Nitrogen in Germanium I

Nitrogen in germanium I. 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

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[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 JOURNAL OF APPLIED PHYSICS VOLUME 84, NUMBER 1 1 JULY 1998

Nitrogen in germanium I. Chambouleyrona) Instituto de Fı´sica ‘‘Gleb Wataghin’’, Universidade Estadual de Campinas, P.O. Box 6165, 13083-970 Campinas, SP, Brazil A. R. Zanatta Instituto 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 1998͒ The known properties of nitrogen as an impurity in, and as an alloy element of, the germanium network are reviewed in this article. Amorphous and crystalline germanium–nitrogen alloys are interesting materials with potential applications for protective coatings and window layers for solar conversion devices. They may also act as effective diffusion masks for III-V electronic devices. The existing data are compared with similar properties of other group IV nitrides, in particular with silicon nitride. To a certain extent, the general picture mirrors the one found in Si–N systems, as expected from the similar valence structure of both elemental semiconductors. However, important differences appear in the deposition methods and alloy composition, the optical properties of as grown films, and the electrical behavior of nitrogen-doped amorphous layers. Structural studies are reviewed, including band structure calculations and the energies of nitrogen-related defects, which are compared with experimental data. Many important aspects of the electronic structure of Ge–N alloys are not yet completely understood and deserve a more careful investigation, in particular the structure of defects associated with N inclusion. The N doping of the a-Ge:H network appears to be very effective, the activation energy of the most effectively doped samples becoming around 120 meV. This is not the case with N-doped a-Si:H, the reasons for the difference remaining an open question. The lack of data on stoichiometric ␤-Ge3N4 prevents any reasonable assessment on the possible uses of the alloy in electronic and ceramic applications. © 1998 American Institute of Physics. ͓S0021-8979͑98͒07213-2͔

TABLE OF CONTENTS E. The Ge–N bond: randomness, charge transfer and electronegativity...... 14 I. INTRODUCTION...... 2 IV. OPTICAL PROPERTIES OF a-Ge ALLOYS.... 16 II. PREPARATION OF COLUMN IV NITRIDES. . . 3 A. Optical properties...... 16 A. Thin film deposition...... 3 B. Optical absorption in amorphous B. Column IV nitrides...... 3 1. Silicon nitride...... 3 semiconductors...... 16 2. Germanium nitride...... 4 1. Optical band gap ͑E04 and ETauc͒...... 16 3. Carbon nitride...... 4 2. Band-gap widening...... 17 4. Tin nitride...... 4 C. Electronic versus structural disorder in C. Other metal and semiconductor nitrides...... 4 amorphous semiconductors...... 17 D. Deposition conditions, particle bombardment, 1. Optical absorption and Urbach edge...... 17 hydrogenation and structure of group IV 2. Structural disorder...... 19 amorphous thin films...... 5 3. Composition, structural disorder and III. STRUCTURE OF a-Ge͑N͒ ALLOYS...... 7 A. Structural studies by EXAFS...... 7 optical band gap...... 19 B. Electronic structure...... 8 V. TRANSPORT PROPERTIES...... 20 1. Theoretical approaches and coordination A. H-free Ge1ϪxNx alloys: crystalline and defects...... 8 amorphous...... 20 2. Experimental reports...... 8 B. Hydrogenated a-Ge1ϪxNx films (xу0.01)..... 21 C. Structural studies by optical techniques...... 10 C. Nitrogen as an impurity in Ge (xр0.01)...... 22 1. Infrared spectroscopy...... 10 1. N in crystalline germanium...... 22 2. Raman spectroscopy...... 12 D. Hydrogenation, structure and stability of 2. N as an active dopant in amorphous Ge a-GeN films...... 13 ͑and Si͒...... 23 D. Photoconductivity in N-doped a-Ge:H films. . . 26 a͒Electronic mail: [email protected] VI. CONCLUDING REMARKS...... 27

0021-8979/98/84(1)/1/30/$15.00 1 © 1998 American Institute of Physics [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 2 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

lengths going from the red well into the ultraviolet, in the same way as the highly successful As-based and P-based Group III alloys used in the infrared, red and green regions of the spectrum. Of particular relevance to the present review is the role of nitrogen, from doping concentrations to the alloy phase, in Group IV elements: tetrahedral C, Si, Ge and ␣-Sn. The driving idea behind this presentation is to review all reported properties of nitrogen in germanium, either crystalline or amorphous. As the understanding of a rather new material improves when its properties are discussed in the context of those of similar ͑and better known͒ systems, we will fre- quently refer to nitrogen in silicon and, whenever appropriate, to nitrogen in tetrahedral carbon and ␣-Sn. By far the most studied among the above systems is 7 silicon nitride. The research on Si3N4 has been largely fu- eled by its use in microelectronics technology, where it acts FIG. 1. Optical gap of column III and IV nitrides as a function of bond length ͑see Refs. 2–6͒. Note the different bond length and gap energy scales. as an effective insulating material and a diffusion mask for The nitrides of heavy isocore elements display similar band-gap variations. impurities. These applications require the etching of the films, which proceeds isotropically for amorphous and highly microcrystalline insulator compounds. Etchants of I. INTRODUCTION technical importance are ammonium fluoride, buffered hy- 2 2 drofluoric acid at room temperature, hot 85% phosphoric The nitrogen atom, 1s 2s 2px2py2pz , can complete its valence shell in one of the following ways:1 acid for pattern etching with metal masks, and miscellaneous other etchants, usual strong mineral acid or bases.8 Table I Ϫ ͑i͒ electron gain to form the nitride ion N3 ; indicates the etching rate of different solutions on a-Si and ͑ii͒ electron-pair bonds ͑single or multiple͒, like in mo- a-Ge based nitrides. The electronic applications of silicon lecular N2; nitride have been complemented with ceramic uses, the im- Ϫ 9 ͑iii͒ electron-pair bonds with electron gain, as in NH2 ; portance of which has not ceased to grow in recent years. At ͑iv͒ electron-pair bonds with electron loss, as in tetrahe- the other end of the research efforts on Group IV nitrides lies dral ammonium. the SnN system, the properties of which are poorly understood at present.10 The molecules NR3 are pyramidal; the bonding involves sp3 orbitals so that the lone pair occupies the fourth position. Nitrogen in diamondlike carbon and carbon nitride CN The propensity of nitrogen, like carbon, to form p␲ – p␲ alloys are the subject of intense research. The interest stems multiple bonds is a feature that distinguishes it from phos- from at least two reasons. First, the possibility of preparing phorus and the other Group V B elements. Thus nitrogen as ␤-C3N4, a hypothetical material predicted to have a hardness 5 the element is dinitrogen, N2, with a very high bond strength comparable to that of diamond. Up to now, however, the and a short internuclear distance ͑1.094 Å͒, whereas phos- efforts made to synthesize ␤-C3N4 have not been very 11 phorus forms P4 molecules or infinite layer structures in successful. Second, nitrogen is a possible donor in dia- which there are only single bonds.1 mond. Substitutional nitrogen gives, besides a variety of The numerous covalent nitrides and their properties, as other not yet fully understood nitrogen related defects, a do- well as their potential applications, vary greatly depending nor level at 1.7 eV below the conduction band.12 Controlling on the element with which nitrogen is combined. For ex- the conductivity of diamondlike films by chemical doping ample, in recent decades an increasing interest developed in would open the possibility for high temperature solid-state the wurtzite polytypes of Group III nitrides: GaN, AlN, and electronics, a matter of interest for military and other appli- InN, which form a continuous alloy system, the band gaps cations. ranging from ϳ2 eV for InN, to ϳ3.5 eV for GaN, to Germanium nitride films13,14 were prepared in the 1960s ϳ6.5 eV for AlN ͑Fig. 1͒.2–6 Thus Group III N alloys could and 1970s but the subject did not develop consistently. potentially be fabricated into active optical devices at wave- In 198515 the Laboratory of Photovoltaic Research at

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

Etch rate at RT Deposition or

Material anneal temperature HF BHF H3PO4 H2SO4 NaOH

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

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UNICAMP, Campinas, started a program aimed at a system- a sputtering system is employed͒. It is to be expected, and atic study of the structural, compositional, electrical and op- indeed observed, that electrical, mechanical, and chemical tical properties of GeN alloys, and of N as an impurity in the properties vary with composition and deposition conditions. a-Ge:H network. Other research groups, particularly in Japan Both the dc- and rf-driven CVD methods depend largely and Germany, have contributed to unravel many puzzling on the plasma chemistry. As a result, the significant param- questions. In the present review, we report results on N in eters that can affect the characteristics and properties of a Ge, with particular emphasis on our own research at Campi- deposited film are more varied than in evaporated or sput- nas. tered material. The literature on a-SiN-based materials deposited by CVD-like plasma-assisted methods is quite ample and very interesting reviews can be easily found.20–23 Re- II. PREPARATION OF COLUMN IV NITRIDES ports on CVD-like deposited a-GeN materials, on the other A. Thin film deposition hand, are relatively scarce.24–28 The chemical vapor deposition ͑CVD͒ of thin films has In the physical vapor deposition ͑PVD͒ of thin layers, become one of the most important methods of film formation more energetic phenomena usually occur. Within the various and now constitutes a powerful and versatile tool in modern PVD-like methods, the sputtering technique is by far the technologies, such as those employed in solid-state electron- most widely used. Sputtering is the plasma-assisted vapor- ization of a material ͑called target͒ by bombarding it with ics. The reasons for the rapidly growing importance of CVD ϩ in the last few years lay primarily in its versatility for depos- high energy particles ͑normally Ar ions͒. Due to these en- iting a very large variety of elements and compounds at rela- ergetic collisions atoms or fragments are ejected from the tive low temperatures, in the form of both vitreous and crys- target’s surface and contribute to film formation. In addition talline layers having a high degree of perfection and purity. to the particles from the target, most of the gaseous species Other unique advantage of CVD over other methods of film present in the plasma undergo physical/chemical reactions formation is the relative ease of creating materials of a wide and determine important properties of the thin films being range of accurately controllable composition and layer struc- deposited. In current practice sputtering discharges are tures that are difficult or impossible to attain by other driven by high frequency ͑13.56 MHz͒ power supplies, and techniques.16 strictly speaking, there is no cathode or anode since the net Roughly speaking, CVD can be defined as a material flow of charged species to each electrode is zero. Due to the synthesis method in which the constituents ͑in the vapor very different mobility of electrons and ions and difference phase͒ react to form a solid film at some surface.17 Thus the in electrode area, however, a negative dc bias develops on occurrence of chemical reactions is an essential characteristic the powered electrode. The election for rf-instead of dc- of CVD. In order to understand CVD processes, one must driven systems is determined by the convenience of the low know which chemical reactions occur in the reactor and to pressure at which rf plasmas operate as well as by the kind of what extent. Furthermore, the effects of process variables material ͑e.g., insulator͒ to be produced. This advantage such as temperature, pressure, input concentrations, and flow combined with capacitive coupling allows the reactive depo- rates on these reactions must be understood. The basic con- sition of insulating and semiconducting films ͑a-SiN and figurations of CVD systems have evolved along the years, a-GeN, for example͒ under highly controlled conditions. Most of the pioneering work on the reactive sputtering of the most important ones being the so-called plasma-assisted 29–36 a-SiN:͑H͒ thin films used gaseous ArϩN2 mixtures in- CVD method. 37 stead of ArϩNH3. The sputter deposition of a-GeN:͑H͒ The uniqueness of the plasma to generate chemically 15,38–40 reactive species at low temperatures is due to the nonequi- films started just a few years ago. librium nature of the plasma state. By nonequilibrium, we mean a gas plasma typically sustained at ϳ0.1 to several B. Column IV nitrides Torr, which exhibits temperatures of the free electrons of tens of thousands of degrees Kelvin, while the temperature 1. Silicon nitride of the translational and rotational modes of free atoms, radi- Thin Group IV nitride films can be deposited by CVD cals, or molecules will be only hundreds of degrees Kelvin. from a variety of precursor gases. In thermal processes, Illuminating discussions about the fundamental aspects and Si3N4 films have been prepared at high substrate tempera- characteristics of low temperature plasma species can be tures ͑1000–1200 K͒ from silane SiH4 and ammonia NH3 in 8,18,19 41,42 found in the specialized literature. a hydrogen atmosphere and from dichlorosilane SiCl4 One of the prime motivating factors in utilizing plasma and ammonia at low pressure.43,44 Plasma deposition of deposition processes is that the substrate temperature can be Si3N4 films from SiH4,N2and/or NH3 has been studied ex- kept relatively low, typically 500 K or lower. Conventional tensively and is now at the stage of production applications CVD processes require temperatures substantially higher that in semiconductor device manufacture, mainly for the passi- may be inappropriate for certain substrate materials or device vation of device surfaces. Si3N4 films 3000–10 000 Å thick structures. The films that are deposited by plasma reactions are excellent diffusion masks for alkali ion contaminants and are usually amorphous in nature, with short-range structural other impurities. The recognition of the potential of low- ordering only. The composition of the films can be varied in temperature deposited nitrides in semiconductor applications a controlled way by changing some key plasma parameters, is not new, and the delay in the widespread use of this ma- such as reactant gas flow ratios ͑and/or target composition if terial has resulted, in part, from the lack of equipment ad- [This article is copyrighted as indicated in the article. 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13,65 64 equate to the film uniformity and production requirements of and NH3 at 650–850 K, from GeH4 and NH3 at 900 K, 66 the industry. The chemistry of plasma Si3N4 deposition is and GeH4 and hydrazine at 600 K. Plasma-assisted CVD extremely complex and the detailed reaction kinetics is not has also been used to deposit GeN films at substrate tempera- 27,28 fully understood, largely because such deposition reactions tures of 500–650 K using GeH4 /NH3 or GeH4 /N2. How- are difficult to diagnose, as opposed to processes taking ever, most of the present understanding of GeN alloys de- place exclusively in the gas phase. Plasma-assisted silicon rives from films deposited by the rf sputtering technique. nitride produced via plasma is different from silicon nitride produced by conventional CVD or PVD techniques. In the 3. Carbon nitride former, the composition can be controlled varying the ratio The interest in hydrogenated diamondlike carbon films of the reactant gas flow, the power level, the substrate tem- (a-C:H͒ stems from their unique electrical, chemical and perature during deposition, and the pressure in the reactor mechanical properties. In the past few years the effect of the vessel. Remote and direct plasma-enhanced CVD techniques incorporation of dopants in these films has been studied. In use SiH and either NH or N as the precursor gases to 4 3 2 particular, since it was proposed that the bulk modulus of the deposit silicon nitride.45–49 Disilanyl amine50 and cyclopen- hypothetical material ␤-C N , structurally analogous to tadienyl substituted silanes ͑Cp SiH and Cp SiH ͒ in com- 3 4 3 2 2 ␤-Si N , may be similar to that of diamond,67 the incorpo- bination with N /NH have also been employed.51 Finally, 3 4 2 3 ration of nitrogen into carbon films has received special at- electron cyclotron resonance ͑ECR͒ plasmas ͑which produce tention. After an early study on rf sputtered CN films, carried a higher degree of ionization than normal rf or microwave out by Cuomo et al. in 1979,68 attempts to incorporate nitro- plasmas͒ have recently been used to deposit silicon nitride. gen into carbon films were made using several techniques: rf In ECR deposition processes, SiH and N /NH or 4 2 3 or dc sputtering,69,70 laser ablation,71 and ion-beam assisted Si͑NMe ͒ H and N are used as precursors and films are 2 3 2 deposition.72 These films are amorphous, and for some par- deposited at substrate temperatures of typically 350–600 K ticular deposition conditions the presence of ␤-C N nanoc- with growth rates of 50– 200 Å minϪ1.52–54 3 4 rystals embedded into the amorphous matrix was determined.70,71 The incorporation of nitrogen improves the 2. Germanium nitride mechanical properties of diamondlike a-C films, in particular 73 Crystalline germanium nitride layers produced by ther- the tribological ones like wear and friction. Nitrogen incorporation up to 20 at % in a-C:H films was achieved through mally activated chemical methods were found to contain 74–78 55 plasma-assisted CVD of different gases mixtures. It was Ge3N2 and Ge3N4 as constituents. Ge3N2 is obtained by the following chemical reactions:56 shown that a-CN:H films deposited by rf plasma decompo- sition of CH –N mixtures can be as hard as a-C:H films, NH3 4 2 and that the incorporation of nitrogen causes a reduction of GeI2 —— ͑GeNH͒nϩNH4I, ͑1͒ → the internal stress without any noticeable change in their and hardness.79,80 A possible explanation for this effect is the 550– 600 K reduction of the mean atomic coordination number, and thus 3͑GeNH͒ —— Ge N ϩNH . ͑2͒ → 3 2 3 a reduced over-constraint due to the presence of threefold N Ge3N2 can also be formed when germane is decomposed by substituting fourfold coordinated C atoms. active nitrogen:57 400– 600 K 4. Tin nitride 3GeH ϩN* —— Ge N ϩ6H . ͑3͒ 4 2 → 3 2 2 There has been little mention of tin nitride films in the Similarly, crystalline Ge3N4 can be prepared from el- literature. Remy and Hantzpergue used reactive cathodic 10 emental Ge reacting with NH3 or starting from Ge͑NH͒2 sputtering to prepare the first nitride films. Crystalline SnN groups by one of the following processes:58,59 films have been recently prepared by reactive sputtering of 81 1000 K 700 K Sn in a N2 plasma, and by magnetron sputtering of a pure GeϩNH3 —— Ge3N4 ͓Ge͑NH͒2͔n ͑4͒ Sn target in a gas mixture of Ar and N2 at room → ← temperature.82 The only CVD SnN reported to date is the in which case, c-Ge3N4 has the hexagonal structure of 58,60 deposition of polycrystalline material by atmospheric pres- phenacite (␤-Ge3N4), the lattice parameters being: a sure CVD from Sn͑NMe2͒4 and NH3 at substrate tempera- ϭ0.8038 nm; cϭ0.3074 nm. There is also a rhombic struc- 83 61 tures of 500–700 K. ture attributed to c-Ge3N4 with lattice parameters a ϭ1.384, bϭ0.906 and cϭ0.818 nm. The density14,62 of Ϫ3 63 C. Other metal and semiconductor nitrides c-Ge3N4 is ␳ϳ5.3gcm and its heat of formation Ϫ⌬Hϳ65 kJ molϪ1. Other possibilities for the preparation Nitrides of a number of metals are of technical impor- of nonstoichiometric GeN thin films, using thermally acti- tance, particularly those of Ti, Zr, Hf, Nb, Ta, and Be ͑and 64 vated CVD methods, are: GeH4 and NH3 ͑at ϳ800 K͒; more recently B, Al, and V͒, which form hard and highly GeCl4 and NH3 ͑with temperatures in the 700–900 K stable refractory materials with very high melting points. 13,65 66 range͒ and GeH4 and N2H4 ͑at ϳ700 K͒, to mention Widely used, the applications of several of these coatings just a few of them. ͑especially TiN͒ include the improvement of the wear resis- 84–86 Stoichiometric (Ge3N4) germanium nitride films have tance of cemented carbide tools, and protective coatings. been mostly prepared by thermally activated CVD of GeCl4 The chemical systems for producing CVD refractory nitrides [This article is copyrighted as indicated in the article. 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TABLE II. Particle bombardment in the formation of thin ﬁlms according to different authors. Depending on the adopted deposition conditions either electrons, charged or neutral particles can be associated to the bombardment of the thin ﬁlm surface during growth.

Author Event Particle Consequences Observation

power delivered: a Ϫ2 Brodie general rf electrons substrate heating Wionϳ30 mW cm Ϫ2 sputtering (sp) WeϪϳ500 mW cm

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

ϩ PAr determines Ar ions are Rossc rf sp a-Si:H Arϩ ions structural accelerated by the features ﬂoating potential

PAr determines electron bombardment Moustakasd rf sp a-Si:H electrons or Arϩ substrate leads to high quality bombardment ﬁlms

Ar and Si species hydrogen atoms can Tardye 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.

6 are based on reacting volatile metal halides with N2, with or applications. In addition to the above mentioned thin film without H2, at high temperatures, i.e., 2700–3000 K for the nitrides and respective methods of deposition, a great effort ZrCl4–N2 or HfCl4–N2 systems, and 1300–2100 K for the is made nowadays to prepare N-based new materials with 87 6 TiCl4–N2–H2 system. Temperatures in the range of 1100– methods other than CVD. 1400 K, however, can be used to deposit TiN films by the latter reaction.88 Films of hafnium nitride HfN, from the re- D. Deposition conditions, particle bombardment, action of chlorides of Hf with N2–H2, have been deposited on W, also at relatively low temperatures in the range of hydrogenation and structure of group IV amorphous thin films 1200–1600 K, and thermodynamic and kinetic studies showed that the process is mass transport limited at low H2 Regardless of the chemical/physical processes involved, concentrations.89 surface adatom mobility ranks among the most important Boron nitride ͑BN͒, prepared under optimal CVD condi- deposition parameters in plasma-assisted deposition meth- tions, is an excellent insulator and p-type dopant source for ods. Besides the thermal energy associated with the substrate junction formation in Si device fabrication. Several CVD temperature, extra energy can be supplied to the growing methods have been investigated for depositing thin BN films film by bombarding it with energetic particles. The process of both microcrystalline and amorphous structure.90,91 Clear enhances adatoms mobility and may induce the removal of vitreous BN films are deposited on a variety of substrates at loosely bonded particles, producing denser films with no co- 90 98 850–1250 K by reacting B2H6 and NH3. The electrical lumnar microstructure. In other words, under certain con- conductivity of these films shows, however, that BN depos- ditions, particle bombardment may cause film etching, re- ited at 1050 K on Si is similar to Si3N4 but is not as good a moving preferentially weakly bonded atoms and/or Naϩ barrier and not as stable chemically.90 protrusions that develop into columns. Ion, neutral species Another refractory nitride of interest is aluminum nitride and/or electron bombardment of the growing surface add en- ͑AlN͒, which has potential as a good dielectric for active and ergy to adatoms and also heat the substrate. The precise iden- passive components in semiconductor devices because of its tification and the role played by these different bombarding large energy gap and high thermal stability.92 Aluminum ni- species is not an easy task and depends very much on the tride films have been deposited by reacting NH3 with AlCl3 deposition method and conditions. Table II reviews briefly 92–95 or Al͑CH3͒3. the ongoing debate on the nature and importance of the bom- Films of gallium nitride ͑GaN͒ have been grown origi- barding species affecting most of the structural properties of nally by the ammonolysis of gallium monochloride GaCl.96 rf sputtered a-Si:H films in an Ar atmosphere.99–103 97 Pyrolysis of the GaBr3–NH3 complex, and ammonolysis of A systematic study on the effects of particle bombard- 95 Ga͑CH3͒3 have also been employed. As for AlN, the main ment on the properties of rf sputtered a-Ge:H does not exist. interest in GaN stems from the possibility of high- However, the importance of particle bombardment to the temperature electronics and short-wavelength optical structural and electronic properties of this semiconducting [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 6 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

FIG. 2. Optical gap (E04)vs͑a͒the target bias and ͑b͒ the substrate temperature of rf sputtered a-GeN:H thin ﬁlms. As seen, the optical gap ͑or nitrogen content͒ is inﬂuenced by both the target bias and substrate temperature ͑see Ref. 108͒.

alloy has been known for a decade.104–107 In particular rf sputtering was one of the ﬁrst methods used to produce a-Ge:H ﬁlms of improved quality.104 Good quality a-Ge:H

samples are currently deposited under varied PVD-like FIG. 3. Optical gaps (E04) of rf sputtered a-Si:H and a-Ge:H thin films conditions.106,107 ͑deposited under similar conditions and in the same deposition system͒ as a Just to illustrate the importance of the adatom mobility function of hydrogen content. Note the band-gap widening mechanism induced by hydrogenation. on the optical properties of the a-GeN:H films consider Figs. 2͑a͒ and 2͑b͒, where the effects of the substrate temperature and of the rf power delivered to the plasma during deposition and temperature, for example͒ while others reflect chemical 108 are presented. The influence of the dc bias and substrate aspects. temperatures ͑adatom mobility͒ on the E04 optical gap ͑pho- Closing this brief introduction on plasma-assisted depo- ton energy at which the absorption coefficient reaches sition of N-containing thin films, it is worth mentioning that 4 Ϫ1 10 cm ͒ of the films is apparent in the figure. device quality a-Si:H ͑and alloy͒ films are best produced by The electronic properties of a-semiconductors depend using CVD-like or soft deposition methods,109 like the an- largely on hydrogenation. Device quality a-Si:H ͑a-Ge:H͒ odic glow-discharge of the SiH4 technique. The sputtering films have a density of dangling bond on the order of technique has been much less used to prepare a-Si:H, al- 1015 cmϪ3 (1017 cmϪ3), as determined from electron spin resonance ͑ESR͒ data. In nonhydrogenated a-Si and a-Ge this density is ϳ1019 cmϪ3. The role of hydrogen, however, is not only to passivate dangling bonds, but also to reduce the strain of the Si ͑Ge͒ network improving the optoelectronic properties. Bonded hydrogen in these a-structures also determines the optical gap because the valence band recedes with hydrogenation. Roughly speaking, the optical gap scales with the hydrogenation. The method/condition of deposition may determine some upper limits ͑optical gap or hydrogen content values, for example͒. Figures 3͑a͒ and 3͑b͒ show the optical gap E04 for a-Si:H and a-Ge:H samples deposited in the same rf sputtering system, in an ArϩH2 atmosphere. Both sample series were deposited under similar conditions but possess quite different degrees of hydrogenation. The incorporation of nitrogen in a-Si and a-Ge occurs in a way similar to that of hydrogen. Most samples considered in this review were prepared by the rf sputtering method. In them the nitrogen concentration is proportional to the N2 or NH3 partial pressure during deposition. Figure 4 shows the nitrogen content of a-GeN and a-SiN alloys, as determined from a deuteron induced nuclear reaction analysis 14 15 FIG. 4. Nitrogen concentration ͑as determined from nuclear reaction analy- ͓ N(d,p) N͔, vs the N2 or NH3 partial pressure during sis͒ in rf sputtered group IV amorphous thin films vs the N2 or NH3 partial deposition. Most of the features present in Fig. 4 are associ- pressure. Within the studied range the nitrogen concentration scales linearly

ated with the deposition conditions ͑substrate bombardment with the N2 and NH3 partial pressures. [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 7

FIG. 5. A basal projection of the phenacite structure. ␤-phase germanium

nitride is hexagonal with two molecular units (2Ge3N4) in the unit cell. The FIG. 6. Magnitude of the Fourier transforms of EXAFS for a-Ge1ϪxNx :H space group is P63 /m. The fractions indicate the distance to the basal plane and c-Ge3N4 samples ͑Ref. 121͒. The arrows indicate expected positions of in c-axis units. features due to possible first- and second-shell configurations, once phase shifts 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, though a-Si:H films of electronic quality have been also from lower to higher interatomic distance. The right-hand side panel shows reported.110–112 Contrary to the case of a-Si:H films, PVD- the four possible second-shell configurations in a-Ge1ϪxNx :H and the interatomic distances expected. like or hard deposition methods involving more energetic processes during deposition, like sputtering deposition techniques and cathodic glow-discharge, proved to be the most appropriate for the deposition of optimized a-Ge:H. The rea- (Ge1ϪxNx) has been retained in the present review. For the sons for this difference are not yet understood, but it is clear sake of clarity, the data from the literature expressed as by now that they may be at the origin of the difficulties 0рxр1.33 will be transformed into the 0рxр0.57 nota- found in the preparation of optimized a-SiGe:H alloys. Yet, tion. Electron-diffraction measurements on a-Si1ϪxNx :H ob- it is important to keep in mind that each deposition method 116 and condition produces thin films with quite specific charac- tained by the glow discharge of SiH4 and NH3 mixtures teristics, some of them being directly correlated to the physi- yield radial distribution functions ͑RDFs͒, the peaks of cal processes occurring during growth. which are a combination of those expected for the first- and second-nearest neighbors in c-Si3N4 and Si. More structural information is obtained by x-ray diffraction117 and neutron III. STRUCTURE OF a-GE N ALLOYS „ … scattering118 on nearly stoichiometric samples deposited by It is convenient to report the structural properties of new CVD. The RDF of the a-phase was found to resemble that of binary alloys in the broad context of well characterized the ␤-phase c-Si3N4 and bond angles of ϳ109.8° and model systems of known stoichiometry and similar valence ϳ121° were found around Si and N, respectively.116 More- coordination. For GeN alloys, the obvious choice is c-Si3N4, over, analysis of small-angle scattering pointed to the exis- an intensively studied electronic material. Two structures of tence of voids, the presence of which reduces slightly the 113,114 c-Si3N4 have been identified by x-ray diffraction. They coordination numbers of Si and N. Extended x-ray absorp- have been named ␣ and ␤ phase, both having a hexagonal tion fine structure ͑EXAFS͒ studies indicate that off- crystal lattice. The coordination numbers for Si and N are 4 stoichiometric a-Si1ϪxNx compounds are chemically and 3, respectively. Figure 5 shows the phenacite structure, ordered.119 as it is called, of crystalline ␤-Si N . It involves covalent 3 4 A. Structural studies by EXAFS bonds between planar bonded N and tetrahedral bonded Si.115 This local structure is consistent with Si sp3 hybrid The structure of GeN compounds has been much less orbitals while N bonding is explained in terms of a linear studied than that of SiN alloys. As expected from the similar combination of p orbitals, the planar geometry being given valence structure of Ge and Si, the structure of c-Ge3N4 is by a strong repulsion of nonbonded Si atoms. According to similar to c-Si3N4. Stoichiometric ␤-Ge3N4 has a density ␳ this picture the s electrons of N do not participate in bonding ϭ5.3gcmϪ3 and a nitrogen concentration of 22 Ϫ3 14,120 121 and the top of the valence band of ␤-Si3N4 is filled by non- 4.6ϫ10 atoms cm . Boscherini et al. studied by bonding N pz electrons. EXAFS ͑Ge K-edge͒ the local order of c-Ge3N4 and of five Off-stoichiometric SiN alloys can be prepared by plasma a-Ge1ϪxNx :H samples with N concentrations ranging from assisted chemical vapor deposition ͑PACVD͒ adjusting the xϭ0uptoxϭ37 at. %. Figure 6 shows the Fourier trans- partial pressure of N2 or NH3 in the reaction chamber. The form of EXAFS signals corresponding to a-Ge:H; composition of off-stoichiometric SiN alloys is given in the a-Ge1ϪxNx :H and c-Ge3N4 samples. The arrows at the bot- literature with two different notations. Some authors prefer tom of Fig. 6 indicate the expected position of features due to use Si1ϪxNx , whereas others indicate composition using to first and second shell configurations, once phase shifts are SiNx . In the former notation 0рxр0.57 (4/7), whereas in taken into account. Quantitative data analyses performed in k the latter 0рxр1.33 (4/3). The notation Si1ϪxNx space indicate that the full EXAFS signal in a-Ge1ϪxNx :H [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 8 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

can be completely simulated by a first shell component due across the network. In a-Si3N4, the silicon dangling bond to Ge and N and a second shell component of Ge in a Ge–N (SiBD) gives rise to states 3 eV above the VB edge, whereas configuration, the relative importance of these contributions the nitrogen dangling bond (NDB) produces a strongly local- varying according to the N concentration. An interesting ized state close to the VB edge. Large N concentrations in- conclusion of this work is that the presence of Ge–N–N can crease the disorder and broaden the VB tail, which is due to be excluded, an indication that when N is inserted in the variations in second neighbor N–N interactions.123 a-Ge:H network its bonding to Ge is determined mostly by There is a clear difference between the behavior of de- Ge–N interactions.121 The constant value found for the fects of Si-rich and of N-rich a-Si1ϪxNx :H alloys. In Si-rich Ge–Ge second shell distance, equal to the average value in material the neutral SiDB produces a half-filled level near c-Ge N , is a further confirmation of this picture. Another 3 4 mid-gap, as in a-Si:H, whereas the energy level of NDB is important feature is that, irrespective of the amount of N in embedded deep in the VB. In c-ora-Si3N4, however, a the samples, interatomic distances remain constant: highly localized p-like NDB appears just above the VB maxi- RGe–Nϭ1.835 Å, RGe–Geϭ2.44 Å, and RGe–N–Geϭ3.19 Å. mum. The roughly sp3 hybridized neutral dangling bond of a The root-mean-square N bond-angle fluctuation was found to trivalent Si site ͑Si surrounded by 3 N, called K0 center͒ is be Ͻ6° in all samples, a value smaller than the minimum located around 3 eV above the VB maximum.130 The simul- fluctuation for bond angles on Ge ϳ9°, suggesting that taneous presence of both types of DBs should cause a charge bonds on N are more rigid than on Ge. Finally, the fact that 0 0 ϩ Ϫ transfer from SiDB to NDB giving: K ϩN ϭK ϩN .At the second shell EXAFS signal can be fitted with Ge in a room temperature a small K0 signal is measured by ESR Ge–N–Ge configuration gives further evidence that when N suggesting that SiDBs are in excess. In Si-rich alloys enters the a-Ge network it exclusively form bonds with Ge (xϾ0.47) the K0 center is more stable in its diamagnetic and thus induces a totally chemically ordered network, in 122 configuration and there is strong experimental evidence in close agreement with the findings on a-Si1ϪxNx :H. favor of the K0 center having a negative correlation energy.131 B. Electronic structure Makler et al.126 calculated the electronic structure of a-Ge1ϪxNx :H alloys, in particular the structure of the VB, 1. Theoretical approaches and coordination defects the band-gap widening mechanism for increasing N concen- The electronic structure of stoichiometric H-free trations, and the energy levels of gap states due to GeDB and a-Si1ϪxNx and a-Si1ϪxNx :H alloys has been studied exten- to NDB. Besides the expected smaller band gap for similar N sively from a theoretical and from an experimental point of concentration, the calculations on a-Ge1ϪxNx :H indicate no view.123–125 The only calculation on the electronic structure major qualitative differences with respect to the electronic 126 126 of germanium nitride has been done by Makler et al. structure of a-Si1ϪxNx :H. Tight binding calculations pre- Compared to a-Si1ϪxNx , only a small experimental research dict a preferential attachment of H to N in a-Ge1ϪxNx :H effort has been made on the electronic structure of alloys, in agreement with experimental evidence.132,133 Re- 127–129 a-Ge1ϪxNx alloys. The available results indicate that garding the energy levels of coordination defects Makler 126 the gross features of the electronic structure of both elemen- et al. find the GeDB at the center of the a-Ge:H gap, where tal semiconductor nitrides is very much alike, as expected it remains for N concentration increasing up to xϭ0.4. from the similar valence configuration of Si and Ge. How- Above xϭ0.4 the GeDB disappears from the band gap. The ever, some differences appear concerning the band gap wid- calculations do not find NDB states in the gap of 126 ening mechanism with increasing nitrogen, the nature and a-Ge1ϪxNx :H. the energy level of defects, and the structure of the valence ESR is a powerful tool to obtain structural information band. on coordination defects. The strength of the ESR absorption Let us first summarize the electronic structure of band is proportional to the density of paramagnetic electrons 123–125 a-Si1ϪxNx . Tight binding calculations of c-Si3N4 in- and the ESR signal structure gives information about local dicate that the top of the valence band ͑VB͒ consists of N bonding. The most useful information comes from the g fac- 2p␲ nonbonding states, the conduction band ͑CB͒ being tor of the paramagnetic center and the hyperfine interaction, given by Si–N anti-bonding orbitals.123,126 The main valence which depend on the nuclear magnetic properties of the at- band consists of three peaks separated by a gap from a deep oms around the center. The neutral ͑paramagnetic͒ dangling fourth peak of N s states. The three main peaks of the VB are bonds in hydrogenated amorphous Si and Ge have assigned ͑increasing binding energy͒ to N p␲, a mixture of gϭ2.0055 and gϳ2.019 signatures, respectively.134,135 N p and Si p bonding states, and N p and Si s states, re- The nature of coordination defects in H-free and hydro- spectively. The results of calculations are in good overall genated a-Ge1ϪxNx alloys has been studied by ESR and agreement with the photoemisison data of Ka¨rcher et al.124 light-induced ESR in rf magnetron sputtered samples.136–138 Theory predicts that the band gap of a-Si1ϪxNx should in- It has been found that both the g value and the linewidth of crease slowly with increasing N up to xϭ0.42, with the band the GeDB decrease with increasing N concentration, as shown edges opening almost symmetrically about mid-gap, and in Fig. 7. The effects have been attributed to the large an- then rapidly until xϭ0.57 ͑stoichiometric composition͒. Hy- isotropy of the g value for Ge in H-free samples,137 and in drogenation is found effective in widening the band gap in hydrogenated samples to the existence of separated Ge-rich Si-rich alloys (xϽ0.42) only, i.e., the N concentration at and N-rich phases in films with large x.138 Note that this last which Si–Si bonds fail to form continuous percolation paths interpretation is at odds with EXAFS data which indicate a [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 9

FIG. 7. g value and linewidth of the resonant ESR broad line of a-Ge1ϪxNx alloys as a function of increasing N2 partial pressure in the deposition chamber ͑Ref. 137͒.

chemically ordered ͑stricto sensu͒ material.121 The origin of FIG. 8. Ultraviolet photoelectron spectra of pure a-Ge and a-Ge1ϪxNx samples ͑xϭ0.11, 0.31, and 0.36͒. The spectra have been obtained using He the satellites in the resonance spectra is due to hyperfine I ͑21.2 eV͒ and He II ͑40.8 eV͒ photons. A recession of the VB maximum interaction with N nuclei, as confirmed by the different hy- and new electronic states are apparent for samples with large N contents perfine structure of GeN samples made using 15N ͑nuclear ͑Ref. 129͒. spin Iϭ1/2͒ and 14N(Iϭ1).137 These preliminary ESR reports indicate the need for more detailed investigations, in order to gain a deeper understanding of coordination defects tron spectra of a-Ge and a-Ge1ϪxNx ͑xϭ0.11, 0.31 and of a-Ge1ϪxNx alloys. Let us consider here now the main differences between 0.36͒ samples. A recession of the VB maximum is apparent the electronic structure of both elemental semiconductor ni- in Fig. 8 ͑left͒. Two N-related features can be recognized trides. clearly in the figure. The peak at binding energy ͑BE͒ 126 ϳ5.2 eV is close to the value determined for the nonbonding ͑a͒ Makler et al. find that the band gap Eg of 124 a-Ge N remains unaltered up to N concentrations of N2p␲lone pair observed for a-Si1ϪxNx alloys, and has 1Ϫx x probably the same origin. As discussed above, two additional xϳ0.5, and that for xϾ0.5, Eg grows quickly to reach the peaks have been reported for a-Si1ϪxNx , one with a BE value Egϭ4.7 eV (xϭ0.57). This is qualitatively different 123 ϳ7.5 eV and the other at energies ranging between 11.4 and from Robertson’s slow linear increase of Eg in a-Si1ϪxNx 12.4 eV, which were attributed to bonding N 2px,y orbitals for 0ϽxϽ0.4, and a faster increase rate at larger x values. 124 Note that neither for a-Ge N , nor in the case of having some contributions from Si 3s and 3p orbitals. 1Ϫx x Following this interpretation the feature at BE ϳ10.9 eV in a-Si1ϪxNx samples does the calculated gap opening rate cor- 124,129 Fig. 8 is probably due to bonding N 2px,y states with contri- respond to the variations measured by photoemission. 129 ͑b͒ Nitrogen is a group V element and, as an impurity, it butions from either Ge 4s or Ge 4p states. Comedi et al., may be an active dopant in the a-Si:H and a-Ge:H networks. however, do not provide a conclusive interpretation of this 126 experimental finding which does not correspond to the three Makler et al. find that a N impurity is a deep donor in 126 a-Si:H, the anti-bonding state laying at 0.45 eV below the VB peaks predicted theoretically by Makler et al. The Fermi energy and the VB maximum have been bottom of the CB. The calculation performed for N impurity 129 in the a-Ge:H network gives a too small energy difference found to experience a small upward shift in energy with increasing x up to xϳ0.25 ͑see Fig. 9͒, a behavior different between the bottom of the CB and the N donor level and the 124 authors cannot conclude whether N impurity is a shallow from that reported for a-Si1ϪxNx , in which the VB maxi- 130 ϩ mum recedes linearly with increasing x. The reason for this donor or not. Robertson’s calculation locates the N4 state just below the CB minimum. The calculations, however, do difference is not yet understood. The sudden asymmetrical not agree with the experimental results on N-doped a-Si:H widening of the band gap in a-Ge1ϪxNx samples for 139–141 xϾ0.22 deduced from Fig. 9 is similar to that reported for and a-Ge:H. We discuss the problem in more detail in 124,125 the section on electronic transport and N doping. silicon nitride compounds. In both elemental semiconductor nitrides the band-gap opening starts at a N concentration when N 2p␲ states dominate the VB maximum. It was 2. Experimental reports found experimentally that, similar to a-Si:H, hydrogenation The structure of the VB and its evolution as a function of does not control the optical gap of a-Ge1ϪxNx :H alloys N content (0ϽxϽ0.36) was investigated in H-free above a well defined N concentration ͑see Fig. 9͒. Note, 129 a-Ge1ϪxNx films by ultraviolet photoelectron spectros- however, that the N concentration at which the sudden open- copy ͑UPS͒.142 Photoelectrons were analyzed with an energy ing of the band gap occurs is significantly lower in GeN than resolution of 0.2 eV, the light source being He I ͑21.2 eV͒ in SiN alloys, i.e., xϳ0.25 and xϳ0.5, respectively. This and He II ͑40.8 eV͒ photons. Figure 8 shows the photoelec- different behavior, related to the nature of the dominant or- [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 10 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

FIG. 10. Absorption IR bands vs wave number of a-Ge1ϪxNx samples. Note the appearance of new absorption bands as the N content in the film in- FIG. 9. ͑a͒ Energy position of the valence band maximum EV , the Fermi creases ͑see Ref. 39͒. energy EF and the conduction band edge EC ͑resulting from the addition of E04 to the VBM͒ as a function of N content in a-Ge1ϪxNx ͑see Ref. 129͒. ͑b͒ Same kind of representation for a-Si1ϪxNx ͑see Ref. 124͒. ͑a͒ The integrated absorption A(␻st)ϭ͐st͓␣(␻)/␻͔d␻ is proportional to the N concentration, i.e., ͓N͔ϭKGe–N 18 Ϫ2 bital at the VB maximum, has not yet been explained. To A(␻st), with KGe–Nϭ5ϫ10 atoms cm . A deuteron- 14 15 some extent, it might be the consequence of different meth- induced nuclear reaction ͓ N(d,p) N͔ was used by the 148 ods used to estimate the concentration of nitrogen. present authors to determine the absolute N concentration in a series of a-Ge1ϪxNx samples. The method allows the determination of N concentrations as low as C. Structural studies by optical techniques 1019 atoms cmϪ3 in samples of typical thickness 10Ϫ4 cm ͑see Fig. 11͒. The existence of a calibration constant facili- 1. Infrared spectroscopy tates the rapid determination of alloy composition from op- The analysis of lattice or network vibration modes has tical measurements. proved to be an efficient tool in structural studies. ͑b͒ The presence of back impurity atoms of a different The structure of a-Ge1ϪxNx and a-Ge1ϪxNx :H alloys electronegativity induces a shift of the peak energy of dipole has been studied with infrared ͑IR͒ and Raman vibrations.149,150 It is worth mentioning here that the absence spectroscopies.24–26,39,40,143–146 Next we discuss the IR spec- of back atom͑s͓͒or dangling bond͑s͔͒ also influences the tra of H-free a-Ge1ϪxNx films of varying composition. charge distribution around the Ge–N dipole. In a-Ge1ϪxNx Crystalline Si and Ge lattices have no first-order IR ab- alloys the frequency shifts are associated with changes in the sorption. In amorphous Si and Ge, however, the lack of long- Ge–N interatomic distance and their analysis sheds light on range order relaxes the crystal momentum and the symmetry the structural modifications produced by the incorporation of rules prohibiting first-order absorption and, in principle, all varying amounts of N, or of other light impurities, like hy- vibration modes can contribute to the first-order absorption process. In other words, within a disordered network all vibration modes can be active in the IR. In the case under study, the introduction of N atoms into the a-Ge network enhances the disorder and induces a considerable charge transfer from Ge to N producing stronger IR absorption bands related to the Ge–N bond structure. The group formed by a planar bonded N atom and its three Ge neighbors is the skeletal Ge3N group. Its normal vibration modes are:147 a breathing mode, an out-of-plane stretching mode, the symmetric and asymmetric in-plane stretching mode, and an in-plane bending mode. The asymmetric stretching vibration involves the displacement of the N atom and of the three Ge neighbors and is strongly IR active. Its strength has been found to be proportional to N concentration.39,148 Figure 10 shows the IR absorption spectra of six a-GeN samples as a function of wave number in FIG. 11. Integrated absorption of the main in-plane Ge–N stretching vibra- Ϫ1 tion mode vs N content for a-Ge1ϪxNx samples ͑see Ref. 39͒. The slope of the 200– 1600 cm range. The main absorption band peak- 18 Ϫ2 Ϫ1 the straight-line fit is KGe–Nϭ5.0ϫ10 cm . The filled circles indicate the ing at ϳ700 cm is the in-plane asymmetric stretching Ϫ1 132,133 integrated area of the absorption bands centered at about 690 cm . The vibration. The analysis of this absorption band as a open triangles represent the integrated area of all stretching vibrations ͑690, function of N concentration indicates that: 870, and 1100 cmϪ1͒. [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 11

FIG. 12. Different chemical environments of Ge atoms suggested to induce FIG. 13. Simplified skeletal representation for the trigermylamine ͑GeH3͒3N frequency shifts of the main Ge–N dipole vibration ͑see Ref. 39͒. molecule ͑top͒. The possible vibrations of the Ge3N group in the a-Ge host are: ͑a͒ symmetric in-plane stretching vibration mode (ϳ370 cmϪ1); ͑b͒ asymmetric in-plane stretching vibration mode (ϳ850 cmϪ1); ͑c͒ symmetric ͑pyramidal configuration͒ stretching mode; and ͑d͒ symmetric near in- drogen or deuterium ͑to be discussed in a coming section͒. plane ͑pyramidal configuration͒ stretching mode. Vibrations ͑c͒ and ͑d͒ are Inductive effects due to increasing nitrogen are detected in characteristic of gaseous trigermylamine ͑see Refs. 39 and 152͒. the main stretching vibration mode. The appearance of new absorption bands at ϳ870 and ϳ1100 cmϪ1 ͑see Fig. 10͒ have been ascribed to Ge–N stretching vibrations in an environment including N atoms and dangling bonds in different trigermylamine molecule possesses three H atoms bonded to bonding configurations. Figure 12 sketches the different con- each Ge. A charge redistribution around the Ge atoms is figurations which, according to our interpretation,39 give rise expected because H is more electronegative than Ge.153 As a to the three bands just discussed. consequence, the stretching vibrations should be more ener- ͑c͒ The absorption features appearing at ϳ450 and getic in the molecule than in an isolated Ge–N dipole. As the Ϫ1 300 cm ͑see Fig. 10͒ have been associated with the sym- N concentration increases in a-Ge1ϪxNx , however, the pres- metric stretching and breathing modes involving the disor- ence of N back atoms leads to a charge redistribution around dered a-Ge3N skeletal group, respectively. Ge atoms which goes in the direction of the trigermylamine The above mentioned vibration modes were identified structure. This fact explains the reasonable agreement of comparing the IR features of a-Ge1ϪxNx alloys with those of peak energies of the asymmetric in-plane stretching vibra- well known gaseous compounds and investigating their tions in both the gaseous and the solid phases ͑850 and Ϫ1 strength and energy position dependence on alloying. A 870 cm , respectively͒ of the Ge3N skeletal group. compound that allows us to establish a useful analogy with The absorption band at ϳ450 cmϪ1 associated with the the a-Ge3N skeleton of GeN alloys is trigermylamine in-plane symmetric stretching mode of a-Ge1ϪxNx is found Ϫ1 ͑GeH3͒3N, an extremely unstable and reactive gaseous at 370 cm in trigermylamine. The energy difference may 151,152 species. The structure of ͑GeH3͒3N is shown in Fig. 13, originate from the nonplanar configuration in ͑GeH3͒3N which also displays its main vibration modes and the corre- which involves the out-of-plane motion of N. Finally, the sponding energy peaks. Note that, besides the absence of a dependence of the absorption strength of the 300 cmϪ1 fea- connective network, the main difference between the Ge3N ture on N concentration in a-Ge1ϪxNx lead us to attribute the Ϫ1 skeleton of a-Ge1ϪxNx films and ͑GeH3͒3N is the presence 300 cm vibration to a disorder-induced breathing mode of of H atoms as Ge bond terminators. Briefly, the vibration the a-Ge network, absent in trigermylamine. properties of trigermylamine indicate that: ͑i͒ the Ge atoms A linear relationship between absorption strength and are probably noncoplanar with N, a valence force field composition with no extra bands at high alloy ratios, as well calculation151 giving a bond angle of 116° at N. ͑Note that as a constant band shape, means that just the density of di- the vibrations of the pyramidal site for a threefold- pole oscillators is varying with composition. Similarly, an coordinated N are qualitatively similar to the vibrations of N equal shift of the absorption peak energy with composition in a planar configuration, both having four characteristic vi- for two different absorption bands, or a linear relationship bration modes involving the displacement of N and its three between maximum absorption coefficients for a couple of Ge neighbors͒; ͑ii͒ there are two IR absorption bands in bands, indicate that both have a common origin. Figure 14 151,152 ͑GeH3͒3N which do not move on deuteration: a very shows the dependence of the absorption strength on alloying strong and broad band at ϳ850 cmϪ1, attributed to the asym- of the asymmetric in-plane Ge–N stretching mode Ϫ1 metric stretching mode of the Ge3N structure ͓Fig. 13͑b͔͒ (ϳ700 cm ), the Ge–N symmetric stretching mode and a weak band at ϳ370 cmϪ1 originating from the sym- (ϳ450 cmϪ1) and the Ge–N breathing (ϳ300 cmϪ1) metric stretching mode of trigermylamine ͓Fig. 13͑a͔͒. modes. As can be seen, the overall absorption increases on The vibration frequencies of the skeletal Ge3N structure alloying, an indication that N is associated with the three in solid a-GeN and gaseous ͑GeH3͒3N differ because the modes. [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 12 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

FIG. 14. Absorption coefﬁcient at peak energy vs N concentration in the

solid phase of the stretching and breathing modes of a-Ge1ϪxNx alloys. Note the dependence of all vibration modes on nitrogen concentration ͑see FIG. 15. Reduced Raman spectra of some a-Ge͑N͒ samples with increasing Ref. 39͒. N concentrations. The features corresponding to the TA- (80 cmϪ1) and TO-like (270 cmϪ1) vibration modes are indicated ͑see Ref. 145͒.

All the inductive effects measured in IR absorption originate from charge transferred from Ge to N atoms. X-ray reaches 0.37, i.e., just above the maximum N concentration photoelectron spectroscopy ͑XPS͒ measurements performed of the samples reported in Ref. 145, the Ge–Ge TO mode on the same a-Ge N samples128,154 corroborate the over- 1Ϫx x disappears while the Ge–N absorption band becomes very all picture of charge transfer deduced from IR spectra. strong. There is wide agreement on the fact that the full width at 2. Raman spectroscopy half maximum ͑FWHM͒ height of the TO-like vibration Raman spectroscopy is a well developed probe tech- mode (⌫TO) and the ITA /ITO amplitude ratio are useful indi- nique to study the structural features of a-semiconductor al- cators of structural disorder. Changes of the measured ⌫TO loys. In a Raman process the incident light frequency is correspond to modifications of the short-range order, particu- shifted by an amount equal to the natural frequency of the larly of the bond angle distribution.156 In the study of Zanatta vibration mode involved in the inelastic scattering. The ab- et al.,145 the laser energy ͑2.6 eV͒ used to record the Raman sence of long-range order favors the observation of all vibra- spectra is coupled mainly to valence and conduction states tions usually forbidden in crystalline lattices by selection associated with Ge–Ge bonds, the Ge–N bonds having a rules. As the short-range order in a-semiconductors is similar much higher energy separation ͑the band gap widens as ͓N͔ to that in the crystalline parents, the phonon density of states increases, as discussed in the following section͒. Thus the is expected to be similar to that of the crystalline phase with changes in frequency and bandwidth of the TO-like mode some broadening due to disorder in bond length and bond reflect basically the distortion of the Ge–Ge bonds with in- angle.155 In addition to structural information, Raman scat- creasing N concentration. For small nitrogen contents, tering measurements in a-semiconductors can give insight ͓N͔Ͻ1022 atoms cmϪ3, the symmetry of the bonds is deter- into the bond type and the nature of disorder. A Raman study mined mostly by the sp3-like character of the Ge–Ge bond. of the structural changes produced by different N concentra- The changes of ⌫TO become more important for tions in the a-Ge network was recently reported.145 Figure 15 ͓N͔Ͼ1022 atoms cmϪ3 and are attributed to the drastic bond shows the reduced Raman spectra of pure a-Ge and of two angle distortions due to changes in the character of the domi- 3 a-Ge1ϪxNx samples. In the case of pure a-Ge the spectrum nant bond, from essentially sp -like ͑Ge–Ge͒ to some com- is dominated by the contributions of the transverse acoustic bination between sp3- and sp2- ͑or p-͒ like Ge–N orbitals. Ϫ1 ͑TA͒-like mode at ϳ80 cm and from the transverse optic The structural changes in the a-Ge1ϪxNx network as the N ͑TO͒-like mode at ϳ270 cmϪ1. The incorporation of N up to content increases were also probed by photothermal deflec- around 1022 atoms cmϪ3 does not affect essentially the shape tion spectroscopy ͑PDS͒, which indicates a linear relation- of the spectrum, except for a slight broadening of the TO- ship between electronic and structural disorder. ͑The point like mode, related to increased topological disorder. For will be discussed in detail in Sec. IV.͒ The interpretation has larger N concentrations, however, radical changes are de- been confirmed also by photoelectron emission studies made tected. The TA- and the TO-like modes broaden consider- on the same samples.129 ably and are also shifted to higher energies. In addition, a Figure 16 shows the shifts of the IR vibration and Ra- broad structure appears at ϳ700 cmϪ1 corresponding to the man scattering signal induced by nitrogen. In this figure, in-plane asymmetric stretching vibration referred to in the both ␻IR and ␻Raman are represented as a function of the E04 previous paragraph. Xu et al.26 reported Raman spectra of optical gap, which scales with N content. The changes can be a-Ge1ϪxNx :H films deposited by plasma-enhanced chemical understood on the basis of a lightening and a stiffening of the vapor deposition ͑PECVD͒. These authors found that when x network. Roughly, ␻IR and ␻Raman are proportional to [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 13

FIG. 16. Frequency shifts of the Ge–N in-plane asymmetric stretching ␻IR and of the TO-like (␻Raman) vibration modes as a function of the optical parameter E04 ͑which depends on N concentration͒͑see Ref. 145͒. The asymmetric stretching vibration and TO-like modes display a similar behavior. Also represented are theoretical data from Barrio et al. ͑see Ref. 146͒ showing the expected increase of the TO-like mode frequency. Large N contents are responsible for a lightening and a stiffening of the network. The dashed straight lines are linear regression fits of experimental data. FIG. 17. ͑a͒ Intrinsic compressive stress vs optical gap E04 ͑ϰN content͒ of a-Ge1ϪxNx :H films deposited by rf sputtering, and ͑b͒ thermal expansion coefficient vs optical gap E04 of some of the a-Ge1ϪxNx :H films of ͑a͒͑see Ref. 158͒. (k/␮)1/2, where k is an effective force constant associated with Ge–Ge ͑and Ge–N͒ bonds, and ␮ is the corresponding reduced mass. As N is added to the a-Ge host, heavy Ge atoms are substituted by light N. In addition, Ge–N bonds X-ray inspection shows that as-deposited hydrogenated 133 are stronger than Ge–Ge bonds ͑EGe–Nϳ2.65 eV and Ge1ϪxNx films are amorphous. Honma et al. investigated 1 EGe–Geϳ1.94 eV͒, producing a stiffening of the network. the structural stability of a-Ge1ϪxNx :H films by studying Note that the interpretation is consistent with results of crystallization processes by differential thermal analysis EXAFS,121 which indicate that root mean square N bond ͑DTA͒ and differential thermal gravimetry ͑DTG͒. Heating angle fluctuations are smaller than the minimum fluctuation up the films under controlled conditions induces the crystal- for bond angles on Ge, an indication that Ge–N bonds are lization of the alloy, revealed by the Ge͑111͒ diffraction more rigid than Ge–Ge. peak. The crystallization temperature Tc increases with N concentration and varies from ϳ635 K for a-Ge:H up to D. Hydrogenation, structure and stability of a-GeN above 925 K for a film containing ϳ40 at. %N. There is a films weight decrease on heating, attributed to the out-diffusion of Hydrogen atoms reduce significantly the density of dan- H and ͑probably͒ Ar, always present in rf-sputtered material. gling bonds of group IV a-semiconductors and, thus, play a Similar qualitative results were obtained by Xu et al.26 on fundamental role in their structural and electronic properties. a-Ge1ϪxNx :H films deposited by PACVD using GeH4 and Another effect of H is to break weak bonds during the depo- NH3 gas mixtures. Thermal annealing of the films up to 875 sition process. Thus, from the structural point of view, the K shows that the inclusion of N improves the structural sta- inclusion of monovalent H reduces the coordination of Si or bility and increases the crystallization temperature. For films Ge and helps in relaxing the over-coordinated network. To a corresponding to xϭ0.44 no crystallization signal can be de- lesser extent, a similar relaxation process is induced by ni- tected by Raman spectroscopy after 30 min anneal at 875 trogen, N having a coordination number smaller than Ge. K.26 Summarizing, nitrogen improves considerably the ther- Hydrogenation of a-Ge1ϪxNx has several consequences, ex- mal stability of a-Ge1ϪxNx films. tending from growth rate to structure and electronic trans- The mechanical properties of a-Ge1ϪxNx :H films, in- port. Of particular relevance to the present review is the cluding stress, thermal expansion, and elasticity as a function bonding of H to N40,132,133 and the likely influence of N–H of temperature and N concentration, have been investigated bonds on deposition rate, network relaxation, and electronic by Marques et al.158 Figure 17͑a͒ indicates that as-deposited transport. The absorption bands associated with hydrogen films have compressive stress ranging from approximately 40 and deuterium bonded to Ge and N have been identified. ϳ1.0 to ϳ7.0 kbar as the E04 optical gap increases from These results suggest that ‘‘preferential attachment’’ 157 of H about 1.0 to 3.0 eV ͓͑N͔ϳ30 at. %, by NRA͒. The compres- 132,133 to N takes place when H atoms are adsorbed on the sive stress, as found in sputtered a-Ge1ϪxNx :H films, is as- growing surface, an indication that the different bonding en- sociated with the presence of impurities, to hydrogenation, ergies of N–H ͑4.5 eV͒ and Ge–H ͑3.8 eV͒ measured in and to Ar inclusion. It gives stable and compact films, in molecules hold qualitatively in the a-network. agreement with the studies of DTA.133 Tensile stress, on the [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 14 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

FIG. 18. Ge 3d core level signals of samples containing increasing amounts of nitrogen x, as indicated in the figure. The centroid of the main contribution 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 levels FIG. 19. Experimental photoelectron spectra of the Ge 3d core level associated with Ge–N bonds experience an apparent chemical shift as N ͑ ͑pluses͒ and the corresponding fitting with Gaussian curves representing the increases , the signal associated with Ge–O bonds of the surface oxide does ͒ different Ge–N components ͑solid lines͒ for a couple of samples ͑see Ref. not shift significantly ͑see Ref. 127͒. 127͒. ͑a͒ Pure a-Ge with a native oxide overlay. The deconvolution represents the Ge 3d5/2 (EBϳ29.2 eV) and the Ge 3d3/2 (EBϳ29.8 eV) core levels corresponding to bulk Ge–Ge bonds, and the contribution of native Ox Ge–O bonds (EB ϳ33 eV). ͑b͒ The deconvolution of the a-Ge0.89N0.11 other hand, is characteristic of unstable films with structural sample displays the main Ge–N and Ge–O bonding components. Neither in defects, such as columnar structure, voids, etc. The thermal ͑a͒ nor in ͑b͒ was the core level signal associated with the Ge oxide de- expansion coefficient decreases with nitrogen ͓see Fig. convoluted in its spin orbit split components. 17͑b͔͒, a consequence of the structural changes of the alloy, which evolves from a diamondlike toward a hexagonal struc- X-ray photoelectron spectroscopy ͑XPS͒ is a very sensi- ture ͑phenacite͒ of smaller overall coordination. tive probe of the atomic environment of a selected element, the binding energy of its core electrons depending on the atomic surroundings. The chemical shifts induced by N in E. The Ge–N bond: randomness, charge transfer and the Ge 3d core levels have been studied by XPS and by electronegativity x-ray excited Auger electron spectroscopy ͑XAES͒.128,154 Charge transfer is an important effect to be considered in The surface analysis techniques probe the outermost atomic SiN and GeN alloys. Electronegativity was originally pro- layers and provide useful information on the atomic environ- posed by Pauling, as a parameter that would allow prediction ment of a given species and its influence on the macroscopic of the approximate polarity of a partially covalent bond.153 properties.142 Consider Fig. 18 where the Ge 3d core level The organic chemistry of Ge shows many similarities to that signals of H-free a-Ge1ϪxNx samples containing different of Si, in some chemical reactions Ge appearing more elec- amounts of N are represented.127 On these ex situ samples no tronegative than Si. Nitrogen has an electron affinity and an chemical or physical cleaning processes were used in order electronegativity greater than Si and Ge.153 Therefore a N to avoid possible preferential sputter etching of species. The atom in the Si ͑Ge͒ network gets an extra electronic charge signal associated with a surface oxide Ge–O, apparent in coming from Si ͑Ge͒ neighbors, the charge of which will be Fig. 18, does not shift as the N concentration increases in the less than 4 by an amount which depends on the number of N alloy. neighbors. Since Pauling’s original proposal a number of The bands associated with the presence of N are a maxi- new definitions of electronegativity and the ways to evaluate mum of five, corresponding to Ge units having 0, 1, 2, 3, and them have appeared. Sanderson’s electronegativity scale,159 4 bonded N atoms. The additivity of the chemical shifts has which besides the atom’s ‘‘ability’’ to hold their own elec- been assumed, the line shapes being the same for each com- trons tightly together taking into account the electronic con- ponent. The fitting procedure is illustrated in Fig. 19 for 127 figuration, gives: ␹Hϭ3.55; ␹Cϭ3.79; ␹Nϭ4.49; a-Ge and a-Ge0.89N0.11. The present authors suggested ␹Siϭ2.84; ␹Geϭ3.59. These figures indicate that Ge is the that it was possible to analyze the subsurface information of element having the closest electronegativity to that of H. As the Ge 3d despite the presence of a thin oxide surface layer. a consequence, the Ge–H bond is less polar than the Si–H Similar to the case of the parent SiN alloys160,161 the analysis bond, which is consistent with the experimentally found of the Ge–N bond distribution indicates a fair agreement preferential attachment of H to Si in a-SiGe:H alloys.157 with the theoretical prediction of the random bonding [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 15

FIG. 21. XPS shift of Si 2p and Ge 3d core levels as a function of the corresponding modified Auger parameter ⌬␣Ј. The possibility of calibrating XPS shifts in terms of the modified Auger parameter is apparent from the figure ͑see Ref. 128͒.

concept.159 Among the various electronegativity definitions and scales being reported in the literature, the one promoted by Sanderson,159 which has been successful in explaining the FIG. 20. Ge 3d and Si 2p core level binding energies and full width at half IR transmission spectra of a-Si alloys,150 was retained. The maximum ͑FWHM͒ height of the XPS signal as a function of N concentra- preceding considerations suggest the possible use of ⌬␣Ј to tion in a-Si1ϪxNx and a-Ge1ϪxNx samples. The data of FWHM correspond to the whole signal. Open squares Ref. 127; filled triangles Ref. 124; define an electronegativity scale. In order to test this idea, ͑ ͒ ͑ ͒ 128 ͑open triangles͒ Ref. 161. Note that both EB and FWHM exhibit similar Zanatta et al. plot ͑as in Fig. 21͒ the XPS shift versus ⌬␣Ј behavior in Ge and Si hosts, except for the values of the maximum FWHM, of some a-Si- and a-Ge-based alloys. The resulting linear and the N concentration corresponding to this maximum ͑from Ref. 127͒. correlation indicates first, the possibility of calibrating XPS shifts in terms of partial charges and second, that ⌬␣Ј might

162 Ϫ1 furnish a reasonable scale of electronegativities in such non- model. A chemical shift per bond ⌬Eb bond ϳ0.5 stoichiometric compounds. Ϯ0.1 eV for the Ge–N bond is found, a value intermediate Finally, an interesting correlation is found considering between the shifts induced by H and O on Ge. A similar UPS the chemical shift induced in the elemental semiconductor 129 study by Comedi et al. gives a slightly smaller chemical core levels by different impurities. Figure 22 displays the shift per bond ⌬E bondϪ1ϳ0.30Ϯ0.08 eV. These values Ϫ1 b chemical shift per bond (⌬EB bond ) for Si and Ge have to be compared with the binding energy shift of hosts154 as a function of the electronegativity of the foreign 159 ϳ0.6 eV measured per Si–N bond in a-Si1ϪxNx atom ͑Sanderson’s scale͒. Note the important direct corre- 124,163 alloys. spondence found between transferred charge and electrone- Figure 20 compares the position and shape of the Si 2p gativity for a collection of quite different systems and from and Ge 3d core levels of a-SiN and a-GeN alloys as a function of nitrogen concentration. The shifts reported in the figure correspond to the centroid of the core level signal and scale in a way similar for both Si and Ge alloys. Figure 20 also displays the FWHM of the XPS signal of the Si 2p and Ge 3d core level signals. Note the asymmetric bell-like shape in both cases, which merely reflects the Si–N and Ge–N bonding distribution. However, two differences appear: ͑i͒ the maximum FWHM is smaller for a-GeN than for a-SiN alloys and, ͑ii͒ this maximum occurs at different N concentrations. A different FWHM maximum may be the consequence of the different chemical shift per bond of Si–N and Ge–N bonds, but why the maximum FWHM occurs at different composition is not yet understood. Using experimental data of XPS and XAES core level 128 shifts, Zanatta et al. determined the modified Auger pa- FIG. 22. Chemical shift per bond (⌬E bondϪ1) as a function of the foreign 164–166 B rameter shift ⌬␣Ј, which is exempt from problems atom’s electronegativity ͑Sanderson’s scale͒. The error bars are associated inherent to the interpretation of XPS and XAES shifts. ⌬␣Ј with some dispersion of reported data and impurity content. The straight values were used to estimate the charge transferred on alloy- lines are linear regression fits of the available data. ͑a͒ SiH, see Ref. 252, ͑b͒ SiN, see Ref. 124, ͑c͒ SiN, see Ref. 253, ͑d͒ SiN, see Ref. 161, ͑e͒ SiO, see ing (⌬nGe). A proportionality between ⌬␣Ј and ⌬nGe was Ref. 173, ͑f͒ SiO, see Ref. 254, ͑g͒ SiF, see Ref. 255, ͑h͒ GeH, see Ref. 256, found, reminiscent of the idea of the local electronegativity ͑i͒ GeN, see Ref. 154, ͑j͒ GeO, see Ref. 257, ͑k͒ GeO ͑see Ref. 258͒. [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 16 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

FIG. 24. ͑a͒ Photothermal deﬂection spectroscopy ͑PDS͒ curves of some a-Ge-based thin ﬁlms ͑the same represented in Fig. 23͒. The optical band gap E04 and the Urbach energy E0 are indicated for each sample. ͑b͒ (␣nE)1/2 vs photon energy for a series of a-Ge-based compounds.

FIG. 23. Transmittance spectra of some a-Ge-based thin films in the NIR- VIS and mid-IR energy regions. Note that both the presence ͑and concen- tration͒ of hydrogen and nitrogen atoms drastically affect the optical trans- of localized states extending into the pseudo-gap. The width parency of these compounds. of the tails determines the most important optoelectronic characteristics of an a-semiconductor and depends on the degree of disorder as well as on the bonding character of the data of different authors. The figure suggests the possibility master states. of a new electronegativity scale, with electronegativity The optical absorption in a-semiconductors is usually 154 Ϫ1 values of Si and Ge ͑corresponding to ⌬EB bond ϭ0͒ of reported in three different energy regions: ͑i͒ a high photon ␹Siϳ2.6 and ␹Geϳ3.1 in agreement with the idea that, in energy region involving transitions between extended elec- some cases, Ge atoms are more electronegative than Si. tron states, ͑ii͒ an intermediate photon energy region corresponding to transitions from or to localized states, and ex- IV. OPTICAL PROPERTIES OF a-Ge ALLOYS hibiting an exponential-like dependence on photon energy. ͓The tailing of these localized states ͑or band edges͒ is char- A. Optical properties 168 acterized by a parameter called the Urbach energy E0 , The introduction of N atoms in a group IV a-network usually determined from PDS data169͔, and ͑iii͒ a sub-gap leads to a new class of materials having quite new optical, absorption region which is associated with: ͑1͒ transitions electronic and structural properties with potential technologi- between deep defect states and the VB or CB, ͑2͒ local vi- cal applications. Large impurity contents ͑alloying range͒ in- brations involving lighter atoms; and ͑3͒ the resonant modes duce important changes of the optical, electronic and struc- of the host network. tural properties of group IV amorphous semiconductors. Hydrogen, for example, relaxes the network and provokes a band-gap widening, a consequence of the recession of the 1. Optical band gap (E and E ) top of the VB.167 On the other hand, small concentrations of 04 Tauc column III and V impurities ͑doping rangeϽ1at%͒provoke The presence of localized states at energies between va- important changes in the electronic properties of group IV lence and conduction bands, makes the optical gap Eg an a-semiconductors. Figure 23 shows the transmission spectra ill-defined parameter in a-semiconductors and, consequently, in the near infrared-visible ͑NIR-VIS͒ and mid-infrared several ways are currently in use to define Eg . The simplest ͑mid-IR͒ energy ranges of some a-GeN thin films with dif- one is to consider Eg as the energy corresponding to an ab- 4 Ϫ1 ferent N and H contents. As seen increasing amounts of N in sorption coefficient ␣ϭ10 cm (E04). Another usual defi- the a-Ge network widen the pseudo-gap. The effects of N nition of the optical gap in an a-semiconductor is the so- 170 doping on the transport properties of a-Ge:H and a-Si:H are called Tauc’s gap ETauc , given by the energy where the discussed in Sec. V. (␣nE)1/2 or (␣E)1/2 vs E plot goes to zero, where n is the index of refraction and E the photon energy, i.e.: 1/2 1/2 B. Optical absorption in amorphous semiconductors ͑␣nE͒ ϭB ͑EϪETauc͒. ͑5͒ The breakdown of the electron k wave vector selection The B1/2 Tauc parameter includes information on the convo- rule for optical transitions in amorphous solids entails an lution of the VB and CB states, and on the matrix element of optical response which is free of fine features, like van Hove optical transition, which reflects not only the relaxed k se- singularities and well-defined band edges. Moreover, with lection rule but also the disorder-induced spatial correlation disorder the abrupt band edges of crystals broaden into tails of optical transitions between the VB and CB. Formally, B1/2 [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 17

1/2 FIG. 25. Optical band gaps ͑E04 and ETauc͒ and Tauc parameter B of a series of a-GeN alloys. The values corresponding to an a-Ge:H sample has FIG. 26. Optical band gap (ETauc) vs the nitrogen content of a-GeN:͑H͒ been also represented. Note the band-gap widening apparent for nitrogen ͑Refs. 39 and 108͒ and a-SiN:͑H͒͑Ref. 259͒ thin ﬁlms. Qualitatively concentrations higher than ϳ1022 atoms cmϪ3 ͑see Ref. 39͒. samples exhibit the same behavior, i.e., a sharp optical band-gap widening at larger N contents.

depends on the product of the oscillator strength of the opti- ally associated with static and thermal structural disorder, is cal transition, the deformation potential, and the mean devia- 174–179 tion of the atomic coordinates.171 not yet well understood despite extensive experimental and theoretical174,180–198 investigations. Part of the problem, Figures 24͑a͒ and 24͑b͒ display the optical gaps E04 and in fact, lies in the precise identification furnished by the dif- ETauc of a-GeN compounds. The optical gap and the Urbach ferent experimental techniques used to probe it. Cody energy depend on the presence ͑or not͒ of H as well as on N. 176 1/2 et al. have proposed a model for a-Si:H assuming the The optical gaps E04 and ETauc and the B Tauc parameter for a-GeN samples with different N contents are represented equivalence of thermal and structural disorder. For the al- in Fig. 25. As seen, important changes in the optical data loys, however, the characteristic energy of the exponential occur for N concentrations higher than ϳ1022 atoms cmϪ3. absorption edge increases to large values and additional terms, which take into account the compositional disorder The optical gap of a-GeN alloys does not depend linearly on 199,200 the nitrogen concentration. due to alloying and hydrogenation, must be considered. Today, the absorption tail is considered to reflect the VB and 2. Band-gap widening CB edge joint density of states ͑DOS͒. It represents then, the The mechanism of band-gap widening of group IV disorder-induced broadening of the bands. The physical ori- a-semiconductors due to alloying with lighter atoms ͑H, C, gin of localized states in group IV a-semiconductors is at- N, O, etc.͒ is well reported in the literature, with tributed to the existence of strained bonds, always present in particular emphasis on a-SiN- and a-SiO-based over-coordinated materials. According to the picture, the alloys.20,147,124,130,172,173 a-network must be strained locally in order to accommodate The overall effect of N atoms on the optical band gap of the atoms in a nonperiodic array. To these strained bonds a-Ge:͑H͒ and a-Si:͑H͒ compounds are represented in Fig. correspond electronic states near the band edges, the density 26. According to the figure, there is a threshold N concen- of which fall off exponentially with energy. Optical transi- tration ͑which is different for Si and Ge a-hosts͒ above tions to and from localized tail states would originate the which important changes of the optical band gap occur. The Urbach edge in the absorption spectrum. In spite of the ap- phenomenon was considered when discussing the electronic peal of this qualitative description, it is difficult to make structure of a-Ge͑N͒ alloys and is related to the substitution predictions based solely upon it. of Si–Si ͑or Ge–Ge͒ bonds by stronger Si–N ͑or Ge–N͒ 1. Optical absorption and Urbach edge bonds. As the N content increases, a N lone-pair band devel- 201 ops and dominates the VB maximum as stoichiometry is In a recent publication the present authors discuss the relationship between the characteristic energy of the expo- approached. The largest optical gap is obtained for the sto- 1/2 ichiometric compound, as expected. On the contrary, for nential absorption edge E0 and the Tauc parameter B of a-SiN- and a-GeN-based alloys ͑Fig. 27͒. No correspon- smaller N contents Si–Si ͑or Ge–Ge͒ bonds dominate the 1/2 VB maximum. A similar optical gap widening mechanism is dence was experimentally found between B and E0 for found in Si and Ge carbides and oxides. small N concentrations ͑less than a few at. %͒. Note that, in all cases, chemical doping provokes a noticeable broadening of the Urbach energy. In contrast, in the GeN and SiN alloy C. Electronic versus structural disorder in regime a correlation between B1/2 and E is found. This ex- amorphous semiconductors 0 perimental finding reflects the structural changes induced by All disordered solids display an exponential-like absorp- elements of different atomic coordination, i.e., the character tion edge, i.e., ln ␣ϰE ͑photon energy͒. The microscopic ori- of the bonding orbital at the top of the VB as the N content gin of the optical absorption tail in a-semiconductors, usu- increases. A connection between the exponential absorption [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 18 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

1/2 FIG. 27. Urbach E0 energy vs the B Tauc parameter for a-Si- ͑Ref. 260͒ and a-Ge-based ͑Ref. 201͒ thin films under different impurities concentra- FIG. 28. Urbach energy E as a function of the FWHM of the TO-like 1/2 0 tions. There is no clear relationship between E0 and B at low impurities vibration mode ⌫ ͑Raman scattering͒ of N-based a-Si ͑Refs. 207–209͒ 1/2 TO concentration. In contrast, E0 and B scale linearly in the alloying regime. and a-Ge ͑Ref. 145͒ thin films. Even for the lowest N concentration ͑doping regime͒ electronic disorder is present in these a-hosts. At higher N contents E0 and ⌫TO increase. tail and the short-range potential fluctuations produced by charged impurities is suggested in Fig. 27. Photoemission spectra, much more sensitive than optical techniques al- consequently, the absorption coefficient depend on both the though not susceptible to selection rules, indicate that this is joint DOS and the transition matrix elements: the case: N concentrations higher than ϳ10 at. % give rise 2 to important changes in the VB extrema of a-SiN and a-GeN ⑀2͑E͒ϰR ͑E͒͑DOSJ͒, ͑7͒ compounds corroborating the above picture.124,129 and The energy region probed by photons around the absorp- ␣͑E͒ϰ⑀ ͑E͒/n, ͑8͒ tion edge corresponds to band-to-band and tail-to-band ͑or 2 band-to-tail͒ electronic transitions in which the probability of where R2(E) is the normalized dipole matrix element a photon of energy E being absorbed is proportional to the squared averaged over all transitions separated by E, and n product of the initial ͑valence͒ and final ͑conduction͒ states is the index of refraction. Jackson et al.202 have reported separated by that energy. This is the joint DOS defined as measurements on a-Si:H indicating that R2(E)ϳ10 Å2 over the photon energy range 1.5 eVϽEϽ3.0 eV, and slightly 2 DOSJ͑E͒ϭ NV͑EЈ͒NC͑EЈϩE͒dEЈ, ͑6͒ greater than 10 Å for 0.6 eVϽEϽ1.5 eV. Strictly speaking, ͵ however, in the sub-gap region R2(E) should depend on the where NV and NC are the VB and CB density of states func- nature of defects and, in particular, on the distribution of tions, respectively. The initial and final states, however, have charge around it, which modifies the dipole moment of the a wave nature and transitions between them are possible only bound electron. So the constancy of R2(E) measured in the to the extent that the corresponding wave functions overlap. band tail energy range of intrinsic a-Si:H is not expected to In crystals, the periodicity of the lattice is reflected in the hold when different impurities are added or defects of a new periodicity of the electron wave functions which leads to the kind created. Potential fluctuations produced by impurities quantization of momentum. In a-materials, the phase of the having an electronegativity different than the atoms of the electron wave function is random, i.e., not well defined with host network, like N in a-Si or a-Ge, for example, are ex- respect to atomic position. Momentum is not a good quan- pected to induce an increased absorption via changes of the tum number and allowed transitions occur between any two dipole moment of localized electrons and holes. The distance states for which energy conservation applies and wave func- to where the perturbation extends will depend on free carrier tion overlapping exists. The matrix elements for optical tran- screening. In singly doped materials the concentration of free sitions between any two such states is the dipole matrix ele- carriers increases rapidly with doping, the potential fluctua- ment R. The imaginary part of the dielectric constant and, tions arising from random charged defects will be effectively [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 19

FIG. 29. Urbach E0 energy and TO-like FWHM Raman signal ⌫TO as a function of the nitrogen content of a-Si- and a-Ge-based compounds.

Whereas E0 ͑electronic disorder͒ is largely influenced by N concentrations FIG. 30. Urbach energy (E0) vs optical band gap (Eg) of various amor- 19 Ϫ3 as small as 10 atoms cm , the ⌫TO values ͑structural disorder͒ do present phous thin films. At low impurities content ͑doping regime͒ there is no 21 Ϫ3 important changes only for ͓N͔у5ϫ10 atoms cm . change in Eg while the electronic disorder increases for increasing amounts of impurities. In the alloying regime both E0 and Eg scale according to the N concentration. screened, inducing important dipole moment changes. In compensated a-semiconductors the carrier concentration is 145,207–209 very low and charges are very poorly screened. The compen- structural and/or electronic disorder. In the small N sated material is therefore anticipated to have long-range po- concentration range ͑doping regime͒, however, both a-Si and tential fluctuations. The optical transition is spatially local- a-Ge hosts do not present appreciable changes in their struc- ized and the absorption modified when the period of the ture, as probed by Raman spectroscopy. fluctuation is comparable with the wave function extent of The dependence of both Eo and ⌫TO on the N concen- the electron and the hole. Much less impressive changes in tration can be better appreciated in Fig. 29 where a series of the absorption edge of compensated a-semiconductors are N-doped a-Ge:H samples ͑using either N2,orNH3as doping expected, in agreement with experimental evidence.109 sources͒ and a-GeN alloys are represented. According to the The picture suggests that the Urbach slope changes mea- figure, while Eo experiences important changes ͑as deter- sured in singly doped samples may originate from a broad- mined from PDS͒ in the whole range of ͓N͔ the Raman ⌫TO 21 3 ening of the joint DOS ͑increased disorder͒ and/or from signal is affected just for ͓N͔у5ϫ10 atoms/cm . Remem- changes of the dipole matrix element of the optical transi- ber that, at this concentration level, considerable structural tion, induced by the dopant species.203 and optoelectronic modifications are detected. Another interesting feature of Fig. 29 relates to the use of N2 and NH3 as 2. Structural disorder doping sources. As can be seen, while the structure is essentially the same in both series of N-doped samples the use of The relationship between Raman linewidth and struc- NH3 leads to a broader Urbach tail, a result linked to the tural disorder in a-semiconductors is abundantly reported in existence of NH complexes which might originate from 156,204–206 n the literature. There is agreement that the FWHM large electrostatic fluctuations in the a-Ge network at void of the TO-like vibration mode (⌫TO) is a useful indicator of surfaces without affecting significantly the surroundings of structural disorder: the larger the disorder the larger the ⌫TO most host atoms. value. Roughly speaking, changes in ⌫TO correspond to modifications of the short-range order ͑SRO͒, particularly of the bond angle distribution. Figure 28 shows the Urbach 3. Composition, structural disorder and optical band gap slope Eo as a function of the FWHM of the TO-like Raman vibration mode ⌫TO measured on different N-containing a-Si Impurities may influence the properties of tetrahedrally and a-Ge samples. Since ⌫TO is associated with the bond coordinated a-semiconductors in a number of ways. Figure angle distribution ͑or SRO͒ a relationship is expected in the 30͑a͒ shows the Urbach slope of N-containing a-Si:H films alloy regime, and experimentally found ͑Fig. 28͒. At high N as a function of their optical gap. It can be seen that in both concentrations ⌫TO increases indicating the occurrence of doping and alloying regimes, the characteristic energy of the [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 20 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

Urbach tail increases with the introduction of nitrogen in the host network. The same phenomenon is observed in a-Ge-based alloys ͓Fig. 30͑b͔͒. The optical gap changes originate from long-range order modifications occurring only at large nitrogen concentrations. These optical gap changes are always accompanied by a broadening of the Urbach tail, the converse not being always true. The Urbach energy may increase considerably without any measurable changes of the optical gap, as in the case of chemical doping for example. A brief summary of the above mentioned optical properties in a-semiconductors follows: ͑i͒ in group IV a-semiconductors, short-range order potential fluctuations induced by chemical doping seem to contribute to the broadening of the exponential absorption tails; and ͑ii͒ in both FIG. 31. Logarithm of the dark conductivity of some selected a-Ge1ϪxNx a-Si:H and a-Ge:H samples no structural disorder associated samples vs TϪ1/4. Note the straight-line behavior in all cases, an indication with doping levels could be observed by Raman spectros- of variable range hopping transport ͑Ref. 39͒. copy. These findings suggest that the increased sub gap absorption of N-doped samples does not originate only from changes of the joint density of states, and may also reflect the data leads us to believe that most of these changes resulted from film contamination upon exposure to the changes of the dipole matrix element for optical transitions. 214 The valence structure and the electronegativity of the impu- atmosphere. The large scatter of the conductivity data, the rities, as well as the existence of impurity clusters, may thus appearance of the Ge–O related absorption band ͑ϳ800 Ϫ1 play a role in the Urbach tail broadening process. cm ) in the N-free films, and the very low room temperature dark conductivity ͓␴d(RT)͔ values for H-free material, indicate that Takano et al.213 were in fact studying Ge oxyni- V. TRANSPORT PROPERTIES trides, as those reported a few years later by Bagratishvili et al.215 We come back to the subject when discussing the Crystalline and amorphous ␤-Ge3N4 are insulating wide dopant properties of N in the a-Ge network. band-gap semiconductors with potential technological Figure 31 shows the dependence on temperature of the 14,133 applications. The band gap of germanium nitride is dc dark conductivity (T) of some a-Ge N samples.39 13,14,65 ͓␴d ͔ 1Ϫx x ϳ4.5 eV and its refractive index 2.05Ϯ0.05. It is apparent from the figure that the addition of N affects 210 Yashiro used Ge3N4 as a passivating layer and a diffusion the conductivity of a-Ge films. As expected for H-free amor- mask to make low leakage Ge p-n junctions. As c-Ge3N4 is phous material having a large density of defects in the an insulator, the discussion of the electronic transport in GeN pseudo-gap, the conductivity is dominated by a variable alloys will refer to: ͑a͒ the conductivity of H-free and hydro- range hopping mechanism,216 with a characteristic genated off-stoichiometric amorphous GeN alloys and, ͑b͒ N 1/4 ␴d(T)ϭ␴0* exp[Ϫ(T0 /T )] behavior. Figure 31 also shows as an impurity in the c-Ge lattice and in the a-Ge network. the conductivity of an intrinsic a-Ge film deposited under the Along the presentation these topics will be compared with same conditions as the alloys. Table III gives the electrical similar studies on N in Si, particularly when notably differ- characteristics of the above series of samples for N varying ences appear between the Ge and the Si host. in the 0ϽxϽ0.36 range.39 The N content of these samples

A. H-free Ge1؊xNx alloys: Crystalline and amorphous

The insulating properties of Si and Ge nitrides deterio- TABLE III. Compositional and electrical characteristics of a-Ge1ϪxNx a rate notably as the N concentration decreases to below samples deposited by the rf sputtering technique. Note that ␴RT does not xϳ0.57. To the present authors’ knowledge there are no re- change with nitrogen up to a concentration of nearly 4 at %. For larger ͓N͔ the band gap widens and ␴ drops very fast. ports on the electronic properties of off-stoichiometric RT

c-Ge1ϪxNx . In contrast, the transport properties of H-free Sample ͓N͔ x ␴RT N(EF) and hydrogenated off-stoichiometric amorphous GeN alloys No. (cmϪ3) ͑%͒ (⍀ cm)Ϫ1 (eVϪ1 cmϪ3) as a function of N content have been addressed by several 1 Ͻ1018 — 9.0ϫ10Ϫ3 1.7ϫ1018 15,26,39,113,211,212 authors. 2nab — 1.5ϫ10Ϫ2 2.3ϫ1018 One of the ﬁrst reports on the effects of N on the opto- 3 na — 1.8ϫ10Ϫ2 2.4ϫ1018 electronic properties of a-Ge was the study by Takano 4 1.6ϫ1020 0.002 1.8ϫ10Ϫ2 1.6ϫ1018 20 Ϫ2 18 et al.213 on rf sputtered H-free a-Ge N samples deposited 5 6.0ϫ10 0.007 1.7ϫ10 1.7ϫ10 1Ϫx x 6 9.1ϫ1020 0.011 1.5ϫ10Ϫ2 1.1ϫ1018 at ϳ350 K. The conductivity was found to depend on the 7 3.0ϫ1021 0.037 1.1ϫ10Ϫ2 1.2ϫ1018 temperature and on the N2 partial pressure during growth. 8 8.8ϫ1021 0.108 3.5ϫ10Ϫ3 — Concomitant changes of the optical properties of the ﬁlms 9 1.8ϫ1022 0.221 2.5ϫ10Ϫ5 — were detected simultaneously and attributed by Takano 10 2.5ϫ1022 0.306 Ͻ10Ϫ7 — 22 Ϫ9 et al.213 to the incorporation of N in the a-Ge network. 11 2.9ϫ10 0.355 Ͻ10 — Whereas some of the changes measured by Takano et al. aSee Ref. 39. could be due to nitrogen incorporation, a detailed analysis of bNot available. [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 21

was determined from a deuteron-induced nuclear reaction, and the concentration evaluated by interpolating between pure Ge and stoichiometric ␤-Ge3N4, which has 4.6ϫ1022 N atoms cmϪ3 and a density of 5.3 g cmϪ3.14,120 Two facts at least emerge from Fig. 31 and Table III: ͑i͒ the absolute value of ␴d and its dependence on T do not change appreciably up to Nϳ3ϫ1021 cmϪ3 (xϳ0.04) and ͑ii͒ higher N concentrations decrease ␴d at all temperatures and the samples display a stronger dependence on T. The relative independence of ␴d on ͓N͔ for xϽ0.04 is the consequence of the large density of deep defects (GeDB) around the Fermi energy N(EF) which dominate the transport. Assuming the wave function localization parameter ␥ϳ0.1 ÅϪ1 determined by Knotek et al.217 for a-Ge, and the conductivity data of Table III, the density of states for xϽ0.04 samples at the Fermi energy N(EF) is of the order of 1018 eVϪ1 cmϪ3 or higher39 ͑see Table III͒, a defect density compatible with transitions between states 100 Å or less far apart. Hence, in Ge-rich material (xϽ0.04) EF is pinned by GeDBs and hopping between them dominates conduction at all temperatures. Honma et al.133 have also reported the conductivity of an a-Ge N film (xϭ0.1) against tempera- 1Ϫx x FIG. 32. Dark conductivity of a-Ge1ϪxNx :H films prepared by rf sputtering ture, and found that log ␴d becomes a straight line when a c-Ge target in an ArϩNH3 atmosphere as a function of inverse tempera- plotted against TϪ1/4, an independent indication that variable ture. The conductivity of an intrinsic a-Ge:H sample is also shown for range hopping prevails in highly defective materials, as for comparison. Note the active doping effects of small NH3 partial pressures in the deposition chamber and, for high NH3 partial pressures, a decreasing samples of Fig. 31. conductivity due to the band-gap widening mechanism induced by nitrogen. As the N concentration into the samples increases The dark conductivity is thermally activated for all sample series ͑see Refs. ͑xϾ0.04, Table III͒ a progressive decrease of the conductiv- 211 and 212͒. ity is expected because nitrogen provokes the widening of the pseudo-gap, inducing a stronger localization of deep defects ͑see electronic structure Sec. III B, and optical proper- tering of a c-Ge target in a gaseous mixture of ArϩN2ϩH2, ties, Sec. IV B͒. As a consequence, ␴d is expected not only were reported by Honma et al.143 and by Chambouleyron to decrease with increasing N, but to have a stronger depen- et al.132 These preliminary reports established the role of H dence on temperature, in agreement with experimental re- in cleaning the pseudo-gap via the passivation of GeDB’s and sults ͑Fig. 31͒. The lack of reliable data on the defect struc- the preference of H to bond to N instead of to Ge atoms. ture in N-rich a-Ge1ϪxNx samples prevents a quantitative More extensive research was done by Marques,211 who stud- analysis along the line used for the Ge-rich alloy. ied several series of samples deposited either by the rf sput- The optical and transport properties of a-Ge0.76N0.17O0.07 tering technique ͑ArϩN2ϩH2,ArϩNH3 and ArϩNH3ϩH2 films formed by nitriding Ge in hydrazine N2H4 vapor at 218 gaseous mixtures͒, or by the rf glow discharge ͑GD͒ of Tϳ1000 K have been reported by Bagratishvili et al. Any 170 39 GeH4ϩNH3. The Tauc optical gap of these sample series comparison with the present authors results becomes diffi- covered the 1.0–3.5 eV range, the wider gap samples being cult due to the large quantity of oxygen present in Bagratish- those prepared by rf GD.211,212 The transport properties of rf vili et al.218 samples ͑the hydrazine contained 1%–2% impu- 170 glow discharge deposited a-Ge1ϪxNx :H samples have also rity water͒. For H-free GeN alloys, xϳ0.17, a Tauc gap of 133 26 39 been reported by Honma et al. and by Xu et al. approximately 0.9 eV is expected, whereas ETaucϭ1.2 eV 218 The most striking effect of hydrogenation on the elec- is reported for the samples containing oxygen, a clear in- tronic transport of amorphous semiconductors is the appear- dication of the important effects of oxygen contamination on ance of a thermally activated conductivity. Figure 32 shows the optical properties. From PDS data of Ge oxynitride the logarithm of the dc dark conductivity of a series of samples, Bagratishvili et al.218 estimated a defect density at 18 Ϫ1 Ϫ3 a-Ge1ϪxNx :H alloys, rf sputtered in an ArϩNH3 atmo- the Fermi energy N(EF)ϳ1.4ϫ10 eV cm , a value sphere, against inverse temperature.211,212 The conductivity comparable to those reported in Table III for xϽ0.04 O-free of a nominally undoped a-Ge:H sample deposited under the samples. Contrary to the case of a-Ge1ϪxNx reported in same conditions is also shown in Fig. 32 for comparison. Table III, a temperature activated dark conductivity at RT The activation energy Ea is obtained fitting the data points and above is found for Ge oxynitride samples.218 with the expression ␴d(T)ϭ␴0 exp(ϪEa /kT), where k is the Boltzmann constant. The figure shows that the conductivity depends strongly on N content. Small NH partial pressures B. Hydrogenated a-Ge N films xу0.01 3 1؊x x „ … in the chamber cause an increase in the dc conductivity and a Hydrogen atoms play a fundamental role on transport simultaneous reduction of the activation energy. This effect properties. a-Ge1ϪxNx :H films, prepared by rf reactive sput- was interpreted as originating from N active chemical doping [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 22 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

FIG. 34. Energy defect levels associated with N in c-Si as found by various FIG. 33. Room-temperature dark conductivity of a-Ge1ϪxNx :H samples as a function of N content. ͑Dots͒ samples from a plasma enhanced chemical characterization and implantation techniques. The donor levels produced by substitutional phosphorus and arsenic are also shown for comparison. a W. vapor deposition system using diluted GeH ͑5% in H ͒ and NH gaseous ͑ ͒ 4 2 3 Kohn, Ref. 219; ͑b͒ See Ref. 264; ͑c͒ See Ref. 265; ͑d͒ A. H. Clark, J. D. mixtures. ͑Ref. 26͒͑Squares͒ rf sputtered samples in an ArϩNH ϩH at- 3 2 MacDougall, K. E. Manchester, P. E. Roughan, and F. W. Anderson, Bull. mosphere ͑see Ref. 261͒. Am. Phys. Soc. 13, 376 ͑1968͒; ͑e͒ A. G. Milnes, in Deep Impurities in Semiconductors ͑J. Wiley, New York, 1973͒; ͑f͒ W. J. Kleinfelder, Techn. Rep. ͑Stanford Electronics Labs, 1967͒; ͑g͒ Y. Tokumaro, H. Okushi, T. of the a-Ge:H network, a subject we will discuss in detail in Masui, and T. Abe, Jpn. J. Appl. Phys. 2 21, L443 ͑1982͒; ͑h͒ S. T. Pan- a coming section. Increasing the partial pressure of NH3 telides and C. T. Sah, Phys. Rev. B 10, 638 ͑1974͒; ͑i͒ K. L. Brower, Ref. above 5ϫ10Ϫ4 mbar during deposition reverses the trend, 225; ͑j͒ K. Nauka, M. S. Goorsky, H. C. Gatos, and J. Lagowski, Appl. Phys. Lett. 47, 1341 ͑1985͒; ͑k͒ L. C. Kimerling and J. L. Benton, Appl. i.e., Ea increases again and ␴d decreases very fast, reaching Ϫ12 Ϫ1 Phys. Lett. 39, 410 ͑1981͒; ͑l͒ D. Wruck and P. Gaworzewski, Phys. Status ␴ϳ10 (⍀ cm) at room temperature for the xϳ0.37 Solidi ͑A͒ 56, 557 ͑1979͒. Singly and double charged Si divacancies give sample. The trend toward an insulating material is concomi- rise to levels at 0.42 and 0.56 eV below the CB, respectively ͓J. Stein, in tant with the band-gap opening mechanism induced by trigo- Radiation Effects Semiconductors, edited by J. Corbett and G. Watkins nal N bonding, referred to in the previous sections on the ͑Gordon & Breach, New York, 1971͔͒. structural and the optical properties of the alloys. Similarly to the crystalline case, activated transport in ductors Si and Ge. The simple model of chemical doping a-semiconductors corresponds to thermally excited electrons assumes that four electrons are needed to fill the bonding in the neighborhood of the CB mobility edge. It dominates in orbitals of the tetrahedral neighbors. The extra electron do- the samples under analysis because hydrogenation has re- ͑ nor case is bound to the impurity by a Coulomb-like attrac- duced drastically the density of deep defects, preventing any ͒ tive potential produced by the extra opposite charge of the significant contribution from phonon assisted tunneling be- impurity atom. The attractive force is modified by the dielec- tween localized Ge s. DB tric constant of the crystal, the lattice potential being taken The conductivity of a series of a-Ge N :H samples 1Ϫx x into account through the use of an appropriate effective prepared by the rf GD of 5% H diluted (GeH ϩNH ) gas- 2 4 3 mass.219 A similar argument may be used to give account for eous mixtures was recently reported by Xu et al.26 Figure 33 the electric activity of shallow acceptor states. shows the room temperature conductivity of Xu et al.’s Within this context, N appears to be distinctly different samples as a function of N content ͑as determined from from other group-V elements, and its electrical activity in the XPS͒ and those of Marques et al.211,212 deposited by rf sput- crystalline lattice of Si and Ge is still a matter of study. Let tering. Note the different ␴ values between the two series. RT us note at this point that the properties of N in diamond are The difference may be due to the use of different deposition very different than in Si or Ge, N being highly soluble in methods, for it is well known that rf sputtered a-Ge:H pos- diamond220 but having a very low solubility in Si and Ge.221 sesses a smaller density of deep defects than anodic glow The difficulty of studying the electrical activity of substitu- discharge samples104 ͑see Fig. 33͒. The large DOS in the tional N in crystalline elemental semiconductors partly de- pseudo-gap of GD deposited samples may hinder the effects rives from:222 i the rather high dissociation energy of the of nitrogenation at small NH partial pressures. Other ͑ ͒ 3 N molecule which prevents the normal mechanism of sources of disagreement might stem from different methods 2 atomic diffusion and ii the fact that the N molecule may used to estimate the N content of the two series, as well as ͑ ͒ 2 enter the tetrahedral interstice of Si and Ge. Moreover, by the total amount of bonded hydrogen atoms. cooling melted Si or Ge in an atmosphere of pure N2 or NH3 the reaction with the liquid phase forms the corresponding C. Nitrogen as an impurity in Ge xр0.01 „ … nitride phase, with planar bonded N and tetrahedral bonded 1. N in crystalline germanium Si or Ge. With the exception of N, the elements of groups III and As a consequence of the above, the manifestation of do- V of the Periodic Table are known to produce localized elec- nor properties of substitutional N in c-Si and c-Ge can be tronic levels near the band edges of the elemental semicon- determined only if it is introduced in the atomic state at [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 23

dissociated by an rf or a dc plasma. Spear and LeComber232 established for the first time that, in tetrahedrally coordinated a-semiconductors, the doping properties of group III and V elements are similar to those found in the crystalline semiconductor parents. The doping process, however, is different in an a-semiconductor compared with a crystal, in the sense that in the crystal constraints of symmetry force an impurity atom of a different valence to have the same coordination number as its host in a substitutional site. In a-networks the local minimization of free energy may also be attained with the impurity atom having a first coordination number equal to its chemical valence ͑self compensated site͒. In this case, the first coordination number is determined by purely chemi- FIG. 35. The antiparallel model of the N pair defect in c-Si and c-Ge ͑see cal parameters and not by the impurity’s 216,233,234 Ref. 227͒. Filled circles: Ge atoms; open circles: N atoms. environment. Both active and inactive doping sites coexist in a-networks and their relative stability results from the combined short-range chemical contributions and the lat- relatively low doses. Ion implantation has been used for such tice strain term, for it is known that tetrahedrally coordinated 223–227 a purpose. The discussion around experimental find- amorphous networks are topologically constrained.235 ings involves the subtleties of annealing processes, always The electrical activity of N impurity in a-Si:H and necessary to relax the radiation damage produced by the ion a-Ge:H has been a matter of debate for many years. It has bombardment. Figure 34 shows schematically the energy been argued theoretically that over-coordination of nitrogen levels associated with N and with N-induced defects reported ϩ ͑N4 ϵpositively charged, tetrahedrally coordinated N͒ should for c-Si, together with the donor levels of substitutional P be impeded in the a-Si and a-Ge networks by its small size, and As. and should only be possible if one or more neighbors are Probably due to nitrogen’s tendency to form only three hydrogen.123 Another proposal236 to explain the large density bonds and leave two electrons in a lone pair, substitutional N of charged Si dangling bonds normally found in N-doped in diamond and Si is distorted in the ͑111͒ direction. This 228,229 a-Si:H considers the existence of N charged dangling bonds gives rise to a deep level in the gap. A study of N in Ϫ ϩ 224 (N2 ) and positively charged SiDBs(Si3). This configuration, c-Ge by Campbell et al. reports that ϳ85% of implanted however, does not give rise to n-type doping, disagreeing N is located in nonsubstitutional sites, and that N does not with the electrical properties found experimentally in N- outdiffuse from c-Ge for anneals as high as 1000 K. More- ϳ doped a-Si:H and a-Ge:H networks. over, electrical measurements indicate that substitutional N The nature of these defects was analyzed in detail by is not a donor in c-Ge. The dominant center observed in Shimizu et al.237 by means of several spectroscopic tech- N-doped c-Ge or c-Si is not substitutional N but a nitrogen niques. Measurements of ESR, light-induced electron spin pair,226,227 which consists of two neighboring ͑100͒-oriented resonance ͑LESR͒, and sub-gap optical absorption using the Ge͑Si͒–N split interstitials, arranged in an antiparallel con- constant photocurrent method ͑CPM͒, were carried out for figuration and with the four bonds forming a square lying on a-Si:H, a-Si C :H, a-Si O :H and a-Si N :H ͕011͖, as shown in Fig. 35. 1Ϫx x 1Ϫx x 1Ϫx x films as a function of film thickness. From the combined The defect structure of N as a substitutional impurity in information provided by these different characterization crystalline elemental semiconductors has also been ad- techniques and from samples of different thickness, both the dressed theoretically230,231 with two different approaches: surface density of dangling bonds in the disordered surface perturbation methods and cluster methods, which stress the layer and the density of dangling bonds in the bulk region local environment of the defect. Atomic relaxation around were obtained, discriminating neutral from charged defects. the defect and electronic correlation are seldom considered. The main conclusions of Shimizu et al.,237 agreeing with our As both can influence calculations by a factor as large as the microscopic picture of N doping in a-Ge:H films,139,238 can gap energy, the defect levels with respect to band edges are be summarized as follows: ͑i͒ the density of dangling bonds not reliably located. deduced from CPM agrees fairly well with the sum of neu- Summarizing, N is a poor donor in crystalline group-IV tral and negatively charged dangling bonds deduced from semiconductors, partly because of its low solubility, and ESR and LESR, both neutral and negatively charged dan- partly because of the tendency of N to form pairs in im- gling bonds contributing to CPM signal; ͑ii͒ the energy po- planted as well as in grown materials. sition and the width of the defect level distribution are almost unchanged regardless of the fraction of charged defects, 2. N as an active dopant in amorphous Ge (and Si) suggesting that both positively and negatively charged de- The problems related to the inclusion of substitutional N fects are paired through a Coulombic interaction, ͑iii͒ many atoms in the crystalline Si and Ge lattices may be partially charged dangling bonds exist in the bulk of a-Si:H, their overcome in a-networks. The deposition conditions ͑fast density being more than five times that of the neutral dan- condensation from the vapor phase͒ allow the incorporation gling bond. Although some charged dangling bonds exist of atomic components, or radicals of complex molecules, even in highly pure a-Si:H films, the origin of these charged [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 24 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

FIG. 37. The doping efficiency of nitrogen atoms vs the concentration of N FIG. 36. Logarithm of the dark conductivity of N-doped a-Ge:H as a func- in the solid phase for N-doped rf sputtered a-Ge:H samples ͑solid circles͒ as tion of inverse temperature. Intrinsic and lightly N-doped samples possess a compared with P-doped glow discharge a-Si:H films ͑see Refs. 238 and single thermally activated conduction regime down to low temperatures. In 262͒. heavier N-doped a-Ge:H films, thermally activated hopping through localized states near the conduction-band edge dominates the conductivity ͑see Ref. 139͒. atoms which are electrically active, was quantified by the present authors238 and found to be of the same order of magnitude as the doping efficiency of P in a-Si:H films, as dangling bonds appears to be unintentionally incorporated N; shown in Fig. 37. It is interesting at this point to compare the ͑iv͒ finally, the incorporation of N largely increases the den- above results with experimental data of N-doped a-Si:H sity of dangling bonds. The charged dangling bonds are ϩ Ϫ Ϫ ϩ films. In what follows we show that important differences likely to be intimate pairs of N4 and Si3 rather than N2 ϩSi3 appear between both materials. 236 pairs as suggested by Robertson and Powell. It is worth One of the first studies on N as a dopant of a-Si:H mentioning here that a recent calculation of defect levels of reported244 an eight orders of magnitude increase in conduc- ϩ 130 N impurity in a-Si finds the N4 center just below Ec . tivity upon incorporation of N, with an activation energy as The active doping properties of N in rf sputtered a-Ge:H small as 20 meV for the most effectively doped sample. films using N2 as a doping source were independently stud- These results, however, could never be reproduced by other 139,140,211,212,238–241 ied at Campinas ͑see Fig. 32, for example, research groups. Most studies on N-doped 211 242,243 for the first evidence͒ and at Kaiserslauten. Figure a-Si:H140,141,245,246 report the following overall picture. N- 36 shows the changes of the dark conductivity of a-Ge:H doped a-Si:H films possess a thermally activated conductiv- samples when small amounts of N2 are allowed to flow into ity, the activation energy E decreases down to a minimum 139 a the reaction chamber during deposition. Experimental data value of around 0.5 eV with doping and a simultaneous 3–5 indicate that the temperature dependence of ␴ can be well d orders of magnitude increase of ␴d is measured ͑see Fig. 38͒. represented by either one exponential function, or by the sum The experimental data suggest that, either the doping effi- of two exponential functions. The single activation energy ciency of N in a-Si:H is very small, or the energy level of the Ea which dominates the electric transport of undoped and defect center giving rise to n-type doping is relatively deep lightly doped a-Ge:H in the 170–400 K temperature range in the pseudo-gap, or both. Meaningful comparisons are of- corresponds, under the assumption of a linear temperature dependence of the pseudo-gap, to the energy difference, at Tϭ0 K, between the Fermi level and the conduction band mobility edge. A second exponential, which appears only in samples deposited under N2 partial pressures higher than 10Ϫ5 mbar, dominates the transport at low temperatures. The analysis of ␴d(T) indicates that transport through band tail localized states ͑hopping͒ dominates the conductivity at low temperatures. The most effectively N-doped a-Ge:H sample possesses a high temperature activation energy Eaϳ0.12 eV, a remarkably low value for a doped a-semiconductor, and a low temperature activation energy EHϳ10 meV, consistent with a thermally activated nearest neighbor hopping at the Fermi energy ͑see Fig. 36, sample with ͓N͔ϳ0.1 at. %͒. These experimental findings indicate that active N doping is a very efficient process in a-Ge:H samples of improved quality. Qualitative similar results were FIG. 38. Room-temperature dark conductivity vs activation energy for N- 242 doped a-Si:H films. The experimental data labeled by a-Si:H͑N͒ correspond obtained by Dru¨sedau et al. The doping efficiency of N in to rf sputtered samples while the other points were obtained from rf glow

a-Ge:H, deﬁned as the fraction of all incorporated nitrogen discharge (SiH4ϩN2) deposited samples ͑see Refs. 141, 246, and 263͒. [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 25

FIG. 40. Dark conductivity at room temperature and activation energy of FIG. 39. Nitrogen incorporation in a-Si:H and a-Ge:H thin films as a func- N-doped a-Si:H films as a function of N2 partial pressure in the chamber tion of N2 partial pressure during deposition. The N concentration in a-Si during deposition. ͑a͒ data from rf sputtered samples and ͑b͒ data from rf was determined from IR absorption data, whereas the impurity content in GD-deposited films. Note the important changes of ␴RT and of Ea induced a-Ge was measured by nuclear reaction techniques ͑NRA͒͑see Ref. 140͒. by small N2 partial pressures ͑see Refs. 140 and 141͒.

ten difficult ͑or impossible͒ to establish between samples deposited in different systems under different conditions or three orders of magnitude smaller than that of P. We believe methods, or deposited using different gaseous dopant this conclusion to be erroneous because the dissociation en- sources. In spite of the difficulty, let us speculate on the ergy of PH3 and N2 are very different. Under the same dis- nature of the N donor impurity under the light of available charge conditions the dissociation of PH3 molecules will be experimental results. accomplished to a much larger extent than for N2. The results of Zhou et al.141 and those of the present authors A comparative study of N doping in rf sputtered a-Si:H 140 and a-Ge:H has been made by the present authors.140,240 Fig- suggest, on the contrary, that the N doping efficiency ure 39 shows the incorporation of N atoms in a-Si:H and might not necessarily be smaller in a-Si:H than in a-Ge:H. a-Ge:H samples prepared in the same deposition system as a Instead, experimental results suggest that the donor level function of the N partial pressure (P ) in the chamber. produced by N in a-Si:H may be deeper than that of N in 2 N2 a-Ge:H, as suggested by the calculations of Makler et al.126 Figure 39 indicates that for similar N partial pressures; ͑i͒ 2 Note that, because of a much larger Ge dielectric constant, the rate of N incorporation with varying P is the same for N2 this is the case for all group III and V elements in c-Si as both elements and, ͑ii͒ under similar PN almost a one order 2 compared with c-Ge. The limited energy range swept by Ea of magnitude larger fraction of N is incorporated in a-Si:H. might then be the consequence of a deep donor level and not The comparison only considers samples having the same de- of a poor doping efficiency. To the present authors knowl- gree of hydrogenation and containing at most 1 at % N at- edge, the doping efficiency of N in a-Si:H has not yet been oms, which corresponds to a true doping concentration clearly established. range. rf sputtering, however, does not lead to device quality Why is N a highly efficient dopant in a-Ge? The size of a-Si:H and its DOS is one order of magnitude larger than ϩ the nitrogen atom would hinder the N4 configuration, unless that of a-Ge:H. a considerable atomic relaxation occurs around the impurity 139,239 Figure 40͑a͒ shows ␴RT and Ea for N-doped a-Si:H as a atom, as suggested by the present authors. Following function of P for the sample series under consideration. 242 N2 Robertson and Powell’s ideas Dru¨sedau et al. suggest that Important conductivity changes appear upon N incorpora- the N-related donor state originates in the presence of hydro- tion, a clear indication of chemical doping. Note that the gen complexes, such as Ge–H –Ge for example, and → 2 ← measured changes occur at relatively low N2 partial pres- not from an over-coordinated N atom. The reason behind the sures, a consequence of the relatively large rf power fed into H-complex configuration should be, according to the Kaiser- the plasma during the sputtering process, which dissociates a slautern group, the preferential attachment of hydrogen to 242 considerable fraction of N2 molecules. Compare this picture nitrogen. To investigate the likely role of H in the doping 141 with that obtained by Zhou et al. ͓Fig. 40͑b͔͒ who found process, N-doped rf sputtered a-Ge:H samples using NH3 an almost identical dependence of Ea and ␴RT on N-doped and N2 as a gaseous dopant source were compared at a-Si:H samples deposited by the rf glow discharge method Campinas.139,239 The analysis of the transport data indicate using a (SiH4ϩN2) gaseous mixture. In the case of GD that, although the doping effect looks similar with the use of samples, however, a much higher N2 pressure is required in either of the two gases, equal amounts of N in the network the chamber ͑see Figs. 1 and 2 of Ref. 141͒ to detect any do not induce equal activation energy and conductivity conductivity change, a consequence of the low rf power fed changes in a-Ge:H. Figure 41 shows the variations of the into the GD plasma, which is insufficient to break a sufficient activation energy in N2- and NH3-doped a-Ge:H as a func- 141 239 fraction of N2 molecules. Zhou et al. compare N doping tion of N concentration in the solid phase. When NH3 is with P doping of a-Si:H under the same deposition condi- used Ea vs ͓N͔ flattens and shifts toward larger N concentrations, and conclude that the doping efficiency of N is at least tions. This is an indication that the N-active doping effi- [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 26 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

determined from EXAFS data of trigonal N in the a-Ge network.121 Let us note here that the above calculation were performed in a H-free a-Ge network. The presence of H may affect considerably the kind of most favorable impurity configuration. In addition, Venezuela and Fazzio249 made ab initio calculations of defect formation energies, the values of which, however, appear to be at odds with chemical intuition and the available experimental data. It is theoretically found, for example, that active electronic doping originates from ϩ ϩ N3 and not from N4 , as expected from the standard theory of chemical doping of a-semiconductors.234 The two most surprising results of the defect energy calculations are: ͑i͒ 0 The neutral N3 center cannot exist because the reaction FIG. 41. High temperature (TϾ300 K) activation energy of the dark con- N0 NϩϩDϪ is exothermic, the donor electron becoming 3→ 3 ductivity vs nitrogen concentration in N-doped a-Ge:H films using N2 localized at deep Ge dangling bond centers. In other words, ͑circles͒ and NH3 ͑squares͒ as doping sources. Both dopant sources induce similar changes of the activation energy. Note, however, the flattening of the the Ge–N bond of the Ge3N skeleton would be shallower 3 curve and its shift to larger N concentrations for samples doped from NH3 than the sp Ge dangling bond. This is at odds with the source ͑see Ref. 239͒. recession of the VB of a-GeN alloys as the N concentration increases, as measured by photoelectron spectroscopies.129 It is also contrary to chemical intuition considering that N is more electronegative than Ge. ͑ii͒ The ab initio calculations ciency is smaller for NH3 than for N2. Moreover, the use of also indicate that the relative concentration of ͓N ͔/͓N ͔ de- NH3 induces a large network disorder ͑see optical properties, 3 4 creases as the N concentration increases in the network. This Sec. IV C͒ due to the incorporation of NH and NH2 groups as Ge dangling bond terminators.239 Tsu et al.247 prepared prediction is not compatible with the known phenacite structure of stoichiometric a-Ge N , in which Ge atoms are tet- silicon nitride films from (SiH4ϩN2) and (SiH4ϩNH3). In 3 4 the latter, they found evidence of the incorporation of NH rahedrally bonded to N atoms, and N is trigonally bonded to Ge in a planar configuration. and NH2 groups into the films. Samples deposited using ND3 Summarizing, the experiments done at Campinas using instead of NH3 indicate that the source of the H in the NH and NH groups is NH and not SiH , suggesting that the N2 and NH3 as dopant sources and the structural calculations 2 3 4 249 donor center induced by nitrogen in a-Si:H is also over- of Venezuela and Fazzio, referring to threefold and four- coordinated N atoms, in agreement with the results of fold coordinated N in a-Ge, give support to the idea that Shimizu et al.237 Infrared studies and chemical bonding active N doping originates from over-coordinated N and not modeling lead Lucovsky et al.248 to propose that fourfold from NH complexes. Yet, there are no reliable theoretical ϩ coordinated N with second neighbor H atoms, as in values for the energy level of the N4 antibonding state. The Nϩ–Si–H linkages, is the active donor configuration of N in question of the different activity of N in the Si and Ge the a-Si:H network. These results, however, should be con- a-networks remains to be answered. The available informa- sidered with caution, in the sense that the N-doped samples tion suggests that the doping mechanisms are pretty much studied by Lucovsky et al.248 contain up to 15 at. % nitrogen, the same in both semiconductors, the less efficient N doping i.e., a N content well above the normal doping concentration of a-Si:H being probably associated with a deep donor level, range and typical of off-stoichiometric silicon nitride. say 300–350 meV below the CB edge. The different energy The question of N doping a-Ge has been recently ad- for the N donor level between a-Ge:H ͑50–60 meV͒ and dressed theoretically by Venezuela and Fazzio.249 The struc- a-Si:H might originate from the smaller dielectric constant ture of a-Ge is generated using the continuous-space Monte of Si and a less effective atomic relaxation around the impu- Carlo method on a large number of atoms. The correspond- rity. ing electronic and structural configurations are obtained D. Photoconductivity in N-doped a-Ge:H films within the framework of the density–functional theory using the local density approximation. The results are illustrative The photoconductivity, defined as the change in conduc- of the importance of atomic relaxation, which appears to be tivity upon illumination, can be expressed as: very large in the a-Ge network. In all N atom inclusion cases ⌬␴ϭe(␮n⌬nϩ␮ p⌬p), with ⌬nϭ␩G␶, where ␮ is the studied by Venezuela et al.249 ͑trigonal or tetrahedral coor- carrier mobility, G the photogeneration rate of electron-hole dinated N͒, a strong inward relaxation of the Ge atoms in the pairs, ␩ the quantum efficiency for free-carrier production vicinity of the impurity is found. The calculations also show and ␶ the lifetime of excess carriers. Depending on recombi- that, of all possible charge states of tetrahedral N, the most nation kinetics, ␶ is usually a function of the generation rate. ϩ ␥ stable is N4 , in agreement with the predictions of the experi- For strongly absorbed photons ⌬␴ϰ(␣F) , where ␣ and F mental studies at Campinas.139,238 As with other group V are the absorption coefficient and the photon flux, respec- 234 ϩ active dopant elements in Si and Ge, the N4 impurity is in tively, ␥ is found to vary between 0.5 and 1.0 ͑bimolecular the center of an almost perfect tetrahedron of Ge atoms, in and monomolecular recombination kinetics, respectively͒ which the Ge–N bond length is approximately 2.1 Å. This and may also depend on the generation rate and the photon figure is to be compared with 1.83 Å, the Ge–N bond length energy. [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta 27

not originate from N donors but from structural randomness caused by the incorporation of nitrogen. VI. CONCLUDING REMARKS The properties of nitrogen impurity in germanium and of Ge–N alloys have been reviewed. Amorphous and crystalline germanium–nitrogen alloys are interesting materials with potential applications for protective coatings and window layers for solar conversion devices. They may also act as effective diffusion masks for III-V electronic devices. The scarce existing information on crystalline ␤-Ge3N4 indicates that its lattice and electronic structure are similar to ␤-Si3N4, by far the most studied column IV nitride. In particular, the band-gap widening mechanism induced by nitrogen closely FIG. 42. ␩␮␶ product and ␥ factor for intrinsic and N-doped a-Ge:H thin resembles both alloys. The top of the valence band is domi- films as a function of N concentration in the solid phase ͑N2 has been used nated in a-Ge3N4 byaN2pnonbonding orbital, the conduc- a dopant source gas͒. ͑See Ref. 250͒. Note the plateau at N concentrations tion band being given by Ge–Ge antibonding orbitals. The between 1019 and 1020 cmϪ3. See Fig. 41 for the corresponding activation mechanical properties of crystalline ␤-Ge3N4 are very poorly energies. known, as well as the possible ceramic uses c-GeN alloys. The review focuses on the experimental and theoretical Marcano et al.250 reported on the secondary photocon- properties of amorphous germanium–nitrogen alloy films, on ductivity of N-doped a-Ge:H samples. Figure 42 shows the which a more abundant literature exists. A general corre- spondence has been found between these and the structural measured ␩␮␶ product and ␥ factor as a function of N doping. Three different behavior regions are apparent from the and optical properties of a-SiN, as expected from the similar figure. The intrinsic and lightly doped samples valence structure of both elements. Some differences appear, however, regarding the nitrogen concentration at which the (NϽ1019 cmϪ3) do not display important changes in the ␩␮␶ band-gap widening induced by nitrogen becomes effective, product and the ␥ exponent remains close to 0.75, an indication that light N doping has not substantially altered the re- the structure of the valence band as determined from electron combination kinetics. It is expected that a fraction of neutral photoemission spectra, and the properties of N impurity in Ϫ the amorphous network. GeDB will be converted to D , i.e., doubly occupied DBs which are ineffective as recombination centers. The E shift Nitrogen is not an active dopant in crystal column IV F elements, but an active dopant in the corresponding amor- ͑see shift of Ea in Fig. 41͒ is not sufficient to transform all D0’s into DϪ’s. The density of remaining D0’s is sufficient phous networks. The doping efficiency of N in a-Ge:H and a-Si:H has been discussed. Some important differences be- to keep the ␩␮␶ product in the 10Ϫ7 cm2 VϪ1 range. The second region corresponds to three effectively tween the two cases have been highlighted. Nitrogen is a very effective active dopant in the a-Ge:H network, inducing doped samples possessing an activation energy E р0.2 eV a a donor level at around 50–60 meV below the CB edge. Its and a ␩␮␶ product distinctly higher than the samples of the doping efficiency is similar to that of phosphorus in a-Ge:H first group (␩␮␶Ͼ10Ϫ6 cm2 VϪ1). The small activation energy for these samples means that most Ge ’s are doubly or in a-Si:H. The conductivity variations of a-Si:H induced DB by N doping are not as impressive. The discussion suggests occupied and the ␩␮␶ product is almost independent of dop- that this poor performance may stem from a N-induced deep ing. The small increase in ␥ may originate from the GeDB density increase with N doping. The remaining samples are donor level, and not from a reduced doping efficiency. The the most N contaminated. The active doping process is very shallowness of the defect level energy in a-Ge:H may be the consequence of a larger atomic relaxation around the impu- ineffective at large N doping concentrations ͑see Fig. 37͒, most N entering in the a-Ge network with trigonal coordina- rity, as found theoretically from a-Ge structural simulations tion. A very different coordination and atom size induce the using a continuous-space Monte Carlo method on a large appearance of a large density of coordination defects. As a number of atoms. More theoretical and experimental work on the nature and energy of N-induced defects would allow consequence, an important ␩␮␶ product drop is expected as well as a pure monomolecular recombination kinetics, as us to understand the role of filled d states in the microscopic measured. origin of N doping efficiency. Masuda et al.251 studied the ESR, light-induced ESR and ACKNOWLEDGMENTS the decay time of the photoconductivity of N-doped a-Si:H thin films. From the behavior of the decay time they con- Most of the work on N in Ge done at UNICAMP was clude that, up to a N concentration of approximately 2 at %, partially financed by the Brazilian agencies Fundaçaõde the conductivity variations are due to the incorporation of Amparo a` Pesquisa do Estado de Saõ Paulo ͑FAPESP͒ and tetrahedral N ͑effective dopant͒. For N concentrations Conselho Nacional de Desenvolvimento Cientıćo e Tecno- Ͼ2 at. % the decay time of the photoconductivity decreases lo´gico ͑CNPq͒. The authors are indebted to their colleagues and the density of deep defects continues to increase, sug- at the Laboratory of Photovoltaic Research for fruitful dis- gesting that most of charged SiDB’s found in the samples do cussions. [This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 143.106.108.149 On: Thu, 18 Jun 2015 18:49:29 28 J. Appl. Phys., Vol. 84, No. 1, 1 July 1998 I. Chambouleyron and A. R. Zanatta

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