.174M. V. Kurik, Phys. Status Solidi A8, 9 ~1971!.175T. Tiedje, J. M. Cebulka, D. L. Morel, and B. Abeles, Phys. Rev. Lett.

46, 1425~1981!.176G. D. Cody, T. Tiedje, B. Abeles, B. Brooks, and Y. Goldstein, Phys.

Rev. Lett.47, 1480~1981!. See also, G. D. Cody in,Semiconductors andSemimetals, edited by J. I. Pankove~Academic, New York, 1984!, Vol.21B, p. 11.

177A. H. Mahan, P. Menna, and R. Tsu, Appl. Phys. Lett.51, 1167~1987!.178S. Aljishi, J. D. Cohen, S. Jin, and L. Ley, Phys. Rev. Lett.64, 2811

~1990!.179J. A. Howard and R. A. Street, Phys. Rev. B44, 7935~1991!.180D. Redfield, Phys. Rev. B130, 914 ~1963!; 140, 2056~1967!.181J. J. Hopfield, Comments Solid State Phys.1, 16 ~1968!.



143.106.108.149 On: Thu, 18 Jun 2015 18:49:29

182J. D. Dow and D. Redfield, Phys. Rev. B1, 3358~1970!.183H. Sumi and Y. Toyozawa, J. Phys. Soc. Jpn.31, 342 ~1972!.184J. D. Dow and D. Redfield, Phys. Rev. B5, 594 ~1972!.185J. Skettrup, Phys. Rev. B18, 2622~1978!.186S. Abe and Y. Toyozawa, J. Phys. Soc. Jpn.50, 2185~1981!.187M. Schreiber and Y. Toyozawa, J. Phys. Soc. Jpn.51, 1528~1982!; 51,

1537 ~1982!; 51, 1544~1982!.188J. Singh and A. Madhukar, Solid State Commun.41, 241 ~1982!.189D. Redfield, Solid State Commun.44, 1347~1982!.190C. M. Soukoulis, M. H. Cohen, and E. Economou, Phys. Rev. Lett.53,

616 ~1984!.191S. John, C. M. Soukoulis, M. H. Cohen, and E. Economou, Phys. Rev.

Lett. 57, 1777~1986!.192Y. Bar-Yam, D. Adler, and J. D. Joannopoulos, Phys. Rev. Lett.57, 467

~1986!.193Z. E. Smith and S. Wagner, Phys. Rev. Lett.59, 688 ~1987!.194C. T. Chan, S. G. Louie, and J. C. Phillips, Phys. Rev. B35, 2744~1987!.195M. Silver, L. Pautmeier, and H. Bassler, Solid State Commun.72, 177

~1989!.196S. John and C. H. Grein, Rev. Solid State Sci.4, 1 ~1990!.197H. M. Branz and M. Silver, Phys. Rev. B42, 7420~1990!.198M. Kemp and M. Silver, Appl. Phys. Lett.62, 1487~1993!.199A. Frova and A. Selloni, inTetrahedrally Bonded Amorphous Semicon-

ductors, edited by D. Adler and H. Fritzche~Plenum, New York, 1985!,p. 271.

200A. Skumanich, A. Frova, and N. M. Amer, Solid State Commun.54, 597~1985!.

201A. R. Zanatta and I. Chambouleyron, Phys. Rev. B53, 3833~1996!.202W. B. Jackson, S. M. Kelso, C. C. Tsai, J. W. Allen, and S. J. Oh, Phys.

Rev. B31, 5187~1985!.203A. R. Zanatta, M. Mulato, and I. Chambouleyron~unpublished!.204M. H. Brodsky, in Light Scattering in Solids, edited by M. Cardona,

Topics in Applied Physics~Springer, Berlin, 1975!, Vol. 8, 208.205M. Wakag, K. Ogana, and A. Nakano, Phys. Rev. B50, 10666~1994!.206J. Morell, R. S. Katiyar, S. Z. Weisz, H. Jia, J. Shinar, and I. Balberg, J.

Appl. Phys.78, 5120~1995!.207E. Bustarret and E. Morgado, Solid State Commun.63, 581 ~1987!.208E. Bustarret, F. Vaillant, and B. Hepp, Mater. Res. Soc. Symp. Proc.118,

123 ~1988!.209M. B. Schubert, H. D. Mohring, E. Lotter, and G. H. Bauer, IEEE Trans.

Electron Devices36, 2863~1989!.210T. Yashiro, J. Electrochem. Soc.119, 781 ~1972!.211F. C. Marques, PhD thesis, UNICAMP, Brazil, April 1989~unpublished!.212F. C. Marques and I. Chambouleyron, inDefects in Materials, MRS

Symposia Proceedings, edited by P. D. Bristowe, J. E. Epperson, J. E.Griffith, and Z. Lilienthal-Weber~Materials Research Society, Pittsburgh,1991!, Vol. 209, p. 555.

213Y. Takano, T. Sato, N. Kitaoka, and H. Ozaki, J. Non-Cryst. Solids55,325 ~1983!.

214W. Paul, S. J. Jones, and W. A. Turner, Philos. Mag. B63, 247 ~1991!.215G. D. Bagratishvili, Yu. N. Berozashvili, M. B. Dzhanelidze, and R. B.

Dzhanelidze, Sov. Phys. Dokl.32, 652 ~1987!.216N. F. Mott, Adv. Phys.16, 49 ~1967!.217M. L. Knotek, M. Pollak, T. M. Donovan, and H. Kurtzman, Phys. Rev.

Lett. 30, 835 ~1973!.218G. D. Bagratishvili, Yu. N. Berozashvili, M. B. Dzhanelidze, and R. B.

Dzhanelidze, Sov. Phys. Semicond.24, 618 ~1990!.219W. Kohn, in Solid State Physics, edited by F. Seitz and D. Turnbull

~Academic, New York, 1957!, Vol. 5.220The Properties of Natural and Synthetic Diamond, edited by J. E. Field

~Academic, London, 1992!.221Y. Yatsurugi, N. Akiyama, Y. Endo, and T. Nozaki, J. Electrochem. Soc.

B120, 975 ~1973!.222P. V. Pavlov, E. I. Zorin, D. I. Tetelbaum, and A. F. Khokhlov, Phys.

Status Solidi A35, 11 ~1976!.223E. I. Zorin, P. V. Pavlov, and D. I. Telel’baum, Sov. Phys. Semicond.2,

11 ~1968!.224A. B. Campbell, J. B. Mitchell, J. Schewchun, D. A. Thompson, and J. A.

Davies, Can. J. Phys.53, 303 ~1975!.225K. L. Brower, Phys. Rev. B26, 6040~1982!.226H. J. Stein, inOxygen, Carbon, Hydrogen, and Nitrogen in Crystalline

Silicon, MRS Symposia Proceedings, edited by J. C. Mikkelsen, Jr., S. J.Pearton, J. W. Corbett, and S. J. Pennycook~Materials Research Society,Pittsburgh, PA, 1986!, Vol. 59, p. 523.

227F. Berg Rasmussen, R. Jones, and S. O¨ berg, Phys. Rev. B50, 4378~1994!.

228C. A. Ammerlaan, inDefects and Radiation Effects in Semiconductors,Proceeding 59, edited by R. R. Hasiguti~Institute of Physics, London,1980! p. 81.

229G. G. DeLeo, W. B. Fowler, and G. D. Watkins, Phys. Rev. B29, 3193~1984!.

230S. T. Pantelides, Solid State Commun.14, 1255~1974!.231P. A. Schultz and R. P. Messmer, Phys. Rev. B34, 2532~1986!.232W. Spear and P. LeComber, Solid State Commun.17, 1193~1975!; Phi-

los. Mag.33, 935 ~1976!.233J. Robertson, Philos. Mag.40, 31 ~1979!.234R. A. Street, Phys. Rev. Lett.49, 1187~1982!.235J. C. Phillips, Phys. Rev. Lett.42, 1151~1979!.236J. Robertson and M. J. Powell, Appl. Phys. Lett.44, 415 ~1984!.237T. Shimizu, H. Kidoh, A. Morimoto, and M. Kumeda, Jpn. J. Appl.

Phys., Part 128, 586 ~1989!.238I. Chambouleyron and A. R. Zanatta, Appl. Phys. Lett.62, 58 ~1993!.239I. Chambouleyron and R. Campomanes, Phys. Rev. B53, 12566~1996!.240A. R. Zanatta, PhD thesis, Unicamp, May 1995~unpublished!.241R. Campomanes, MSc dissertation, Unicamp, September 1994~unpub-

lished!.242T. Drusedau, B. Schro¨der, and H. Oechsner, Solid State Commun.79,

799 ~1991!.243T. Drusedau, J. Non-Cryst. Solids137&138, 821 ~1991!.244J. Baixeras, D. Mencaraglia, and P. Andro, Philos. Mag. B37, 403

~1978!.245H. Watanabe, K. Katoh, and M. Yasui, Thin Solid Films106, 263~1983!.246B. Dunnett, D. I. Jones, and A. D. Stewart, Philos. Mag. B53, 159

~1986!.247D. V. Tsu, G. Lucovsky, and M. J. Mantini, Phys. Rev. B33, 7069

~1986!.248G. Lucovsky, M. J. Williams, S. S. He, S. M. Cho, Z. Jing, and J. L.

Whitten, MRS Symposia Proceedings, Vol. 336, edited by E. Schiff, M.Hack, A. Madan, M. Powell, and A. Matsuda~Materials Research Soci-ety, Pittsburgh, 1994!, p. 637.

249P. P. M. Venezuela and A. Fazzio, Phys. Rev. Lett.77, 546 ~1996!.250G. Marcano, A. R. Zanatta, and I. Chambouleyron, J. Appl. Phys.75,

4662 ~1994!.251A. Masuda, K. Itoh, M. Kumeda, and T. Shimizu, J. Non-Cryst. Solids

198–200, 395 ~1996!.252J. Reichardt, R. L. Johnson, and L. Ley, Physica B & C 117&118, 877

~1983!.253S. Hasegawa, T. Tsukao, and P. C. Zalm, J. Appl. Phys.61, 2916~1987!.254G. Hollinger and F. J. Himpsel, Appl. Phys. Lett.44, 93 ~1984!.255K. J. Gruntz, L. Ley, and R. L. Johnson, Phys. Rev. B24, 2069

~1981!.256F. Patella, F. Sette, P. Perfetti, C. Quaresima, C. Capasso, M. Capozi, A.

Savoia, and F. Evangelisti, Solid State Commun.49, 749 ~1984!.257D. Schmeisser, R. D. Schnell, A. Bogen, F. G. Himpsel, D. Rieger, G.

Landgren, and J. F. Morar, Surf. Sci.172, 455 ~1986!.258R. A. Riedel, M. Turowski, and G. Margaritondo, J. Appl. Phys.55, 3195

~1984!.259S. Hasegawa, M. Matsuda, and Y. Kurata, Appl. Phys. Lett.58, 741

~1991!.260S. Hasegawa, M. Mutuura, and Y. Kurata, Appl. Phys. Lett.49, 1272

~1986!.261J. Vilcarromero and F. C. Marques~private communication!.262R. A. Street, D. K. Biegelsen, W. B. Jackson, N. M. Johnson, and M.

Stutzmann, Philos. Mag. B52, 235 ~1985!.263A. Marimoto, M. Matsumoto, M. Yoshita, M. Kumeda, and T. Shimizu,

Appl. Phys. Lett.59, 2130~1991!.264J. B. Mitchell, J. Shewchun, and D. A. Thompson, J. Appl. Phys.46, 335

~1975!.265T. B. Karashev, R. M. Aranovich, A. A. Vaino, and A. A. Tali,Radiation

Physics of Nonmetallic Crystals, edited by N. N. Sirota~Nauka i Tekh-nika, Minsk, 1970! ~in Russian!, p. 174.



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