INTRINSIC OR EXTRINSIC ORIGIN OF THE RECOMBINATION AT EXTENDED DEFECTS

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INTRINSIC OR EXTRINSIC ORIGIN OF THE RECOMBINATION AT EXTENDED DEFECTS

B. Sieber

To cite this version:

B. Sieber. INTRINSIC OR EXTRINSIC ORIGIN OF THE RECOMBINATION AT EXTENDED DEFECTS. Journal de Physique Colloques, 1989, 50 (C6), pp.C6-47-C6-56.

�10.1051/jphyscol:1989604�. �jpa-00229635�

REVUE DE PHYSIQUE

APPLIQUEE

Colloque C6, Supplbment au n06, Tome 24, Juin 1989

INTRINSIC OR EXTRINSIC ORIGIN OF THE RECOMBINATION AT EXTENDED DEFECTS

B. SIEBER

Laboratoire de Propriett5s et Structure de 1'Etat Solide, CNRS UA-234, Bat. C6, Universite des Sciences et Techniques de Lille, F-59655 Villeneuve d'Ascq Cedex, France

R6sum6 Cet article resume les principaux resultats obtenus par les techniques EBIC et Cathodoluminescence (CL) sur l'origine de la recombinaison des porteurs de charge aux dislocations dans le silicium et les semiconducteurs compos6s, ainsi qu'aux joints de grain symetriques dans le silicium. Pour ces deux types de defauts seront presentees des experiences qualitatives reliant 11activit6 Qlectrique du defaut avec sa structure et la presence de precipites. Quelques experiences quantitatives de contraste sur les

dislocations, analysges theoriquement, ont dgalement permis de deduire l'origine de la recombinaison. Celle ci n'est toujours pas Qtablie de facon certaine pour les

dislocations, alors que sa nature extrinsgque semble dorenavant incontestable pour les joints de grain.

Abstract This paper reviews the major EBIC and CL experiments on dislocations in silicon and compound semiconductors and on symmetric grain-boundaries in silicon, from which the origin of the minority carrier recombination could be inferred. Qualitative experiments were often combined with transmission electron microscopy in order to correlate the defect electrical activity with its structure and with the presence of precipitates.

Quantitative contrast experiments were sometimes theoretically analysed to deduce the nature of the recombination. The conclusions are not still well-established for dislocations, whereas the extrinsic origin of the recombination at symmetric grain- boundaries in silicon has been already ascertained.

I-INTRODUCTION

The intrinsic or extrinsic nature of the recombination of minority carriers at extended defects has been an open question over now many years. Local methods such as Electron Beam Induced Current (EBIC) and Cathodoluminescence (CL) appeared very attractive, powerful and sensitive in this research area, as they both allowed the characterization of the electrical activity of an isolqted defect in well-defined surroundings.

i/- intrinsic recombination occurs via the core of the defect when deep and/or shallow levels are present in the forbidden band gap of the semiconductor; they are associated either with dangling bonds, or with a reconstructed core or with imperfections (such as solitons on dislocations [I], or with kinks [2]). The presence of such defects can be predicted by microscopic experiments such as High Resolution Electron Microscopy (HREM) [3-51 and can be associated with energy levels by theoretical calculations [6-81. These levels can also be evidenced by macroscopic experiments such as Hall effect [9,10], photoluminescence 110-131, electron spin resonance (EPR) 1121, deep level transient spectroscopy (DLTS) [13-151,

...

^But,

in these experiments, it is not always straightforward to detect the participation of the defect to the signal, and to assign precisely its energy level(s); this is mainly due to the usual presence of surrounding impurities and point defects, and also to the averaged results over a large number of dislocation types

.

ii/- extrinsic recombination of carriers may occur via the electrically active impurities and/or point defects which either decorate the core of the defect, or which are segregated around it. They can be present in the form of a Cottrell atmosphere, or as precipitates.

It can be innnediately noticed that in many cases, both types of recombination can be expected, the preponderance of one type on the other being related to the defect charge and to the ionization of the surrounding point defects. It seems therefore quite obvious that recohination will depend on the character of the defect, on the type of semiconductor (elemental or compound), on the experimental conditions under which it has been introduced in the crystal, on the semiconductor doping type and level, as well as on the presence of residual impurities. Each material and defect can be conqidered to a certain extent as unique from the point of view of recombination.

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1989604

C6-48 R E W E DE PHYSIQUE APPLIQUEE

Dislocations and particularly grain-boundaries (GBs') have been mainly characterized in silicon, because of its technological importance. As the luminescence efficiency of silicon is quite low, a major part of the literature on recombination deals with EBIC experiments.

This technique has also been used for compound semiconductors, as it allows an easier quantitative characterization of the bulk (diffusion length for instance), and of the contrast of defects, than CL does. But, nevertheless, dislocations in compounds have been predominantly studied by CL (even in semiconductor with an indirect band gap as Gap) because no specimen preparation is required, and also because local spectroscopic investigiitions are possible.

The first important studies on the recombination at defects have been made by observations of the same defects with either a Scanning Electron Microscope (SEM) [16-311, with a Scanning Transmission Electron Microscope (STEM) or a TEM [32-341 in the EBIC and/or CL modes, combined with Transmission Electron Microscopy (TEM) analysis to assess the character of the defect [16-341, its degree of dissociation (in the case of dislocations) [20], the presence of precipitates [18,19,20,22,27-29,351.

The influence of heat treatments on the contrast of extended defects has also been much studied, and has brought many information [21,24,27,36-381, as did the influence of the diffusion of specific species at GBst[29,35,38,39].

It can also be noticed that a quantitative analysis of EBIC contrast values, in order to derive the electrical activity of the defect has been -and still is- more frequent for dislocations [40-431 than for GBs'. CL contrast values of dislocations have been also theoretically analysed, taking advantage of the possiblity to perform experiments and compare results on specimens with very different doping levels [44,45].

Experiments such as the variation of EBIC contrast with specimen temperature and electron beam current, which bring many information about the recombination at the defect have been made already few years ago on dislocations [41-43,46-491, but only very recently on GBs'in germanium [50].

This paper will not be an exhaustive review, as have been previous ones [51-531, but it will give and discuss the most important results obtained on defects in silicon and 111-V compounds.

11- RECOMBINATION AT DISLOCATIONS

Grown-in and diffusion-induced dislocations have been mainly characterized in silicon, by collecting the EBIC signal with p-n junctions. Schottky barriers configuration has been preferred in the case of deformation-introduced dislocations in silicon ~ o m p y y d s , ~

The doping levels, of either n or p type silicon, range from 10Bdto 10 cm

.

^The

dislocations are often parallel to the surface, but located at different depths from one specimen to another. Futhermore, they have been observed with different experimental confitions ( ccelerating voltage Eo in the range 5-35 kV, beam current I ranging from 10 to lo-' A; the detection limit of the amplifier can vary from 0.P to 1%); very often the bulk diffusion length L, or minority carrier lifetime is not reported. Few authors have taken into account some of these parameters (mainly Eo, dislocation depth, L, surface recombination velocity v and once I ) within EBIC or CL theories in order to either confirm experimental observation; [40] or to assess the nature of the recombination from experimental b contrast values [41,43,441. Comparison between dislocations of different character located at various depths under the surface could also be made by use of an appropriate theory [54].

In other cases, when quantitative analysis was undertaken, the contrast of dislocations was compared in one specimen after it was checked, by a careful1 TEM analysis, that they were located at the same depth. Temperature dependence of contrast was used to derived information about the nature of the recombination [41,43], to propose a mechanism of recombinations [40]

and also to locate the energy levels in the band gap associated with the dislocation [43,47,48].

A comparison of contrast values found in different specimens will not be undertaken in this paper. They will be sometimes noticed when necessary, for instance when they are very different from one author to another.

11-1 Extrinsic recombination

By combined SEM(EBIC)/TEM investigations, Kittler and co- workers [24,25,40] could not observe any EBIC contrast from grown-in Shockley dislocations (b=a/6 <112>), Frank partial dislocations (b=a/6 <Ill>), stair-rod partial dislpgatio-y (b=a/6<110>) and complete dislocations (b=a/2 <110>) in phosphorus-doped (n=5.10 cm ) as-grown silicon epilayers.

After one hour time annealing at 900 K in an inert gas atmosphere, or after a storage for few months at room temperature, they could show that 1% of stair rod, 55% of Shockley partials and all the complete dislocations, as well as all the Frank partials bounding the stacking

faults, exhibited an EBIC contrast. As no precipitates could be detected by TEM [25], deep impurities (heavy metals) gettered at dislocations with various efficiencies, were thought to be responsible for the observed electrical activity. A theoretical analysis of part of these results has been made in the case of large diffusion lengths L which supports these conclusions [40]. The model of Kittler and .Seifert relies on an estimation of the minimum detectable line density of states of a dislocation, e which is found to be greater than that really produced by a 'clean' dislocation a126%#; them to conclude that such a dislocation could not be detected with their EBIC system. The model is as following: in the charge-collection Schottky barrier configuration, the contrast of dislocations is given by the relation [55,56]

c =

r

function'(R, L, defect geometry) (1)

where R is the electron penetration depth; T, which is the dislocation strength, characterizes its electrical activity. If the.dislocation is described as a pipe reduced diffusion length Ld, of dimension 1, and of diameter" small as compared to L,

r

can be written as:

When (L /L)L <<1 (case of large L), (2) becomes d

2 2 r = 1 / L d

The main assumption of the model is that the minority carrier lifetime T inside the dislocation is controlled by non interactive deep centers, of kind d, occupie$ by majority carriers. This allows them to write

I

T =

N o v d d th

where ND is the density of recombination centres homogeneously distributed inside the defect volume their capture cross-section and v the minority carrier thermal velocity. Using L

= (DT)'"~(D: minority carrier diffu~ivity),~k~uation (3) becomes:

therefore,e (cm -1 ), the line density of deep centers, can be related to the contrast c (equation

(b)

^{) :}

F is a geometrical correction factor.

In order to find ed m. , Kittler and Seifert estimate the minimum measurab-1% co-qrast c . to be 0.5 %; assumin$ #Rat the hole capture cro? section o is equal to 10 cm

,

as g%& by various authors, a value of about 2.10 pm '-is 4 found Qor e

.

This is about 7 times greater that the line density of states at a Shockley partial

%fh%

has two dangling bonds per lattice constant; this confirms that these dislocations cannot be detected by conventionnal EBIC systems, unless they are decorat~p~by-ypurities. As a matter of fact, assuming, this time, a capture cross-section of 10 for impurity deep centers (as found for gold and copper in silicon), the authors find that a line density of impurities 120 times smaller than that of dangling bonds woultilgrodu e the same contrast. It has to be mentioned that the capture cross-sections of 10 cm-' has been derived for deformation- introduced dislocations, and that Kittler and Seifert have studied grown-in dislocations.

The same conclusion has been drawn by Menniger et a1 [21] who performed identical combined experiments on pure edge grown-in dislocations (b =a/2 <110> and <loo>) parallel to the glide directions, in n and p type Si specimens. The enhancement of the EBIC contrast after an annealing in the range 750-1100 K was coupled to an increase of the bulk diffusion

length. This was considered as a proof that the presence of a Cottrell atmosphere (probably of copper in that case) around the dislocation was responsible for the recombination.

All the previously described observations on grown-in dislocations are in good agreement with those (by EBIC only) of Bodarenko et a1 [37] and of Castellani et a1 [36] on deformation-introduced dislocations in phosphorus and boron doped silicon. These authors have reported a higher EBIC activity of 90° dislocations than that of screw dislocations, as being due to a greater contamination of 90°.

In all these experiments, gettering of impurities has been either supposed or theoretically confirmed, but has not be detected. This is perhaps due to a not so efficient segregation. Gettering is of course better evidenced when the 'dot and halo' contrast (grown-in dislocations perpendicular to the surface in Te doped GaAs for instance [26,57]) or when the 'line and halo' contrast (90" diffusion-induced dislocations parallel to the surface in boron doped silicon [19]) is observed. In that cases, it seems obvious that impurities must play a predominant role, even if the dislocation core is active. So, recording of the EBIC or CL contrast of such defects, and comparison with those obtained on cleaner dislocations _could appear very attractive. Unfortunately, it seems that in few cases, Te doped GaAs specimens for instance, the dot does not correspond to a single dislocation, but to an irregular tangled dislocation with associated large loops on which impurities (heavy metal atoms) have segregated [26]. This means that a detailed TEM analysis of the dislocations under studied must be carried out, if possible, if one does want to avoid erroneous conclusions.

Such precautions have been taken by Blumtritt and coworkers [18,19,22,231 who studied the dependence of the EBIC contrast with the character of the dislocation. Grown-in Frank partials were observed as being electrically active in boron-doped silicon epilayers. By quantitative EBIC contrast measurements, 90" dislocations (b =a/2 <110>) were found to be more active than 60° dislocations, and much more active than screw dislocations. Precipitates on dislocations gave also rise to a greater EBIC contrast. These results were again explained in terms of a stronger gettering of edge dislocations. An obvious proof of the direct correlation of the electrical activity with the degree of decoration, was that a strongly decorated 60" dislocation, with the line and halo contrast, exhibited a higher contrast than an adjacent 60" dislocation not surrounded by a Qalo [23].

- 3 In extreme cases, such as dislocations in n doped silicon (phosphorus = lo2' cm ), the correlation between a high EBIC contrast and the presence of impurities and precipitates (futhermore detected by TEM) at the dislocation, as reported by Lesniak and Holt [46] seems the most probable.

The history of the material under study, the nature of the intrinsic and extrinsic point defects, as well as the way the dislocations have been introduced seem very important in the control of their electrical activity, and thus in the nature of the recombination. This means that the comparison between dislocations in various types of specimens can be made only if one has many more data on 1/ the bulk parameters (diffusion length for instance)

2/ the contrast of the defect on which has already been made a depth correction 3/ the nature of the impurities and point-defects which are aggregated at the dislocation.

-3 Deformation-introduced screw dislocations in phosphorus doped silicon samples (n=9.10 15 cm ) were shown, by Ourmazd et a1 [48] to be more active than 60" dislocations in the presence of point defects -produced by the deformation and detected by EPR-. In their absence, 60" dislocations were more active than screw dislocations.

From variation of the contrast of these dislocations with the electron beam chopping frequency, this influence of point defects on the recombination at dislocations was easily evidenced [47].

At first sight, these results could appear in contradiction with those of Blumtritt and coworkers [18,19,22,23]. But, it has to be kept in mind that the dislocations were not introduced at the same temperature in both experiments. Futhermore, the assumed point defects are impurities in Blumtritt et a1 experimements, whereas they are effectively point defects in Ourmazd et a1 experiments.

This points out the fact that great care has to be taken if one does want to avoid misleading interpretations of one set of experiments. Of course it would be better to know the particular nature of the point defects and impurities, as well as the way they segregate at the disloc&ions. This can be made in compound semiconductors by spectroscopic CL investigations[58] and in all semiconductors by scanning DLTS [59]. Another way is to develop a model which takes into account the band structure of the dislocation, and to compare theoretical EBIC contrast values with experimental ones.

Such an EBIC theory has been proposed by Wilshaw and Booker [43]; it describes the recombination of electron-hole pairs at a negatively charged dislocation in n type silicon.

To calculate its recombination strength

r

as defined by Donolato [55,56], to which the EBIC contrast was found to be proportional, they assume that i) the dislocation is screened by a space charge region of depleted majority carriers where holes have a reduced carrier lifetime

-id;

i i ) the free holes are captured by the attractive potential around the dislocation into

bound hole states at the top of the valence band, iii) the capture of the majoritr, carriers is dependent on the height of the barrier around the dislocation iv) as the holes are not re-emitted in the valence band, rd depends on their cascade capture. The dislocation strength i" was thus found to be a function of the capture of majority carriers and of the recombination of minority carriers at the dislocation level. 'I was dependent on dislocation parameters such as the number of states Nd per unit dislocation length, the position of the energy level, the temperature T, as well as on experimental conditions such as the electron beam current Ib. Experimental curves of the variation of EBIC contrast versus Ib and also versus T were made on straight screw and 60° dislocations. Wilshaw and Booker, after fitting of theoretical and experimental curves, suggested that the electrical acJivi_ty could be due to impurities at the core of the dislocation, N being too low (1.6~10 pm ) to ensure a recombination via the dislocation states. d

The amount of results on the recombination at dislocations is much lower in compound semiconductors than in silicon.

A detailed analysis of the CL contrast of deformation-introduced dislocafions s be mafig by IUamatsu et a1 1441 in InP specimens of various doping levels (n-10

,

1 0 , ''01

cm ) . Microhardness tests at 490K on a (001) face allowed them to distinguish 60°a (P core)

dislocations from 60°$ (In core) ones inclined at 45O degrees from the surface. a and f3 dislocations were in the glide configuration, and both had the same CL contrast, whatever was the doping level.

In their theoretical analysis, Akamatsu et a1 describe the dislocation as a cylinder of radius X 136,601 X being the Debye-Hiickel screening length, inside which the minority carriers have a reduced lifetime rd; the bulk lifetime r was determined from CL intensity versus beam voltage curves. As the electrical activity of the dislocations was characterized by the ratio -rd/r, a theoretical analysis of CL contrast, based on the resolution of the three-dimensionnal continuity equation in the presence of a defect was undertaken. Comparison of theoretical and experimental contrast values, rd/r being an adjustable parameter, allowed them to notice that the recombination efficiency of the dislocations increases with the dopifg lev_q. Equation (4) of Kittler and Seifert gavClg rec_Pination centre density N

d Of 2.10 cm

,

assuming a capture cross-section of 10 cm

.

As this corresponds to

age

occupied centre among 10 sites along the disloca&+on -lne, for a doping level of 1015 cm , and to more than one trap for each site for n=10 cm

,

it was believed that the impurities segregated at the dislocations over a distance larger than the Debye-Huckel length, were at the origin of the recombination, at least in doped materials.

Another experimental and theoretical way to explore the nature of the recombination by cathodoluminescence, was developped by Dimitriadis [45] on dislocations in vapor phase epitaxy (VPE) Gap layers. The dislocation was theoretically treated as an internal cylinder surface of radius r with an infinite non-radiative surface recombination. The minority carriers were assume8 to recombine by diffusion to the dislocation. By solving the transient continuity equation in cylindrical coordinates, with judicious boundary conditions, the dislocation efficiency, E, could be calculated as a function of r

,

and of the bulk diffusion length L. E was then related to the lifetime rd corresponding Po the recombination at the dislocation by:

1 1 1

with - - - -

+

-

T' T

rd

r' was the lifetime when the electron beam was positioned at the dislocation, r was the bulk lifetime.

By measuring the decay of the green CL at and away from the dislocation, E was found to be equal to 10% and 32% in VPE undoped and VPE doped but Cr diffused (1200K for 24 H) layers respectively. The bulk diffusion length was decreasing by Cr diffusion. By fitting the experimental values of E with the theoretical E versus L curves,, a radius r

<

0.1 nm and r =0.02 pm were respectively assigned to the dislocations in VPE undoped and ?n Cr diffused layers. Thus, the nature of the recombination at dislocations was found to change from 0

intrinsic (core effects) to extrinsic (Cottrell atmosphere of impurities) when increasing the impurity concentration. Futhermore, spectroscopic CL investigations of the defects segregated at the dislocations, could evidence an enhancement of a vacancy-donor complex which, associated with impurities, could explain the observed reduced lifetime [61].

Combined with a theoretical analysis of the CL contrast, this kind of experiment could give very useful information about the recombination behaviour of dislocations in compounds.

Nevertheless, only very few data of the modification of the CL spectra brought by

electrically active dislocations are available in the literature. Petroff et a1 1331 have observed by STEM/CL a small shift in the emission energy of the 'bound exciton" or electron at some molecular beam epitaxy (MBE) growth-induced dislocations in undoped GaAs. The CL analysis was made at temperatures 16K-400K, with a beam voltage of 100 kV. The structural properties were assessed by STEM and TEM analysis. The bound exciton could be correlated with the presence of single isolated dislocations or dislocation clusters.

11-2 Intrinsic recombination

The contrast of dislocations is not always related to the presence of extrinsic centres at the dislocation, as it is often suggested when the extrinsic origin of the recombination is asserted. This point has been evidenced by low temperature STEM/CL investigations, made by Myhajlenko et a1 [34] on grown-in dislocations in undoped n type as-grown and heat-treated

(T=1023 K) LEC InP specimens. They could observe the quenching of the exciton luminescence by some screw dislocations, and not by dipoles, which they attributed either to the presence of dangling bonds, if.the screw are dissociated, or to their strain field inducing deformation potential and piezoelectric effects.

Petroff et a1 [32] argued that a possible reconstruction of the core of an edge Lomer-Cottrell dislocation (b=a/6<110>) in Ga

AlxAsl-yP heterostructures could be at the origin of its electrical inactivity. Core recok'fructlon A u l d then be produced by a complete elimination of dangling bonds and kinks.

Lowering of the total core energy by reconstruction of dangling bonds seems less favourable in compound than in elemental semiconductors, because bonding of atoms of the same chemical nature is not so likely to occur 1621.

The increase of the EBIC contrast of grown-in edge dislocations with their degree of dissociation, as okp;m~5) by . combined SEM(EBIC)/TEM(Weak-Beam) experiments in phosphorus doped silicon (n.10 IS, following Ourmazd and Booker [20] due to new recombination centres introduced by the core of the Shockley 60° partials in which the perfect edge dislocation is dissociated. Electrical inactivity of the intrinsic stacking fault (SF) between them was assumed, as all the extrinsic SF never gave rise to a contrast higher than 0.1% (detection limit). Ourmazd and Booker also observed an increase in the electrical activity of Frank dislocations when their direction changed from <112> to <110>. This was associated with a modification of their core structure, as their Burgers vector and character remained equivalent.

The measured EBIC contrast in Ourmazd and Booker's experiments was in the range 0.1-0.6%, with a peak around 0.2-0.3%. These values are lower than the detection limit of Kittler and Seifert's EBIC system (0.5%) who claimed that clean dislocations could not be detected. This has to be kept in mind, even if, as argued by these authors, the results are difficult to compare because of the different geometries used. The same remark can be made for Blumtritt et a1 experiments, who associated electrical inactivity of dislocations with the absence of impurities (detection limit -1%).

Different recombination activities of neighbouring 60° dislocations, evidenced by Heydenreich et a1 1231 and by Pasemann el a1 [ 5 4 ] , have been also related to different states of dissociation into 30° and 90° partials, stressing the fact that not only impurities are involved in the recombination process.

HREM observations [3,4] and theoretical calculations [6,7] have shown that core reconstruction of 30° and 90° partials in the glide configuration (30°G, 90°G) is energetically more favourable than that of 60°G partial. The 60°G partial dislocation only is associated with energy levels in the band gap, due to its highly distorded reconstructed bonds. Therefore, the higher electrical activity 'of 90° perfect dislocations observed by Blurntritt and coworkers [18,19,22,23] could have also an intrinsic origin, due to the 60°

partials, in contrast to the suggestion of the authors.

Core reconstruction is not a priori constant all over the partial length. This is especially true for the 30° partial which could have a core containing G sites as well as shuffle interstitial S. and shuffle vacancy Sv sites, as proposed by Louchet and Thibault-Desseaux [62].

A.

and S sites are obtained from G sites by emission or absorption of a row of point defects? 30°G gore is easily reconstructed, even if solitons with a single dangling bond can remain; reconstruction of 30°Si core is unlikely to occur, and that of 30;Sv core could only be partial. On the opposite, 90° partial seems to have a more or less unlform glide core.

These considerations cannot explain why, in many experiments previously mentionned, the 60° perfect dislocations (made of 30° and 90° partials), has been observed to be more active than the perfect screw (made of two 30° partials). Therefore, the extrinsic origin of the recombination, as suggested by the authors seems, up to now, the most reasonnable and straightforward; it has just to be noticed that the crystallographic structure of the partial dislocations (core reconstruction, altenate glide and shuffle segments, . . . ) was derived from HREM images taken on ingots different from those used for EBIC experiments.

Therefore, in all the experiments described in this paper, the electrical activity of dislocations cannot be surely correlated with their crystallographic and band structures.

~t can seem surprising that Ourmazd [41] has excluded the participation of impurities to the recombination atlgrown-'n 90° perfect and dissociated dislocations when the specimen was

-

5

heavily doped (n=10 cm ) . First of all, it was argued that any precipitate on the dislocations could be detected. Secondly, Ourmazd correlated the EBjLC contrast to the number N of recombination centres within the generation volume equal to 8b :

where od, vth and D h e th ir usual meanings (cf 5 11-1-a). b=l p.

By taklng o =lo-" cm-', N was found to be equal to about 100. This value was much lower than the number of dangling bonds along the segment of an undissociated shuffle edge d dislocation (*5000), and therefore consistent with recombination centres at the dislocation core itself.

Controversial capture cross-sections in silicon appear in the EBIC literature.

ThBk

tak n by Ourmazd is 100 times greater than that used by Kittler and Seifert [40]. od=10 cm-' is the capture cross-section of shallow levels in silicon. The justification of such a high o comes from a similar cascade capture process at J&se i'vels and at the dislocation, as first shown by Ourmazd [41]. Another value of o =10 d cm is also available; for the first time it has been deduced by fitting experimental and theoretical variations of the d minority carrier lifetime obtained by EBIC experiments, with the deformation-introduced dislocation density 1631. It is now usual, in EBIC contrast theories, to correlate the reduction of the minority carrier density with a reduced minority carrier lifetime T~ inside the volume of the defect. This concept is valid for independent recombination centres located outside the action of any electric field; but, unfortunately, this is not the case for dislocations. Therefore, a new description of the contrast, which is not phenomenological, and which involves the band structure of the dislocation, seems appropriate to go further in the understanding of the recombination process. In the model of a two stage recombination developped by Ourmazd, the capture of electrons by the dislocation is not rate controlling.

At high temperatures, the thermal reemission of holes captured at the hole band is predominant, whereas their capture by the deep level will control the recombination at low temperatures. A detailed description of this model is beyond the scope of this paper, but let us just notice that these assumptions are opposed to those of Wilshaw and Booker [43].

Fitting of theoretical and experimental EBIC contrast variation with,temperature, in the range 120K-300K, confirmed Ourmazd's previous assumptions that the energy level of a dissociated 90° dislocations was lower-lying than that of a undissociated 90°.

EBIC contrast versus temperature curves have, following Ourmazd, to be carefully interpreted with the help of a recombination theory. He assumed, due to the high doping level of his specimen, that the Fermi level was constant over the temperature range, and that only the variation of the controlling processes with T had to be considered.

In contrast, Kimerling and co-workers [47] have derived the energy levels of stacking faults and Frank partials by interpreting their EBIC contrast versus T curves (T=80&220KJ only in terms of Fermi level variation. The specimen had a low doping level of 5.10 cm and was first heat treated at 700°C for 2 hours before a second annealing at 1200°C for 16 hours in an H atmosphere which developed the stacking faults. The contrast of the SF was

2 .

decreasing when increasing temperature, from about 10% at 80K to 2% at 220K, with a large transition of 10KT; it was associated either with a high spatial density of interdependent states, producing a broad energy level, or with more than one level. Although these values are much higher than those observed by Ourmazd and Booker [20], Kimerling et a1 claimed that the electrical activity of the SF was associated with the fault itself, because they did not detect, by TEM, any precipitation. They could locate the SF energy level in the upper half of the band gap; that of the Frank partial bounding the SF was lying deeper in the gap, as its contrast value, independent on its direction, in contrast to Ourmazd and Booker's results

[20], was constant with T (~~14%).

To be able to characterize the electrical activity of a dislocation, any contrast measurements have to be made within a well defined experimental procedure. As a matter of fact, a saturation of the recombination process, for instance, can occur even for electron beam currents which correspond to a low injection level [43]; futhermore, contrast versus T curves can be drhstically modified by changing the electron beam power, as eviqenced by Lesniak and Holt [46].

C6-54 REVUE DE PHYSIQUE APPLIQUEE

Thus, only a theory in the sense of those developped by Ourmazd [41] and by Wilshaw and Booker [43] which takes into account: - the experimental conditions (accelerating voltage Eo, electron beam current Ib) .

-

the junctlon geometry and the bulk parameters (diffusion length, temperature)

-

the band structure of the dislocation and

-

the recombination processes,

will be able, after comparison with experimental results, to bring more valuable information about the intrinsic and/or extrinsic nature of the recombination. Futhermore, in addition to the electric field produced by a charged dislocation, strain fields as well as piezoelectric fields in compound semiconductors should be considered.

111- RECOMBINATION AT GRAIN-BOUNDARIES IN SILICON

All the works undertaken on the recombination at GBs in silicon have led to the conclusion that its origin was extrinsic.

Such an unambiguous conclusion could have been drawn because, in contrast to dislocations, coherent planned experiments have been made. From the results of the French group 'ARC silicium polycrystallin', the origin of the recombination at the most important GBs in silicon is now well-established.

The growth of specific crystals, essentially bicrystals, has been fundamental for the determination of the nature of the recombination. The electrical activity of GBs, the microscopic observations of their structure, by HREM, as well as their band structures, by means of DLTS, have been characterized in specimens coming from the same ingots; theoretical calculations of the energy levels have also been performed on well-known structures. The EBIC electrical activity of the GB has currently be correlated with its structure, by conventional TEM diffraction apalysis, as well as with the presence of impurities, precipitates by microanalytical microscopy.

It is now usual to characterize a GB by its coincidence site lattice (CSL); in this model, the C index is the reciprocal density of sites in coincidence. In the following, only the results obtained on symmetric GBs will be described, as they have been mainly studied. It can just be worth noticing that aspetric GBs (C=9 {lll)//{115} or I=27 {111)//{1 11 11) for instance), exhibit a stronger EBIC contrast than that of symmetric GBs. This could be explained by the presence of a high density of defects such as dislocation and precipitates.

Symmetric GBs were mostly issued from bicrystal ingots. Their interface plane is a mirror plane. They exhibit a lower electrical activity when their rotation axis is [Oll]

(X=3 {lll}, X=9 {122)) than when it is [001] (X=13 {510), C=25 (710)). The two former ones, C=3 and C=9, have been shown to be reconstructed without any dangling bonds 1641, the bonds in X=9 being a little distorded. As for these two GBs no energy levels are introduced in the band gap, the presence of impurities was evoked to explain their recombination activity. The same statement was made for X=25 (710) and X=13 {510), even if any theoretical calculations have been performed on these GBs, to assess the positions of their energy levels. But convergent observations by HREM and DLTS [15,65] indicate that there are no dangling bonds at GBs C=25 and X=13 cores, and no energy levels in the gap. Many EBIC experiments have evidenced that all these four GBs are electrically active only after a heat treatment [27-291.

The decoration bv impurities of primary dislocations in GBs with a small deviation to

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coincidence qives rise to an enhancement of their electrical activity in comparison with the same GBs with an exact coincidence relation [28].

Battistela observed that annealing of a bicrystal under nitrogen at 750°C for 2 hours gave rise to an heterogeneous EBIC contrast at a C=9 (122) GB [27]. By chemical etching,,the dark spots were found to correspond to precipitates with a high local concentration of oxygen; they perhaps also contain heavier atoms, such as copper [28] which is the major contaminating impurity in silicon. When the annealing time was increased up to 24 hours, the activity of the GB X=9 became homogeneous, due to a higher density of precipitates at the

interface [271.

The GBs X=13 {510} and X=25 1710) behave similarly with respect to annealing.

Futhermore, a gettering effect at the GB C=25 C7103, detectable by the appearance of a 'line and halo' EBIC contrast, was observed after annealing at 750°C for 24 hours [27].

The carrier recombination at GBs is also annealing temperature dependent: the EBIC contrast of a C=13 {510) GB has been observed to increase with temperature 1271.

A detailed DLTS study by Broniatowski [15] evidenced the strong dependence of the density of states of the GB C=25 (710) on the annealing temperature and atmosphere. The exact origin of this effect is not yet fully understood, even if the participation of oxygen and carbon, the main residual impurities in silicon, could be suggested; but, it emphasizes the

fact that segregation of different impurities can create different recombination centres at GBs, This phenomenom has also been observed by Chari and co-workers 1381 who performed hydrogenation by microwave plasma annealing at 320°C for 6 hours on polycrystals pre-annealed at various temperatures under high static vaccuum for 24 hours. The total passivation of the GB C=25 (710) occured when the pre-annealing temperature was 450°C; but hydrogen did not interact with the recombination centres created at 750° and 900°C.

A direct proof of the extrinsic nature of the recombination at GBs was given recently by Bary and Nouet [31] for the GBs C=9 (1223 and C=3 (111). By undertaking combined SEM(EBIC)/TEM experiments on polycrystalline silicon, they could show that a GB C=9 {122) was electrically active when located between two grains, whereas no activity could be detected when it was inside one grain. It was argued that the impurity concentration being higher at the interface than in the grain, because of growth, such observations could be explained in terms of impurity segregation. These authors also correlated that lack of activity of the GB C=3 {Ill), whatever was its position in the crystal, to its perfectly reconstructed structure, which turns to be a less efficient sink for impurities than the GB Z=9 {I223 with its distorted bonds.

The electyical activity of GBs C=3 {Ill} has been found to be correlated with impurity decoration [29] or with the presence of a periodic network of Frank partials which accommodate the deviation to exact coincidence. But, even in that case, dislocations have been described to play a role in the recombination not because of their electrical core properties, but because of their attractive strain field for impurity aggregation.

Therefore, the recombination at symmetric coincidence GBs is not related to dangling bonds, but is only dependent on the GB- impurities interaction which still needs further experiments to be clarified. As a matter of fact, some impurities enhance the electrical activity of the GBs, whereas other ones passivate them; to the first group belong oxygen, carbon, copper, to which can be added alurnino-silicate precipitates on C=3 {Ill}, as identified by Maurice [35]; to the second group belongs hydrogen and aluminum.

The structure of the GB remains of course very important; for instance, a facetted Z=3 'GB has been observed to have specific recombination sites located at the intersection of the (111) and {112} planes; such a phenomenom could not occur in straight GBs [30].

ACKNOWLEDGMENTS

The author would like to thank J.L. Farvacque for his valuable comments on the manuscript as well as J.L. Maurice for discussions about the recombination at GBs.

Heggie,M.I. and Jones,R., J. de Physique C4 (1983) 43 Hirsch,P.B., J. de Physique C6 (1979) 27

Bourret,A., Desseaux-Thibault J. and Lancon F., J. de Physique C4 (1983) 15 Hutchison,J.L., J. de Physique C4 (1983) 3

Bourret,A. and Bacmann,J.J., Revue Phys. Appl. 22 (1987) 563 Marklund,S., J. de Physique C4 (1983) 25

Teichler,H. and Veth,H., J. de Physique C4 (1983) 93

Mauger,A., Bourgoin,J.C., Allan,G., Lannoo,M., Bourret,A. and Billard,L., Revue Phys.

Appl. 22 (1987) 579

FerrB,D., ThSse d'Etat Lille (1987)

Schroter,W., Scheibe,E. and Schoen,H., J. of Microscopy 118 (1980) 23

Depraetere,E., Vignaud,D. and Farvacque,J.L., Solid State Comm. 6 4 (1987) 1465 Weber,E.R. and Alexander,H., J. de Physique C4 (1983) 319

Hamilton,B., Peaker,A.R. and Wight,D.R., J. de Physique C4 (1983) 233 Gelsdorf,F. and Schroter,W., Philos. Mag. A 49 (1984) L35

Broniatowski,A., Revue Phys. Appl. 22 (1987) 585

Titchmarsh,J.M., Booker,G.R., Harding,W., and Wight,D.R., J.Mat.Sci. 12 (1977) 341 Ourmazd,A., Darby,D.B. and Booker,G.R., Inst. Phys. Conf. Ser 36 (1977) 251 Blumtritt,H. and Gleichmann,R., Ultramicroscopy 2 (1977) 405

Blumtritt,H., Gleichmann,R., Heydenreich,J. and Johansen, H., Phys. Stat. Sol. (a) 55 (1979) 611

Ourmazd,A. and Booker,G.R., Phys. Stat. Sol.(a) 55 (1979) 771

Menniger,H., Raidt,H. and Gleichmann,R., Phys. Stat. Sol.(a) 58 (1980) 173

Heydenreich,J., Blumtritt,H., Gleichmann,R. and Johansen,H., Cryst. Res. and Technol.

1 6 (1981) 133

C6-56 REVUE DE PHYSIQUE APPLIQUEE

ueydenreich,J., Blurntritt,H., Gleichmann R. and Johansen,H., SEM I (1981) 351 Kittler,M. and Bugiel,E., Crystal Res. and Technol. 17 (1982) 79

Bugiel,E., Kittler,M., Borchardt,A. and Richter,H., Phys. Stat. Sol.(a) 84 (1984) 143 Chu,Y.M., Darby,D.B. and Booker,G.R., Inst. Phys. Conf. Ser. 60 (1981) 331

Battistela,F., These Toulouse 1985 and Battistella,F. and Rocher,A., Semicond. Sci.

Technol. 2 (1987) 226

Rocher,A., Revue Phys. Appl. 22 (1987) 591 Maurice,J.L., Revue Phys. Appl. 22 (1987) 613

Laval,J.Y., Maurice,J.L. and Cabanel,C., Revue Phys. Appl. 22 (1987) 623 Bary,A. and Nouet,G., J. Appl. Phys. 63 (1988) 435

Petroff,P.M., Logan,R.A. and Savage,A., J. of Microscopy 118 (1980) 255

Petroff,P.M., Weisbuch,C., Dingle,R., Gossard,A.C. and Weigrnann,W., Appl. Phys. Lett.

38 (1981) 965

Myhajlenko,S., Batstone,J.L., Hutchinson,H.J.and Steeds,J.W., J. Phys. C 7 (1984) 6477 Maurice, J.L. These 3O cycle. Paris. 1982

Castellani,L., Gondi,P., Patuelli,C. and Berti,R., Phys. Stat. Sol. 69 (1982) 677 Bondarenko,I.E., Yakimov,E.B. and Yarykin,N.A., Acta Physica Polonica A69 (1986) 393 Chari,A., de Mierry,P., Menikh,A. and Aucouturier,M., Revue Phys-Appl. 22 (1987) 655 Martinuzzi,S., Revue Phys. Appl. 22 (1987) 637

Kitt1er.M. and Seifert,W., Phys. Stat. Sol.(a) 66 (1981) 573 Ourmazd,A., Crystal Res. and Technol. 16 (1981) 137

Ourmazd,A., Wilshaw,P.R. and Booker,G.R., J. de Physique C4 (1983) 289 Wi1shaw.P.R. and Booker,G.R., Inst. Phys. Conf. Ser. 76 (1985) 329

Akamatsu,B., Jean-Louis,A.M., Papadopou10,A.C. and Henoc,P., Phys. Stat. Sol.(a) 95 (1986) 557

Dimitriadis,C.A., Solid State Comm. 49 (1984) 1111

Lesniak,M. and Holt,D.B., Inst. Phys. Conf. Ser. 76 (1985) 25

Kimerling,L.C., Leamy.H.J. and Patel,J.R., Appl. Phys. Lett. 30 (1977) 217 Ourmazd, A., Wilshaw, P.R. and Cripps, R.M., Inst. Phys. Conf. Ser. 6 1 (1981) 519 Jakubowicz,A. Habermeier,H.U., Eisenbeiss,A. and Kass,D., Phys. Stat. Sol. (a) 104 (1987) 635

Tabet,N. and Monty,C., this issue Dupuy,M., J. de Physique C4 (1983) 277

Yacobi,B.G. and Holt,D.B., J. Appl. Phys. 59 (1986) R1 Booker,G.R., Inst. Phys. Conf. Ser. 60 (1981) 203

Pasemann,L., Blumtritt,H. and Gleichmann,R., Phys. Stat. Sol. (a) 70 (1982) 197 Donolato,C., Optik 52 (1978/79) 19

Donolato,C., Appl. Phys. Lett. 34 (1979) 80

Boulou,M. and Schiller,C., J. Microsc. Spectrosc. Electron. 6 (1981) 39 Steeds,J.W., this issue

Breitenstein,O., this issue

Sieber,B., Philos. Mag. B 55 (1987) 585

Davidson,S.M. and Dimitriadis,C.A., J. of Microscopy 118 (1980) 275 Louchet,F. and Thibault-Desseaux,J., Revue Phys. Appl. 22 (1987) 207

Casta$fiini,A., Cavalcoli,D. and Cavallini,A., to be published in the proceedings of the 8 Int. Summer School 'Defects in Crystals', Sczczyrk, May 1988

Bourret,A. and dlAnterroches,C., Jt de Physique 43 (1982) C1-1 Maurice,J.L., private communication